Search Documents

By F. S. Pettit

The oxidation of Ni-3 to 25 wt pd Al alloys has been studied in 0. 1 atm of oxygen at temperatures between 900° and 1300°C. These alloys have been found to oxidize by three different mechanisms which depend on the temperature of- oxidation and the alloy composition. Two of the three mechanisms do not permit a continuous layer of Al,0, to be formed on the alloy surface and the oxidation rates are greater than that for pure nickel. The third mechanism results in the formation of an external A1203 scale and lke oxidation rates are about three orders of magnitude smaller than those for pure nickel. The minimum amount of aluminum required for the formation of external scales of A L,0, has been determined. NICKEL-base alloys are currently the main source of materials for use at elevated temperatures in gas turbine engines. These alloys are usually coated to obtain oxidation resistance. Coatings on nickel-base alloys are frequently formed by reaction of the alloy with aluminum whereby alloyed nickel aluminides are formed. The alloyed nickel aluminides provide protection to the nickel base alloy because external scales of A1203 (alumina) are formed during oxidation and mass transport through A120, is slow in comparison to mass transport through most other oxides. In view of the protective properties of A1203, it is important to know how much aluminum is required in these alloys in order to form external scales of AlzO,. The present paper is concerned with the oxidation kinetics and the oxidation mechanisms of Ni-A1 alloys and the minimum amount of aluminum required in these alloys for the formation of external scales of Alz03. THEORETICAL CONSIDERATIONS When a Ni-A1 alloy is heated in oxygen at elevated temperatures, the following reactions can take place on the surface of the alloy where the oxide phases are assumed to be virtually pure: These oxide phases are in the form of nuclei scattered over the surface of the alloy and, in view of their rapid formation, they need not be in equilibrium with the alloy. As the oxidation process continues, equilibrium between the alloy surface and the oxide phases is approached and the stability of the oxide nuclei is determined by the composition of the alloy at the alloy/oxide interface because of the following reactions: 3NiA1204 + 2A1 (alloy) = 4AlzO3 + 3Ni (alloy) [41 4Ni0 + 2Al (alloy) = NiA120, + 3Ni (alloy) [ 5 1 Application of the mass-action law to Eqs. [4J and [5J yields the following equilibrium conditions for these reactions: where aA1 and aNi are the activities of aluminum and nickel, are the standard free energies of formation of NiO, A1203, and NA1204, respectively, R is the gas constant, and T is the absolute temperature. If the composition of the alloy at the alloy/oxide interface is such that (akl/ahi) is greater than the equilibrium values defined by Eqs. [6] and [7], then Reactions (41 and [5] will go to the right as written. Conversely, if the alloy composition is such that the activity ratio (aLl,/aki) at the alloy/oxide interface is less than the equilibrium values, then Reactions [4] and [5] will proceed to the left. The equilibrium activity ratios in Eqs. [6] and 171 can be calculated since values for the standard free energies of formation of the oxide phases are available. Standard free energies of formation for NiO and A1203 have been tabulated by Elliott and ~leiser.' The standard free energy of formation for NiA1204 can be obtained from the data of Tretjakow and Schmalzried.' The results of these calculations are tabulated in Table I. Table I shows that the following inequality is valid over the temperature interval 900" to 1300°C: (Reaction [5]) (Reaction [4]) « 1and therefore the aluminum activity for these compositions can be taken as equal to the square root of the activity ratios (i.e., aNi = 1). If equilibrium is estab-

Jan 1, 1968

By S. Feuerstein, L. Rice

The effect of low pressures on the flow and fracture behavior of molybdenum is described. For poly crystalline samples, room-temperature tensile tests indicate greater ductility under 10 Torr than under intermediate pressures up to and including atmospheric pressure (760 Torr). In addition, tests conducted at 760 Torr under atmospheres of air, dry nitrogen, and purified argon exhibited no apparent difference in mechanical properties. Critical tests involving baking in situ as well as those involving single-crystal deformation further imply that the ductility effect is a pressure-dependent phenomenon related only to the fracture process. This dependency is discussed in terms of adsorption and diffusion contributions. THE effect of very low pressures on material properties has heretofore been presumed to be important only for substances possessing relatively high vapor pressures at ambient temperatures. Research has therefore been concentrated primarily on organic solids and liquids, and in some instances on metals such as zinc and cadmium. Most vacuum-effect studies' on the mechanical behavior of metals have been performed under conditions of either cyclic loading or creep rupture at elevated temperatures, i.e., over extended time periods. These studies were not restricted to high vapor pressure materials but also encompassed such metals as gold, copper, and nickel. Very little concern, however, was placed upon the importance of a vacuum environment on the mechanical behavior of metals subjected to a simple unidirectional deformation at ambient temperatures. A tension test is generally of short duration as compared to a creep test, and at room temperature vacuum effects if any would be expected to be surface-limited. In early 1963, Kramer and podlaseck2 reported a change in the bulk flow behavior of aluminum single crystals during room-temperature tension tests. The deformations were performed under pressure conditions of 760 to 3.4 X 10-8 Torr and indicated for the first time a vacuum surface effect contributing to the bulk tensile behavior of metal specimens. As a consequence, an experimental program was initiated in this Laboratory to study the effects of ultrahigh-vacuum conditions on the mechanical behavior of metals. The results of a preliminary study on poly-crystalline molybdenum3 revealed, unlike Kramer's observations of changes in the stress-strain behavior, only an increased ductility under ultrahigh vacuum. Flow behaviors were nearly identical for all tests re- gardless of pressure. This paper presents comprehensive results obtained in this area of research. 1) EXPERIMENTAL PROCEDURE Three material categories were used in this study: sintered and are-cast polycrystalline molybdenum of nominal purity 99.93+ pct and single-pass electron-beam zone-refined molybdenum single crystals having a nominal purity level of 99.99+ pct. The interstitial levels (weight percent) as determined by the Materials Testing Laboratories, Division of Magnaflux Corp., were as follows: sintered molybdenum (C— 0.005, H—0.0004, O—0.015, N—0.008); and arc-cast molybdenum (C-0.0038, H-0.0003, O-0.015, N-0.023). Single-crystal molybdenum obtained from Materials Research Corp. had a typical interstitial analysis of C-0.0015, H-0.00007, O—0.00045, and N-0.0001. Tensile specimens having a 5 mm diam by 50.8 mm length were prepared from these materials. An average grain diameter of 0.059 mm was obtained for the sintered specimens following a 4-hr, 1600°C heat treatment. Grain sizes from 0.019 to 0.149 mm were obtained in the arc-cast specimens following heat treatments from 1100° to 1600°C for 1 hr. This series of specimens was used exclusively for the grain-size effect studies. All samples were electrolytically polished in 97 pct sulfuric acid solution prior to testing. Experiments were performed at room temperature in an ion-pumped ultrahigh-vacuum system positioned in an Instron tensile machine, Fig. 1. A constant strain rate of 4.2 x 10-4 sec-1 as derived from crosshead displacement was assumed for the deformations. Starting vacuums ranged from 2 to 0.5 X 10-10 Torr. These pressure measurements were made using corrected values4 of an NRC Redhead gage. Comparative readings were also made against a G.E. triggered dis-

