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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 J. L. Jorstad

The hypereutectic Al-Si alloys 390 and A390 have wear characteristics superior to any of the more common aluminum casting alloys. This excellent wear resistance, coupled with good mechanical properties, high hardness, and low coefficient of thermal expansion, has made these alloys candidates for the all-aluminum internal combustion engine, the application for which they were primarily developed. A390 alloy is intended for sand and permanent mold casting, and requires a lower iron ccxltent (0.5 pct maximum) than its die casting companion 390 alloy (0.6 to 1.1 pct Fe). The 17 pct nominal silicon in these alloys provides sufficient quantities of the hard primary silicon phase to assure a high degree of wear resistance, yet little enough of this phase, to minimize the casting and m achining problems associated in the past with the hypereutectic alloys. The mechanical and physical properties of 390 and A390 alloys are competitive with the best of the common aluminum casting alloys with the exception of low ductility. This low ductility is not considered a detriment, since in many applications these alloys replace cast iron. Corrosion resistance is similar to the standard alloys 380 and 333. Numerous parts, of varying degrees of complexity, have been cast of alloys 390 and A390, including engine blocks and heads. The deviations from normal good handling and casting practices that are required for the alloys are those associated with the presence of the primary silicon. To obtain optimum strength and machinability, the molten alloys should be treated with phosphorus to refine the primary silicon phase. Care must be exercised to properly control the pouring temperature and rate in order to minimize primary silicon growth and segregation. AUTOMOTIVE engineers have for years recognized the advantages of lighter weight and better heat transfer that aluminum could offer as material of construction for engine blocks. These advantages alone, however, are insufficient justification for acceptance of aluminum engines by the automotive industry. To be acceptable, an aluminum engine first must pass all tests to which cast-iron engines are subjected, and, in addition, must represent no cost penalty in manufacturing. None of the conventional aluminum casting alloys have sufficient wear resistance to withstand the tests to which cast iron is subjected as a cylinder bore material, and recent attempts to manufacture aluminum engines with cast-in-place iron sleeves could not compete costwise with the equivalent cast-iron engine. The prerequisite to an acceptable aluminum engine seemed to be a casting alloy with a high degree of wear resistance that would eliminate the necessity of cast-in iron sleeves or liners. The hypereutectic A1-Si alloys 390 and A390 have met this challenge.' Combined with modified pistons and a patented cylinder bore surface finish,2 the performance of these alloys in more than 1300 cold start and cold scuff tests, over 3500 hr of dynamometer wear and endurance tests, and more than 350,000 miles of road tests leaves no doubt as to the feasibility of running on bare aluminum cylinder bores. ALLOY DEVELOPMENT The desirable characteristics of the hypereutectic A1-Si alloys have been recognized for some time. Their excellent wear resistance and low coefficient of thermal expansions have been used to advantage in the production of diesel-type pistons in Europe and to a limited extent in this country. The major contributor to the good wear resistance of these alloys is the extremely hard primary silicon phase in the microstructure. Although the percent silicon used experimentally has been much higher, most commercial alloys have ranged from about 14 to 25 pct. Generally, as the silicon content is increased, wear resistance increases, and the coefficient of thermal expansion decreases. With increases in silicon, however, machinability and cast-ability tend to suffer. The development of a satisfactory aluminum engine block material at Reynolds involved the testing and evaluation of many alloys, covering a wide range of silicon contents and numerous variations of iron, copper, magnesium, manganese, and other alloying additions. Castability, mechanical and physical prop-

