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By J. H. Bechtold

Tensile properties of annealed molybdenum were investigated from 1000" to—200°C. In the vicinity of room temperature a well-defined transition in tensile properties occurs. Reduction in area decreases from over 80 pct to 0 pct and, simultaneously, the yield strength increases about 60,000 psi. MOLYBDENUM, tungsten, iron, and many other metals and alloys with the body-centered cubic and hexagonal close-packed crystal structures change from ductile to brittle behavior over some relatively narrow temperature range. The temperature at which this transition in ductility occurs depends on chemical composition, microstructure, and method of testing. When fine grained, completely recrystallized molybdenum of commercial purity is tested in uniaxial tension at a slow strain rate, the transition occurs at slightly below room temperature.' Iron2 and tungsten3 tested under similar conditions become brittle at temperatures close to the boiling point of nitrogen, — 195°C and at about 200°C, respectively. A sharp change in ductility is uncommon in metals with the face-centered cubic crystal structure, however, a relatively sharp transi- tion in notched bar impact properties has been observed below room temperature in a face-centered cubic copper-antimony alloy.4 A transition from ductile to brittle behavior is usually explained by the concepts originally developed by Ludwik,5 and discussed and elaborated on by others including Orowan,6 Hollomon and Jaffe,7 and Gensamer8 In brief, there are two aspects to strength, first the resistance to flow and second the resistance to fracture. Both are functions of temperature, strain, strain rate, stress system, composition, and microstructure. With decreased test temperature the flow stress required to initiate plastic deformation increases more rapidly than the fracture strength and brittle fracture occurs at that temperature where the yield strength equals or exceeds a brittle fracture strength. This investigation was conducted to determine the effects of temperature, strain, and rate of straining on the flow and fracture strengths of molybdenum above, through, and below the ductile-to-brittle transition. Other variables were held constant during testing. This information is believed to show

Jan 1, 1954

By J. A. Snyder, W. L. Finlay

SEVERAL notable studies of the effects of sub- stitutional solutes have been reported. Outstanding among these for its elucidation of general principles, is the summarized work of Hume-Rothery and coworkers.1 For a systematic study of the effects of substitutional solutes on the mechanical properties of the solvent lattice, the work of Brick, Martin, and Angier² on copper; of Frye and coworkers³,4,5 on copper and silver; and of Lacy and Gensamer6 and of Austin7 on iron may be mentioned. Titanium is one of the few elements offering the possibility of a similar study of interstitial solutes since, unlike most metallic elements, it exhibits extended solid solubility for most of the interstitial solutes—hydrogen, boron, nitrogen, oxygen, and carbon. This paper reports a survey of the effects of dilute solutions of nitrogen, oxygen, and carbon on the mechanical properties of high-purity alpha titanium. Previous Work: Van Arkel and De Boer showed" that high purity titanium was quite ductile and De Boer, Burgers and Fast established".' that not only were oxygen and nitrogen very soluble in solid titanium but that, long before the limit of solid solubility of either had been reached, these interstitial solutes rendered the metal brittle. Their electrical resistance data as a function of

Jan 1, 1951

By Dell P. Williams, Howard G. Nelson

The tensile behavior of magnesium single crystals at a temperature of 26º ± 2ºC was investigated at varying pressure levels from 760 to 8 X 10-8 tow. For crystals deformed at a constant linear strain rate of 0.061 min-1. there was no apparent effect of the vacuum environment down to 10-8 torr. However, from a vacuum of 10-5 torr down to the lowest pressure attained (8 x 10-8 tor,), the effect of the vacuum environment was to increase the strain at the end of the first linear (or easy-glide) stage and to decrease the slope of this stage; the magnitude of this effect was found to be dependent on pressure. The slope of the second linear stage was also decreased at the lower pressure levels; however, this effect appears to be caused primarity by the vacuum effect on the first linear stage. The magnitude of the vacuum-environmrnt effect was also found to be dependent upon strain rate. The reduction in flow stress at 5 x 10-8 torr, as compared to the stress in air, was obserced to be about 10 pct for a strain rate of 0.0250 min-' and 70 pct for a strain rate of 0.67 min-1. No systematic effect of pressure on the critical resolvled shear stress was observed. The data interpreted in terms of a modified form of Mott's work-hardening mechanism for hexagonal metals. It has long been known that surface conditions influence the mechanical behavior of materials. Roscoe 1 demonstrated the importance of surface conditions on mechanical behavior by testing cadmium crystals with oxide layers. He found that the "apparent" critical resolved shear stress for cadmium was increased approximately 2.5 times by the presence of an oxide layer 1000 atoms thick. The magnitude of this effect was dependent on film thickness with crystals having films less than 20 atoms thick still showing a definite increase. This observed strengthening could not be attributed to any strength of the oxide film because of 1) the thinness of the films and 2) an observed increase in the flow stress as deformation proceeded. Harper and cottrel12 obtained similar results on zinc crystals which had been etched and steamed to give artificially thick oxide films. However, in the course of this study, they also carefully examined the very early stages of plastic flow for zinc crystals which in one case had been electrolytically polished and in the other had been lightly etched to produce a thin film. They then found that the critical resolved shear stress for both surface treatments of the crystals was actually identical but that the etched crystal exhibited a much more rapid initial hardening rate. This had not been noted previously. This result indicated that, in fact, the film caused an increase in the stress for continuation of slip, rather than for the initiation of slip. Lipsett and King3 investigated the effect of evaporated gold films on the plastic behavior of cadmium single crystals. For large angles (=75 deg) between the specimen axis and the normal to the slip plane, they found that the stress-strain curve of the coated specimen always lies above that for a clean specimen. The shear strain at the end of the first linear stage was greater in clean specimens than in coated specimens and the slope of the second linear stage was