Jan 1, 1967

By B. H. Kear, S. M. Copley

Stress-stvain curves are presented for Ni3Al (y') cvystals in several ovientations, deformed in tension and compression at constant displacement rate, at temperatures from 70° to 2000°F. Both the yield stress and wovk havdening increase with temperature, with the magnilude of the effect dependent on ovientation. The yield stress maximum occurs at 1500°F in [001], at 1400°F in [011], and at 1200°F in [111]. The suppression of the yield stress peak in [011] and [111] orientatiorzs is due to the onset of cube slip, rather than octahedval slip, a1 elerated temperatures. The temperature and orientalion dependence of work hardening in nominal single slip orientations corvelates with changes in the CRSS ratio for octahedral slip and cube slip, in agreement with a model for work hardening based on pinning of screw dislocalions by cross slip from an octahedral plane into a cube plane. It is concluded that the unique plastic properties of y' have a decisive influence on both ductility and strength characleristics of y +y' nickel-base superalloys. It has long been recognized that the strength of y + y' nickel-base superalloys depends on the precipitation of the y' phase [basically Ni3(Al,Ti)]. This paper presents new data on the strength characteristics of single crystals of simple y' (Ni3Al), and forms the first step in a systematic program aimed at elucidating the mechanisms of hardening in the complex commercial alloys. 1) EXPERIMENTAL PROCEDURE Single crystals of off-stoichiometric NiA1 were grown from the melt under vacuum by a modified Bridgman method. The melt was poured into a preheated alumina mold and crystal growth was promoted from one end by gradient cooling. The crystals were homogenized by annealing in hydrogen at 2400° F for 72 hr followed by furnace cooling. Chemical analysis of samples taken from several crystals gave an average composition of 88.2 wt pct Ni. Spectrographic analysis gave as the principal impurities in weight percent— Si (0.02), Mg (0.01), Fe (0.02), Ti (0.002), Cu (0.008), and Co (0.03). Tensile specimens with specifications as in Fig. 1 were prepared by a series of operations involving electrical discharge machining, precision grinding, and electropolishing. Compression specimens with dimensions 1/4 by 1/4 by -3/4 in. were prepared in a similar manner, except for the initial shaping operation using a precision cut-off wheel/two-circle goniometer unit to give selected crystal orientations. Specimens were deformed in a Baldwin-Wiedemann testing machine with furnace attachment, using a strain rate of 5 x 10-4 sec-1. The strain measuring device consisted of extension arms attached to the tensile grips (or compression plattens) at one end and leading out of the furnace to an LVDT at the other. Temperature was controlled by a thermocouple placed in contact with the specimen. According to the known phase diagram for the Ni/A1 system,' when an alloy of the specified composition is cooled from the melt, primary y (nickel solid solution) dendrites grow at the expense of the liquid phase, which becomes enriched in aluminum. At the eutectic composition the remaining liquid freezes as a two-phase mixture of y + y' (Ni3Al). Upon further cooling, a solid state transformation occurs, involving the precipitation of y' in the primary dendrites, and the complete transformation of the eutectic mixture to massive y'. The as-cast structure consists, therefore, of y + y' dendrites with interdendritic regions of massive y', i.e., transformed eutectic, Fig. 2. 2) DISCUSSION OF RESULTS 2.1) Stress-Strain Curves. Fig. 3 shows tensile stress-strain curves for crystals in orientations close to [001] deformed at temperatures from 70° to 2000°F. Both the yield stress and total elongation are strongly temperature-dependent. At 70°F, easy glide is absent, and the work hardening coefficient BIT - G/300. At T > 1500°F, the negative slope of the stress-strain curves is due to pronounced necking in the crystals. The yield stress maximum at 1500° F corresponds with a minimum in the ductility, Fig. 4. The extensive duc-

Jan 1, 1968

By C. B. Sclar and

A graphical solution is presented for the equation v =v &-cU where vb is the volume of liquid b of density 6 that must be added to liquid a of volume va and density d, in order to obtain a heavy-liquid solution of preferred density dm for mineral fractionation. The equation is valid only for pairs of liquids whose volumes are additive, but empirical density-composi-tion data show that this condition is met over wide compositional ranges by all of the heavy solutions that are commonly used for mineralogical analysis. The chart is a nomograph which consists of two horizontal volume scales ad a vertical density scale arranged so that one volume scale occurs on each side of the density scale. All the scales are linear. For two liquids of density d, and &, respectively, each volume scale has an independent level on the density scale, ad the graphical solution for vb on the nomograph is obtained by construction of one straight line. The chart can be prepared easily to cover any density range with any desired accuracy by proper selection of the scales. Heavy-liquid separation of particulate samples is an indispensable analytical procedure in the mineralogical and mineral-process ing laboratory. Heavy liquids may be used 1) to isolate specific minerals in the form of high purity products or to fractionate samples into several products with density limits for petrographic, chemical, spectrographic, and X-ray diffraction analysis,'-' 2) to determine the density of minerals,'-l4 3) to facilitate the recognition and identification of optically similar associated minerals,15 and 4) to determine quantitatively the degree of liberation of ore minerals in ores and mill products at various limiting sizes." The preparation and recovery of heavy liquids suitable for mineralogical analysis are discussed fully in the literature and operational procedures and special equipment for heavy-liquid separation are described in numerous references.'747837' 13,18,25-42 Useful tabulations of heavy liquids for mineralogical analysis are given by Tickell, Mil-ner,' Twenhofel and Tyler,' and Lange and a very complete up-to-date list of minerals arranged according to increasing density is given by Mursky and Thompson. USEFULNESS OF THE DENSITY CHART For any particulate sample, the limiting densities selected for its fractionation by means of heavy- liquid separation depend on the respective densities of the constituent minerals and the specific objectives of the separation. In many instances, the densities of readily available and practicable heavy liquids do not coincide with those that are required, and suitable liquids of the proper density are obtained by dilution with miscible liquids. In the general case, one may wish to mix either a pure liquid (A) or a solution of two miscible liquids (AP,,) of known density with either a pure liquid (B) or a different solution of the two miscible liquids (A,B,) of higher or lower density in order to obtain a final liquid of the preferred density. In conventional practice, the latter is reached empirically by trial-and-error addition of increments of one liquid to the other. The density of the final solution may be determined either by accurately weighing a known volume of the final liquid,'' 532;4;40, or by means of I)a Westphal balance, 2)a hydrometer, a refractometer in the case of binary liquid systems for which density-refractive index data are available,15 4) natural or synthetic .indicators of known density, a pycno-meter. Like all trial-and-error methods, these procedures are time consuming and many require special equipment which may not be available, The purpose of this paper is to show that the volume of a liquid of known density which is to be added to another liquid of known volume and density in order to obtain a solution with an intermediate preferred density may be determined rapidly and accurately by a graphical method provided that the two liquids form ideal or quasi-ideal solutions whose volumes are additive. Empirical data47'8~'8~4B'49 show that solutions prepared from pairs of miscible

Jan 1, 1961

By R. W. Armstrong, P. C. Jindal

The hardness dependence on grain size for polycrys-talline cartridge brass and a leaded brass has been measured by Brine11 and Rockwell B testing. In each case, the hardness, H, depends on the average grain diameter, 1, according to: H =Ho + kHl-1/2 where Ho and kH are experimental constants. Diamond pyramid hardness values have also been measured as a function of the indentation size and grain size to give additional information on the nature of the hardness test and the dependence of hardness on micro-structure. The hardness of polycrystalline brass depends on its grain size. Bassett and Davis' demonstrated this as early as 1919 by making Brinell hardness measurements on cartridge brass. Since then, the hardness of this type of material has been measured as a function of grain size by making Rockwell,2'3 Vickers,4 and Brinell5 tests. he hardness dependence on grain size has also been measured for other materials. Angus and summers6 investigated the grain size dependence of the Brinell hardness of polycrystalline copper and a Cu-4.5 pct Sn bronze. In other studies, nickel,? Armco iron,Big an Fe-0.07 pct C alloy,I0 and an 0.39 pct C-12.45 pct Cr stainless steel" have been investigated. In some of the preceding cases, the hardness results have been analyzed to show that the hardness varies with the average grain diameter, 1, according to an l-l\4, l-1/4 or I-2 dependence,11-13 The studies of the influence of grain size on hardness have not been based on any theoretical model. This may be because the hardness of a material is itself a complicated property. However, attempts have been made to correlate, experimentally and theoretically, the hardness of a material with its unidirectional stress-strain behavior.14-l6 On this basis, Hall" proposed that the polycrystal hardness dependence on grain size might follow directly from the Hall-Petch18,19 relation for the grain size dependence of the yield stress. Thus, the hardness-grain size relation was given as: H = Ho + kHl-1/2 [1] where Ho and kH were taken as experimental constants. The relation was applied to the measurements on brass,' copper,6 bronze,= and Armco iron.' More recently, this relation was shown by Armstrong and jindal20 to adequately describe the measurements on cartridge brass made by Bassett and Davis' and Babyak and Rhines.5 In this case, the relationship was taken a step further by independently relating the values of Ho and kH to the values of oyand ky, previously reported by Armstrong, Codd, Douthwaite, and petch21 from measurements of the yield stress dependence on grain size for this type of material. In the present investigation, new Brinell and Rockwell B hardness measurements have been made as a function of grain size for a cartridge brass and a leaded brass. In addition, diamond pyramid hardness values were measured as a function of the indentation size. All these results are applied to a further analysis of the hardness dependence on grain size. MATERIALS AND EXPERIMENTAL METHODS Cartridge brass and a leaded brass were selected for this investigation for two main reasons: it was anticipated 1) that these materials could be cold-worked and recrystallized to a wide range in grain size and 2) that the results to be obtained on these typical industrial materials could be usefully compared with previous investigations. The chemical analyses of the actual materials which were employed are given in Table I. The as-received 1/2- and 3/4-in.-thick plates were given various reductions in thickness by cold rolling. The rolled material was heat-treated at various temperatures between 330" and 850°C for differing time periods from 5 min to 9 hr to achieve a variation in the average grain diameter between 0.0339 and 0.000543 cm.22 During heat treatment, the brass was protected from zinc loss by packing it in chips or foils of the same composition material. Reasonably equi-axed grain structures were obtained in each case. The metallurgical grain sizes of the specimens were determined from measurements of the average linear intercept on a random line. Annealing twin interfaces were not counted along with grain boundaries. The