Jan 1, 1969

By D. W. Lillie

Property changes produced by irradiation of Al-0.4 wt pet Li alloys at 270°C to a burnup of 0.155 pct of all atoms are described. Metallographic evidence of the formation of internal pores and the concentration of gases at grain boundaries is shown. Some conclusions are drawn concerning desirable property characteristics for resistance to radiation swelling. RADIATION induced gases are of great practical significance in nuclear engineering since they affect the behavior of fissionable fuels and control rod materials. In fuels the gases xenon and krypton are the major offenders, producing serious em-brittlement at low irradiation temperatures and gross swelling at high irradiation temperatures. In control rod materials containing boron, the production of helium from the B 10 (n, 0) Li7 reaction produces similar reactions. Since the trend in reactor design is for operation at increasingly high temperatures, the phenomenon of high-temperature swelling by these induced gases has been the subject of considerable recent study. The objectives of such work have been to obtain engineering information and quantitative data on the effects in specific fuels and control rods and to obtain further understanding of the mechanism of gas damage so that more intelligent corrective steps can be made. Analogue systems are often useful for the latter objective and in this case alloys containing Li are suitable in view of the reaction Li6 + n - He4 + He 4. Both products are gases and no serious radiation problem is produced to complicate post irradiation handling. Considerable caution is, of course, necessary in any operations in which tritium release is possible. The present program was initiated to provide a detailed study of the effects of He + H3 generated by irradiation in Li-bearing alloys. The results on Mg-Li solid solution alloys have been reported elsewhere.' This paper presents information on irradiation of an A1-0.4 pct Li alloy in which the lithium was present as highly enriched Li6 in order to maximize the radiation burnup. PROCEDURE The original intent in the selection of the alloy content was to have Li in solid solution at all temperatures. It became apparent, however, that the solubility of Li in A1 was substantially lower than that reported in the original reference consulted.' This was substantiated by the work of Nowak3 who reported solubility of less than 0.01 pct Li at 100°C. The alloy used (0.4 pct Li) was therefore two phase at room temperature. The irradiation temperature was estimated to be about 270° C, and at that temperature all or almost all of the Li would be soluble under equilibrium conditions. The precise state of the alloy under irradiation is not known. Aluminum used was 99.996 pct purity (max. impurities Si 0.002 pct, Mg 0.001 pct, Fe 0.0005 pct, and Cu 0.0005 pct). The lithium was obtained from Oak Ridge National Laboratory; 96.1 + 0.1 pct of the lithium was ~i~, but chemical impurities were 0.25 pct Ca, 0.05 pct Fe, 0.02 pct Na, 0.02 pct Cu, 0.01 pct Sr, and 0.01 pct Ba. The two metals were melted in an alumina crucible in a resistance heated furnace in an atmosphere of helium. A l-in. ingot was cast in an iron mold, machined to %-in. round, hot swaged (350°C) to 0.080-in. round, cold swaged to 0.060-in. round, and finally cold drawn to 0.040-in. round. Four-inch wire lengths were then cut and annealed at 250°C for 4 hr in helium. Irradiations were carried out in a VT hole in the Argonne CP-5 reactor. The wires were held in an aluminum jig, sixteen wires per capsule, and were maintained in a helium atmosphere for maximum heat transfer. Eight wires of pure aluminum and eight of Al-Li alloy were placed in each capsule. Argome Laboratory statements were accepted as establishing flux and exposure levels. These were based on previous experience and flux map-