Jan 1, 1965

By Leonard R. Weisberg

HE procedure of directional freezing by the Bridgman technique1 is frequently used in crystal preparation. On those occasions where the crystal is regrown, it can be advantageous to Crop part of the ingot before regrowth to achieve purification. This note describes some simple, yet interesting, observations concerning the effects of cropping an ingot between successive freezes. First, it is not effective to Crop the front end of an ingot, whether or not the distribution coefficient, k , of the impurities is greater or less than unity. This can be shown as follows. The concentration C, of an impurity in a directionally frozen ingot of length d with a uniform cross section is given by where x is the distance along the ingot, and C0 is the initial concentration per unit length. It is assumed in deriving this equation that diffusion is negligible in the solid, is complete in the melt, and that k is constant. Consider that a section of the ingot Ad is removed from the front end. Then the average impurity concentration in the remaining (d- Ad) section of the ingot can be shown by integration to be given by When the ingot is remelted and refrozen, the new impurity distribution is given by where xl is the distance along the refrozen ingot. Substituting the value of c into Eq. [3], one finds This is identical to Eq. [11 when x = (x, Ad), which represents the original impurity distribution in the (d - Ad) section of the ingot after the first directional freeze. Therefore, the cropping and re-freezing accomplished nothing. This can also be seen physically by realizing that the average impurity content in the melt after the first Ad section is frozen, is identical to that after the resulting ingot has been cropped and remelted. Hence, it is assumed in the following sections that only the tail end of the ingot is cropped. However, it must be realized that for impurities with k greater than one, the pieces cropped from the tail end are purer, and are therefore saved. After cropping an amount Ad from the tail end of an ingot, the average impurity concentration of impurities in the remaining portion of the ingot is given by wheref= d/d. For the case of k < 1, the purification is defined as c,/c. It is obvious from Eq. [5] that the purification increases asf increases, but the increase in purification is not large. For example, iff is doubled from 0.1 to 0.2, the purification increases by only 10 pct for k = 0.5, by only 24 pct for k = 0.1, and by 32 pct for k =0.01. Much smaller values off than 0.1 are not desirable since insufficient purification is achieved. These figures indicate that much larger values off do not lead to substantial improvement, especially considering that much of the original ingot is lost. Therefore, it is generally useful to Crop about 1/10 of the ingot to achieve purification. As compared to doubling the value off, it is much more helpful to remelt, refreeze, and crop the ingot twice. Using Eq. [51 reiteratively, and assuming f= 0.1, it is found that the purification after this process is found to be increased over the purification after a single freeze and cropping by 33 pct for k= 0.5, 450 pct for k = 0.1, and 4100 pct for k = 0.01. This indicates that considerable purification can be obtained by this procedure. However, if purification is the main goal, one should resort to zone refining,3

Jan 1, 1962

By J. F. Smith

ELASTIC constants of yttrium and eleven of the rare earth elements have been measured. This has been accomplished by measuring the propagation velocities of ultrasonic pulses. The velocity measurements were made by noting the elapsed time between an applied pulse and the appearance of successive echoes. The pulse-echo technique has become relatively common, and the basis of the measurement is outlined in several standard tests, e.g., Heuter and Bolt.1 Two important assumptions are made when using this method for the determination of elastic moduli. First, the pulse is assumed to propagate as a plane wave. To justify this assumption, the cross section of the test specimen should be large compared to the wave length of the transmitted pulse. For these measurements on the rare earth elements the minimum cross section wave length ratio was about 20. A second assumption which must be made when using polycrystalline specimens is that the grains are sufficiently numerous and randomly oriented so that the specimen may be approximated as elasti-cally isotopic. The grain size in the as-cast specimens which were used is believed to be sufficiently fine for this approximation, and the degree of preferred orientation is believed to be small. Supporting this viewpoint is the fact that elasticity measurements by the pulse-echo technique on as-cast specimens of other metals give results which are in satisfactory agreement with values obtained by other techniques. Hence, it is believed that the measured elasticity values for the rare earth elements are not apt to be seriously in error due to the assumption of elastic isotopy. The apparatus used in this determination was a modified Sperry Reflectoscope with a precisely calibrated time base circuit. Piezoelectric transducers were used both to generate the pulse and to detect the echoes. X-cut quartz crystals were used for the longitudinal modes, and Y-cut quartz crystals were used for the shear modes. Frequencies of 2.25 and 5 megacycles per sec were employed. Most of the specimens were 2.5 to 5.0 cm in length. Results Table I is a tabulation of the measured propagation velocities of longitudinal waves and of shear waves, together with densities obtained by Sped- ding, Daane, and Herrmann2 from lattice parameter measurements. The shear modulus, Poisson&apos;s ratio, and Young&apos;s modulus were calculated from the data in Table I through use of the following relationships: µ = pv32 [1] s=2-(vb/vs)2/2[1-(vb/vs)2] [2] and Y= 2µ(l+s) [3] where s is the shear modulus, Y is the Young&apos;s modulus, s is Poisson&apos;s ratio, p is the density, V0 is the longitudinal wave velocity, and v, is the shear wave velocity. The calculated quantities are shown in Table 11. Compressibility values can be computed from the relationship and the computed values may be compared with measured values obtained by Bridgman.8 Table III lists the data for such a comparison. The agreement is satisfactory, being particularly good for neodymi- um and gadolinium, poor for samarium, and intermediate for the other elements. Some variation is expected since Bridgman&apos;s measurements are isothermal, whereas the present measurements are adiabatic. In addition there may be some effect due to small variations in purity. Though most of the rare earth metals measured by Bridgman and all of the metals used in the present measurements were prepared at this laboratory, minor purity variations might exist since the samples came from different lots. Also, any imperfections such as dislocations, vacancies, voids, grain boundaries, microcracks, etc., should contribute to the volume change of a com-

