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By R. F. Mehl, C. S. Barrett, A. H. Geisler

THE studies of the mechanism of precipitation and of the resulting property changes in aluminum-silver alloys1-3 have presented some new concepts of the aging reaction-concepts that may be fundamental in establishing a complete quantitative theory of age-hardening. With both aluminum-copper and aluminum-silver alloys, thin platelike aggregates form parallel to the Widmanstätten planes of precipitation.1,4 These platelets grow in lateral dimensions and in thickness with aging time until they are thick enough to produce an X-ray diffraction line pattern. Their structure is that of a transition lattice which is coherent with the matrix, since the atomic spacings in the matrix and the precipitate are the same on each side of their common interface.2,4 The equilibrium precipitate forms later from and at the expense of the transition lattice. Since age-hardening occurs during the formation and growth of the transition lattice, the major contributor to the hardening no doubt is the stress condition in the matrix resulting from the matrix holding the precipitate in the coherent condition. The development of these stresses also appears to be responsible for an increase in the electrical resistance superimposed on the normal decrease resulting from a change in the concentration of the matrix during aging.3 The two consecutive reactions constituting the aging process are accompanied by characteristic microstructural alterations. The initial precipitation process results in the formation of a Widmanstätten figure, while the later transformation of the transition lattice to the equilibrium lattice is accompanied by a coalescence, in the case of aluminum-copper alloys,5 and by the appearance of light areas at the grain boundaries with aluminum-silver alloys showing a structural rearrangement of matrix and precipitate.3 This latter reaction is similar microscopically to the so-called "discontinuous precipitation." Possibly discontinuous precipitation is a process that often if not always follows true precipitation of a coherent lattice and during which the transformation of the coherent transition lattice to the equilibrium lattice takes place accompanied by matrix recrystallization.3 EXPERIMENTAL PROCEDURE The alloys were made using 99.97 per cent Al or distilled magnesium, with

Jan 1, 1943

By J. Handin

Triaxial compression tests with independently applied external confining pressures and internal pore pressures show that the ultimate compressive strengths of representative oil well cements are nearly linear functions of effective pressure — the difference between external and internal pressures on the jacketed cylindrical specimens (to 15,000 psi). The strengths are little affected by the test temperature to 350F (not to be confused with the curing temperature). At an effective pressure of 15,000 psi, strengths are in the rage of 30,000 to 50,000 psi, comparable to those of sedimentary rocks under similar conditions. The cements become very ductile even under low effective pressures; permanent shortenings of 30 per cent or more are attainable without rupture. INTRODUCTION Since the pioneering work of Richart, Brandtzaeg and Brown1 on the failure of cement under combined compressive stresses, it has been recognized that ultimate compressive strength is greatly enhanced by the application of confining pressure. More recently, Mc Henry2 showed that the strength of concrete was a linear function of the effective pressure (the difference between the external confining pressure on a jacketed specimen and the internal fluid pore pressure), at least for a range of 0 to 1,500 psi. The effect of temperature had not been investigated, and no previous systematic triaxial compression testing of naterials used for oilwell cementing seems to have been done. The present work was suggested by the late J. hi. Bugbee3 who stated that "consideration of the common application of high-pressure hydraulic fracturing to the initial completion or recompletion of wells, and the large pressure drawdowns in some producing wells, particularly those in abnormally high-pressure gas-condensate reservoirs, raises the question of what is a suitable cement strength for various completions. The intuitive belief exists that cement strength need be no greater than formation strength. Tests should, however, be conducted at downhole conditions." The ultimate compressive strengths of rocks penetrated by the borehole must rise several fold with increasing depth.4 This marked enhancement of strength is due to the influence of the effective pressure, the total weight per unit area of the overburden less the hydrostatic pore pressure. (The effect of temperature due to the geothermal gradients is relatively small for depths to 30,000 ft.) .A significant comparison of the strengths of rocks and cements at downhole conditions requires knowledge of the confined compressive strengths of cements as well. EXPERIMENTAL PROCEDURES The theory and technique of triaxial compression testing are fully discussed in earlier reports.4-6 Briefly, cylindrical specimens 1-in. long and 0.5-in. in diameter are jacketed in thin copper tubes of negligible strength, placed in the test chamber and subjected to an external confining pressure of kerosene and loaded axially by the piston at a strain rate of 1 per cent per minute. Pore pressures of water (or kerosene) are applied independently through the hollow piston and are maintained constant during the shortening of the specimen, Tests at sensibly 0 pore pressure are arranged so that any free water in the cement can escape to the atmosphere. (If egress of water were denied, pore pressure would rapidly attain the value of the external confining pressure because of reduction of pore space.) The test chamber can be heated for high-temperature experiments. Unless otherwise noted, the cement samples were air dried for about a week. Recorded during a test are pore and confining pressures, shortening and axial differential force (total force less the product of the confining pressure and the area of the piston). True stress-strain curves are constructed by dividing the measured differential force by the actual cross-sectional area of the specimen at a given

