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

Constant-load creep-rupture tests at 1100°, 1300° and 1500°F were made on a Type-316, 18 Cr-8 Ni-ZMo, austenitic stainless steel to determine the relationship between ruptzire life and other aspects of creep behavior, The test program was designed on a statistical basis to check the variability of various creep properties and to determine emprically the dependence of minimutn creep rate and ruptzire life on initial stress, The dependence of the rupture life, tn, on the initial stress, a,, is found to be related to the stress dependence of the wrinimum creep rate, i,, and secondary creep strain, The instabilities or breaks found in the conventionual log - logt, plots can be traced to either a change in the linear dependence of log i, on log a, or to a change in secondary creep strain. In the alloy tested, rupture is found to be predowmanantly of the inter crystalline type, For this material, the secondary creep strain is -found to depend on the nature of the grain boundary precipitate, which affects grain boundary migration. It has been shown for a variety of metals and al tested at elevated temperatures under creep conditions that the empirical relationship exists between rupture life, tr, and minimum creep rate, i,. This relationship, in which C and a are in some cases independent of stress or temperature, shows in a general way that creep rupture depends on creep behavior prior to tertiary creep stage. To determine more fully the factors which control creep rupture requires, therefore, more knowledge of the relationship between rupture and creep behavior. To obtain such information, constant-load creep-rupture tests have been made on a Type-316, 18 Cr-8 Ni-2 Mo, austenitic stainless steel at 11000, 1300°, and 1500°F. These tests were designed on a statistical basis to determine the variability of the various creep properties measured and to determine empirically the stress dependence of minimum creep rate and rupture life. As a result of this work, various empirical relations have been established defining more closely the factor, C, in the relation between rupture life and minimum creep rate. This factor is found to be proportional to the amount of strain during secondary creep. The variation of secondary creep strain with rupture life and temperature is determined and its effect on rupture behavior discussed. The variability in minimum creep rate and rupture life is also dis- cussed and the dependence of these quantities on initial stress is interpreted. TEST MATERIAL AND TEST PROCEDURES The material tested is an austenitic stainless steel, AISI Type 316, of the following composition: C, 0.07 pct; Mn, 1.94 pct; P, 0.01 pct; S, 0.021 pct; Si, 0.38 pct; Cr, 18 pct; Ni, 11.4 pct; Mo, 2.15 pct; All 0.003 pct and N, 0.043 pct. The material was received as a hot-rolled 0.5-in. diam bar which was sectioned into 3.25 in. long blanks. Each blank was heat treated by holding 1/2 hr at 2000° F followed by water quenching. A 3/8-in. coupon was removed from each heat-treated blank and the microstructure was examined, the grain size determined, and the hardness measured. The microstructure was found to be uniform from blank to blank and the grain size was found to range from 4 to 6 ASTM numbers. Hardness, measured using a 20-kg load, varied between 127 and 148 Vpn. Tensile creep-rupture specimens having a 0.25 in. diam within a 1.5-in. reduced section were machined from the heat-treated blanks. Constant-load tests were made on these specimens at 1100°, 1300°, and 1500 F. Six specimens were tested at each stress level. Three creep-rupture machines were employed for tests at each temperature, two specimens being tested in each machine for each stress level. To minimize any effect due to local inhomogeneities along the length of the as-received bar the specimens were randomized before test. Each specimen was tested to rupture and in most cases an autographic extension-time curve was obtained. For very low stress levels at 1500°F no auto-

Jan 1, 1962

By A. Phillips, D. H. Killpatrick, V. Kerlins

The tensile failure of two dispersion-strengthened Ni-20 Cr alloys was studied and compared to the fracture of a similar alloy with no dispersoid. The fracture characteristics were studied using electron fruc-tography and transmission electron microscopy. In all three cases, the mode of failure was found to be microvoid coalescence. The failure in the dispersion-strengthened alloys was found to have initiated at the particles. The size of the dimples in the fracto-graphs was found to be related to the spacing of the particles, but not to the total elongation before failure. The elongation before failure was found to be related only to the amount of dispersed phase. These results are compared to those predicted by a theoretical model of ductile failure. THE continually increasing strength requirements for creep-resistant materials capable of long life at service temperatures above 1600"F (875"C) have developed considerable interest in dispersion-strengthened nickel-base alloys. The high yield strength of these alloys results from the presence of a dispersion of fine particles of thoria which act to impede the normal motion of dislocations. Many theories have been presented to explain the role of the dispersion in the strengthening of these alloys. A good review of these theories is presented by Ansell.' The ultimate strength of such alloys is not only related to mechanisms raising the yield strength, but also to the amount of work-hardening which occurs. This work-hardening is determined by the work-hardening rate and the amount of plastic strain before failure. Since fracture limits the amount of plastic strain, a study was undertaken to gain a better understanding of the fracture mechanisms in these alloys. Electron fractography and transmission electron microscopy were used to study the fracture characteristics in two dispersion-strengthened alloys and one similar alloy containing no dispersoid. These results are related to the tensile properties of the dispersion-strengthened alloys at room temperature and at 2000°F (1095°C). The results are also related to a theory for ductile fracture. 1) PROCEDURE Standard tensile specimens were made from 0.020-in. (0.051-cm) sheets of three different Ni-Cr alloys. The alloys had nominal compositions of Ni-20Cr, Ni-20Cr-2Th0,, and Ni-2OCr-4Th0,. All alloys were supplied in an annealed condition. The specimens were fractured in tension at room temperature to determine the effect of the dispersed thoria on the fracture appearance. The tensile properties were determined from the average of a minimum of three ten- sile specimens for each condition tested. The fracture surface of all of the alloys was examined by electron microscopy using standard two-stage plastic-carbon replica techniques. Thin foils of the dispersion-strengthened alloys were used to investigate the size and spacings of the dispersion particles, and to investigate sections of the 2 pct thoria alloy taken from areas of the specimen adjacent to the fracture surface. The thin foils were mechanically ground to 0.010 in. (0.025 cm) and then chemically polished to approximately 0.001 in. (0.003 cm) in a solution of 29 g of ferric chloride and 10 ml hydrochloric acid in a liter of water. The polishing solution was maintained at 150"F (65°C) during the thinning operation. The final thinning of the foil was done electrolytically using a solution of 700 ml ethanol, 100 ml 2-butoxy eth-anol, 120 ml distilled water, and 78 ml perchloric acid (70 pct). The potential was maintained at 15 v and the bath temperature at -20°~ (-29°C) during the thinning operation. The final thickness of the foil was approximately l000A as determined from the width of twins boundaries observed in many of the foils. 2) RESULTS AND DISCUSSION The electron fractographs of the fracture surfaces of the alloys are shown in Figs. 1, 2, and 3. The normal interpretation of this type of fractographz is that, as a result of the difference between elastic and plastic properties of the matrix and the particles or other inhomogeneities in the alloy, microvoids are formed