Jan 1, 1958

By D. Argyriades, G. Derge, G. M. Pound

The electrical conductance of molten FeS was studied as a function of temperature and composition. It was found that stoi-chiometric FeS (36.5 pct S) shows a minimum specific conductance of 400 ohm-1 cm-1. For the Fe-rich melts the conductivity increases with increasing amount of iron to about 4850 ohm-1&apos; cm-I at 26 pct sulfur. For the sulfur-rich melts the conductivity increases with increasing amount of sulfur until it reaches a maximum of 1300 ohm1 cm-1 for 37.4 pct S and decreases to 800 ohm-1 cm-1 for 38.3 pct sulfur. It is concluded that molten stoichiometric FeS is an intrinsic semiconductor in which the forbidden gap in the electron energy level diagram is quite narrow. Excess sulfur acts as an acceptor impurity in providing a component of positive hole conduction, while excess iron acts as a donor impurity to increase the n-type conduction. K. Bornemann and G. von Ftauschenplatl used a four-terminal, direct-current cell to study the electrical conductance of molten copper sulfide. Their measurements showed a high conductivity, of the order of 100 ohm-1 cm-" and a positive temperature coefficient. From electrolysis studies, Savelsberg2 concluded that molten Cu2S and molten FeS are electronic conductors. Further. he found that molten FeS has a high specific conductance, of the order of 1500 ohm-&apos; cm-&apos;, and a small negative temperature coefficient. Pound, Derge, and Osuch,3 using an ac Kelvin circuit and a four-terminal cell, measured the specific conductance of molten Cu,S, molten FeS, and various mixtures. They found that the specific conductance of the solutions followed a roughly additive rule of mixtures. Yang, Pound, and Derge4 showed from electrolysis measurements that the mechanism of electrical conduction in molten Cu2S, FeS, and mattes is primarily electronic. Further, they observed that small additions of CuCl suppress the electronic conduction whereupon the system behaves ionically.

Jan 1, 1960

By J. L. Wyatt

The electrical resistance of titanium as a function of purity and temperature was measured from —325" to 2800°F. Two points of inflection in the data plots were found, and an increase in resistance with increase in temperature above the transformation was observed. IN 1948 Greiner and Ellis&apos; reported on the thermal and electrical properties of titanium, including a study of the effect of temperature on the electrical resistance. The materials used in the research were obtained from the United States Bureau of Mines, where powder metallurgy techniques were being employed for consolidation purposes. These authors found a decreasing resistance with increasing temperature above 1625°F, and concluded that the latter temperature represented the point of transformation from hexagonal close-packed to body-centered cubic structure on heating. They reported a specific resistivity value of 55x10-6 ohm-cm, and showed that dp/dt was positive, decreasing with increasing temperature above ambient. In one of his early texts Hume-Rothery2 mentioned that certain metals, including titanium, exhibited minima in plots of resistance vs temperature. He indicated that this occurred at 212°F. Reporting many years ago on the effect of temperature on the electrical resistance of zirconium, de Boer and Fast³ showed that phase transformation in this sister metal occurred over a range of temperatures, the extent varying with its purity. Graphs of resistance vs temperature, showing transformation effects over a range of several hundred degrees, were presented for relatively impure metal. In the latter case the resistance decreased with increasing temperature throughout the region of partial phase transformation, thereafter increasing as predicted for normal metals by modern theories of the solid state. Since the quality of metal used in previous studies was substandard compared with iodide-type titanium, and in view of the fact that there seemed to be some rather unusual observations in the reported literature, it was decided to undertake a more extensive study of the thermal effects using high purity titanium, and to include temperature ranges beyond those which had been studied in order to ascertain whether or not the characteristics of zirconium would be duplicated. Experimental Materials Materials used in this investigation included commercial quality titanium (Ti-75A produced by the Titanium Metals Corp. of America) in the form of ¼ in. diam rods, and high purity iodide titanium (produced by the Foote Mineral Co.) in the form of crystal bar of irregular size. The latter was cold-swaged into rod form and machined to the desired size. Lead wires to all specimens consisted of 18 gage Ti-75A. Chemical analyses of the stock materials are given in Table I. Experimental Procedures Specimens 2½ to 3 in. in length were prepared with diameters nominally 0.15 to 0.20 in. Two holes,

Jan 1, 1954

By F. W. Glaser

Iron-silicon alloys have had a great influence, in many ways, in modern industry. Silicon steels have been used almost exclusively for the construction of electrical machinery, but have also become an important material for structures which benefit by the high tensile and yield strengths of this material. In elaborating on the early history of work done on iron-silicon alloys, Greiner, Marsh and Stoughton1 went back as far as 1808, and pointed to Berzelius as the first producer of ferrosilicon. During the last twenty years some very important contributions in the form of studies of the iron-silicon system, as well as of the physical properties of these alloys, have been published. In consideration of the brit-tleness and poor machinability of alloys containing more than 5-6 pct Si, commercial interests have stimulated the research done on iron-silicon alloys, but mainly in the range of 0 to 6 pct Si. Yensen2 has shown the influence of im- purities, such as carbon, sulphur and phosphorus on the magnetic properties of 2, 4, and 6 pct Si alloys. Corson3 extended his investigation of physical properties to approximately 32 pct Si, and attributed the brittleness of the high-silicon material to the particular constitutional features of the alloys. Though mentioned earlier in the literature, the effect of nitrogen on crystal growth and magnetic permeability of iron-silicon alloys has been demonstrated recently by Morrill4 and Rostoker,5 respectively. In all the above quoted investigations, except that of Rostoker,5 sand or chill casting processes were employed for the preparation of samples. A major advantage of the powder metallurgy process is that it does away with time-consuming machining in many applications. The powder metal industry is always interested in pro- ducing a finished product employing a material which has presented considerably greater difficulties in manufacture by other methods. Steinit26 described preparation and testing of various Fe-Si alloys which were produced by the powder metallurgy process for the manufacture of armatures for small electric motors or generators. While Steinitz investigated alloys up to 6 pct Si, Rostoker5 went as high as 9 pct Si. The author7 previously discussed the diffusion of iron-silicon compacts containing 6 and 10 pct silicon. It is the object of this paper to study the progress of diffusion occurring in iron-silicon alloys by electrical resistivity measurements. All samples were prepared by pressing and sintering of compacts. Iron, silicon, and Fe-Si powders of various mesh size and compositions were employed to obtain test samples. It was assumed that complete diffusion had been obtained when the electrical resistivity reached a constant value characteristic for that particular alloy. To this end, electrical resistance measurements were used to follow the progress of alloying under various conditions.