Jan 1, 1966

By P. Ramachandrarao, T. R. Anantharaman

EVER since the technique of quenching metals and alloys from the melt (splat cooling) was perfected a decade ago, it has been recognized that the grain size of products solidified by this technique may be extremely small.' The further observation that fcc metals quenched from the liquid state contain very few dislocations has led to the inference2 that metals are subjected to negligible or no stresses during the rapid solidification characteristic of the "gun" or "piston-and-anvil" technique. Evidence for the incidence of appreciable densities of stacking faults has, however, been obtained in case of some splat-cooled fcc and hcp alloys,3 although not for pure metals. In the light of these earlier observations it was considered desirable to study X-ray line-broadening effects, if any, in fcc metals rapidly cooled from the melt. In the present work pure silver (>99.99 pct), aluminum (>99.99 pct) and lead (>99.9 pct) were quenched from the liquid state from temperatures about 50°C above the melting point by the "gun technique" and the resulting foils were subjected to X-ray examination in a Philips Diffractometer. The quenched foils (up to -10 u thick) did not generally stick to the substrate surface and could be easily transferred to the Diffractometer without introducing any plastic deformation. The profiles of the first five reflections from the foils were recorded in each case with Cu Ka, radiation at the slowest available scanning speed of 1/8 deg per min. To correct for instrumental broadening, profiles were also recorded from the metals annealed in vacuo at suitable temperatures. The integral breadths of the X-ray reflections were arrived at by a procedure described earlier.' There was a distinct suggestion of preferred orientation in the recorded intensities of reflections from aluminum and lead foils. Such an effect was not observed in case of silver. In addition, the integral breadths of X-ray reflections from splat-cooled aluminum and lead were not significantly different from those recorded for the annealed metals. The analysis was therefore continued only for silver where the X-ray line broadening was appreciable. The pure diffraction broadening, B, was evaluated for each X-ray reflection (hkl) from silver from the observed, B, and instrumental, b, breadths with the aid of each of the three equations due to Scherrer,5 Anantharaman and Christian,6 and Warren and Biscoe,7 respectively: Bs= B-b BAC=B- b2/B Table I gives the values of particle size, 71, the lattice strain, E, arrived at by the use of the following well-known relations and on the assumption that all observed diffraction broadening could be attributed to lattice strain or particle size, respectively: E = 1/4 cos ? n= B cos ? where A is the wavelength of X-radiation and 0 is the Bragg angle. As no significant peak shifts or asymmetry could be detected in the profiles from the foils, the possibility of any significant contribution due to twins or stacking faults was ruled out. The absence of faults is by no means surprising since pure silver is known to develop stacking faults only on severe deformation and the stacking fault densities recorded so far for even silver filings have been extremely low.' The very low values for percentage mean deviation from the mean value for the particle size in Table I strongly suggest that all observed broadening in splat-cooled silver can be attributed only to small particle size. This conclusion receives further support from the lowest mean deviation recorded for data computed from the Scherrer equation based on Cauchy profiles that are considered characteristic of particle size broadening. Further analysis for separation of particle size and lattice strain effects was considered unnecessary in view of the very large mean deviations obtained for strain values and also the earlier results suggesting absence of even detectable strain in metals and alloys quenched from the melt. It is to be stressed in this connection that the particles are actually grains and not cells formed by walls of high dislocation density usually encountered in deformed samples. As such, the absence of strains is not surprising. The present results are probably the first to record X-ray line broadening due only to small particle

Jan 1, 1970

By D. E. Menzie, D. L. Dauben

This paper discusses the physical parameters involved in the slow flow of high molecular weight polymer solutions in porous media. The interacting effects of polymer properties and porous media properties on flow performance are considered. Experiments were conducted with the polyethylene oxides, with molecular weights ranging from 200,000 to over 5,000,000. Frontal advance velocities ranged from I to 30 ft/day. The porous matrix consisted of a flow cell packed with glass beads. Polymer solutions were characterized by viscosity and normal stress measurements. Under certain conditions, unexpectedly high flow resistance was observed. This behavior was observed to be a function of flow rate, pore size, polymer molecular weight and concentration. The polymer solutions exhibit "dilatant" flow behavior in porour media in contrast with the pseudo-plastic behavior in simple flow systems. A theoretical ex-planation of such behavior is presented. INTRODUCTION The use of polymers in the injected water of a waterflood increases oil displacement efficiency by reducing the mobility (k,/pu) of the driving phase. A reduced driving phase mobility results in improvements in the areal sweep efficiency and in the vertical coverage in stratified reservoirs. This mobility reduction may be achieved by a permeability reduction, a viscosity increase or by a combination of the two. Early attempts at increasing the injected water viscosity were not successful because of the poor economics involved. The use of such materials as glycerin, sugar or glycols to increase water viscosity was not economically feasible. Attempts in using certain naturally occurring polymers were not too successful because of the high polymer losses to the rock. A high molecular weight, partially hydrolyzed polyacry-lamide was introduced 3 years ago as a waterflood additive.',' Initial work by Pye' indicated that the presence of these polymers in dilute concentrations decreases the water mobility 5 to 20 times more than would be expected from measurements of the solution viscosity. Such an effect would be of obvious economic value since only a small polymer concentration would be required to accomplish a large reduction in water mobility. This research was designed to study the basic flow mech anisms of polymer solutions in porous media. The interacting effects of polymer properties and porous media properties on flow performance are considered. This study indicates some of the conditions under which these interactions can lead to significant mobility effects. is valid only for Newtonian fluids. For polymer solutions and other non-Newtonian fluids the equation must be modified to consider that viscosity is a variable quantity. Modification of the Darcy equation to include non-Newtonian effects has been the subject of several recent investigations."" The modifications generally adopt a rheological model. such as an Ellis or Dower law model. to Dorous media by defining some characteristic channel radius. Most of these studies showed the porous media flow behavior to be predictable from viscometric data. Investigations involving the flow of high molecular weight polyacrylamide solutions through cores have generally encountered the high flow resistances reported previously by Pye. This high flow resistance has generally been attributed to an in-depth permeability reduction, as evidenced by a reduced permeability to water which has displaced polymer solution from the porous media. Marshall and Metzner" report high flow resistances, which they attribute to viscoelastic effects. In this paper, the effects of polymer molecular weight, pore size, flow rate and concentration are considered. The polymers used are the polyethylene oxides known as Poly-ox. This class of polymers was used because of its availability in a wide range of molecular weights and because of its known ability to propagate well through porous media. THEORY PHYSICAL PROPERTIES OF POLYMER SOLUTIONS The polymer molecule dissolves in water by means of hydrogen bonding, but retains some of its own structural identity while in solution. Nonionic polymers, such as polyethylene oxide, are generally considered to have a random coiling configration. This type of molecule has the ability to "sequester" or hold a large volume of solvent within its coils in a manner similar to that of a sponge. The random coil is easily deformed and under an applied stress