Jan 1, 1970

By R. E. Delavault, H. V. Warren

In normal soils there are usually 10 to 50 parts of copper in every million parts of .soil. Only 0.2 to .5 pct of this copper can be found by any simple cold chemical attack. Now, with rubeanic mid reagent paper, a prospector or field geologist can detect as little us 4 ppm of readily available copper ill soil. This degree, of sensiticity is enough to determine the presence. of copper anamalous an as and, ecentually, to discouer copper mineralization. Circumstances determine whether it is better to make analyses in the field or in a permanent laboratory. The rubeanic acid test described in this article has been designed primarily for field use: it is simple and virtually foolproof, and it requires a minimum of field kit." It is sensitive, easily de- • Ed. Note: Persons Interested in purchasing kits suitable for rubeanic acid prospecting can obtain information by writing Eldrico Geophysical Sales Ltd., 633 Hornby Street. Vancouver 1, B.C. The University of British Columbia does not produce these kits for sale and has no financial interest in their production tecting 4 ppm of readily extractable copper in a soil. This is by no means a quantitative test, but it is accurate enough to provide a valuable indicator of copper anomalous areas for both prospectors and field geologists. The easiest method for detecting metal deposits that do not produce visible float or stains is to make a simple chemical test for the metal in overlying soil, or in the silt of a stream that may have picked up metal farther upstream. In Brief: Testing for copper may be done easily by shaking a soil sample with strong acetic solution in a small test tube and pouring the mud into a small filter, the tip of which rests upon a strip of reagent paper impregnated with rubeanic acid (di-thio-oxamide). When copper is present—and only when it is—a blue spot develops. The more copper, the darker the spot. If the copper content is merely the small amount present everywhere, there is a pale blue or hardly visible spot; if it is abnormally high, the spot will be dark. There are, of course, intermediate cases where the experienced geochemist cannot tell offhand whether a medium-strength spot represents rich agricultural soil, weak copper mineralization, or distant rich copper mineralization. Reagents and material are inexpensive; the test may be readily done on the spot with a simple kit easy to pack and handle. Anyone interested in general problems of soil sampling as applied to prospecting may refer to an article recently presented to the AIME. In exploration work it is the contrast between the metal content of anomalous and background areas that is important; absolute values become of greater interest when an anomalous area is being investigated in detail. With specific reference to copper, it has been the authors' experience that the amounts of metal extracted from anomalous and normal soils with buffer solutions of decreasing pH show better contrast if an acid reagent is used. This contrast tends to increase with increasing acidity until 3 to 4 pH is reached. Using a short cold attack on unheated soil, it has been found that further increases in acidity do not produce better results, and only increase the hazards involved in carrying strong acids. An acidity of about pH 4 is satisfactory for direct determination of copper by dithizone. But dithizone itself introduces some problems: it must be made up fresh at frequent intervals, and with some soils, notably those with much ferric iron, oxidation mag take place before all the copper has reacted with the dithizone. Rubeanic acid keeps its strength unimpaired for long periods, is unaffected by oxidation, and is practically specific for copper at pH 4. Consequently it seems an ideal reagent to use in prospecting for copper. History and Background: Rubeanic acid (systematic name: dithio-oxamide (SC-NH2) has long been known as a spot test reagent for some heavy metals with which it gives a number of compounds. Only copper and some metals of the platinum family are believed capable of providing any ru-beanate compounds under conditions of moderate

Jan 1, 1959

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' 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,' Rosenberg and Cadoff,9 and Bilello and cadoff' 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'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-' 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 W. D. Ludemann, J. E. Dor, L. A. Shepard

Recovery of the creep resistance of 99.99 pct pure Al was studied at temperatures 540°, 573°, 600°, and 611°K. Poly-crystalline specimens crept under a stress of 950 psi to a strain of 5.5 pct were allowed to recover for periods of from 1 min to 16 days under a residual stress of 4.4 psi. Increased creep rates upon reapplication of the 950 psi stress evidenced softening of the material. The activation energy for the recovery process was found to be 64,000 cal-per mol. THREE major observations have clearly revealed that the creep of polycrystalline aluminum above about 510°K is controlled by the rate of climb of jogged edge components of dislocations past the barriers impeding their motion: 1) Extensive polygonization, characteristic of climb, takes place during creep.'9' 2) The activation energy for creep equals the estimated activation for self-diffusion. 3) The stress law for creep coincides with theoretical predictions based on the climb mechanism.4,5 The decreasing rate during the primary stage of creep must be ascribed to the introduction of additional barriers to the glide of dislocations. During secondary creep the density of barriers must remain constant, the rate at which new barriers are formed being equal to the rate at which they recover. For this reason the nature of the barriers, and their rates of formation and recovery, are significant to a complete understanding of creep. It is proposed here to study the kinetics of the recovery of barriers to creep of high-purity polycrystalline aluminum in the climb range. The technique will consist of creeping a series of tensile specimens under a prescribed stress to given state following which the specimens will be recovered for various times and temperatures under approximately zero stress. The amount of recovery will be determined by comparing creep curves following the recovery. EXPERIMENTAL PROCEDURE The material used for this investigation was a 99.99 pct pure Al supplied by the Aluminum Co. of America in sheets of 0.001-in. thickness having an H-18 temper. The spectroscopic analysis of impurities gave Cu-0.004 wt pct, Fe-0.002 pct, Si-0.001 pct, others-0.000 pct. Creep specimens were machined with their tensile axes in the rolling direction of the sheets. Annealing in a molten potassium nitrate-potassium nitrite bath for 1 hr at 686K, followed by air-cooling, resulted in a grain diameter of from 0.22 to 0.27 mm. Creep machines were equipped with constant stress lever arms of the type suggested by Fullman, Carreker, and Fishher,' which maintained the applied stress to within 0.04 pct of the reported value of 950 psi. Creep strains were calculated from extensions over a 6-in. gage section to the nearest 3 X 10"5. The temperatures of testing, 540°, 573", 600°, and 611°K, were obtained by immersing the specimen and extensometer assembly in a molten KNO2-KNO3 bath. Temperatures during the periods of creeping could be maintained constant to within better than ± 1°K. Correlation of creep data at different temperatures of testing demanded the use of an appropriate temperature-compensated time 8. Previous investigations have shown that the creep of pure metals at temperatures above about 0.55 of their melting temperatures can be correlated by the functional relationship3" e=f(0), s= const. [1] where e = total plastic strain during creep f = a function that depends on stress a = stress 6 = te-Q/RT = a temperature-compensated time t = time under test Q = 35,500 cal per mol = activation energy for creep R = gas constant, and T = absolute temperature In the present series of tests all of the creep was conducted under the same stress of 950 psi. To provide identical initial conditions for the recovery all specimens were first crept to a temperature-compensated time of = 47 X 10-14 min which resulted in a creep strain of 0.055 + 0.005 regardless of the creep temperature. During recovery, a small stress (4.4 psi) was permitted to remain upon the specimens to maintain tautness in the pulling assembly and to prevent mechanical damage to the soft specimen. After each recovery period, the full stress was reapplied and