Jan 1, 1950

By W. R. McMillan, E. A. Gulbransen, K. F. Andrew

Annealed and electrolytically polished pure iron was oxidized between 650° and 850°C at oxygen pressures of 0.1 to 2 microns Hg. Electron optical studies showed that oxidation occurs discontinuously over the surface and orients crystallographically with the substructure of the metal grains. IN previous papers&apos;, " the authors described the formation of oxide films on high purity iron at norrnal pressures and at temperatures below 300°C. Electron optical studies showed that the oxide film consisted of many individual oxide crystallites having nearly random orientation on each metal grain. The size of the oxide crystallites depended upon the temperature and upon the extent of oxidation. In these studies the crystallite size varied from 250 to 1800Å. Recently, Bardolle and Benard3 studied the crystal habit of the oxide crystallites formed on Armco iron between 650" and 850°C in what might be defined as a "poor vacuum" atmosphere in which the pressure could be varied between 10-1 to 10-3 mm of Hg. Their results showed that, at 850°C and a vacuum of l0-2 to 10-3 mm of Hg, a few well oriented nuclei of oxide were formed on the metal grains. At pressures between l0-1 to 10-2 mm of Hg, many oriented nuclei of oxide were formed. At 750°C the metal surface was bright although an oxide was present. Unfortunately, the extent of oxidation, the oxygen pressure, and the influence of carbon in the metal were not determined, and it was, therefore, impossible to give a theoretical interpretation to the very beautiful light micrographs. Bardolle and Benard used the light microscope to determine the crystal habit of the oxide and X-ray diffraction to determine the orientation relationships. It was believed that the use of controlled amounts of oxidation at high temperature is an important method for the study of the initial stages of oxidation of metals at high temperatures. This paper presents the first results of an electron optical study of this problem for iron. Since the electron microscope and electron-diffraction methods using stripped oxide films and surface replicas yielded finer details of the crystal habit and structure of oxide film, these methods were used in this study. Two types of iron having different carbon contents were used: Puron and Armco* iron. A com- parison of the results for these materials made possible a study of the influence of the surface oxide-carbon reaction on the initial stages of oxidation. Before discussing the experimental results, the thermodynamic equilibria of 1—the oxidation reaction, 2—the surface oxide-carbon reaction, and 3—the solubility of oxygen in iron, will be considered. In addition, it is important to consider the kinetic-theory predictions on the rate of collision of oxygen with the iron surface. Thermodynamic and Kinetic-Theory Calculations Only one type of reaction is usually considered in the oxidation of iron at normal pressures at high temperatures: Fe(a) + 1/2O2(g) = FeO(s) 3Fe(a) + 2O2(g) = Fe3O4(s) 2Fe(a) +3/2O2(g) = a-Fe2O3(s). [11 However, two other reactions may occur: FeO(s) + C (solid solution in Fe) + Fe(a) + CO(g) [2] and Fe(a) + O2(g) e Fe(a) (solid solution of 0,). [ 3 ] In addition, iron will react with the higher oxide at certain temperatures to form FeO. For a complete interpretation of the oxidation of iron at low pres-