Jan 1, 1969

By K. E. Beu

IN the measurement of retained austenite concentrations in steels using the integrated intensity method,1 Averbach has pointed out" that the absorption factor A(0) for a flat sample making a glancing angle 4 with the incident X-ray beam can be combined with his constant factor R to obtain our constant factor" G provided that "1—There is no preferred orientation in the sample, and 2—The geometric requirements [for the sample, film, and X-ray beam] have been met precisely."' If A(0) can be combined with R to give G according to the equation G = R . A(0) [14] then the possibility for making non-compensating errors in the austenite determination has been eliminated. This has been described previously in a technical note." Because of the brevity of the previous note,3 it was impossible to emphasize the fact that the two conditions on preferred orientation and geometry can be easily met experimentally, contrary to Aver-bach's statement that: ". . . the necessary conditions must be tested experimentally for each determination, and this is done most easily by observing whether the apparent absorption has the form of Eq. 2 [the theoretical equation for A(0)]."" he way in which these two conditions can be met experimentally will be discussed briefly to help clear up this point. In addition to these two conditions, other factors such as sample shape, homogeneity, and grain size which also affect A will be included in this discussion. A (0) depends on sample shape. The theoretical function for A was derived originally for a flat surface.' In general we have found that a flat sample surface is readily obtainable." If such a surface can * The effect of surfaces which are not flat on the measurement of retained austenite will be discussed later. be obtained, it has the following advantages: 1—it is easily reproducible from sample to sample, 2—it is the form required for metallurgical examination —this type of examination being frequently desirable for this work, and 3—it is an efficient shape for diffraction purposes. Perhaps the most important of these features is that a flat sample surface is easily reproducible; hence, this requirement on the repro-ducibility of A from sample to sample is met for all such samples. Fig. 1—Schematic diagram of the quartz crystal monochromator diffraction unit. The centerline of the main beam is at 6' to the target face. The tangent to the crystal face is at 16.8' to the moin beam. The centerline of the monochromatic beam is at 33.6 to the main beam. The sample surface can be rotated in its own plane. The angle of the sample surface and the monochromatic beam can be adjusted by rotating about the vertical axis, The film holder can also be rotated about B so that the film can be exposed over the desired angular range. For Fe K, X = 1.932A. 1011 planes of quartz have d = 3.35A. A(0) depends on homogeneity and grain size. If the sample is badly segregated or the grain size is large, the effect of micro-absorption and primary extinction1. ' must be considered. It has been shown, however, that for most plain carbon or low alloy hardened steels, neither micro-absorption nor primary extinction effects are present.' A (0) depends on camera geometry. For a given angle 0, A remains theoretically constant from a geometrical viewpoint only if the following factors are kept constant: 1—the angle of inclination .+ of the sample to the X-ray beam, and 2—the centering of the sample with respect to the film cylinder. These are mechanical problems which can be solved readily if the facilities of a good machine shop are available. The arrangement used to insure that the angle 4 remains constant and the centering of the sample is reproducible is indicated schematically in Fig. 1. The sample is clamped against a thin plate with a hole in it by means of a spring-loaded pres-