Jan 1, 1961

By T. H. Alden

BaCKOFEN, Avery, and Turner have shown"2 that the large tensile elongation in superplastic metals is correlated with a high strain-rate sensitivity of the flow stress. At present, the reported superplastic materials are complex eutectic or eutectoid systems with a large solubility range in at least one of the terminal phases. This note describes the behavior of a solid-solution alloy (at 22°C)5,= Sn-1 wt pct Bi, which has been processed in order to produce a fine grain size.2,3,6 The chill-cast, 1/2-in.-diam ingot was homogenized for 1 week at 130°C, extruded at O°C to 0.083 in. diam, quenched within 15 sec, and stored at -196°C. Samples were warmed to room temperature 16 hr prior to mechanical testing or metallography. During tension testing at room temperature, it was found that the flow stress reached a steady-state value which depended on the imposed elongation rate and was independent of strain to an error of about ±5 pct. In Fig. 1 is plotted log a vs log e where a is the measured steady-state flow stress and i the imposed strain rate. These quantities were calculated using "instantaneous" area and length to an accuracy in the sample diameter of ±5 x 10-4 in. The log a-log 1 curve is sigmoidal. The strain-rate sensitivity, which may be given by d log old log 6,' has a maximum value of about 0.48 between i = 5 x 10-4 and 5 x 10-3 per min. The value is well above frequently measured values for metals at creep temperatures (0.1 to 0.3) and comparable to the rate sensitivity for Pb-Sn solder: which has been shown to be superplastic.2,3 Note, however, that below or above the strain rate of 10-3 per min the rate sensitivity of Sn-1 pct Bi decreases and at high rates is more nearly like that of conventional metals. The grain size of this material is relatively uniform, Fig. 2, and may be specified by an average linear intercept i.' Grain growth occurs at room temperature but after 16 hr, L = 5. Tensile elongation was measured on finer-grain-size specimens, which had been at room temperature 15 min prior to testing. After 15 min, the samples are fully recrystallized. Three strain rates, 5 x 10-3, 10-2, and 2.5 x 10-2 per min, were employed for which the measured elongations varied from 425 to 500 pct. These results would seem to justify the term superplastic, applied to this alloy, although the ductility is inferior to that of the Sn-Bi eutectic.3 The Sn-1 pct Bi alloy, which is a single-phase material with equiaxed, fine grains, is an excellent one on which to make an estimate of the Nabarro-Herring9 creep rate. According to the model of Avery and Backofen2 a substantial part of the deformation must be accomplished by N-H creep in the strain rate-temperature regime in which d log a/d log 1 = 0.5. The maximum N-H creep rate is given by9 L2kT where v is the atomic volume and D the self-diffusion coefficient. For this test v = 3.O x 10-23 cu cm, D = 1.4 e-23,300/RT = 2.5 x 10-l7 sq cm per sec,7 L = 5x 10+4 cm, T = 295°K, and a = 2000 lb per sq in. = 1.38 x 10' dyne per sq cm. The calculated creep rate is E= 4 x 10-11 per sec which may be compared with the measured creep rate at this stress, Fig. 1, of 6.7 x 10-5 per sec. The N-H creep rate is seen to be negligible. Earlier consideration2 of this question was in error because incorrect diffusion coefficients were used. The assistance of R. R. Russell in metallography is acknowledged gratefully.

Jan 1, 1967

By Byoung Whie Lee

In order to investigate the change in surface energy as a function of cold work ad environment, the contact angles between a liquid and the surface of elec-tropolislzed aluminum with various degrees of cold work were measured. Specinlens cut from aluminum rod (99.99 pct Al) were annealed at 600°C for 24 hr and then cold compressed parallel to the drawing di-rectzon of the rod in increments of approximately 10 pcl from 0 to 70 pct. The sections, taken transversely to the drawing direction, were mechanically and electro chemically polished. Contact angles between the sessile drop of distilled water, glycerin, and glycol and the specimen surface were measured optically. The differences in interfacial energy were calculated from the contact angle data. The effects of chemical washing and surface roughness of the specimen on the contact angle were also studied. The liquid contact angle was a linear function of the degree of cold working (D) ad changed markedly with differed liquids. THE plastic deformation of polycrystalline metals induces a preferred orientation in crystal structure. A preferred orientation changes the atomic density of a metal surface. This change in surface atomic density brings about the changes in surface interactions of the metals."' Also, the plastic deformation of metals raises the internal energy by increasing the stored energy.' Consequently, the total Helmholtz free energy of the metal increases and, in turn, the surface free energy of the metal. In order to study the effect of plastic deformation on solid surface phenomena, the liquid contact angle of a sessile drop made on cold compressed and elec-trochemically polished aluminum was measured. The preferred orientation in aluminum, due to uniaxial compression, is known to be [110] direction parallel to the compression axis. However, this description is only a first approximation of a rather complex orientation. In actuality, half the crystallites are more than 10 deg from this orientation, regardless of the amount of compression." For randomly oriented fcc crystals, the average surface atomic density can be calculated as follows: The number of atoms in a unit volume is 4/a3, where a is a lattice constant. Hence, the number of atoms per unit surface area is about (4/a3)2/3 = 2.52/a2 However, when the preferred orientation took place along [110] direction, the surface atomic density is as follows: The area of (110) plane of a unit cell, considering ridges and valleys, is BYOUNG WHlE LEE, Member AIME, is Research Metallurgist, Atomic Energy Research Institute of the Republic of Korea, Seoul, Korea. Manuscript submitted May 29, 1968. IMD 2 x J2- a x Ga = 6a2 Since this plane contains 4 atoms, the surface atomic density is 4i6a2 = 2.3/a2 Upon changing from random to preferred orientations, the surface atomic density decreases about 10 pct. The increase in the stored energy and the change in surface atomic density due to plastic deformation would certainly alter the surface tension of a metal. EXPERIMENT High purity aluminum rod of 6-mm diam (99.99 pct Al, obtained from Johnson, Matthey & Co., Ltd, London, England) was cut into sections approximately 1 cm long. The transverse cross section was mechanically polished. Then, the specimens were annealed in air at 600°C for 24 hr. Thus, eight annealed specimens were uniaxially compressed at room temperature from 0 to 70 pct of cold working with an increment of 10 pct. The extent of cold working was determined by the percentage of reduction in axial length to the original length. The specimens were further polished mechanically and, then, electrochemically with the Modern Battelle Type 1 electropolishing solution with a current density of 0.3 to 0.4 amp per sq cm and an electrode potential of 8 to 12 v at the solution temperature of 78" to 82°C for approxirr~ately 5 min. The Modern Whlte Diffusing Screen \t/ --------------- re Goniometer for Mercury Lamp ^-^ Sessile Drop ¦=• Condenser c=> Microscope l_*~_i bJ Mlcromster Filer with Protmcter Fig. 1—Schematic diagram of apparatus set-up for liquid contact angle measurement.