Jan 1, 1955

By H. A. Domian, H. Biloni, G. F. Bolling

Jan 1, 1965

By V. A. Phillips

Transmission-electron micrographs of electro-thinned samples of bulk-aged Cu-3.1 pet Co alloy show an aging sequence, supersaturated solid solution — coherent particles — quasi -coherent particles — noncoherent particles. Hardening is due to precipitation of coherent spherical fee coball-rich particles showing coherency strain fields, which are resolved at between 15 and 30A diameter. Loss of- full coherency did not occur until well into the overaged region, even with the assistance of deformation after aging. Different average particle diameters of 123, 92, and 149 ± 10Å were observed in samples aged to peak yield strength at 600°, 650°, and 700°C, respectively, indicating that there is no critical size for peak hardening. Noncoherent particles tended to develop (111) faces and became octahedral in shape. Dislocations tended to nucleate spherical coherent particles which eventually grew together forming large elongated particles. The surface energy of a noncoherent (low-angle) inter-phase boundary is estimated to he about 50 ergs per sq cm. A number of particle lining-up phenomena were observed. Overaging is principally attributed to increase in particle spacing, progressive loss of coherency, and increase in amount of discontinzdous precipitation. COPPER dissolves about 5.6 at. pet (5.2 wt pet) of cobalt at 1110oC1 and the solubility decreases to 0.75 at. petl (0.54 at. pet)2 at 650°C and to 0.1 at. pet or less at lower temperature.&apos; It has been known for many years3-5 that Cu-Co alloys are capable of age hardening. Since cobalt is fee above 417°C and its atom size is only about 2 pet smaller than that of copper, precipitation of coherent particles would be expected. The equilibrium phase precipitated at 700°C and below contains about 10 pet Cu in solution which tends to stabilize the fee structure, lowering the transformation temperature to 340oc.l The alloy is known to undergo discontinuous precipitation in addition to general precipitation; while the former can be seen with an optical microscope, the latter precipitates are not visible except in the grosly overaged condition.5, 6 Extensive use has therefore been made of the ferromagnetic properties of the precipitate in order to follow the course of aging, and it has proved possible to measure the average particle size, spacing, approximate shape, and volume fraction and to determine that the particles are coherent without ever seeing a particle (see for example Refs. 2, 7, and 8). The magnetic measurements of particle size are limited to diameters below about 120Å.7 The present study was undertaken using the techniques of transmission-electron microscopy in order to check the above conclusions, to extend the previous magnetic work to larger particle sizes, and to attempt a more detailed correlation of properties and structure. A portion of this work has already been published.9-11 The present paper is concerned with the metallographic features of precipitation in relation to aging curves. Bonar and Kelly12&apos;13 have published preliminary results of a similar study on single crystals of Cu-2 at. pet Co. EXPERIMENTAL Preparation of Alloy. A Cu-Co alloy, containing 3.12 wt pet (3.36 at. pet) Co by analysis, was prepared from 99.999 pet purity oxygen-free copper and electrolytic-grade cobalt. The alloy was melted and cast in vacuo in a high-frequency furnace using a graphite crucible and mold: Analysis showed chat 0.004 pet C was picked up during melting. The 1-1/2-lb ingot was homogenized in hydrogen for 24 hr at 1000°C. Slices were cold-rolled to 0.005 or 0.003 in. thickness, with an intermediate 650°C anneal in hydrogen at 0.080 in. thickness. Batches of six to ten strips were solution-treated in sealed-off quartz tubes in high vacuum in a vertical furnace and quenched by dropping into iced brine containing a device which snapped off the nose of the tube. Solution treatment consisted of 1 hr at 990°C or 2 hr at 965°C. The latter was employed for all mechanical-property studies, since a tendency was noted for the higher temperature to give porous material. Strips were usually aged individually in a horizontal vacuum furnace, inserting into the hot zone and withdrawing into a cold zone without breaking the vacuum. This method gave a rapid heating rate, permitting the use of short aging times. In some cases, particularly for the longer aging times at the higher temperatures, samples were sealed individually in quartz tubes in high

Jan 1, 1964

By J. L. Brimhall, R. A. Huggins

The structure of deformed internally oxidized alloys of siluer- cotztcrining magnesium and copper-containing aluminum in was studied by thin-film transrrzissiotz electron microscopy. With low to moderate amounts of stvain, the deformation structure in t11e dispersion-hardened alloys was found to consist of a high density of dislocations arranged it a random distribution or in a very fine, but poorly defined, cell stucture, depending on the fitness of the dbsbersion and amount of deformation. 112 contrarst, both pure copper and silver show relative1y Large, distinct cell structures. The particular structures in the alloys are believed to result from particles acting as both dislocaliot~ sources and deflector-s. The Inck of a pvorzolrr~ced deformation texture tr~ regiotzs qf&apos; seoere loctll plastic curvature after heavy deformation can also be traced to the formation of clislocntions on many slip systemes and a lack 01 local concentrations of dislocations of predominantly one sign. THE presence of fine dispersions of a second phase has long been known to have a strong influence upon the mechanical properties of solids. Until recently, the influence of particles on the deformation mechanisms has been assumed to be solely that of barriers to dislocations. This is seen in the classic models of edge dislocations being restrained by the particles.1 4 There is, however, considerable evidence that this is not their primary effect, and that the particles can have a large influence on the defect structure generated by plastic deformation. The presence of particles causes changes in the surface topography observed after deformation. In overaged precipitation-hardened materials containing hard, noncoherent precipitates, fine intersecting slip lines are found after deformation of single crystals. indicating the activation of many slip systems. This has been extensively observed in the A1-Cu systen;, both by optical microscopy and replica electron microscopy.&apos;"13 In some of the observations. the slip lines were seen to extend between the particles. Similar observations of extremely fine slip lines or the apparent absence of normal slip lines have also been found in A~-M~."

Jan 1, 1965

By E. J. Whittenberger

OVER the past two decades, the Cr-Ni stainless steels, popularly termed 18-8 steels, have been used in ever increasing amounts in the aircraft, automotive, chemical, transportation, and building industries.&apos; The contribution of nickel to these steels is associated with its strong austenite forming tendencies which assure good hot and cold forming characteristics, and with its beneficial influence upon corrosion resistance, especially in oxidizing acids. The adjustment of the chromium and nickel contents has been used advantageously in controlling the rollability and the response of the mechanical properties of these steels to cold reduction. With the advent of the jet age and the government stockpiling of nickel, considerable emphasis has been placed upon the development of satisfactory low nickel or nickel-free substitute austenitic stainless steels, especially for nonmilitary applications.2 Fer-ritic 17 pct Cr (AISI Type 430) stainless steel has been effectively substituted for the austenitic grades in applications with mildly corrosive exposure. However, the limited formability and ferromagnetic behavior of this grade do not permit its satisfactory use for many applications. To achieve better formability in a nenmagnetic steel and conserve nickel, a number of low nickel substitute austenitic grades of stainless steel have been developed in recent years. In most of these steels, manganese and nitrogen have been substituted for a portion of the nickel content, for example, 17 pct Cr-4 pct Ni-6 pct Mn (AISI Type 201), 18 pct Cr-5 pct Ni-8 pct Mn (AISI Type 202), and 16 pct Cr-1 pct Ni-17 pct Mn, all of which contain approximately 0.15 pct N.1 The precipitation of ferrite when large ingots of these Cr-Ni-Mn-N steels are heated to temperatures high enough for conversion to slabs on conventional universal slabbing mills has seriously affected the hot workability of these steels. This has been especially true in the case of the 16 pct Cr-1 pct Ni-17 pct Mn (sometimes referred to as 15 pct Cr-1 pct Ni-15 pct Mn) composition." Because it is possible that little, if any, nickel will be available for nonmilitary stainless steel applications in the event of a national emergency, a need has existed for a nickel-free austenitic stainless steel containing hot and cold forming characteristics similar to 18-8 steels and corrosion resistance at least comparable to the 17 pct Cr ferritic steel. The aforementioned deleterious effect of ferrite upon the hot working characteristics of the 16 pct Cr-1 pct Ni-17 pct Mn, steel indicates the importance of establishing the austenite-austenite plus ferrite phase boundaries as the first step in the development of nickel-free or low nickel substitute steels. This paper presents the data collected in an elevated temperature constitutional study of the Cr-Ni-Mn-N system at the South Works of the United States Steel Corp. Experimental Procedure The 300 compositions used in the investigation were melted in a 26 Ib induction furnace and cast in 3 14 in. diam by 18 in tall hot-tcpped ingot mo1ds.