By J. F. Pierce, S. M. Enterline

Since January 1959 a zinc refiner of novel design has been in operation at Blackwell, Okla., producing 99.995+ zinc from the output of the Blackwell horizontal retort smelter. The refiner is a continuously operated, pyro-metallurgical unit with principal dimensions horizonhl permitting one floor operation. It is very ruggedly constructed. Furthermore it is not subjected to deterioration from high concentrations of iron in the feed metal because the boiling units are separate from the rectification columns and constructed with silicon carbide brick arches through which the heat is transferred radiantly to the zinc. A zinc refiner of novel design has been in operation since January 1959 at the Blackwell, Okla. plant of the Blackwell Zinc Co., Inc., a wholly owned subsidiary of American Metal Climax, Inc. Developed by the company's technical and operating staffs, the new refiner produces special high-grade zinc (nominally 99.995+ pct Zn) from the prime western1 slab output of the Blackwell horizontal retort smelter. The refiner is a continuously operated, pyromet-allurgical, high-capacity unit employing fractional distillation to separate zinc from its impurities, consisting principally of lead, cadmium, and iron. In prime western slab, these elements may vary in concentration from 0.005 to 1.6 pct. Table I shows the impurities usually encountered in the Blackwell prime western metal fed to the refiner together with their boiling points and vapor pressures at 907oC. The commercial pyrometallurgical smelting processes for extracting zinc from ore include distillation from retorts and a limited degree of refining can be accomplished by redistillation in similar retorts. To achieve special high-grade purity, an impractical number of batch redistillations would be required, however, with an accompanying loss of much metal. A continuously operated and integrated unit is required in which fractional vaporization and condensation can take place in a large number of stages. The first commercial continuously operated unit was developed by the New Jersey Zinc Co.4 In the separation of any two substances the degree of separation is limited if a mixture exists having a constant boiling point. The Zn-cd5 and Zn-pb6 systems contain no constant boiling points and hence these metals may be separated to any degree of purity. Basically, the pyrometallurgical refining of zinc involves three steps: a) iron, copper, and other metals having no appreciable vapor pressures at 907ºC concentrate in the bath when the impure zinc is boiled, b) lead, bismuth, antimony, and other metals having boiling points above but closer to that of zinc are selectively condensed from the vapor boiled from the bath, and c) cadmium and other metals having lower boiling points than zinc are fractionally distilled by reboiling and subsequent condensation of the zinc from the vapor. The boiling and condensation of zinc must be conducted in vessels constructed of materials which are refractory at the temperature involved (907 ºC) and it is highly desirable that they also be resistant to attack by the various metals. This requirement dictates construction from relatively small, simply shaped pieces which limits the pressures which may be employed to avoid leakage through the numerous joints. Stresses induced must be limited essentially to compressive ones and suitable provision must be made to permit thermal expansion of the components in order to minimize these stresses. Where heat is to be added for boiling or removed for condensation, refractory material through which the heat flows must have a high thermal conductivity. Silicon carbide is highly conductive but is subject to attack by iron and is expensive compared to other commonly employed refractories. A hard, dense aluminous fireclay is well suited to contain these metals where high thermal conductivity is not required. Zinc will penetrate the pores of fireclay brick but only to a limited depth in an exterior wall and has little effect other than an increase in weight and conductivity. The design objectives of the Amax refiner were: a) to achieve a long life by heavy, rugged construction and by employing silicon carbide refractories only where required for heat conductivity and not in contact with liquids of high iron concentrations, b) to simplify operation, manning and control by essentially horizontal construction, and c) to obtain high capacity in a single unit. DESCRIPTION Flow of metal through the Amax refiner is shown diagrammatically in Fig. 1 and schematically in Fig. 2.

Jan 1, 1963

By C. W. Thompson

Paul D. Suloff (Goodyear Tire and Rubber Co., Inc., Akron, Ohio)—I would like first to comment on problems of the conveyor belt discussed in Mr. Thompson's excellent paper, since that is what we hope we know most about. Twists in relatively wide conveyor belt unavoidably produce a lateral maldistribution of tension, raising tension at belt edges and reducing it at the center. They also produce a lateral collapsing force on the belt at the center of the twist owing to the inherent tendency of all the longitudinal elements of the belt to try to pass through a point at the twist center. Calculation of the twist geometry by the methods shown in Mr. Thompson's paper keeps these extraordinary forces within limits which the belt designer can tolerate. No reduction in belt life due to twisting need be contemplated when this geometry is maintained. There is a minor exception that belts of extreme lateral flexibility will tend to curl laterally at the center of the twist. However, any ordinary fabric construction will perform satisfactorily in this respect. These twists are always made in regions of low tension in the conveyor so that even in the edges of the twist, belt tension does not exceed the average tension found in highly stressed regions of the conveyor. Offsetting these out-of-ordinary belt stresses is the advantage that Mr. Thompson has brought out of getting the return run up out of the dirt where it can be seen. This not only makes it easier to train, but also, in the event that it is not properly trained, frees it of the normal return run edge wear hazard. It is well known that return run edge wear is a prominent cause of belt mortality underground. An interesting aspect of this two-way conveyor is that the belt may be made what is known as a Mobius Strip. A Mobius Strip is obtained by splicing a belt after turning one end of it 180" about its longitudinal axis. In other words, one end is turned upside down before splicing. A belt spliced in this fashion turns itself upside down every time it comes around, but the twist which has been put in the splicing, of course, stays at one location on the conveyor, in this case one of the twist sections at the end. Turning the belt over every revolution might have advantages in some cases. Belts could be made with equal covers and the two sides worn equally and simultaneously. In this case there would be no problem of getting belts on upside down by mistake. However, the two-way conveyor does not have to be a Mobius Strip. It can be twisted in such fashion that the same side is up on both runs. It is simply a question of which way the final 90" twist is made before joining the ends. Another interesting aspect of the two-way conveyor is the problem of operating two-way conveyors in series. Here the sequencing of starting brings up some new problems. It will be recognized, although not always at first glance, that if the starting sequence is planned for one run of the conveyor the reverse will result on the other run. With the two runs carrying bulk material in both directions a reverse sequence on one run would be intolerable. In this situation the only solution appears to be a simultaneous starting of all conveyors in the series. However, with the coal in one direction and intermittent supplies in the other it would be entirely practical to sequence the conveyors for the coal run and accept a reverse sequence on the supply run. The two-way conveyor also lends itself to new driving possibilities. First, it is quite possible to drive at the head end of each run, which of course, means a drive at each end of the two-way conveyor. Driving in this way a given belt can be extended to substantially greater lengths than a conventional conveyor with drive at one end only. In addition to this, under certain conditions the conveyor can be extended to extreme length by driving at one end and at some intermediate point on the most heavily loaded run. As a particular case, a belt carrying coal downgrade and supplies back upgrade could be extended to extreme lengths by driving at the head of the coal run and at an intermediate point of the supply run. Mr. Thompson has been a pioneer in belt conveyor transportation underground and his accomplishment here with the first two-way conveyor of any consequence is another notable addition to the art.

Jan 1, 1954

By John V. Beall

Let's start at the beginning, Mr. Turner. Where and when were you born? In Lansing, Mich., on July 31, 1880. And what was your education? I went to the University of Michigan, where I got an A.B. degree in 1902; then to the Michigan College of Mines at Houghton, graduating from there in 1904 with a B.S. and an E.M. What was your first job in the mining industry, Mr. Turner, and what were you paid? In the summer of 1902, 1 was a field assistant for the U. S. Geological Survey in southwestern Idaho and southeastern Oregon, working under Israel C. Russell, professor of geology at the University of Michigan, who incidentally, had just returned from an expedition to study the effects of the eruption of Mt. Pelee, on Martinique. I was paid $60 a month and found; in addition to doing geologic and hydrographic work, I cooked many of the meals for a party of four, and wrangled four horses. I slept for four months under the stars, without a tent.