Jan 1, 1970

By G. C. Rauch, T. H. Courtney, J. Wulff

The superconducting behavior of Nb-Ti alloys containing 40 and higher wt pct of Nb and variable oxygen content was studied as a function of thermomechanical processing. Critical current density (Jc) us applied transverse magnetic field (H) data were obtained for 0.010 in. diam wires at 4.2° Kin steady fields up to 120 kOe. The precipitate responsible for increase in Jc in each of the alloys studied was related to niobium and oxygen content. In alloys containing 40 to 60 wt pct Nb with less than about 1600 ppm of oxygen, w precipitation in the range 350° to 450°C appears to be responsible for the maximum value of Jc observed. In alloys of higher oxygen content (-2500 ppm), the oxygen-enriched a phase which precipitates in the temperature range 450° to 550°C is a more effective fluxoid pinner. Maximum Jc in low oxygen (350 ppm) 80 wt pct Nb alloys was achieved by 600°C aging heat treatments. Measurements of resistive critical field also reflect the changes in composition accompanying precipitation in these alloys. It is well known that the superconducting critical current density, Jc, of solid solution superconducting alloys can be increased by precipitation heat treatments.1-4 The precipitate, to serve as an effective superconzlucting fluxoid pinner and increase Jc, must be of proper size and spacing.' Its electronic properties must also be different from those of the matrix.6 As in age-hardenable systems, optimum temperature and time of aging are dictated by the composition and history of the particular solid-solution alloy under consideration. The alloy diagram of Ti-Nb7-9 shown in Fig. 1 is typical of ß-stabilized Ti alloys. Precipitation to achieve high Jc in Ti-Nb and Ti-Ta alloys is usually carried out in what is believed to be a two-phase a + ß region.1,2,4,10 Previous work on Ti-Nb4,11 and on other ß-stabilized systems such as Ti-V, Ti-Mo,12-14 has shown, however, that the first precipitate to appear is not necessarily (Y. Under specific conditions the meta-stable w phase is precipitated instead. A further complication arises in such alloys from the strong influence of contained oxygen on the phase stability and on the kinetics of precipitation.15,16 The superconducting critical current density of Nb-Ti alloy containing less than 36 wt pct Nb has been studied by Vetrano and Boom1 and Kramer and Rhodes.4 The most effective aging temperature was found to be about 425°C; fluxoid pinning was attributed to the precipita- tion of a by Vetrano and Boom and of w by Kramer and Rhodes. Although the latter workers used recrystal-lized material and the former cold-worked, studies carried out in our laboratory lead us to believe that the type of precipitate responsible for increased Jc is unaffected by prior cold-work.17 That other types of precipitates may also be effective, especially in higher oxygen-containing alloys, has been demonstrated by Reuter, et al.2 in this laboratory with the aid of electron microscopy. Reuter et al. found that at an aging temperature of 600°C, fee-TiO is precipitated, resulting in high values of Jc. Even higher values of Jc were observed after heat treatment at 500°C, but the precipitate could not be indexed as fee-TiO. Whether the precipitate was oxygen-rich w or a could not be clearly distinguished. The work nevertheless indicates that appreciable amounts of oxygen alter the nature of the phases precipitated at different temperatures, at least in alloys containing more than 36 wt pct Nb. Since such compositions are more readily worked into suitable wire than lesser niobium content alloys, it was decided to study the response to aging of the alloy compositions reported in the present paper with the hope of clarifying the effect of composition on the precipitation processes responsible for high values of Jc. I. EXPERIMENTAL Nb-Ti alloys containing 40, 50, 60, and 80 wt pct Nb were prepared by are-melting 35 g ingots of iodide crystal bar titanium and electron-beam melted niobium.

Jan 1, 1969

By R. E. Pray, M. C. Fuerstenau, J. D. Miller, B. F. Perinne

Quartz cannot be floated with potassium amyl xanthate as collector at any pH. Complete flotation is achieved with certain minimal additions of amyl xanthate and Pb from pH 5.8 to 8.5 and with amyl xanthate and Zn from pH 7.5 to 8.1. The active species of these metal ions responsible for activation are shown to be PbOH and ZnOH No flotation was effected when Cu and Mg were added as activators. Previous work1,2,3,4,5,6,7 has shown that quartz may be floated in the presence of sulfonate or soap when the following two criteria are met: 1) some portion of the added metal ion has hydro-lyzed to its first hydroxy complex, and 2) a precipitate of the metal-collector has formed. Since sulfonate and soap, two substantially different anionic collectors, conform to these criteria, this phenomenon may well be expected to be of a general nature. In this view then, it would appear that any anionic collector, including xanthates, should also float quartz when these two criteria are met. The object of this paper is to demonstrate the conditions under which quartz will respond to flotation in the presence of potassium amyl xanthate. EXPERIMENTAL MATERIALS AND TECHNIQUES Quartz was prepared by leaching the sized sample (48 x 150 mesh) with HC1 until no iron could be detected in the leach liquor. Conductivity water, made by passing distilled water through an ion exchange column, was used in the investigation. Potassium amyl xanthate was selected as collector. Purification of this reagent was effected with the following method: 1) a laboratory sample of potassium amyl xanthate was dissolved to saturation in hot ethyl alcohol and filtered several times while still hot, 2) the filtrate was cooled, during which time potassium amyl xanthate crystallized out of solution; the xanthate was then separated from the solution by filtration, 3) the amyl xanthate was dissolved in fresh hot ethyl alcohol and filtered, 4) step (2) was repeated, 5) the crystalline potassium amyl xanthate was washed with ethyl ether five times, and 6) the xanthate was dried for 24 hours under continuous vacuum. Inorganic impurities are insoluble in hot ethyl alcohol whereas dixanthogen is soluble in ethyl ether. Five-gram charges of 48x150 mesh quartz were conditioned for 3 min in 125 cc of solution containing the reagents and were floated with the small-scale cell and procedure described in previous papers.478 Reagent grade HCl and NaOH were used for pH adjustment, while reagent n-amyl alcohol was added as frother. Frother addition is necessary with this technique, but it should be mentioned that data, similar to those shown in Fig. 2 for 1 x10-4 mole per liter Pb++ and 1 X 10-4 mole per liter AX-, were also obtained with a Hallimond cell in the absence of frother. EXPERIMENTAL RESULTS Clean quartz could not be floated with amyl xanthate at any pH. Flotation was achieved under certain conditions, however, when various metal ions were added to the system. The activating ions used were lead, zinc, copper and magnesium, added as reagent grade chloride salts. Lead as Activator: The effect of Pb++ as activator was examined at five different additions of amyl xanthate (Fig. 1, 2 and 3). When 1x 10-5 mole per liter xanthate was added in the presence of 1 x 10"4 mole per liter Pb ++, no flotation was obtained at pH 6.5 or below, while a recovery of about 60% was effected at pH 7. When the pH was increased to 7.8, no flotation was again obtained, Complete depression was noted from pH 7.8 to about 10, at which point a recovery of about 30% was achieved. At about pH 11, the precipitate of lead amyl xanthate which was present below this pH dissolved and complete depression occurred. When 2.5x 10-5 mole per liter xanthate was added with the same lead addition, no flotation was obtained at pH 5.8 and below, whereas a recovery of

Jan 1, 1965

By J. J. Gilman, T. A. Read

S URFACE effects in the cleavage of brittle crystals have been known for some oftime,1, 2 but our knowledge of surface effects in the plastic deformation of crystals is of relatively recent origin. In 1936, Roscoe9 eported his observations comparing oxidized and clean cadmium crystals. He found that cadmium crystals 1/2 mm in diam when covered by an oxide film less than 1200 atoms thick had critical shear stresses about twice those for clean crystals. This effect is anomalous in that it cannot be caused by the inherent shear strength of a normal oxide film. Either an interaction between the film and the crystal causes the large increase in strength, or the oxide film has an abnormal mechanical strength. Several authors since Roscoe have studied surface effects in crystal plasticity. A complete review of the literature will not be presented, but most of the results may be summarized as follows: 1—The critical shear stress for slip is increased markedly by very thin oxide films on zinc4 and cadmium crystals;" ' also, the entire stress-strain curve is raised by an oxide coating;3,4 both effects increase with decreasing specimen radius. 2—Surface active substances such as oleic acid destroy the hardening effect of oxide films.4, 6-8 The time required for the reduction of the hardening effect to begin increases with the viscosity of the solvent4~ 3—There has been dispute in recent years about the effects of surface active substances themselves on plastic deformation. Rehbinder, Lichtmann, and Maslenikov9 state that surface active substances reduce the strength of metal monocrystals markedly and increase the number of slip lines per unit length at a given strain. The effect varies with orientation for tin crystals being smallest for x, = 0, largest for x, = 45", and intermediate as x, + 90°.10 Several workers claim to have disproved the results of Rehbinder')"." but Rehbinder and Lichtmann9 have defended themselves and Masing with his collaborators" has confirmed some of Rehbinder's results. 4—Immersion of metal specimens in electrolytes such as sodium chloride reduces their strength12,13' and subsequent cathodic polarization increases the strength reduction. For noble metals, the strength reduction occurs for both anodic and cathodic polarization," increasing with increasing polarizing voltage. 5—Electroplated copper coatings effect the creep rates of zinc monocrystals. Poly crystals are not affected." 6—Immersion in dilute acid destroys the hardening effect of an hydroxide film. When an acid solution is applied to a steadily creeping crystal, the creep rate jumps to a very high value and then drops gradually to a steady value which is higher than the initial steady rate.*' " If the hydroxide film to be removed is very thin (about 100A thick), the creep rate does not change measurably, but a small rapid strain-increment is observed when dilute acid is applied." 7—Crystal plasticity depends on specimen dimensions. This was observed by Lichtmann and Ven-stroml" for monocrystalline tin and by Andrade and Kennedy" for polycrystalline lead. 8—Immerson of copper crystals in mercury has no effect on their stress-strain curves.'" 9—Bombardment of creeping cadmium crystals with a particles which penetrate the surface to only 0.005 mm may increase their creep rates as much as five times.''' The experiments which will be described in this paper pertain to: (a) the effect of a crystal's shape on its plastic deformation, and (b) the effect of metallic surface films on plastic deformation. For the work pertaining to (a), crystals of zinc, tin, and lead were grown which had various shapes. These were tested in simple tension. Macroscopic rotations about the tension axis which are unpredicted by the classical theory of slip were observed, and serrated edges were observed in tin crystals. For the work pertaining to (b), the effects of evaporated and electrodeposited films of copper, nickel, gold, and zinc on the mechanical properties of zinc, tin, and lead crystals were determined. The films influenced both slip and twinning of the crystals in creep and tensile tests. Experimental Details Chemically pure zinc, tin, and lead were used in this investigation. The zinc was 99.999t pct pure;