Jan 1, 1954

By C. G. Dunn, P. K. Koh

Microstructure, magnetic torque, and texture data before and after grain growth were obtained on two 3.25 pet Si-Fe specimens having initially the same cold-rolled textures and the same primary recrystallization textures. The latter textures consisted of four components, two only of which were strong; the strong components were near (120) [001] and (320) [001]. The textures of the two specimens obtained by grain growth were different. In one specimen the two weak components were converted into strong components by a secondary recrystallization process; in the other specimen the primary recrystallization texture was retained by a normal grain-growth process. The difference in behavior of the two specimens was interpreted in terms of a difference in the relative grain sizes between weak and strong components; i.e., in terms of a geometrical factor that would alter the early growth rates of grains belonging to the two weak components of the texture. Specifically the growth rate was considered as the product of two terms: a driving force and a boundary mobility. The texture changes observed are considered to favor the oriented-nucleation growth-selectivity theory more than the oriented-growth theory. WHEN primary recrystallization and grain-growth textures form as a result of annealing deformed metals or alloys, the fundamental processes involved are nucleation and growth. Knowledge, therefore, about nuclei and the factors influencing their growth is required not only for the understanding of texture development but also for its control. In the oriented-growth theory of texture formation1, 2, 3 the problem of nuclei is generally disposed of by the assumption that nuclei are available in all orientations; the theory then has to explain the final texture from the initial texture and the way boundary mobility depends on orientation. In the oriented nucleation theory of texture formation"' nuclei are considered to be present only in certain orientations but the importance of growth selectivity is still recognized. For example, the mobility of boundaries of approximately the same orientation and of the coherent type in twins is believed to be low while that of high-angle boundaries, in general, is high. Because of this the theory may be described as an oriented-nucleation growth-selectivity theory. The main feature of the oriented-nucleation theory, however, is the importance given to nuclei; i.e., how they form in specific orientations and how they grow. In the present investigation two cold-rolled single crystals were used to study the effects of growth after primary recrystallization. The cold-rolled textures and the primary recrystallization textures were determined from portions of these specimens and reported earlier.",' The two cold-rolled crystals were in the (111) [112] stable end orientation, the recrystallization textures were nearly identical and consisted of two strong components, designated M and M', which were near (120) [00l] and (320) [00l], respectively, and also some specific weak components near (111) [110] and (111) [110]. It is well recognized that a single strong component in the primary recrystallization texture may be needed for secondary recrystallization3, 5, 8 but unfortunately when this is the case the amount of material left in deviating orientation may be too small either for growth or for positive identification. The present samples proved to be almost ideal for the study of minor components and their influence on texture changes produced by growth after primary recrystallization. The results obtained are interpreted in terms of present knowledge of grain-growth processes. Experimental Procedure Two lots of silicon-iron alloy of essentially the same composition, namely, 3.25 pet Si, 0.004 C, 0.009 P, 0.010 S, 0.035 Mn, 0.070 Ni, 0.090 Cu, 0.009 Sn, with traces of A1 and Cr and the balance Fe, were converted into single crystals 0.025 in. thick and 1.25 in. wide. The orientations were predetermined by the method of reorienting a seed crystal as described elsewhere.3 pecimen 1 from one lot had the (335) [556] orientation while specimen 2 from the second lot had the (111) [112] orientation.* Both crystals were cold rolled to a reduction in thickness of 70 pet while widening approximately 20 pet. Samples of each were selected for determining 1) the cold-rolled texture, 2) the time required at

Jan 1, 1959

By F. F. Schmidt, H. R. Ogden, E. S. Bartlett

Continuing tantalum alloy development studies have been concerned with a more detailed investigation of promising binary, ternary, and more complex tantalum alloys containing Groups IV-A, V-A, VI -A. VII-A, and VII-A metal additions. A primary objective of these studies was to improve the elevated temperature strength of tantalum with minimum degradation of fabricability and low-temperature ductility by alloying. Several alloys were found to exhibit hot tensile strengths in the excess of 25, 15, and 10 ksi at 2700°, 3000°, and 3500°F, respectively, combined with excellent cryogenic toughness and ductility. Particularly outstanding in this respect are alloys in the system Ta-Mo-W. Ternary parameter curves describing hot-strength, low-temperature ductility, recrystallization temperature, weld ductility, and fabricability are presented for the Ta-Cb-V, Ta-Hf-W, Ta-Mo-Hf, Ta-Mo-V, Ta-Mo-W, and Ta-V-W systems. At present, only four metals qualify for large scale structural applications at temperatures above about 1800°F. These are tungsten, tantalum, molybdenum, and columbium. Research and development work on these refractory metals has progressed to the point where efforts are now primarily directed toward the improvement of high- and low-temperature mechanical properties through alloying. High melting point combined with excellent low-temperature ductility, fabricability, toughness, weldability, and high solubility and tolerance for both substitutional and interstitial solutes are important attributes of tantalum. Conversely, its high density and relative scarcity compared with other refractory metals are the major deterrents to the use of tantalum. Early studies in the Battelle alloy development program established base line data on the mechanical and metallurgical properties of high-purity tantalum.' These studies were extended to include a preliminary investigation of the effects of various alloying elements on the established base properties.2,3 Recent studies involved a more detailed investigation of promising binary alloys containing Cb, Hf, Mo, Os, Re, Ru, and W, and ternary Ta-Cb-V, Ta-Hf-W, Ta-Mo-Hf, Ta-Mo-V, Ta-Mo-W, and Ta-V-W alloys. Results of major importance obtained are given in this article. EXPERIMENTAL PROCEDURES Electron-beam-melted tantalum (Wah Chang) and high-purity alloying elements were consolidated as 150-g ingots by multiple nonconsumable arc melting under a partial atmosphere of helium. Homogeneity of the button ingots was checked radiographically. Alloy buttons measured about 1-1/2 to 1-3/4 in. in diam by about 0.35 and 0.50 in. thick. Nominal alloy compositions were checked on each alloy button by weighing after arc melting. In most instances, weight losses were within 3 pct of the total weight charged. Subsequent chemical analyses on a number of selected alloys showed interstitial contents to range as shown below: Range of Interstitial Content, ppm 0 C N H 25-75 50-100 10-50 1-5 Metallic alloying additions were found to agree well with nominal compositions based on weight change data. Fabrication directly from the cast button to sheet was conducted at low (75° to 900°F), intermediate (1800° to 2000°F), or high (3000°F) temperatures, depending upon alloy content and hardness. Evacuated stainless steel or molybdenum packs were used for fabrication at intermediate and high temperatures, respectively. The temperature required to complete recrystal-lization in one hour was determined by hardness measurements and metallographic observations (longitudinal test sections) after vacuum annealing at 180°F intervals in the range 2010° to 2910°F. All mechanical property evaluations were conducted on recrystallized sheet, nominally 0.04 in. thick. Weldability of the various alloys was screened using manual TIG welding techniques. Bend ductility (bend axis perpendicular to weld axis) was the criterion used for evaluation. Tensile specimens employed 0.200 by 1 in. gage lengths, and bend test specimens measured about 1 by 1/4 in. Low-temperature tensile tests were conducted using a crosshead speed of 0.02 in. per min up to the point of yielding, and 0.05 in. per min to fracture. High-temperature tensile tests were conducted in vacuum using a crosshead speed of 0.01 in. per min