Jan 1, 1958

By W. H. Class

The embrittling effects of oxygen, ozone, nitrogen, air, and surface residues, on NaCl has been investigated. The embrittlement by ozone and oxygen was found to be associated with the formation of a NaClO3 surface compound. In these cases the initial crack that was responsible for fracture (in a bend test) always nucleated at the corners between the tension and side faces. The behavior of air was very erratic and on certain days did not produce enzbrittlement. During these periods, crystals that had become embrittled by the ozone treatment completely recovered their ductility after a short exposure to the ambient atmosphere, It was established many years ago1 that considerable ductility could be obtained in NaCl single-crystal specimens if the crystal surfaces were dissolved in water either during or immediately prior to the test. The original interpretation of this effect by Joffe attributed the enhanced ductility to the removal of surface microcracks by dissolution. Later investigations2&apos;3 have suggested that the exclusion of air from the specimen surface is the criterion for extensive plastic flow prior to fracture. The air em-brittlement in this later work was attributed to the diffusion of gaseous atoms into the surface layers of the crystal, thereby impeding the movement of dislocations. This model satisfactorily accounts for the reembrittlement observed after further air exposure subsequent to the water dissolution treatment. However, the situation has recently become more complex by the observations in several laboratories4-t that under certain conditions air exposure does not impair the ductility of NaC1. It has also been recognized5 that improper drying operations after water dissolution can leave surface precipitates that lead to embrittlement. Cleavage defects on as-cleaved crystals can often be another source of embrittlement. In the present work the effect of the gaseous atmospheres nitrogen, argon, air, oxygen, and ozone, on the ductility of rock salt was studied extensively. The embrittlement resulting from oxygen and ozone exposures was found to be associated with the formation of a NaC1O3 surface film. It is suggested that certain atmospheres, one of which often can be ambient air, which inhibit the formation or favor the decomposition of this compound, can promote ductility. Thus one aspect of the Joffe effect is certainly related to the removal of surface compounds or complexes by water dissolution. The effect of surface precipitates that remain after drying operations and of cleavage defects were also studied. In neither of the latter cases was the embrittlement as severe as that found with a NaClO3 surface layer. PROCEDURE AND SPECIMEN PREPARATION The nature of the embrittlement produced by the agents mentioned above was studied by means of microscopy, mechanical testing, and X-ray diffraction. Specimens were cleaved from large crystals of optical quality sodium chloride obtained from the Harshaw Chemical Co., and, except for those tested in the as-cleaved condition, were given a 15- to 20-sec immersion in distilled water followed by a rinse in absolute methyl alcohol. The specimens were then blotted on a soft, absorbent paper, and dried by a few seconds exposure to a stream of warm, dry air. Such a procedure was found to give a control surface which was microscopically free of residues. (A few crystals were intentionally painted with a concentrated NaCl solution in order to investigate the effect of surface residues). All specimens were of 0.140 sq in. cross-section. Crystals prepared in the above manner were immediately placed in a gas train where they could be exposed to the desired gases for preselected periods of time. For the oxygen and nitrogen exposures, pure reagent-grade gases were employed. The ozone was provided in the form of an ozone-oxygen mixture (approximately 10 pct ozone) prepared by passing commercial grade oxygen over a strong ultraviolet light source. All gases were dried prior to their introduction into the train. Since argon was found to be completely inert in its behavior (i.e., residue-free specimens that were exposed to argon were not embrittled), it was periodically utilized to check the control specimen surfaces as well as the condition of the gas train used for aging the specimens. After exposure to the gaseous media in question, the crystals to be used for the measurement of the strain to fracture were transferred from the gas train to a protective oil bath (without further exposure to the atmosphere) where the tests were conducted in three-point bending. The apparatus was so adjusted that the load could be applied at a constant, continuous rate. Other Snecimens from the gas train were deformed