Jan 1, 1949

By J. C. Swartz

New measurements of graphite and cementite solubilities in ferrite have been obtained to resolve disagreements among previous data. To measure graphite solubility, specimens of iron were equilibrated with rarious carburizing atmospheres at temperatures be-twee~z 480" and 1435°C. The carbon activity of each atmosphere relative to graphite was taken from published data. After quenching from each equilibration treatment the concentration of carbon (typically 10 pprn by wt) was determined from the height of the Snoek peak in the internal friction spectrum. A few inclividual analyses by combustion support the reszdts of Smith on the solubility of graphite near the eutectoid temperature. Combination of these combustion and internal friction analyses determines In a, (at 0.01 wt pct) = 24,000/RT - 12.72, where a, is the activtity relative to graphite in ferrite at fixed cornposition. The same specimens were used to determine the solubility of self-stressed cementite precipitates. The results indicate solubilities of 6.3, 23.2, and 68 pprn carbon by wt at respectively 400°, 500". and 600°C. The once large discrepancy between the cementite solubility determined by internal frictiori and that calculated from flee energy data has been almost completely resolved. A recent study1 (hereafter designated I) on the solubility of cementite in a iron did not fully resolve the disagreement among previous data. The remaining disagreement is illustrated by the phase diagram in Fig. L The data of I are represented by the line labeled "stress-free" cementite. They determine a heat of solution of cementite of (14.8 * 0.2) kcal per mole which contrasts with the (20 * 1) kcal per mole derived from Smith's graphite solubility2 and the formation energy of cementite.3 Toward resolving this disagreement we have obtained new data on the solubilities of graphite and cementite. METHODS As in I the Snoek peak in the internal friction spectrum was used to measure the concentration of carbon dissolved in ferrite. This method relies on the relation C=kQ-1 [1] in which C is the concentration of dissolved carbon, Q-I is the Snoek peak height, and k is a proportionality factor which is a function of the peak temperature, polycrystalline texture, and other parameters characteristic of the ferrite lattice. To determine true heats of solution via Eq. [1], it is necessary to establish that k is independent of the temperature from which the specimen is quenched. Some evidence on this point is contained in a separate report.* Additional evidence is shown by the data points in Fig. 1. These represent the Snoek peaks of two specimens carburized to about 10 ppm C, copper plated, and quenched first from 888º then from 588°C. The copper plate was removed before each Snoek peak measurement. These data show that there is no detectable dependence of k on quenching temperature. The solubility of graphite was determined by repeatedly carburizing various specimens of iron wire with gas mixtures of known carbon activity relative to graphite. To cover a wide temperature range within the a phase boundaries it is desirable to use as low a carbon concentration as is compatible with reasonable precision. With the present internal friction apparatus the best compromise of these desires is a concentration near 10 wt ppm. The dashed line in Fig. 1 was used to approximate the activity of carbon* for

Jan 1, 1970

By R. E. Reed-Hill, W. H. Hartt

High-purity magnesium single crystals, oriented with basal plane parallel to stress axis, were deformed in tension at room temperature so as to form second-order (1011)- (1012) twins. Investigation by optical and replica electron microscopy revealed that these twins can support very large internal strains approaching 1000 pct and that fracture follows as a direct consequence of these strains. At low strains the basal plane within (10i1)- (10i2) twins rotates with increasing deformation until it becomes approximately parallel to the twin boundaries. This rotation is not consistent with deformation by basal slip in the double twin, but may be rationalized in terms of nonbasal slip. Observations also indicate that basal slip in the primary (1011) twin, prior to second-order twinning, may contribute to this rotation. Deformation markings that are difficult to interpret in terms of slip on previously observed systems were noted in the twin bands. A possible means of rationalizing these markings is to relate them to grain boundary shear. The fracture associated with these twins is thought to initiate by formation of voids at the twin boundaries. These grow into microcracks, with ultimate fracture occurring by tearing of the interconnecting regions. THE significance of deformation and fracture along second-order (10i1)-(10i2) twins in polycrystalline magnesium1-' and many magnesium alloys4 was discussed in a previous paper.5 It was noted that this double twinning mode effectively controls the room-temperature deformation and fracture of magnesium single crystals, oriented with the basal plane parallel to the tensile axis. Reed-Hill and Robertson6 observed that, although these crystals may fracture at a macroscopic strain of less than 1 pct, shear strains up to 1000 pct could occur within (1011)- (10i2) twin bands. They concluded that the fracture was "ductile" and resulted from concentrating the deformation into a very small volume. The present research was undertaken to obtain a better understanding of the deformation and fracture associated with these twins. EXPERIMENTAL TECHNIQUES The experimental procedure involved straining high-purity magnesium single crystals in tension at room temperature. Specimens were oriented so that the applied stress direction was within 2 deg of the [1010] crystal axis. The specimen cross section was rectangular with faces closely parallel to (1210) and (0002). The specimen preparation and testing procedures have been described in detail elsewhere.5 Second-order (1011)- (1072) twins nucleated during straining were studied using optical and replica electron microscopy techniques. For the latter, cellulose acetate-carbon double replicas were employed. Observations were made on the (1210) crystal surface, which is the plane of shear for twins of this type. In addition to the single crystals, some investigations were carried out on longitudinal polycrystalline specimens. These were obtained from high-purity magnesium plate with a texture in which the basal planes of the grains tended to be nearly parallel to the specimen axis. All. polycrystalline specimens were annealed prior to testing in order to produce a coarse grain structure, permitting Laue back-reflection X-ray photographs to be obtained from grains of interest. EXPERIMENTAL RESULTS a) Deformation Within ( 101 l )-( 1072) Twins. Fig. 1 is an optical micrograph demonstrating the extent of the plastic deformation that can occur in (1011)-(10i2) twin bands. Notice at the lower left corner of the photograph the large displacement of the dark band (a polishing step) where it crosses the twins. The shear strain in the twin band at this position was computed to be 700 pct, using the ratio of step width to twin thickness. It was previously shown5 that (1011)- (1012) twin bands are often composed of small, separate twins aligned close to the macroscopic habit plane. Fig. 2 is an electron micrograph from near the upper end of the twin band in Fig. 1. This illustrates that the twin habit often does not coincide exactly with the macroscopic band habit, but may be inclined at a slightly smaller angle to the matrix basal plane. Figs. 3 through 6 are electron micrographs of the twin band in Fig. 2 illustrating successive steps in the development of deformation inside the band. A twin band such