Jan 1, 1953

By A. H. Daane, R. L. Myklebust

The yttrium-manganese system has been investigated by thermal, metallographic, and X-ray diffraction methods. There are three intermetallic compounds present: YMn2 which melts congruently, YMn4, which undergoes syntectic decomposition, and YMn,, which undergoes peritectic decomposition. The compound YMn4 is ferromagnetic at room temperature with a Curie temperature of 214°C. There are eutectics at 25.2, 60.9, and 82.0 wt pct Mn which melt at 878°, 1100°, and 1075°C, respectively. Crys-tallographic data are given for YMn4, and YMn12 . The terminal solid solubilities are low. In a general program of study of yttrium metal in this laboratory some alloy systems of this metal with elements of the first transition period have been examined. This work was originally instigated by experiences in cladding yttrium with jackets of some protective metals such as Inconel for high-temperature service in air.' In some cases a low-melting phase was observed to form between the Inconel and the yttrium resulting in failure of the samples. In characterizing this reaction, a survey was made of the systems of yttrium with chromium, manganese, iron and nickel,, and it was found that a eutectic was formed between these metals and yttrium on the yttrium-rich side of the system. This present study of the Y-Mn system was carried out to examine in more detail the alloying nature of yttrium, and to correlate trends that have been observed in previous related studies. It was observed by Voge13 that in the systems of lanthanum, cerium, and praseodymium with each of the metals in the first transition series, the tendency to form compounds diminished in the order nickel, cobalt, and iron while no compounds were formed with manganese, chromium, or titanium. Beaudry and Daane4 observed a similar behavior in the systems of yttrium with some members of the first transition series except that the tendency for compound formation was greater. Both the Y-Ti5 and Y-Cra systems consist of simple eutectics, while in the Y-Fe,7 Y-CO, and Y-Ni4 systems, there are four, eight, and nine compounds, respectively. In addition to the above trends, the similar atomic radii and electronegativities of yttrium and thorium invite a comparison between their alloying behaviors with a common element such as manganese. In crys- tallographic studies, Florio et al.' have identified three intermediate phases in this system which are ThMn2, Th6Mn23, and ThMn,,. Gschneidner and Waber10 have examined published information on alloy systems of the rare-earth metals and have correlated this information with current alloying theory. From their study, they predicted that the Y-Mn system would contain one intermetallic compound. On the basis of this prediction, the trend in the alloying behavior of yttrium with the elements of the first transition series and the alloying behavior of thorium with manganese, one might expect from one to three intermetallic compounds to form in the Y-Mn system. A consideration of Hume-Rothery's rules of alloying based on size-factor, electronegativity, and valency suggested a small terminal solubility and possible compound formation. The present study was undertaken to confirm these predictions of low terminal solid solubility and compound formation and to establish the general alloying behavior of yttrium with manganese. EXPERIMENTAL Materials. The manganese used in this investigation was obtained from the Foote Mineral Co. as electrolytic plates of 99.9 pct stated purity; the yttrium metal was prepared in this laboratory. Table I gives the analyses of these materials. For the solubility studies at the yttrium-rich end of the alloy system, distilled yttrium, whose major impurity was 200 ppm Ti, was used. Preparation of Alloys. The alloys were formed by comelting the two metals in an are-melting furnace under an atmosphere of argon. The buttons thus formed were inverted and remelted three to five times to promote homogeneity. Due to the high vapor pressure of manganese, it was assumed that the weight lost during are-melting was all manganese. This assumption was based on the very good agreement observed between calculated compositions and chemical analyses of several alloys. The compositions of the dilute alloys used for solid solu-

Jan 1, 1962

By W. C. Thurber, C. J. McHargue

THE deformation textures of metals have been extensively studied because of both the practical implications in metal fabrication and the fundamental insight into the behavior of the metal during deformation. Most studies have been concerned with pure metals or solid-solution alloys and relatively little attention has been directed toward more complex alloy systems. Barrett' and Brick2 have summarized work on multiphase systems in terms of the relative ease of deformation of the phases. Thus, if the phases are deformed about equally, each develops its usual texture for the particular fabrication process. This behavior has been reported for Ag-Cu and Cd-Zn alloys3 and 62-38 brass.4 A dispersion of hard particles in a softer matrix interferes with the normal reorientation of the matrix lattice during deformation, thus distorting the texture or in some cases preventing its development. Increasing the amount of carbon in steel has been observed to increase the randomness of the ferrite5 and almost random textures have been reported for A1-12 pct Si eutectic alloy wires.6 It has also been noted that if the second phase is platelike, deformation causes the platelets to line up in a common direction. The present work concerns a quantitative study of the deformation textures in aluminum-uranium alloys containing 5 or 13 wt pct U. Alloys in this composition range are characterized by a pure aluminum matrix in which the intermetallic compound UA14, is dispersed. From a fundamental standpoint, this system affords the opportunity for studying the effects of a dispersed phase on texture development in an essentially pure metal matrix. Further, these alloys are commonly used as fuels for nuclear research and testing reactors, and a knowledge of the preferred orientation may be useful in evaluating physical and mechanical properties as well as response to fabrication. The deformation textures in sheets of aluminum (99.996 pct), Al-5 pct U alloy, and Al-13 pct U alloy were determined for reductions in thickness of 90 pct by cold rolling. The texture in an alloy rod of the latter composition which had been extruded at 455°C was also determined. It can be readily ascertained from the aluminum-uranium phase diagram that the 5 pct U alloy is hypoeutectic and contains 7.3 wt pct (3.4 vol. pct) UAl4. The 13 pct U alloy has the eutectic composition and contains 18.9 wt pct (9.4 vol. pct) UAl,; however, because of nonequilibrium solidification conditions, the typical eutectic microstructure was not developed. The alloys were prepared by open-air melting in graphite crucibles. The slab castings were cropped and cold reduced 90 pct in thickness on a two-high mill. The slabs were reversed after each pass. Although the uranium-bearing alloys exhibited edge cracking, they could be cold rolled to the desired thickness. Small coupons 1.5 by 0.75 by 0.1 in. thick were sheared from the center of the rolled plate with a major axis parallel to the rolling direction. To obtain a specimen of sufficient thickness for machining of the X-ray diffraction specimen, three coupons were laminated into a single three-ply sandwich with a suitable adhesive. Adhesives which could be cured at room temperatures were selected in order to prevent recrystallization during bonding. Bondmaster M-648 (Rubber and Asbestos Corp.) was used for the pure aluminum laminate and Armstrong A-6 (Armstrong Products Co.) was used for the alloys. The alloy for extrusion was cast in a cylindrical graphite mold! cropped and machined into a 3-in. diam by 4-in. long billet. This billet was preheated to 455°C, upset to 3.125 in. diam, and extruded on a 700-ton hydraulic press at a speed of 6 fpm through a flat-face, 1-in. diam die. The rod was quenched with a water spray as it emerged from the die. Pole-distribution data were obtained using the diffractometric method described by Jetter and Borie.7 A spherical X-ray diffraction specimen of