Jan 1, 1963

By Margaret V. Doyle, R. W. Powers

INTERNAL friction measurements, carried out as functions of temperature, have been used extensively to obtain data on the mobility of interstitial impurities in the Group V metals, vanadium, colum-bium, and tantalum. For a given frequency of vibration, the internal friction peaks at a temperature which depends principally on the particular metal and interstitial impurity. Relaxation times and corresponding activation energies are determined from a knowledge of such peak temperatures for various vibration frequencies. These experimentally determined relaxation times are related to the average jump frequencies of the interstitial impurity atoms using the model of Snoek1 and can be related to macroscopic diffusion coefficients as shown by Polder.2 In these Group V elements, the impurity whose diffusion rate has been the hardest to pin down by the internal friction technique has been carbon. The experimental difficulties associated with carbon in these metals were not evident in the earliest investigation carried out by Ke on tantalum.3 The carbon peak was thought to occur near 140°C at 1 cps on the basis of the enhancement of the height of this peak measured on specimens which had been coated with colloidal carbon and heated in vacuo at 1200°C. Presumably on the basis of similar evidence, Wert identified a peak also near 140°C at 1 cps in columbium with the diffusion of carbon in that element." Later, the difficulties of admitting carbon into these metals without simultaneously introducing oxygen were pointed out by Ang and Wert- nd by Marx, Baker, and Sivertsen.6 The latter suggested that the peaks attributed to carbon in tantalum and columbium might arise from the presence of traces of oxygen, adventitiously absorbed during the carburization. Such suspicion was confirmed, at least in tantalum, when it was found possible to prepare C-Ta specimens without contamination by oxygen. Specimens sci prepared, aitnough high in carbon content, gave no evidence of a damping peak near 140°C.7 Recently, the authors were able to Show that carbon diffuses in tantalum at a rate only slightly slower than that for nitrogen.8 At 1 cps, the carbon peak lies at 34g °C, only a few degrees above the nitrogen peak at 341 °C. However, the two could be distinguished by a different stability with respect to aging at 400 °C and by a different dependence of peak height on solute concentration. On the basis of this experience with tantalum, it appeared worthwhile to search also for a carbon peak in columbium in the vicinity of the nitrogen peak." In this report, the location of the carbon peak in columbium is made known and its position relative to the nitrogen peak is compared with the authors' previous findings in tantalum. In addition, a second source of internal friction seen only in specimens of very high purity is discussed. Columbium was obtained from the Fansteel Metallurgical Corp. as either 30 or 40 mil diam stock. The manufacturer gives the following analysis as typical of this material: 99.4 pct minimum Cb, 0.10 pct maximum C, 0.5 pct Ta, 0.04 pct Ti, 0.02 pct maximum Si, and 0.01 pct maximum Fe. Wires 13 in. long were degassed near 2200°C for 4 to 16 hr in an ultimate vacuum better than 2.10" mm Hg in an apparatus described previously.8 This treatment reduced the oxygen and nitrogen damping peaks observed in the as-received material to less than 0.0001 in Q-1 above background. The latter value varied somewhat from specimen to specimen in the range 0.0004 to 0.0008, depending chiefly on the wire diameter and the frequency of vibration. The high temperature evacuation left the specimen in a rather soft condition, the Vhn being reduced to 40 from a value near 121 in the as-received conGi~iur~. The gram size was about the diameter of the specimen. After outgassing, specimens were loaded with carbon by absorption of such hydracarbons as methane or ethylene. Carbon was also added to a few speclrnens by heating them at