Jan 1, 1969

By T. E. Scott, D. H. Sherman, C. V. Owen

Mechanical properties and optical metallographic characteristics of vanadium containing 53 ppm (wt) H were investigated from, 77° to 298°K. A sharp ductile to brittle transition induced by the presence of hydrogen was observed between 238° and 223°K and a mar-tensitically formed hydride appeared at 200°K. While there appears to be a subtle relationship between the hydride formation and the ductile-brittle transition, it was concluded that the hydride per se was not sufficietzt to cause the severe em brittletnent. On the basis of the metallographic experiments and the results of a Petch analysis it is suggested that the mechanism of the embrittlement involves a reduction of the true surface energy, or cohesion, of the uanadiunz. THAT hydrogen can severely embrittle vanadium is well-documented;'-5 however, the mechanism by which hydrogen induces the observed embrittlement has not been established. It is not known whether the embrittlement of vanadium is caused by the formation of vanadium hydride or by a more subtle role of hydrogen. Until very recently6>' the solid solubility of hydrogen in vanadium at temperatures in the,range of the embrittlement was not known to a useful accuracy. In general, a severe embrittlement of the type produced by the presence of hydrogen in vanadium can be caused by changes in either of two properties of a metal. Either the yield stress of the metal is raised to the extent that plastic deformation at the tip of a crack cannot occur or the cohesive properties of the metal are reduced, at least locally, to the point at which cleavage is the natural response to applied loads. Of course a combination of these is possible. It is not known which of these properties of vanadium is influenced most by the addition of hydrogen. The purpose of this investigation was to determine the role of hydrogen in the embrittlement of vanadium. Accordingly, the solubility of hydrogen in vanadium in the temperature range where embrittlement is most severe was determined, the appearance of vanadium hydride was examined in relation to the initial ductile to brittle transition temperature, and an attempt was made to evaluate the effect of hydrogen on the yield stress and cohesive properties of vanadium. MATERIAL AND PROCEDURES Sample Preparation. The vanadium used in this investigation was prepared by the iodide refining process described by Carlson and owen.' Three ingots, 3 in. diam by 6 in. long, were cold-swaged to 0.096 in. diam, cut into short lengths, and then arc-melted under purified argon into one ingot. This process was repeated to insure a uniform composition in the final ingot. Chemical analysis of the metal is given in Table I. On the basis of work of Loria et a/ .' the following sequence of treatments was performed in order to obtain a uniform equiaxed ultimate grain structure: a) cold swage q-in.-diam ingot to 0.230 in. diam; b) anneal 6 hr at 1000°C in vacuum; c) by cold swaging reduce 81 pct in area in three steps: 0.230 to 0.187 in. diam, 0.187 to 0.150 in. diam, 0.150 to 0.096 in. diam. This sequence provided the desired grain structure after subsequent annealing treatments. Because annealing in evacuated quartz capsules resulted in contamination by oxygen, all anneals in this investigation were performed in a 304 stainless-steel chamber under a dynamic vacuum of 10- I Torr with the vanadium wrapped in tantalum foil. Prior to all annealing treatments the wires were electropolished for 1 min at 16°C in an 80 pct methanol-20 pct sul-furic acid electrolyte at 12.5 v. Following the final reduction the 0.096-in.-diam wire was cut into 2-in. lengths. A 1-in. gage section was reduced to 0.090 in. diam by electropolishing as above. Five different grain sizes were obtained by annealing the 2-in. samples at different temperatures for 8 hr to insure uniform grain size and residual hydrogen removal. The annealing temperatures used and the resulting grain sizes are given in Table 11. The final step in sample preparation was to charge half of the specimens of each grain size with hydro-

Jan 1, 1969

By R. E. Reed-Hill

The lattice reorientations in (1011) and (1013) twins of pure magnesium have been investigated using polarized light. Both forms (Ire subject 20 almost complete second-order twinning on the (1012) plane of the first-order twin. The observed deviations of composition planes by 6 and 3 deg, respectively, from coherent twinning planes are caused by second-order twinning. The experimental data indicate that (1011) and (1013) are reciprocal twins with a twinning shear 0.136. In a previous paper1 it was shown that the primary fracture mechanism in single crystals of magnesium strained in tension parallel to the basal plane is parting or fracture inside twins. In the temperature interval 25oto286oC, the observed twinning form, dominant in fracture, was a very small twin which formed in bands. Because these twins were very small and appeared in extensively deformed regions they did not lend themselves to X-ray analysis. Lacking Cou-ling and Pearsall&apos;s2 recently reported polarized light technique for magnesium, the twinning plane indices (K1) were originally determined by two surface composition plane measurements. This plane fell in a major zone (axis < 1120>) at an angle of 55.75 deg to the parent crystal basal plane. Because its habit deviated by 6 deg from{l0ll), 61.9 deg, the experimental data per se were not consistent with simple (1011) twinning and, for expedience, the twin was designated with habit plane indices (30341, 54.5 deg, representing the lowest set of indices corresponding to measured angles. Couling, Pashak, and sturkey3 have also found bands of reoriented lattice with a similar habit in polycrystalline specimens of magnesium and certain dilute magnesium alloys. These bands are significant because their formation in large numbers during plastic flow permit specific magnesium alloys to undergo very extensive deformation by cold rolling. Using polarized light2 they showed that the basal plane in the reoriented material was nearly parallel to the band habit and, as a result, proposed a mechanism for band formation involving a double twinning process in which (1011) twins form and then retwin according to (1012). Because of the obvious similarity between the twin bands observed in single crystals and in polycrystalline specimens, it was decided to investigate if the two phenomena were not identical and, if possible, to verify the retwinning hypothesis. The present paper is a report of this investigation and shows that the twins with the (3034) habit correspond closely to retwinned (l011) twins. Evi- dence is also given for an analogous retwinning process in (1013) twins. EXPERIMENTAL PROCEDURE Rectangular single-crystal specimen deformed in tension (stress axis in the basal plane parallel to [1010], as previously described,4 were polished electrolytically by Jacquet&apos;s method5 and etched in acetic picral.2 All twins studied in this investigation belong to a major zone, whose axis, [1210], was perpendicular to both stress axis and two sides of each specimen [(1210) surfaces]. These latter contain the plane of shear of the twins and measurements of the basal plane trace of twins on these surfaces determines the lattice orientations of the twins. The position of the basal plane in each case was determined with a polarizing microscope using a modification of the technique of Couling and Pearsall.2 For each twin the four positions of maximum extinction were measured, the data averaged, and then the pair of extinctions corresponding to the basal plane was determined by rotating the microscope stage and observing the position at which the color changed from orange to blue with a gypsum red, first-order plate inserted between the crossed Nicol prisms. EXPERIMENTAL RESULTS (1013) Twins—In general, especially in speci-mens deformed at 150°C, (1073) twins appear in greater numbers and larger sizes than (1011) twins, and the lattice reorientation inside these twins is therefore easier to study and measure. Fig. 1 shows an area of an unetched and unpolished (1210) crystal surface containing a number of (1073) twins sloping upward to the right at about 29 deg with the basal plane of the original crystal (horizontal direction). Slip line segments may be seen in both the parent crystal (running horizontally) and in the twins (sloping downward to the right at approximately -22 deg with the horizontal). As may be seen in Fig. 2, the slip lines in the twin are not at the proper angle (64 deg from basal plane of the