Jan 1, 1961

By Arnulf Muan, Hobart M. Kraner

Ferromanganese alloys react with aluminu-silica brick in blast furnace hearths and cause the formation of new phases with low refractoriness and consequent failure of the refractory lining. The nature of these reactions is explained on the basis of petrographic observations, thermodynamic data, and Phase equilibrium considerations. Animproued hearth design resulted from these findings. The production of ferromanganese in blast furnaces presents many problems which are not encountered in the production of iron in the same furnaces. One of the most serious of these problems is the rapid failure of refractories caused by their reaction with the ferromanganese metal. The present paper describes such a "hearth attack", its symptoms, causes, and cure. Following an introductory discussion of operating practice in ferromanganese blast furnaces, results are presented of a petrographic examination of samples from a furnace which had to be shut down because of refractory failure. The causes of this failure are then analyzed on the basis of petrographic observations combined with thermodynamic and phase equilibrium data available in the literature. The latter type of data were subsequently used as a guide in recommending improvements in hearth designs which have since proven themselves successful in practical operations. I) FERROMANGANESE BLAST FURNACE PRACTICE It is considered good practice to blow in a furnace on iron and operate it in this way for a period of time preliminary to producing ferrornanganese. During this time, iron generally penetrates the bottom refractories six or more feet below the original working surface of the hearth. Not only does this metal pervade the joints to these depths, but the pores of the refractories are usually also filled by molten iron. The high throughput of the iron furnace provides ample heat to maintain a high temperature in the hearth refractories. This, together with a high ferrostatic pressure, causes porous clay brick of the hearth to shrink. During subsequent use of such a furnace in ferromanganese production, the iron in the pores of hearth brick if; replaced to some extent by ferromanganese. The replacement is usually not complete. This is probably due to the lower metal throughput and lower prevailing temperatures resulting from this in the brick of the hearth bottom. When the ferrornanganese furnace is banked—or when the temperature in the furnace falls due to a furnace delay—-manganese oxidizes even though the atmosphere is largely CO. Considerable expansion accompanies this oxidation of manganese from metal to oxide. The hearth refractories are then subjected to the disintegrating forces of the volume expansion. These forces may in extreme cases be large enough to lift the furnace off its mantle supports as was described in a previous paper.' When cooling takes place slowly, however, and the temperature remains at a fairly high level, the refractories undergo fluxing action by the MnO. This latter reaction and the reduction of SiO2 of the clay brick is the general subject of the present paper. The observations to be described herein were made on a ferromanganese furnace located at the Johnstown plant of Bethlehem Steel Co. The furnace was lined entirely with clay refractories and had operated on iron for 266 days and on ferromanganese for 96 days prior to being taken off because of a hearth wall failure. II) PETROGRAPHIC EXAMINATION AND CHEMICAL ANALYSIS The photograph reproduced in Fig. 1 shws a cross section of the upper part of an 18 by 9 by 4-1/2 in. bottom block removed from the blast furnace. This brick came from a depth of approximately 6 ft below the original working face of the bottom. It will be noted that the 4-1/2 in. dimension is almost the same as when the brick entered the furnace. Three characteristically different zones marked I, II, and III on the print of Fig. 1 are apparent. Results of chemical analyses of the original brick and samples from each of these zones are listed in Table I. Petrographically the various zones were characterized as follows: The central zone marked I was essentially un-reacted brick. The microstructure consisted of fine needles of mullite, and the pores in this entire area were filled with iron. This is characteristic of

Jan 1, 1962

By J. P. Morgan, J. R. Dawson

A review of Jones & Laughlin's Vesta coal preparation plant flowsheet in 1959 indicated that, of the three sizes making up the company's metallurgical product, the -48-mesh fraction offered the greatest potential for improvement. Efforts to improve quality and yield were directed toward the use of froth flotation. After laboratory and plant tests were completed, it was decided that treatment of the entire -@-mesh fraction of the coal mix by flotation was a significant step forward in providing better coal quality and recovery. Reduction of ash in the metallurgical coal from 6.75 to 6.25% provided material savings at the coal washer, coke plant and blast furnace. The Vesta preparation plant of the Jones & Laughlin Steel Corp. is located on the Monongahela River, eight miles south of Brownsville, Pa. This plant is capable of cleaning 2400 tph of a mixture of Pittsburgh seam high volatile coals from the Vesta and Shannopin mines to produce a metallurgical coal of 6.75% ash and a steam coal containing 14% ash. Heavy media vessels are utilized to upgrade the coarse coal while launders and tables handle the -1/4-in. fraction of the raw coal feed. Cleaned -1/4-in. coal goes to boot classifiers where the 48 mesh material settles and is removed by bucket elevators. This 1/4x48 mesh material is then sent to centrifugal dryers for dewatering. As designed, the -48 mesh fraction overflows the boot as a 7 to 8% solids slurry and is recovered by cyclones and a thickener and then dewatered by filtration. The products of the heavy media circuit, the launder-table-boot-dryer circuit and the cyclone-thickener-filter circuit, are combined to constitute the shipped metallurgical coal. The Vesta-Shannopin mix is combined at Jones & Laughlin's Pittsburgh and Aliquippa Works with purchased low and medium volatile coals for the production of metallurgical coke. A review of the preparation plant flowsheet in 1959 indicated that of the three sizes making up the metallurgical product, the -48-mesh fraction, containing 16.7% ash, offered the greatest potential for improvement of coal quality. This fraction was considered low, averaging 55%, while the rejects averaged only 33% ash. It was evident that while the boot overflow fraction represented only 6% of the product mix, a reduction in ash to 6 or 7% would reduce the ash of the metallurgical coal by 0.5%. The original plant flowsheet handling the boot overflow (Fig. 1) consisted of ten 14-in. cyclones treating 70% of the overflow and producing a 48x150-mesh underflow and a -150-mesh overflow. The cyclone overflow was combined with the remaining 30% of the boot overflow to constitute the feed to a 150-ft thickener. The cyclone underflow and the 48x325-mesh thickener underflow were fed to vacuum disk filters for dewatering prior to combination with the coarse metallurgical coal. The thickener overflow stream containing the -325-mesh material was split, 40% being recycled directly to the plant water system and 60% being flocculated and clarified in an 85-ft thickener. The underflow of this thickener was pumped to waste. Efforts to improve quality and yield on the -48-mesh coal were directed toward the use of froth flotation early in the investigation. A review of the fine coal flowsheet indicated that three streams should be considered for treatment by froth flotation. These were the -325-mesh thickener overflow, the -150-mesh cyclone overflow and the -48-mesh boot overflow. Flotation of the thickener overflow was considered basically to improve recovery by upgrading the reject stream so that it could be fed to the filters along with the thickener and cyclone underflows. The overflow from the thickener contained 3.7% solids and had an ash analysis of 33%. Flotation of the cyclone overflow and boot overflow were considered to upgrade coal quality as well as yield. Treatment of the cyclone overflow, a 5.4% solids slurry averaging 27% ash, would eliminate the coal thickener and provide a feed of flotation concentrate and cyclone underflow to the filters. Treatment of the boot overflow would eliminate both the cyclones and thickener and feed only flotation concentrates to the filters. Flotation feed in this