Jan 1, 1958

By H. H. Klepfer, M. E. Snyder

UNTIL very recently, the commercial availability of the rare earths as metals has been very limited. Fabrication of mill products from these metals has not been studied in most cases. This note reports the results of the development of fabrication techniques for thulium. Thulium has a melting point of about 1550°C and a hexagonal close-packed crystal structure. It oxidizes in air to give a black oxide (TmzO,). The procedures for producing thin thulium foil were developed on one ingot weighing about 220 g and were subsequently applied in processing about a pound of metal to foil. Several alternates for the various fabrication steps were investigated and will be discussed. The metal fabricated was in the form of commercial chill-cast ingots 1 in. in diam and 2 in. long weighing approximately 220 g. Impurities in the ingots were reported by the vendor to be 4000 ppni tantalum, 2000 ppm calcium, 200 ppm nickel, 100 to 200 ppm iron, 100 ppm europium, and less than 100 ppm copper, lutetium, and ytterbium. In addition to these impurities, several salt-like inclusions as large as l/8 in. in diam were revealed along the center line of the one ingot sectioned. Preliminary tests indicated that small wafers cut from the as-cast ingot would not fabricate readily by rolling. Forging of copper-jacketed wafers was therefore attempted. At 1550"F forging was satisfactory but an apparent reaction of copper with thulium demanded investigation of lower temperatures. Therefore, the remainder of the test ingot was forged at 1450°F—with only minor cracking. All ingots forged were inserted into copper tubes of 1 in. 1D and 0.125-in. wall thickness. The jackets were sealed by flattening the ends of the tubes and welding under helium. Heating time at 1450°F was 30 min. Press forgings of 1/8 in. per pass were used, followed by 10 min reheats. When the ingots had been squared and reduced to 0.250 in. in thickness, the original copper jacket was stripped off and replaced by a new jacket in preparation for hot rolling. Hot rolling at 1450" F without edge cracking was readily accomplished after forging. Excellent results were obtained with 10 pct reductions of thickness followed by 5 to 10 min reheats. After reduction of thickness from 0.250 to 0.100 in. the copper jacket was removed. It was found, in fact, that hot rolling in air was possible. A tenacious black oxide similar to that seen on zirconium was formed during 3 min reheats at 1450°F. Reduction in air to 0.010 in. foil was possible taking 10 pct reductions per pass. The oxide coat formed during hot rolling in air could best be removed by sand blasting and pickling. Common pickling solutions containing polar solvents were found to attack the metal too rapidly and a concentrated nitric-hydrofluoric acid mixture attacked neither the oxide nor the metal. The most satisfactory pickling solution was 52 vol pct concentrated nitric acid-48 vol pct glacial acetic acid. After forging to 0.250 in., vacuum annealing and cold rolling was found to be another satisfactory alternate to hot rolling in copper jackets. After forging. a hardness of Rockwell B76 was found. Annealing in vacuum (2 X l0-5 mm of Hg) for 1 hr at 1200"F did not alter this value. Annealing at 1470"F for 1 hr brought the hardness down to R;]63. With cold rolling (5 pct per pass) the hardness returned to about R,1:76 after 20 pct reduction and edge cracking became noticeable. However, cold rolling to a total of 40 pct reduction in thickness (RI,,83) was possible before edge crack propagation became serious. Good surface finish was obtained, and the metal loss due to oxidation was minimized by cold rolling and vacuum annealing. Using this procedure the yield of 1.25 in. wide by 0.010 in. thick foil from a 1-in. diam ingot was about 40 pct.

Jan 1, 1961

By L. Stone, H. Margolin

The Ti-C-N and Ti-C-O systems were investigated in the temperature range from 500° to 1400°C and in the composition range up to 2 pct C and 5 pct N or 0. Characteristic isothermal sections at 800°, 900°, 1000°, and 1300°C are presented. The Ti-N-0 system was studied in the temperature range from 900' to 1400°C with alloys containing up to 6 pct total alloying content. Characteristic isothermal sections at 1000° and 140O°C are presented. Melting-point data for all three systems are also included. THIS paper reports on one of a series of investi-gations which have been conducted on the phase diagrams resulting from interstitial alloying with iodide titanium. The other investigations involved delineation of the binary systems with carbon,' nitrogen and boron,h and oxygen. The Ti-0 binary system has also been investigated by Bumps et al.' In varying degrees, each of these interstitial elements has been shown to stabilize the low temperature a modification of titanium1-5 and each forms a face-centered cubic TiX compound (henceforth designated 6). In addition, the Ti-N and Ti-0 systems reveal a low temperature tetragonal phase (6) formed by a peritectoid reaction between a and TiX Experimental Procedure The development of experimental techniques for the study of titanium alloy systems has, to a large extent, become standardized. In this investigation, the equipment and procedures described in detail by Cadoff and Nielsenl have been used. Arc Melting: In general, binary alloys with carbon, nitrogen, and oxygen, prepared in the composition range of interest in this investigation, show negligible composition changes during arc melting. However, the possibility of the formation of some gaseous combination of alloying elements such as CO, CN, or NO during the preparation of these ternary alloys was considered. Calculations showed that the evolution of only 0.05 gram of such a gas would be detectable as a pressure change in the closed system used during preparation of these alloys. Such pressure changes were not observed. Consequently, nominal compositions have been used in plotting the data. The compositions of the materials used in the preparation of the alloys are shown in Table I. After melting for 3 to 5 min at 275 to 350 amp, the alloys were checked for homogeneity by microstruc-tural examination. Alloys containing up to 1 pct C were homogeneous in the presence of less than 3 pct N or 0. At higher alloying contents, some inhomo-geneities in the carbon distribution became evident. Alteration of the melting procedure toward longer times and higher currents did not improve the homogeneity of these alloys. Ti-N-0 alloys were homogeneous in the range to about 3 or 4 pct total alloying addition. Beyond this, almost all of the specimens showed as-cast microstructures consisting only of the phase. Consequently, inhomogene-ities could not be detected by examination of micro-structures. Ten alloys from each of the systems were analyzed for two of the elements present (oxygen being omitted in all cases and titanium being omitted in the Ti-C-N alloys). In all cases the analyses were found not to be sufficiently precise to serve as criteria for the total composition of these alloys. On the basis of phase distribution in heat-treated alloys, however, it appears that carbon is distributed throughout the alloys most uniformly, with oxygen and nitrogen following in that order. Heat Treatment: Specimens for heat treatment were wrapped in titanium sheet before sealing in the argon-filled quartz capsules. Heat-treatment times varied from 100 hr at 800°C to 0.5 hr at 1400 °C. After heat treatment the specimens were quenched by breaking the capsule in water. With the exception of alloys in the low composition region, heat treatment did not have an appreciable effect on the as-cast microstructures. Metallography: Following heat treatment, the specimens were prepared for metallographic examination by grinding on emery paper and electrolytic polishing. For the majority of the specimens a 10-sec etch with Remington "A" agent (25 pct HNO3, 25 pct Hf, and 50 pct glycerin) adequately