Jan 1, 1961

By J. Campbell

Various models are discussed for the evaluation of the negative pressures which may occur in solidifying materials which exhibit various deformation modes: elastic-plastic, Bingham, viscous, or creep flow. The inadequacy of the previously proposed elastic-plastic solution for solidifying metals is revealed by comparison with the more reliable creep results which are given graphically for aluminum, copper, nickel, and iron. The maximum tensions experienced in the liquid phase of solidifying spheres ranging in size from large castings to submicron powders are in the range from —10&apos; to —105 atm for these metals. THERE has been much recent interest in the negative pressures associated with the volume change on solidification and in the possibility of the occurrence of cavitation. Considering the freezing of a highly supercooled liquid, an attempt to evaluate the stresses in the liquid ahead of the rapidly moving solidification front has been made by Horvay1 on a microscale and by Glicksman2 on a macroscale. In a casting of a wide freezing range alloy, the pressure differential due to viscous flow of residual liquid through the pasty zone has been discussed by Piwonka and Flemings,3 In a previous publication4 the author has attempted to estimate the negative pressure occurring in the residual liquid of a spherical casting, employing an elastic-plastic model to describe the collapse of the solidified shell under the internal tension. An earlier model assuming a rigid shell was shown to be inaccurate by many orders of magnitude. The elastic-plastic model is critically reviewed here, and other models are developed which are thought to be more closely related to metals and other materials near their melting points. The spherical geometry (Fig. 1) is chosen because the highest shrinkage pressures would be developed, although the analyses are readily adaptable to cylindrical geometry. A parallel sided casting experiences little internal tension because of the relatively easy dishing inward of the sides. (This commonly observed phenomenon has previously been attributed solely to atmospheric pressure.) Furthermore, small regions of confined liquid in a large solidified volume of a casting approximate reasonably well to spherical geometry. ELASTIC-PLASTIC MODEL The author has shown4 that as solidification proceeds the internal hydrostatic tension builds up until the elastic limit of the shell is exceeded. At this point the internal pressure is closely -2Y/3. Subsequently a plastic zone spreads from the inner surface toward the outer surface of the shell. When the whole casting is deforming plastically a rather more generalized analysis taking account of the externally applied pressure PA + 2y/b gives the internal pressure as: P = Pa + 2y/a + 2ys/b - 2 Y In(b/a) [1] The 2y/a and 2ys/b terms result from the tendency of the liquid-solid and solid-vapor interfaces to shrink, reducing their energy, and thereby helping to collapse the solid phase and compress the liquid phase. The 2y/b term would be important only for powders. The last term arises because of the plastic restraint of the solid, resisting collapse and so effectively expanding the residual liquid. From Eq. [I] it is easily shown that there is a minimum in the pressure at the radius amia= y/Y [2] which is of the order of 103K for the metals aluminum, copper, and iron, and corresponds to the minimum pressure Pmin = 2 Y[l-ln(bY/y [3] The results of a fully worked out elastic-plastic solution are given in a previous reporL4 The main criticism which may be leveled at this analysis when applied to metals at their melting points is the strong dependence of the yield stress on the strain rate. The strain rate varies with both solidification conditions (e.g., whether chill-cast or slowly cooled) and during solidification, as is indicated in the following section. Thus an appropriate choice of Y is very arbitrary. Before proceeding to a discussion of models which are strain-rate-dependent, it is necessary to evaluate the strain rate as a function of the rate of solidification. SOLIDIFICATION RATE Various empirical relations have been deduced5 for the rate of thickening of the solid shell by pour- out tests on partially solidified spheres. These, however, are unsatisfactory for our purposes since they become very inaccurate when the liquid core is very small. A theoretical approach is therefore necessary, and some solutions are set out below. Making the assumptions of constant surface temperature of the casting during freezing, no superheat, and a material freezing at a single temperature, Adams8 deduces the approximate solution: which becomes when b » a: Employing a semiempirical approach vallet6 finds:

Jan 1, 1969

By J. S. Rinehart, W. C. Maurer

The mechanics of impact crater formation in rock, particularly sandstone, has been sutdied, the velocity range being approximately that normally associated with oilwell gun perforators. The bullets were small steel spheres having diameters of 3/16, 9/32 and 7/16 in; impact velocities ranged from 300 to 7,000 ft/sec. The craters have two distinct parts — a cylindrical hole (or burrow) with a diameter the same as that of the impacting sphere, and a wide-angle cup comprising most of the volume of the crater. The burrow is fornred as material in front of the projectile is crushed and pushed aside, forming a cylindrical hole surrounded by a high-density zone. The clip forms as fractures are initiated in front of the projectile and propagate along logarithmic spirals, approximaling maximum shear trajectories, to the free surface of the rock. A most significant observation (made for the first time) was that, below the base of the cup in one type of sandstone, there are a group of similar fractures, not extending to the surface, which are spaced uniformly a few millimeters apart. Each fracture follows roughly the contour of the base of the cup and appears to require a certain threshold impulse to initiate it. These fractures comprise a relatively high fraction of the total, newly exposed surface area. The volume of the material removed by crushing varies as the first power of the impact velocity and the volume removed by fracturing, as the second power of the impact velocity. Penetration varies linearly with the impact velocity and is inversely proportional to the specific acoustic resistance of the target material, the proportionality constant being dependent upon the shape of the projectile. INTRODUCTION Yield of oil from a producing well is frequently enhanced by firing bullets and shaped charges through the well casing into the oil-bearing rock, forming craters and fractures from which oil can flow more readily. The purpose of this investigation has been to develop a better understanding of the mechanics of impact crater formation in rock, particularly sandstone, the velocity range being approximately that normally associated with oilwell gun perforators. FORCES OPERATIVE DURING IMPACT When a projectile moving at considerable velocity strikes a- massive target such as oil-bearing sandstone, intense and complex transient stress situations develop within both the projectile and the rock or sandstone against which it is striking. Usually the struck rock fails, the missile or projectile penetrating into the rock to some depth where it comes to rest or is forcibly ejected from its burrow by expansion of a plug of target material compressed in front of it. When the impact velocity is very high, the projectile itself may fail, breaking apart or becoming distorted; this situation is not considered here, the discussion being limited to nondeforming projectiles. Many experimental studies&apos;.&apos; have been carried out to determine the nature of the mechanics of crater formation and the salient features of the forces coming into play, some of the earliest studies being the French Army experiments performed at Metz between 1835 and 1845.&apos; The stratagem in most instances has been to make a post-mortem examination of the crater, measuring volume and depth of penetration and deducing force relationships from these observations rather than performing the more difficult (usually almost impossible) feat of measuring stresses during penetration. In many materials, the force acting during penetration of the projectile is found to be the sum of two components—(1) a constant force, independent of the velocity, representing some inherent strength of the target material; and (2) a component, proportional to the square of the velocity, representing inertial forces. For such materials, the average force per unit area acting on the projectile at any instant while it is in motion and being decelerated may be written F/A = a + bv2 . . . (1) where v is the velocity of the projectile at that instant, A is the cross-sectional area of the penetrating projectile taken normal to its trajectory, and a and b are constants which are dependent upon the target material and the shape of the projectile. It follows that the total penetration s is given by .........(2) where v, is the velocity of the projectile when it just strikes the target. Values of a and b for spherical projectiles impacting in a loose sand-gravel mixture and compacted earth were obtained in the Metz experiments. For sand-gravel, a and b are 620 psi and 0.0115 (psi) (ft/sec)&apos;, respectively; and for compacted earthworks, a and b are 432 psi and 0.0008 (psi) (ft/sec)&apos;. Figs 1 and 2

By Hans Lundberg

AFTER the war years, great activity has been shown in geophysical exploration for ore. The appreciation by mining and government geologists of geophysical techniques and results is largely responsible for the popularity of these methods. In Canada, plans are being made to study, by geophysical methods, the structure of a portion of the Pre-Cambrian shield in the hope that such studies will not only give a clearer picture of the geology, but will be of direct aid to the mining companies when selecting areas for detailed prospecting. Although geophysics thus has become more popular, some engineers are still reluctant to employ the methods to their full extent. Allotted funds are often rather low, so that the least expensive

Jan 1, 1947

By George Otis Smith

WHAT American can enter this Western empire without his imagination being stirred by the stories of its past-yes, and even more by visions of its future! Whether we travel by rail or by auto, our pathway has its traditions of exploration and its records of early settlement. The mountain passes and the routes used by today&apos;s commerce are lasting memorials of the pathfinders whose unconquerable purpose was to search out a way-Lewis and Clark, Bonneville, Pike, Long, FrBmont, their names still live in your rivers and lakes and peaks. It was the glory of these adventurers that they spied out the land so that others might follow that long pro- cession of pioneers who occupied the new country- fur traders, gold miners, cattlemen, and homesteaders. Each of these pioneers attacked the problem of transportation in his own way, and each sought by his own method to win the resources whose richness had been reported by the pen of the Government explorer or by the swift rumor of the unnamed and unfamed adventurer. Trails were worn deep by these home or fortune seekers as they passed through the mountain barriers in successive waves-and these were waves that did not ebb, for this was an army of occupation. Along these main pathways of the West, as one historian has pointed out, the figure of chief interest was the sad- faced woman sitting on the front of the covered wagon.

Jan 1, 1926

By Augustus Locke

I WISH to urge on this Committee the task of evaluating mining geology. -My motive is as follows: It, is a. duty of the Institute from time to5 time, to establish the social perspective of the professions grouped within it, not simply to promote their use, but to determine with engineering definiteness their true powers and functions. This becomes the delegated duty, therefore, of this committee in respect to geology. Be- cause my own subject is mining, I shall confine my remarks to mining geology. But I believe that the evaluation should be made to include all of that geology termed economic. The perspective for such evaluation is shown in the nature of the public support given to geology. The federal Geological Survey, the leading organization on the economic side, is actuated by the zeal of a small group of devoted professionals who consider its work valuable, &apos; It proceeds, however, on a starvation ration of appropriations, and that implies that neither the corporations, the voters nor the vote getters consider it important. Altogether, in&apos; the public consciousness, geology is negligible. That, at the same moment, .the profession should act with full confidence, while everyone else should act&apos; with opposite motives, is not fitting. For us, as a committee, to remain incurious regarding such a disparity of beliefs is quite impossible.

Jan 1, 1931

By J. J. Jakosky

The strategic importance of the metallic minerals in our industrial economy, and the declining rates of discovery have focused attention on means of exploration for new mineral deposits. A consideration of exploration techniques leads to an attempt to evaluate the effectiveness of geo- physics as a tool for mining exploration. One means of evaluation will be to compare the use of geophysics in petroleum exploration with its possible uses in mining.

Jan 1, 1949

By R. G. Bowman

LEAD is a dull and sober metal, and in times of economic stress it en- joys, or at least occupies, a position which partakes somewhat of its physical lethargy and stability. The past ten years have witnessed in this country a decline in the production of lead, and a shrinkage of ore reserves, following closely the general economic situation. Metal prices have not justified exploration and development of new sources. Momentary rallies in industry have brought no response from lead.