Jan 1, 1961

By E. S. Machlin

The wzodel of Blandin and Diplunt;, generalized to include a phase factor, is applied to the carbon-carbon interaction in iron. Darken&apos;s "energetic" model is generalized to include not only first neighbor interactions but further neighbor interactions as well. On the bases of these generalized models relations are derived for the activity of carbon in both austenite and ferrite in terms of the carbon-carbon Pair interaction energies. A single function then yields the pair interaction energies consistent with the experimental activities of carbon in both ferrite and austenite. Thus, a simple explanation is given for the observation that the nearest-neighbor interaction between carbon is repulsive in austenite and attractive in ferrite. Certain consequences of this approach are explored. OnE object of the present paper is to attempt to take into account the consequences of electrostatic contributions to the carbon-carbon pair interaction energy for carbon as a solute in iron. Friedel&apos; has shown that oscillations in electrostatic potential are to be expected about a solute atom in a metallic solution. Blandin and 6lant6&apos; have shown that such oscillations yield an interaction energy between pairs of solute atoms that obeys the relation: W{ = A cos(2ftFri + 4>)/(kFri)3 [l] where kF = Fermi wave vector, ri = distance between solute atoms comprising the pair7 <p = phase factor dependent only on electronic nature of solute and solvent, A = coefficient dependent only on electronic nature of solute and solvent. Machlin3 found that Eq. [I] accurately described the pair interaction energy derived from short-range order measurements based on field ion microscope observations of dilute alloys of platinum. He also found that the value of the phase factor $ derived from residual resistivity measurements agreed well with that obtained from the analysis of the short-range order data. Harrison and paskin4 were able to predict the long-range ordering energy of 0 brass using Relation [I] and residual resistivity values to predict the value of the phase factor $. Machlin5 has repeated their analysis and applied it to the prediction of the long-range ordering energy in AgZn and AgCd with excellent agreement between prediction and experiment. Both A and $ are independent of the crystal structure. The Fermi wave vector depends uniquely upon the conduction electron concentration per unit volume in the spherical approximation of the Fermi surface. Thus, Eq. [I] is expected to apply to both fer- rite and austenite with only one set of values of A and $. Mossbauer studies6 yield the result that iron has one 4s electron. We shall make an assumption found to hold previously for platinum3&apos;7 and nikel, which is that only the s electrons are involved in shielding the perturbing potential of carbon. With this assumption, kF = 1.35 A-l. Although A and $ may be obtained from certain mdels&apos;&apos;&apos; we shall take A and $ to be empirical constants in the spirit of Kohn and osko.&apos; Thus, Eq. [I] involves two adjustible parameters. Consequently, two independent relations in A and $I are required in order to evaluate them for carbon as a solute in iron. We may use a recent analysis of Aaronson, Domain, and poundg who showed that Darken&apos;s energetic model,1° as well as others, can be used to describe the activity-temperature data for carbon in iron in both the aus-tenitic and ferritic phases. Darken&apos;s model takes into account only first neighbor pair interactions. For our needs, all neighbor pairs need to be taken into account. It is convenient to generalize Darken&apos;s model. The result for the partition function for austenite is: over the temperature range 800" to 1200°C and where the uncertainty corresponds to one standard deviation. Eq. [4] effectively yields only one relation. Another relation is required to obtain unique values for A and $. One property of Eq. [4] is that it is independent of crystal structure. Hence, data for a iron can be used to obtain another relation. To arrive at this relation we must generalize Eqs. [2] and [3] so that they may be applied to the bcc a iron. The result is that:

Jan 1, 1969

By M. C. Flemings, R. Mehrabian, G. E. Nereo

Analytic expressions were developed and applied in two preuious papers to predict effects of solidification variables on macrosegvegation. In this paper, experiments are reported to test the analyses quantitatively and qualitatively. Good quantitative agreement is obtained between predicted and measured compositions for the following types of segregation (all experiments were on laboratory-size ingots of Al-4.5 pct Cu alloy): (a) inuerse segregation, (b) composition distribution in ingots solidified with unidirectional heat and fluid flow, (c) simulated centerline segregation, and (d) segregation resulting fro, change in solidification cross section. In addition, qualitative agreement of prediction of theory with experiment is obtained for (a) segregation resulting from sudden change in heat extraction during solidification, (b) centerline segregation in a bidirectionally solidified ingot, and (c) under-riser positive segregation. IN two previous papers,&apos;&apos;2 it was shown that a variety of apparently different types of macrosegregation result from the same basic mechanism. This mechanism is the flow of solute -rich liquid to feed solidification and thermal contractions. Analytic expressions were derived and examples given of their application to predict macrosegregation. Calculations were for Al-4.5 pct Cu alloy. In this paper, we describe a series of experiments designed to test results of the foregoing calculations. All experiments were on laboratory-size ingots (weighing 10 lb or less), of A1-4.5 pct Cu alloy, nominal composition. APPARATUS AND PROCEDURE A typical mold and chilling arrangement used for unidirectional solidification is shown in Fig. 1. Before casting, this assembly was placed in an electrically heated air furnace, and heated to 680" to 700°C. When the plaster reached the desired temperature (usually 1 to 2 hr after the furnace temperature), the mold was filled through a funnel extending from the outside of the furnace into the cavity. The assembly, with liquid, was held for several minutes to allow convection to subside. The coolant, either air or water (depending on the desired solidification rate), was turned on and the ingot solidified. With the four sides and top insulated, heat removal for solidification was through the bottom surface, resulting in a fully columnar structure. Design of the casting cavity was varied, as discussed below, to produce desired macrosegregation effects. However, for castings solidified with unidirectional heat flow, the top was always open to the furnace atmosphere and temperature. The bottom sur- face rested on a chill through which the heat was extracted, and all other surfaces were insulated with plaster. In one case to be discussed, the mold design was changed to obtain bidirectional solidification. In this case, two vertical chills were used (on two faces of the mold) and the bottom was insulated. In a last case discussed, that dealing with under-riser segregation, experiments were conducted in bottom-chilled sand molds at ambient temperature. Except for this one case, all experiments were conducted in the electrically heated air furnace as described above. In many ingots, thermocouples were used to obtain continuous records of solidification. In all cases, these were chromel-alumel couples with output continuously monitored on a twelve-channel recorder. Where thermocouples were employed, these were embedded in the refractory mold walls during mold-making so their heads extended to the vertical centerline. Six to nine couples were employed, placed at intervals from the chill. METAL-MEL TING PROCEDURE Aim analysis was A1-4.5 pct Cu; alloy was prepared from high-purity aluminum (99.9 pct Al) and A1-Cu master alloy (50 pctOFHC Copper, 50 pct high-purity Al). The A1-4.5 pct Cu alloy was cast in 5- to 7-lb ingots for subsequent remelting for the macrosegregation experiments. Melting for ingot casting was as follows. After preheating the melting crucible for an hour or more, the charge was melted and superheated to a temperature of 690" to 720°C. After stirring, the melt was degassed with chlorine for 10 min. After stirring again for 1 min, a gas sample, under reduced pressure, was taken. When the sample was gas-free, the crucible