Jan 1, 1954

By G. Sachs, C. C. Chow, L. J. Klingler

Most structural parts which are heat treated are designed using strength properties which have been determined in the principal direction of the wrought material. For example, for rolled or drawn materials, properties are given for the rolling or drawing direction. The structures, however, may be loaded so that the critical stress is in some direction other than that for which the properties of the material are known. Investigations of forged products,123 have shown that while the yield strength and tensile strength of carbon steel billets and bars vary little with the direction of the test specimen in relation to the fiber, the contraction in area in tension tests and the impact strength in notched bar impact tests decrease in the transverse direction. The contraction in area and, consequently, the fracture stress of hard aluminum and magnesium alloy forg-ings have also been found to be lower in the transverse direction. An investigation on aluminum alloy plate4 likewise has shown the dependence of the fracturing characteristics upon the direction* of the test specimens, the longitudinal direction being considerably stronger than the transverse direction and the normal direction, with the normal direction being the least strong. The variation of properties with direction has been explained by a type of anisotropy called mechanical anisot-ropy. This anisotropy results from the elongation, in the direction of the principal strain, of certain phases, inclusions, and/or cavities in the metal during working. A mechanical fibering is thus produced which seems to persist through annealing and heat treatment. This investigation was initiated to determine the flow stress and fracture stress, at high hardness levels, at 90° to the rolling direction in a round steel bar. It is this direction which receives the critical stress in drawing dies machined from round bars. Preliminary tests showed a large difference in properties between the 0° and 90° directions. Consequently, it was felt that a more complete investigation, utilizing several types of tests, was warranted to determine the flow and fracture characteristics of a steel at various orientations, for a number of hardness levels. This investigation was conducted on an air hardening nondeform-ing die steel. Material and Procedure The distribution of properties was made on a 3-in. round bar of annealed high-carbon, high chromium steel of the following analysis: Pct Carbon................... 1.53 Manganese................ 0.39 Silicon.................... 0.27 Chromium................ 11.76 Vanadium................. 0.25 Molybdenum.............. 0.81 The 3-in. bar was produced from an 8-in. ingot, which was annealed and forged to a 4-in. square billet. The billet was annealed and rolled to a 3-in. round which was then annealed and straightened. This steel is an air hardening die steel which has very good dimensional stability on hardening; therefore, the residual stresses resulting from hardening would be expected to be low. A hardness survey across the diameter of the annealed bar showed no difference in hardness from the center to the outside. However, the test sections of all the specimens were taken approximately half way between the center of the bar and the surface to avoid any surface effect or possible porosity at the center. Tension, compression and bend tests were made on specimens hardened and tempered at six different temperatures. The tension test specimens, Fig 1, were machined from the bar at orientations of 0, 22.5, 45, 67.5 and 90° from the axis of the steel bar. The specimens were rough machined, heat treated and then ground to size. The test section on each specimen was lapped in a direction parallel to the axis of the specimen to remove any transverse scratches which might act as stress raisers. The specimens were tested in fixtures which insured concentricity of loading of less than 0.001 in.5 The transverse strains were measured with a radial strain gauge,5 the least count of which was 0.0001 in. change in diam. The compression specimens, Fig 1, were machined from the bar at ori-

Jan 1, 1950

By R. F. Campbell, R. F. Miller

SIXTEEN steels containing different combinations of chromium and molybdenum, in amounts up to 5 per cent of each element, were examined for microstructure and hardness after air cooling and after furnace cooling from suitable austenitizing temper- [ ] atures. Diagrams summarizing the structures and hardness resulting from these two cooling rates are presented (Figs. 3, 4, 7).† Data were also obtained on the influence of these two elements, individually and in combination, on the depth 0f decarburization observed after furnace cooling (Fig. 8). COMPOSITION OF MATERIAL The chemical composition of the steels investigated is given in Table I. The plain carbon steel was made in a 25-ton Heroult electric-arc furnace, while the alloy steels were made in a 4.5-ton Ajax-Northrup elec- [ ] tric induction furnace. The cast ingots were hot-rolled to I.75-in. bars. AIR-COOLED SERIES The I.75-in. bars were forged to flat plates, 0.25 in. thick, from which specimens 0.5 in. wide by 2 in. long were cut. The specimens were pretreated to improve their homogeneity by heating to I830°F. (1000°C.) for 20 min. and oil quenching. They were then reheated to I740°F. (950°C.) for one hour and air-cooled, after

Jan 1, 1941

By M. E. Wadsworth, P. H. Yu, C. K. Hansen

A study of the rate of dissc5lution of chalcopyrite (CuFeS2) in acidic solutions under oxygen overpressures was carried out by measuring the rate of formation of cupric ions in solution. Effects of temperature, oxygen partial pressure, surface area, and concentration of sulfuric acid were evaluated. A sized batch of chalcopyrite was leached in the temperature range of 125° to 175°C and in the pressure range of 75 to 400 psi of oxygen. In 0.5 N H2SO4 all products of reaction went into solution except for trace amounts of elemental sulfur. The dissolution of chalcopyrite followed linear kinetics and was essentially independent of hydrogen ion concentration for H2S04 concentrations between 0.2 and 0.5 N. The oxygen dependence indicated adsorption approaching limiting values with increasing oxygen pressure. The linear mechanism was explained in terms of steady-state adsorption of oxygen at the chalcopyrite surface followed by a surface reaction. The enthalpy of activation for adsorption of oxygen was found to be approximately 33 kcal per mole. An activation enthalpy of approximately 9 kcal per mole was observed for the surface reaction. Charge transfer reactions are not rate controlling in the process.