Jan 1, 1962

By Cloyd M. Smith

A few years ago, the Ventilation Committee established the practice of resenting one topic each year for discussion at the annual meeting. The practice has met good response on the part of committee members and I suggest that it be continued. The topic chosen for this year, "Underground Anemometry," is a topic which has bothered me for more than 20 years. It seems to me that the coal industry is content to rely on slipshod methods for measuring the rate of flow of air underground, so I prefaced my discussional charge to committee members with the statement that I regard air measnrements made in the usual way, with hand held anemometer, as no good. Agreements and disagreements came in from more than a dozen engineers, some of whom are with operating companies, coal and metal; some with manufacturers; others with government agencies. The statement was accompanied by a questionnaire on the use of the rotating vane anemometer and by one describing two methods of using a mechanically held anemometer. The questionnaire will be considered first. The questionnaire and statement are as shown on pages 5 and 6, the committee members and respondents are given on page 4, and the general comments of the latter on page 5. Questionnaire 1. Has your company or agency issued written instructions for care and use of anemometers? If so, please enclose a copy with reply. Only one answer, McElroy's, was affirmative. It gave reference to Bureau of Mines publications1'5 which recom- mend the hand held anemometer for rough measurements and indicate that am accuracy of 5 pct can be had if calibration and method factors are used. Mathews said that instructions are principally oral while Maize reported that state inspectors of his department are well trained in use and care of anemometers. 2. Are your anemometers calibrated periodically? If so, by their manufacturer? or by? Are calibralion corrections applied to all observed mean velocity readings? Only one respondent, Lee, answered negatively as to calibration. This means that anemometers are generally calibrated but the questionnaire failed to ask how often this is done. As no one volunteered the information, we have no data on this point. In six cases the instruments are sent to their manufacturers for calibration. but Krickovic reports that his company limits manu-facturers' calibrations to anemometers which are used by operating personnel; those used by the engineering department being calibrated by U. S. Bureau of Standards. The Anaconda Copper Mining Co. has its ventilation engineers calibrate its anemometers. Most of the respondents say that a calibration correction is applied to each mean velocity reading, but Krickovic limits this to surveys made by the engineering department. Since Lee does not calibrate, he has no correction to apply. Maize reports that his department has its anemometers calibrated but does not apply corrections. 3. Do your men hold the anemometer by hand in measuring air flow? for 1 min? or traverse the section? for 1 min? or at? points for 5 sec each? Of 10 replies, 6 were "yes," 3 were "no," and one was "seldom" with respect to holding by hand. Among the six hand holders, four hold in a central position in the measuring section for a minute, except that two of them, Krickovic and Matthews, traverse the section by hand for survey or fan test. Their operating personnel hold by hand, centrally, for routine measurements. McElroy sometimes traverses with hand-held anemometer in rapid survey work. 4. If the anemometer is not held by hand, how is it supported? Augustadt supports the anemometer on an adjustable rod, Condon on "a rod of sufficient length to reach all points with observer standing in one position throughout traverse and at arm's length from plane of traverse." I presume that arm's length must be interpreted liberally enough to allow for arm movement, otherwise it would be impossible to manipulate the anemometer throughout the traverse section. Mancha upholds Condon in this method of traversing. Glanville hangs the anemometer on the end of a 4-ft staff by the hasp at the top of the anemometer frame. McElroy mounts it on the end of a rigid square shaft, 12 in. long, the staff being at right angles to the axis of the instrument. He traverses the section in two halves, holding the anemometer 3 feet from his body.