Jan 1, 1969

By J. C. Sarace, S. M. Sze, C. R. Crowell

Thin films of tungsten 077 n-type germanium, silicon, and gallium arsenide were obtained by reacting tungsten hexafluoride with the semiconductor surface in an argom atmosplrere at temperatures between 325° and 400° C. Capacity-voltage, current-iloltage, and photoelectric measurements were used to investigate the characteristics of the tungsten -semiconductor diodes thus Produced. The junctions are shown to he very close to ideal Schottky barlp/ers with barrier heights measured with respect to the Fermi energy of 0.18, 0.65, and 0.78 1.1 jar W-Ge, W-Si, and W-GaAs, respectively. The electrical properties of the W-Si interface show no deterioration when heated to 1000°C in dry forming gas for 5 min. A theoretical value of the Richardson constant, A, appropriate to the semiconductor-hand structure has been used in evaluating the current-voltage characteristics. ThE W-Si surface-barrier diode was initially proposed for investigation because the eutectic temperature with silicon (1400°C) is much higher than that in the Si-Au system 1370°C).1 This would permit more flexibility in heat treatment and possibly provide greater reliability at elevated temperatures. The lower work function of tungsten (4.54 ev)2 compared with that of gold (4.78 ev)3 also suggested that a lower barrier height would be obtained with tungsten and hence a lower forward bias and lower minority carrier injection ratio for a given current density. The investigation was extended to include the characterization of W-Ge md W-GaAs surface-barrier diodes. The tungsten films have been produced by reacting WF6 with germanium, silicon, and GaAs surfaces in an argon atmosphere at temperatures from 300° to 500°C.4 This process is a very satisfactory alternative to the relatively difficult process of evaporating tungsten films in vacuo. To ensure an adequate electrical characterization of the tungsten-semiconductor interface, three types of barrier-height measurements have been performed. The mutually consistent results obtained lead to the conclusion that the tungsten-semiconductor junctions are indeed of the Schottky type. EXPERIMENTAL PROCEDURE The apparatus used for producing tungsten films is shown schematically in Fig. 1. It consists of an argon carrier gas line to which metered amounts of tungsten hexafluoride can be rapidly added. The mixture passes through a heated reaction tube containing the semiconductor slices and is exhausted to a hood. The argon is purified by passage through a 6-in. column of titanium turnings maintained at 800°C. The tungsten hexafluoride dispensing arrangement was designed by V. C. Garbarini and W. R. Bracht.4 A measured amount of liquid tungsten hexafluoride is injected into the argon stream and vaporizes. The mixture passes through a sodium fluoride absorption cell to remove traces of hydrogen fluoride, then into the nickel reaction tube. The tube is 12 in. long with an inside diameter of 1/2 in. The center section is heated by a furnace of the self-supporting wire-filament type. It was chosen for its rapid heatup and cooling. The wall thickness is 10 mils except for a 3-in. hot zone which is 30 mils thick to reduce thermal gradients along the length. The samples to be coated are placed on a sapphire plate and centered in this section. The tungsten is deposited with the following sequential steps: the loaded reaction tube is flushed with argon at a rate of 500 cu cm per min and heated to 370°C. Then 0.45 g of liquid tungsten hexafluoride is injected into the argon stream. The samples are held at this temperature for 2 min. The tube and samples are then cooled to room temperature and the samples removed. Tungsten films were grown on (111) faces of silicon and germanium polished with Linde A abrasive and lightly etched with HF-HNO3. The GaAs surfaces were (100) faces chemically polished with a H2SO4-H2O2 solution. The films have typical sheet resistances of 0.2, 8, and 15 52/0 when grown on germanium, silicon, and GaAs, respectively. After the tungsten deposition, ohmic contacts were alloyed on the back surfaces of the wafers. Ohmic contacts to the germanium and silicon were obtained by alloying Au-Sb at 370°C, ohmic con-

Jan 1, 1965

By H. M. Beatty, T. L. Joseph, Gust Bitsianes

SUCCESSFUL blast-furnace operation depends upon securing an optimum balance between a number of important variables. This balance will vary somewhat from furnace to furnace in the same plant and with raw materials available in different plants. It is difficult to isolate the effect of any one variable unless it has a rather pronounced effect upon the efficiency of operation. The physical properties of iron ore and coke are known to affect the process, yet it is difficult to set up any definite yardstick by which these differences can be appraised. In recent years, whenever practicable, operators have taken steps to obtain better control over the absolute size and uniformity of size of the ore and coke. The size and physical character of the stone, however, have received comparatively little attention, emphasis being placed upon chemical composition, fluxing efficiency, and action as a desulphurizing agent. Since the stone occupies only about 10 per cent of the volume of the stock column, its influence upon permeability is less than that of the ore and the coke. The large pieces now used have a tendency to roll to the center and thus increase the permeability of this part of the stock column. Smaller pieces would be more effective in promoting the permeability of denser portions of the stock. The size and physical properties of blast- furnace stone are important also for reasons best understood by considering the results obtained by the crushing and the sizing of iron ore. Crushing ore to about 2 in. at Provo, Utah, and to about I in. in Alabama produced material somewhat less permeable to gases, even though sized. In spite of some sacrifice in permeability, the efficiency of the process was improved by crushing and sizing the ore, as indicated by a substantial saving in fuel and increased output. This saving in fuel is due in large measure to more thorough reduction of ore in the upper part of the furnace and to a corresponding decrease in the amount of carbon consumed prematurely above the tuyeres by the reaction CO2 + C = 2CO ? 70,080 B.t.u. [I] The physical properties of stone, its size and rate of calcination are important because they influence the rate at which C02 is released and the completeness of calcination before reaching a level in the furnace where the temperature does not exceed 1050°C. (1920°F.). Briefly, carbon dioxide from the limestone or from the reduction of ore may cause premature com- bustion of carbon above the tuyeres. Assuming that 1000 lb. of limestone is used per ton of pig iron, about 435 lb. of C02 will be liberated during calcination of a good grade of stone. If all this CO2 from the stone were to oxidize carbon prematurely, about 120 lb. would be consumed above the tuyeres in a heat-absorbing reaction. The amount of carbon thus consumed when a

Jan 1, 1942

By J. Hunter

LouisiaNa ranks fifth among the oil-producing states. In 1941 the state-wide production was slightly in excess of 118,000,000 bbl. of crude oil and conden-sate—an increase of 15,000,000 bbl. over the total production for the year 1940. During 1941, south Louisiana produced approximately 90,900,000 bbl. of oil, while north Louisiana produced approximately 27,000,000. South Louisiana, therefore, produced about 3¼ times more oil than north Louisiana. During the past year, 574,800,000 M cu. ft. of natural and casinghead gas were produced in Louisiana; thus 1941 gas production exceeded the 1940 production by 104,900,000 M cu. feet. Completion records for the year 1941 show that 1512 wells were completed in or adjoining producing fields. These wells, which were drilled to extend known producing areas or were located within producing areas, are not considered as wildcat wells. Of the total number of field wells drilled, 1152 wells found production; the remaining 360 were junked or dry and abandoned. This means that 76 per cent of the wells drilled were successful. Considering the precarious nature of oil-field development, this percentage of new producing wells is exceedingly satisfactory. Of the total number of field welis com-pleted as producers, approximately 86 wells found new and distinct producing sands. During the past year, 30 new fields were found. Four parishes—namely, Caldwell, Catahoula, Orleans, and Natchitoches— entered the ranks of oil-and-gas-producing parishes in the state during that period. This brings the total number of oil-and-gas-producing parishes in the state up to 48. Thus, 75 per cent of the parishes in Louisiana now contain one or more oil and/or gas fields. Although sufficient data are not available to list the addition reserves found in Louisiana during 1941, it is felt that such reserves undoubtedly exceed twice the total production for that year, or at least 200 million barrels. In reserves, therefore, Louisiana ranks third among the various oil-producing states. South Louisiana Although 16 new fields were found in south Louisiana, the outstanding developments in that part of the state were not limited to new discoveries. In fact, most of the increase in reserve in that area has been most consistently accumulated through extensions of known producing sands and through the finding of new sands in already productive fields, especially at depths below the presently producing sands. Since data concerning most of the new extensions and sands are too extensive for discussion here, the more important developments arc briefly outlined below: Abbeville, Vermilion Parish.—Production extended ½ mile to the west. Anse La Butte, St. Martin Parish.—New oil sand discovered at 9400 ft., thus extending production to east flank. West flank defined by dry holes. Barataria, Jefferson Parish.—New oil sand discovered on east side of field at 7620 ft. New

Jan 1, 1942