Jan 1, 1973

By W. A. Hoyer, G. A. Lock

Laboratory and field tests have shown that it is feasible to log for copper in a qualitative way using the techniques of neutron activation analysis. Laboratory tests showed that the 0.51 and 1.04 mev gamma rays of 62Cu and 66Cu, respectively, could be diagnostic. However, in the core holes, the detection of copper was limited to 0.3% concentration because of interferences from the 18Si and Al gamma rays. Both stationary and continuous logs were made, but only the stationary data yielded a recognizable copper response. This response consists of a combination of two effects-the direct response from copper at 0.51 mev and an increase in low energy gamma ray absorption due to mineralization. Both low energy chemical (PuBe) and high energy accelerator neutron sources were used for the laboratory and field tests, while the same 2 X 5 in. Nal scintillation detector served for all the tests. Potential improvements in copper detection might be possible with solid state gamma-ray detectors which have better gamma ray resolution but much less detection efficiency. This type of activation logging requires complex equipment and trained personnel to conduct field operations and may find application where coring is difficult or prohibitive.

Jan 1, 1973

By Orville R. Lyons

THE Russellton preparation plant of Republic Steel Corp., located north of Pittsburgh, Pa., prepares 3 1/2x0-in. Thick-Freeport coal by means of a 13-ft 6-in. diam Chance cone and 16 No. 7 Deister tables. The plant feed ranges between 370 and 410 tons per hr, of which approximately 31 pct is ¼x0-in. size, and the plant reject ranges between 14 and 16 pct of the feed. The 3½xY4-in. washed coal, the cone product, is dewatered and desanded by a double-deck Parrish-type shaking screen 46 ft long and 7 ft wide. The 1¼x¼-in. portion is further dewatered by two 4x12-ft single-deck vibrating screens. The ¼x0-in. washed coal from the tables is distributed to three 54x70 in. Bird centrifuges. The dewatered centrifuge product is combined with the product of the de-watering screens prior to loading into railroad cars. During the early part of 1951 complaints were received from officials at Republic's coke ovens concerning the moisture content of Russellton washed coal. The records revealed that during the first half of 1951 the shipped washed coal moisture averaged

Jan 1, 1954

By F. R. Hensel

ACCORDING to statements by Guertler,1 Smith and Hamilton were the first to study the copper-titanium alloys, but owing to the presence of large amounts of impurities their data are inconclusive. M. A. Hunter and I. W. Bacon later investigated the electrical conductivity and the temperature coefficients at room temperature. The present writers reported their results2 in November, 1930. Recently Droll published data on copper-titanium alloys3 which agree closely with those obtained by the authors. MATERIAL TESTED The titanium used was in the form of square sintered rods, the analysis of which runs as follows: titanium, 95.4 per cent; silicon, 0.7; iron,1.8; aluminum, 1.4; copper, 0.1; manganese, 0.05. . The ingots were prepared in a high-frequency induction furnace in air atmosphere. Their analyses are given in Table 1. Up to 5 percent titanium the alloys could be forged. Above that amount they were red short and broke up easily. Owing to their rapid absorption of nitrogen and oxygen from the air, the alloys with the higher amounts of titanium tend to form surface films which greatly increase the casting difficulties, especially when small masses of metal are handled. It was, however, possible to prepare sound ingots up to 25 per cent titanium.

Jan 1, 1931

By Carl Henrich

I have read with particular interest that portion of the discussion by Anton Eilers referring to the high-lime (and also high-alumina) slags made by August Raht in 1881, while smelting the Horn Silver ores at Franklin, Utah. The two analyses of such slags, furnished by Mr. Raht to Mr. Eilers are: (a) Not determined in this analysis. (b) Not given in Mr. Eilers' paper. Oxygen ratio, A1 as base. I. 0 in acid: O in base:: 18.1:18.9:: 0.96:1 11. 0 in acid :O in base:: 18.7:17.5::1.06:1 Average: 18.4:18.15=1.01:l Oxygen ratio, AI as acid. I. O in acid: O in base: :23.2:13.8:: 1.68:1 11. O in acid: O in base::23.1:13.5:: 1.76:1 Average: 23.15:13.4 = 1.73:l By assigning to alumina the the of a base, Mr. Eilers deduces the formula 3Feo.-2SiOa + 4Cao.2SiO². The nearest approach to a simple formula expressing the average composition of these slags, which I can figure, is 9RO.Al2O3.6SiO2. Allowing a ratio of 4Feo : 5Cao, we have for this formula:

Jan 1, 1917

By W. V. Vietti, W. A. Oberlin

Crude petroleum produced in the Texas Panhandle oil field is both an asphalt and a paraffin-base oil and is further characterized by being a high-gravity crude with an extremely high cold test,. An eight-well composite of Panhandle crude showed the following characteristics: Gravity, 36.6° A. P. I. Viscosity: 61 sec. Saybolt at 100° F. 170 sec. Saybolt at 75° F. 1560 sec. Saybolt at 50° F. (stirring) Cold Test, 60° F. Distillation Test: Per Cent. Per Cent. Gasoline........................ 28.8 Wax........................... 19.2 Kerosene........................ 4.8 Asphalt........................ 13 .0 Gas oil......................... 30.7 Sulfur1......................... 0.8 Fuel oil......................... 34.7 Subsurface Conditions The original pressure in the formation in the Panhandle field closely approximated 420 Ib. gage or 435 Ib. absolute pressure. Temperatures taken indicate that the formation temperatures do not exceed 90° F. The temperatures in Table I are typical measurements taken on one lease. * A contribrition from the research department of the Marland Refining Co. of Oklahoma and the production engineering department of the Marland Oil Co. of Texas. 1 Smith [Analysis of Panhandle and Rig Lake Crudes. U. S. Bur. Mines, Circular No. 6014 (December, 1926)] found an average sulfur content of 0.6 per cent.