Jan 1, 1950

By C. Y. Ang, C. Wert

Quench aging of supersaturated solid solutions of nitrogen in columbium takes place in reasonable times in the temperature range 300&apos; to 500°C. Changes in internal friction, hardness, and electrical resistivity are found to accompany this change. THERE is practically no detailed information in the literature concerning interstitial solid solutions in either columbium or tantalum. The fact that both of these metals are used extensively commercially in both pure and alloyed states makes studies of their properties of value. While the plan is to study both these metals and their alloys extensively, so far only a small number of measurements have been made. This paper presents data obtained in the initial investigations of columbium. The principal interstitial impurities of columbium are likely to be hydrogen, carbon, oxygen, and nitrogen. The first of these, hydrogen, is outside the scope of internal friction measurements at present and nothing quantitative is known about the initial hydrogen content of the specimens. However, it is believed that it is completely absent after the preliminary heat treatments are carried out. The carbon content of the original samples is low and can be reduced still further during sample preparation. Oxygen is normally present to some extent in columbium and can also be added or removed easily. Nitrogen is also normally present in small amounts and can also be controlled within limits. Since it is not possible yet to remove all the nitrogen from columbium and since it is possible to remove essentially all the oxygen and carbon, initial measurements were made on the alloy system nitrogen in columbium. In this sense then, the alloys are true binary alloys. Little is known about the effects of nitrogen on the physical and metallurgical properties of columbium, so it was necessary to perform some rather elementary measurements on the specimens. These measurements are described briefly in the next section and are discussed in detail in later sections. Scope of the Work That internal friction is associated with diffusion of nitrogen in columbium has been known for some time. It is similar in nature to that for carbon and nitrogen in a-iron. In these latter cases, the information obtained from internal friction has been useful metallurgically.&apos; and it is not too much to expect useful information from similar studies on columbium containing nitrogen. The technique of internal friction measurements is adequately described by Ke2 At a frequency of 1 cps the internal friction peak for nitrogen in columbium occurs at 300°C and that for oxygen at 160°C." So that all the measurements will be meaningful metallurgically, the internal friction has been calibrated by making a measurement of the damping at the two peaks and a measurement of oxygen and nitrogen content by chemical analysis. A comparison of the two gives the following calibration factor: To go from damping capacity expressed in terms of 1/Q to weight percent of oxygen and nitrogen in solid solution in columbium, multiply the peak values of internal friction by about 4 and 3 respectively." Data from which these factors were calculated are presented in Table I. The increase of electrical resistance of columbium caused by the addition of oxygen and nitrogen as interstitial impurities has been quantitatively determined. At room temperature the percentage increase is small for small additions of impurities (<0.1 wt pct); as the temperature is lowered it becomes an increasingly larger factor of the total resistance. An empirical relation has been established to express the resistivity at — 195°C in terms of weight percent of oxygen and nitrogen in solid solution. It is easy to make supersaturated solid solutions of nitrogen in columbium and to observe aging phenomena. The optimum aging temperature, as in the case for carbon and nitrogen in a-iron, occurs at the temperature for which the time necessary for a single place-change in diffusion is about 1 to 10 millisec. This means for nitrogen in columbium that significant property changes should take place for this alloy at about 400" to 500°C in reasonable time intervals. Results of such aging experiments are