Jan 1, 1940

By I. B. Cadoff, J. C. Bilello, R. Rosenberg

Diffiusion of magnesium into the surface of LiF crystals to controlled depths and subsequent heat treatments provided a wide range of surface zone harahesses and structure, The bend strength of the LiF crystals was increased by as much as an or-dev of magnitude. Ductility was achieved when dislocation generation occurred in the diffusion zone or when dislocations penetrated to the surface from the intevior. A critical surface hardness of 130 to 140 kg per sq mm was found helozu which generation could take place in the diffusion zone and ahoue which the zone was impenetrable, This hardness was obtainable by several methods, among them the aging of quenched MgF2 -LiF solutions to produce MgF, precipitation. Maximum hardness was ohtained in quenched specimens with no visihle evidence of MgF,. Diffusion-zone formation followed a parabolic rate law and an activation energy of 20.9 kcal per mole was obtained for the process. RECENTLY, the properties of ionic crystals as related to surface condition have been receiving much attention, specifically the transitions between ductile and brittle behavior. Originally Joffe 1 showed that NaCl crystals could be made ductile by immersion in water and related this to the elimination of surface microcracks. Aerts and DeKeyser 2 and Gorur 3&apos; have subsequently shown that ionic crystals are inherently ductile and are embrittled through contact with air. Machlin and Murray4 hypothesized that a layer of NaCIO3 produced by contact of ozone with NaCl induced embrittlement by acting as a barrier to outward dislocation flow. westwood,&apos; Rosenberg and Cadoff,9 and Bilello and cadoff&apos; have reported surface strengthening of LiF crystals by coating with a magnesium compound and then heat treating for adherence. The major effect of the coat was to inhibit dislocation-slip lines from reaching the specimen surface. westwoods showed microcrack formation and fracture to be caused by slip-band interactions at the surface. The material presented in this paper is an extension of the work reported earlier by Bilello, Rosenberg, and cadoff&apos;6,7 and illustrates the wide range of surface properties and bulk behavior obtainable by use of heat-treated magnesium-diffused surface regions in LiF crystals. EXPERIMENTAL PROCEDURE The LiF single crystals were obtained from the Harshaw Chemical Co. Some batch to batch variation was observed; therefore all specimens for a given test series were cleaved from the same crystal. The typical dimension used was 1 by 0.1 by 0.40 in. Surface damage resulting from cleavage was removed by chemically polishing in a 2 pct NH4OH solution. Each group of specimens was given a vacuum anneal at 700°C for 4 hr to provide a base standard for measuring comparative effects of various surface treatments. To produce the reacted surface zone, the annealed specimens were immersed in a boiling suspension of MgF, in doubly distilled HzO, agitated slbwly for 30 sec, removed, and dried at room temperature. Uniform coatings of MgF, were deposited with a thickness of approximately 5 mils. It should be noted that this technique can be modified for use with crystals which are soluble in water by using boiling absolute alcohol as the dissolving medium. This was found effective for the coating of NaCl with MgF2. The diffused surface zone was obtained by annealing the coated samples at elevated temperatures in a vacuum of lob4 mm Hg. Penetration depth was controlled by varying the annealing time from 1/2 to 28 hr. After heat treatment, the samples were tested for bend strength and hardness. Load was applied by four-point bending in a hard-beam testing jig. Four-point rather than three-point bending was used to provide a wide area of constant stress and to minimize the effect of localized inhomogenities in the specimen. The deflection rate was 8 x min-&apos; and the distance between knife edges was 1/4 in. Load-time curves were obtained from a chart recorder coupled to the machine and converted to resolved shear stress on the shear plane vs deflection, as plotted in the figures. The unstressed portions of the sample outside of the two outer knife edges were used for the microhardness studies. Microhardness measurements were made with a Bergsman tester attached to a Reichert metallograph. All hardness impres-

Jan 1, 1964

By A. Yoshikawa, T. Sugeno

To detenmine the origin of the sharp fatigue limit in many ferrous metals, S-N curvces were determined in push-pull fatigue at 18.6 kc per sec at room temperature and - 67°C for various kinds of iron. These differed in carbon and nitrogen content and grain size (25 and 2500 grains per sq mm). Sharp knees appeared in the S-N curves of fine-grained specimens with high interstitial content at both temperatures. In coarse-gruined specimens, knees were not sharp .regardless of the interstitial content and the testing temperature. It is concluded from these results that the controlling factov for the appeavance of the sharp fatigue limit is not strain aging but grain size. It has already been mentioned that iron and some iron alloys show fatigue curves characterized by a more or less sharp knee and a fatigue limit.&apos; Explanations of the origin of this characteristic behavior of fatigue in iron appear to fall into two broad classes. One is concerned with interstitial impurities through the aging or locking effect and the other is based on intrinsic plastic properties of iron. Levy and sinclair2 and Lipsit and Horne3 have suggested that the fatigue limit observed generally in iron and steel is due to interstitial solute aging during the course of fatigue. On the other hand, Thompson and wadsworthl have suggested that the fatigue limit may be analogous to the sharp yield point, in the sense that alternating stresses above the fatigue limit are necessary to break down the dislocation locking. However, there are some experimental results which cast doubt on the general validity of these suppositions.4-8 Ferro and Montalenti&apos; have recently suggested that the typical form of the fatigue curve of iron seems to come from an intrinsic plastic characteristic of the iron structure. This suggestion is based on the experimental result that pure iron containing an insufficient amount of interstitial solute to alldw appreciable strain-aging effects still showed a fatigue limit at room temperature. Although systematic investigations on the effect of grain size are scarce, it is thought that this effect is important in relation to the sharp fatigue limit. In fact, sensible differences were observed by Oates and Wilson4 in fatigue behaviors of a low-carbon steel depending on grain size. Ideally, the fatigue limit will be defined as the upper limit of stress amplitudes at which the life is infinite. This ideal fatigue limit can never be determined because of the limited time available for experiments, and moreover scattered data points often leave the construction of S-N curves somewhat arbitrary. And, it is usually stated that a material has a fatigue limit when its statistical S-N curve becomes apparently parallel to the N axis with a more or less sharp knee. The meaning of the fatigue limit and the sharp knee will be discussed in detail in relation to the construction of S-N curves. The origin of this sharp knee was investigated by determining S-N curves of iron for various interstitial contents, testing temperatures, and grain sizes. EXPERTMENTAL PROCEDURE The material as received was a commercial-grade low-carbon steel in the form of 4-mm-diameter rods with the following composition in weight percent: C, 0.09; Si, 0.017; Mn, 0.43; P, 0.006; S, 0.012; Cu, 0.11. Five different groups of specimens were tested. The grain sizes and interstitial contents in each group are given in Table I. The specimens of group C-02 were decarburized by annealing for 100 hr at 700°C, followed by an additional 100 hr at 850°C in wet hydrogen. Specimens of groups N-2 and N-9 nitrided in a mixed atmosphere of ammonia and hydrogen after an initial treatment identical to group C-02. The specimens of group C-9 were prepared in a manner similar to those of group C-02, except heat treatment was carried out in a vacuum. After polishing their surface with emery paper, the specimens of all groups were given a final anneal for 30 min at 930°C and air-cooled in an evacuated quartz tube. For group C*-9 this

Jan 1, 1965