Jan 1, 1928

By A. R. Troiano, E. A. Steigerwald, F. W. Schaller

D. N. Williams(BatteZle Memorial Institute)-The authors have presented an extensive collection of arguments pertaining to the role of stress in hydrogen embrittlement. The basic assumption of these arguments is that stress is essential as a driving force for diffusion of hydrogen. However, in view of the rapid diffusion rate of hydrogen in steel at low temperatures as evidenced by the ability to alloy with hydrogen by electrolytic charging, it seems unlikely to me that stress acts in this manner. Instead, I would like to suggest that stress may be essential for the nucleation of hydrogen in steel. I believe that such an assumption would equally well satisfy the experimental data presented by the authors. I should also like to comment on the interpretation of the data presented in Figs. 5 and 6. Fig. 5 shows the effect of hydrogen content on ductility in a rupture test. The curve as drawn by the authors indicates an abrupt transition from ductile to brittle behavior at about 0.001 amp per sq in. (5 ppm). The data points would equally well support a gradual transition from ductile to brittle behavior extending over the range from 0.0008 to 0.02 amp per sq in. (4 to 8 ppm). A gradual transition would be the ex-

Jan 1, 1962

By R. S. Shoemaker, F. W. McQuiston

EMPEROR GOLD MINES Vatukoula, Fiji 1974 Operating Data to July 1 ORE DESCRIPTION: Ore bodies occur in basalt and associated ash beds overlain by andesite tuffs and sediments. The basalt in much of the ore mined is weathered with alteration of olivine to serpentine and iron oxides, and augite to chorite. The most important host minerals for gold are the tellurides sylvanite and hessite. Next in importance the sulphide s, pyrite, marcasite and arseno-pyrite. A small amount of free gold also occurs. Chalcopyrite, sphalerite, tetrahedrite, bornite, galena and native tellurium are present in small quantities. TYPE OF MINE: Underground and open pit ORE HAULAGE TO MILL: 3.5 ton bottom dump skips from underground. 7 ton trucks from open pit. CAPACITY IN TONS/24 HRS: 1345 short tons ASSAY HEADS (1973): 0.232 oz Au per ton, 0.357 oz Te per ton PRODUCTS: Gold Bullion (400 oz bars) Metallic Tellurium (1 2 lb bars ) RECOVERIES: AU 86.2% Te 88.0% ORE GRINDABILITY: 9.25 Kwh/T TOTAL POWER: 25.7 Kwh/T WATER: Pumped 112 mile to a 318, 000 gallon head tank. Consumption: 2 140 gal per ton, no reclaimed water used. TAILINGS: Pumped 2 /3 mile through 8" pipe. Uphill sections are steel. Downhill sections are fibro asbestos.

Jan 1, 1975

By O. J. Kensok

THE literature contains three previous studies of the chemical activities of Cd-Mg alloys: Trumbore, Wallace, and craigl obtained chemical activities of magnesium at 543" by performing electromotive-force measurements; Borg and irchenall calculated activities at 590°K based on vapor-pressure measurements of cadmium over Mg-Cd alloys by the Knud-sen effusion method; and Terpilouski and arcka measured the chemical activities of magnesium in the liquid state by means of electromotive-force measurements. In this study the vapor pressure of cadmium was determined over nine solid Cd-Mg alloys in the temperature range 571" to 697°K by the Knudsen effusion method. This study was prompted by the lack of precision in the data reported by Borg and Birchenall. Knudsen cells used during this study were fabricated from 0.020-in.-thick tantalum tubing. The tops and bottoms of the cells were cut from 0.010-in.-thick tantalum sheet, press-fitted, and welded in place. The orifice was drilled in the cell top and subsequently reamed to a knife-edge thickness. The diameter of the orifice was measured with a traveling microscope capable of reading 0.0001 cm. The sample was heated by means of a resistance

Jan 1, 1967

By Charles Clayton

A study of the literature shows that the greater part of research work on annealing of cold-worked iron has been for the purpose of studying the effect on grain-size and properties other than hardness. No reference has been found of experimental work of the same nature as that explained in this short paper. The material used throughout was Vismera iron, an iron containing 0.03 carbon and made several years ago by the Inland Steel Co. One-half inch stock was cut into cylinders ¾ in. long. The cylinders were then compressed for 60 sec. in a Reihle testing machine under loads as indicated below: COMPRESSION, AVERAGE LENGTH, SERIES POUNDS INCHES 1 10,000 0.693 2 15,000 .599 3 20,000 .501 4 25,000 .426 5 30,000 .378 6 35,000 .351 7 40,000 .311 Fig. 1 shows these cylinders after compression. Each series con¬sisted of 17 specimens lettered 0, B, C, D, E, E:, G, H, I, J, K, L, M, N, P, Q, and R. One specimen of each series was reserved for study of the properties in the cold-worked condition. The other specimens were annealed in a Hump furnace for ½ hr. at the following temperatures:

Jan 2, 1926