Jan 1, 1954

By R. I. Jaffee, F. C. Holden, H. R. Ogden

The effects of grain size and shape on alloys of titanium with nitrogen and aluminum have been determined. Increasing a grain size decreases strength and hardness and increases impact resistance. Quenching from the ß field produces subgrain markings delineating a plates in Ti-N alloys but not in Ti-AI alloys. This suggests a precipitation from the high nitrogen alloys. ALPHA titanium alloys are single-phase alloys with the hexagonal-close-packed structure of a titanium. If these alloys are worked and annealed within the a field, equiaxed structures result in which strength is derived from the solid solution of the alloying elements present. Strain hardening can be used to increase further the strength of a titanium alloys. This is the only other mode of strengthening the a type of alloy. Quenching from the ß field has been noted by many investigators to have little effect on the strength of iodide titanium and on commercial titanium. Ti-A1 alloys were found to be little different in hardness, whether annealed in the a field or quenched from the ß field.&apos; These facts illustrate a characteristic of a titanium alloys: they are insensitive to heat treatments from the ß field. Underlying this heat-treatment insensitivity are the facts that the a-ß field is generally narrow in a titanium alloys, the transformation-temperature range increases with a-stabilizing content, and the a solubilities of the alloying elements are greater than the ß solubilities. When quenching is done from the ß field, there is no supersaturation of the alloying elements in the acicular transformation structure produced. Welds in a titanium alloys generally are ductile, because no transformation hardening occurs during cooling.1,2 Other desirable features of a alloys are: no thermal-stability problem, excellent toughness, and retention of strength at elevated temperatures better than a-ß or ß alloys of comparable room-temperature strength. The prime factors governing the properties of a titanium alloys are the amount and kind of solute present. This can change the a phase from a soft, tough material into a hard, brittle material with no apparent change in microstructure. Moreover, brit-tleness can exist over a wider alloy-content range, without change of phase, than in most other metals. The solutes for a titanium may be divided into interstitial and substitutional types. Oxygen and nitrogen dissolve interstitially in rather large amounts, extending well through the brittle range, and the transformation-temverature ranges are raised in the process.:&apos; , carbon has a maximum interstitial solubility of about 0.5 pct at the peritectoid temperature, while at lower temperatures the solubility decreases to about 0.2 pct.3,5 Substitutional a solutes include aluminum and tin. Aluminum has a high a solubility of 25 pct Al, and increases the transformation-temperature range markedly.6,7 Tin has a high a solubility of about 20 pct, but has a relatively innocuous effect on the transformation-temperature range.&apos; The choice between interstitial and substitutional solutes generally has been conceded to the substitutional solutes because of a belief that they do not have as adverse an effect on ductility and notch toughness. However, a recent investigation8 on the effect of hydrogen on titanium has shown that this element, which is not under compositional control, has a powerful detrimental effect on notch toughness. Hydrogen was shown to form a hydride which is practically insoluble (about 0.002 pct) in a titanium at room temperature, although the solubility at the eutectoid temperature of about 300°C is relatively high, about 0.16 pct. Limited data were presented on the effect of hydrogen on a alloys, where the same impairment of toughness as in unalloyed titanium is found. The work reported here describes the effects of interstitial and substitutional solutes on 1—the mechanical properties of high purity titanium free of titanium hydride and 2—the influence of structural variables on these properties. Nitrogen was selected as an example of an interstitial solute with high solubility and aluminum as a substitutional solute with high solubility. All compositions were well within the ductile range, since it was not intended to study the brittleness problem here, but to study only the factors influencing optimum mechanical properties. Procedures The alloys were prepared from iodide titanium as 1/2 Ib ingots double melted to insure homogeneity.