Jan 1, 1962

By W. F. Carew, F. A. Crossley

IT has been reported that the Ti-8 pct A1 alloy is ductile as water quenched from 800°C but brittle as annealed at 650 °C." The present, somewhat limited, investigation was undertaken to discover the cause of the embrittlement. In order to evaluate to some extent the continuity of the embrittling phenomenon, alloys of 6, 8, and 10 pct A1 were studied. Subsequent to the completion of most of this work, Ence and Margolin reported that a Ti,A1 phase exists and that two or more unidentified compounds may exist in the Ti-A1 system.&apos; Materials, Equipment, and Techniques The alloys were prepared using Kroll process sponge of 124 Bhn hardness. In addition, some 8 pct A1 material was made with iodide titanium. The sponge-base alloys were prepared as melts of 4 lb or more by double arc melting. The iodide Ti-8 pct A1 alloy was prepared as 200-g buttons. The alloys were forged at temperatures of 980" to 1150°C. Results of chemical analyses are given in Table I. A Baldwin-Southwark 60,000-lb capacity tensile machine, equipped with an autographic stress-strain recorder, was used for the tensile tests. Standard tensile test specimens, 0.252 in. diam, were used. Standard Charwy V-notch impact specimens were tested using a -kiehle machine. specimens were sealed in evacuated Vycor bulbs for heat treatment. Powder samples for Debye-Sherrer anal- ysis were prepared by filing and screening through 325 mesh. The —325 mesh fraction was taken in an effort to concentrate a microstructural constituent assumed to be less ductile than the a phase. The powder samples were stress relieve2 by annealing

Jan 1, 1958

By R. G. Hudson, L. Yang

In electrochemical separation of metals, it is necessary to control the potential applied between the electrodes so that only the desired electrode reactions can occur. A knowledge of the minimum potential needed for a given electrode reaction to take place continuously is therefore pertinent to the success of the process. In molten chloride systems at high temperatures, activation polarization is believed to be small,1,2 so the minimum potential needed for the electrolytic decomposition of a molten metal chloride MC1n (n = valence of the metal ion) is therefore essentially the equilibrium electrode potential E of the metal-chlorine galvanic cell in the solution. If the decomposition products are pure M and chlorine at a pressure of p atm, E*Eo-rt/nf in [i] Since EO (the standard electrode potential of MCln) is a constant at a given temperature and R (the gas constant), and F (the Faraday&apos;s constant) are all constants, the variation of E with the electrolytic bath conditions is largely determined by the effect of the latter on the activity of the metal chloride amcln (standard state = pure MC1,). Measurement of E and study of how amCin varies with temperature and concentration, the nature of the solvent, and that of the coexisting solutes are therefore pertinent in the selection of optimum operating conditions for the electrochemical separation of metals. Because of its low melting point, high decomposition potentials of its components, and the low solubility of heavier metals in it, the LiC1-KC1 eutectic melt is a useful solvent and has been used in the electrodeposition of Mo,3Ce,4 and u5 metals. Senderoff and Brenner3 measured the potentials of Zn/Zn++, Fe/Fe++, Cu/Cu+, Mo/Mo+++, and Ag/Ag+ against an Ag/AgCl (pure) reference electrode at 600°C and a concentration of 4.1 mole pct of each chloride and came to the conclusion that complex ion formation occurred in some of the melts. Laitinen and LiUe gave an electromotive force series at 450°C for a number of {Metal/Metal ions (unit activity)) electrodes in LiC1-KC1 eutectic melt, the {pt/pt++ (unit activity)) electrode being chosen as the reference. Walker and Danly7 studied the ther-modynamic properties of NiC12 in LiC1-KC1 eutectic melt over a wide range of temperatures and concentrations by measuring the electrode potential between a nickel electrode and a chlorine electrode in the melt. Since potential measurements made by using the chlorine reference electrode are of more direct significance both in the study of electrochemical separation and in the evaluation of the thermodynamic properties of the melt, it is therefore decided to measure the electrode potentials of cells of the type {M/MCin (mole fraction = N) in LiC1-KC1 eutectic melt/Cl2 over a range of N values and temperature for a number of metal chlorides. The results should indicate the order by which the various metals plate out on the cathode and how this order is affected by temperature and the concentration of the metal chloride in the solu-

Jan 1, 1960

By R. M. Hudson

Equilibrium constants have been determined for the jeaction C (in steel) + H2O = CO + H2 as a function of carbon content (0.013 to 0.74 ujt pct) and temperature (1200° to 1800°F) by using a flow system in which the water-vapor content of various gas mixtures containing hydrogen, carbon monoxide, and nitrogen was systematically controlled. The equilibrium partial-pressure ratio K1 = PcoPH2/PH2O was obtained for two low-carbon steels and four Fe-C binaries from a series of runs made at each temperature. By comparison of Ki with the analogous ratio for the reaction C (graphite) + H2O = CO + H2, the activity of carbon in steel (ac) was comDuted for each temperature and carbon content studied. From these data values of the equilibrium constants for two other carburization-decarburiza-tion reactions of importance were estimated. The chemical compositions of the two steels and four Fe-C binaries used are given in Table I. Specimens of Steel E were cut from a 0.035-in.-thick commercially box-annealed and temper-rolled lot of a low-carbon low-metalloid steel. Steel F was a commercially hot-rolled rimmed steel that was cold-reduced at the Laboratory to a 0.059-in. thickness. Fe-C binaries A, B, C, and D were vacuum-melted and poured, hot-rolled, pickled, and cold-reduced at the Laboratory to furnish specimens 0.035 in. thick. Sheet specimens were cut from all lots to approximately 3/4 by 3 in. by gage and were vapor-degreased with trichlo-roethylene and weighed before heat treatment. Heat treatments were carried out in a hinged-top Hevi-Duty electric furnace with a Vycor-glass tube. When the furnace was at the desired

Jan 1, 1965

By John J. Gilman

F many years it has been suspected that a correlation existed between pits produced by etching and the density of dislocations in crystals. In 1953, the interest in this correlation was greatly stimulated when Vogel et al&apos; demonstrated a 1:l correlation between etch pits and dislocations in low angle grain boundaries of germanium. About two years ago, the author began attempts to produce etch pits at dislocation sites in zinc mono-crystals. These attempts met with success a year or so ago,&apos; and since that time a survey has been made of the etch pit patterns that are associated with various modes of deformation in zinc monocrystals. Unfortunately, the etching technique that will be described is too delicate for use as an accurate quantitative tool. Nevertheless, it has been of considerable assistance in clarifying the structure of plastically deformed crystals. Technique Many attempts were made to produce etch pits in high purity zinc crystals. It was found that several techniques would produce surface grooves at grain boundaries that had misorientation angles greater than about l°, but no technique was successful for boundaries with smaller angles. Among the things that were tried were: electrolytic CrO, (both etching and polishing concentrations), HNO, in various water and alcohol solutions, dilute HC1, many 0-0,-Na,SO, solutions, Cr0,-HC1 solutions, picral, electrolytic Na,S,O,, and electrolytic NaOH and KOH. A few attempts to use cathodic sputtering and thermal etching were also made.

Jan 1, 1957

By G. B. Craig, G. L. Montgomery

A simple technique has been developed which reveals glide dislocations in zone-refined aluminum as etch pits. The technique has been tested quantitatively by making dislocation counts on aluminum single crystals deformed in pure bending. The method has the advantage of being insensitive to the orientation of the surface under examination. A number of etchants which will reveal dislocations in annealed aluminum have been reported.&apos;-&apos; These etchants apparently depend on decoration of the dislocations by an impurity (probably iron) and cannot, therefore, be used to reveal fresh glide dislocations. In addition, it is unlikely that screw dislocations will be etched because of their small attraction for impurity atoms. Amelinckx6 has reported that glide dislocations located along slip traces can be revealed by fluo-rated aqua regia. This work indicated that dislocations themselves, irrespective of decoration by impurity atoms, are responsible for the etch pits. However, as a result of additional etching experiments, Wyon and Marchin7 concluded that dislocations alone are not sufficient to explain the localization of etch pits on slip lines in aluminum using chemical etch-pitting techniques. In addition, their results show that glide dislocations are revealed only on surfaces approximately parallel to a (100) axis and that not all slip systems are attacked. These observations have been confirmed by the present authors. The technique described in this article is insensitive to the orientation of the surface, and reveals dislocations on all glide systems. Fig. 1 shows a deformed region in the neighborhood of a scratch. Dislocations can be seen along three slip traces. In the present work, single crystals of zone-refined aluminum are polished for from 2 to 5 min in a chemical polishing reagent held at 100°C, as described by Herenguel and second:&apos; H2Sol 60 pct H3PO4 30 pct HNO3 10 pet The samples are then coated with a thin oxide film by annealing in air at 450°C for 1 hr and cooling in the furnace. After deformation, dislocation etch pits are developed by immersing the specimen in the same polishing solution which is now held at a temperature of 50°C. The formation of small hydrogen bubbles indicates that etching has occurred. The sample is immediately removed from the solution. Prolonged etching results in enlargement of the pits and, in regions of high dislocation density, the overlapping and loss of resolution of individual pits. It is the experience of the authors that contact with iron will ruinously contaminate the reagent. With careful etching, dislocations separated by half a micron can be resolved optically and the method has been shown to be sufficiently sensitive to reveal dislocations produced during the pre-easy glide region of deformation, Fig. 2. A few examples of glide-dislocation configurations, as revealed by this technique, are shown in Figs. 1 to 4.

Jan 1, 1964

By J. P. Hirth, G. M. Pound

The flux of evaporation J from {loo) , {111), and (110) planes of silver single crystals and from poly crystalline specimens was measured at 744°C under a high vacuum using a time average weight loss method. The evaporation coefficient a was computed from the measured flux and vapor pressure data through the relationship J = aPe/21rmkT. The evaporation coefficients for the close packed planes and the poly crystalline specimens were found to be about 0.4 providing tentative confirmation of the theoretical limiting law for evaporation of metal crystals, a = 73(P/Pe) + % . The (110) planes faceted to form (111) planes, resulting in a greater evaporating area and larger a . The evaporation morphology of the plane surfaces and the surfaces of spherical single crystals is described in some detail. RECENTLY a detailed theoretical analysis of the mechanism and kinetics of metal crystal evaporation was made.1&apos;2 A mechanism for vaporization involving the recession of monatomic ledges across a metal surface was deduced, and evaporation rate equations were derived by considering the dynamics of these ledges. In most metals, the process of evaporation is one of dissociation of atoms from ledges, surface diffusion, and desorption. The gross vaporization rate is controlled by the surface diffusion and desorption rates. It was shown that the monatomic ledges, under the action of surface diffusion gradients, tend to assume a terminal velocity and spacing. Further, it was shown that crystal edges and grain boundaries are a ready source for monatomic ledges. In the case of evaporation of large, perfect crystals with clean, low-index surfaces, the kinetic treatment1 leads to a limiting law for crystal evaporation: where J is the gross flux of vaporization, a is the evaporation coefficient, Pe the equilibrium vapor pressure, m the atomic mass, k Boltzmann&apos;s constant, T the absolute temperature, and P the actual pressure of vapor over the crystal surface. Imperfections such as grain boundaries, pores, cracks, and macroscopic steps should have the effect of causing positive departures from this limiting law. Such an imperfection is a source for monatomic ledges, which, near their origin, are closely spaced and thus give rise to higher evaporation rates there. However, under many experimental conditions,&apos; this effect should be negligible due to the relatively low concentration of imperfections. In general, a more important cause of positive departure should be screw dislocations, and the effect of these has been considered in some detai1.2,3 Monatomic ledges emanate from points where the surface is intersected by a dislocation with a screw component normal to the surface. The dynamics of ledges arising from a screw dislocation are similar to those of ledges arising from a crystal edge. There is a perturbed region near the origin where the monatomic ledges are, under certain experimental conditions, closely spaced and the

Jan 1, 1960