Jan 1, 1955

By H. L. Marcus, C. D. Turk

Titanium was physically vapor-deposited by electron beam high rate evaporation. Rates of 50,000 and 85,000Å per min at deposition temperatures of 480°, 535°, and 595°C were used. Deposited thickness ranged front 0.004 to 0.010 in. The evaporated titanium had a high degree of basal Plane texture. The specimens were tested in tension at temperatures ranging from 78° to 520°K. From the data, the yield stress, thermal component of the yield stress, and work hardening characteristics were found as a function of temperature, deposition rate, and deposition temperature. It was found that Taylor&apos;s equation, best fit the work hardening data with m = 1/2. MANY investigations have been made on the tensile deformation characteristics of commercial- and high-purity titanium.1"7 The data obtained from these experiments emphasize the fact that the mechanical properties are generally related to the method of production. The correlation of mechanical properties to production history arises from the dependence of impurity concentration on the extraction and melting technique used. Additional effects are derived from crystallographic texturing which can result from fabrication history. Lee and Backofen8 and other investigators9 have shown that crystallographic texturing in titanium sheet has a significant effect on the deformation characteristics. Bunshah and Juntz10 prepared high-purity titanium by high rate physical evaporation techniques. These authors found the deposited material to have a columnar structure with grain sizes ranging from 12 to 104 µ, depending upon the substrate temperature. The room-temperature strengths of the coarse-grained distillate were similar to arc-melted, rolled. and annealed crystal bar titanium. Commercial-purity titanium was also produced experimentally in our laboratory using high rate physical vapor deposition techniques. The commercially pure material has a finer structure than the high-purity distillate,10 and is highly textured. Since the method of preparation and degree of texturing affect the mechanical properties, a research program was initiated to determine the temperature dependence of the yield stress and work hardening characteristics as a function of high rate deposition parameters. This paper relates the results of the investigation. EXPERIMENTAL A) Material Preparation. The titanium was received in the form of commercial-grade 0.096-in.-diam wire which has a composition of 0.016 pct C, 0.0037 pct N, 0.110 pct 0, 0.011 pct H, 0.05 pct Mn, 0.051 pct Fe. balance Ti. The wire was continuously fed into a vacuum evaporation unit as shown schematically in Fig. 1. The deposition chamber was evacuated, purged with argon, and reevacuated to a pressure of 10-5 Torr which was maintained during deposition. A focused electron beam was used to evaporate the titanium from a water-cooled copper crucible. The titanium vapor was condensed onto a preheated titanium substrate. The temperature of the substrate was controlled within ±5°C for the deposition temperatures of 480°, 535°, and 595°C. Control of the input power enabled the use of two deposition rates. 50,000 and 85,000Å per min. The evaporated titanium had a final analysis similar to that of the input titanium, within experimental limits. B) Testing. The evaporated titanium was physically stripped from the titanium substrate and machined into tensile specimens with 1-in. gage lengths. In addition, tensile specimens of as received annealed 35-A titanium sheet with a 20 µ grain size and 45 pct cold-rolled 35-A were also prepared. All specimens were tested in tension in an Instron at a strain rate of 3.3 x 10"4 in. per in. per sec. Proportional limit has been used as a measure of yield strength. Tests were conducted at 78°, 220°, 300°, and 520°K. Metallography, X-ray diffraction, and electron microscopy were used to correlate structure with the deposition parameters and mechanical properties. Room-temperature X-ray measurements were made to determine the phases present. Pole figures and a scans were made to determine deposition texture. Because the vapor deposition was symmetrical about the normal to the substrate, there was no preferred di-

Jan 1, 1969

By H. H. Podgurski, F. N. Davis

In silver alloys containing less than 0.2 wt pet Cu. the reaction 9 + 1/2 0, = CuO(s) was found to proceed to equilibrium between 700o and 808oC. From measurements of the equilibrium dissociation pressures of the CuO at several temperatures, the differential heat of solution and the activity coeflicients for copper in silver were calculated. These values were found to be in reasonable agreement with those calculated from data appearing in the literature. A metastable "copper oxide" with an atom ratio of oxygen to copper as high as 1.7 was formed by internal oxidation at 300°C of these same dilute Ag-Cu alloys. No anomalous behavior was noted in the temperature dependence of oxygen solubility in silver. The solubility minima between 300" and 800°C reported several years ago can be accounted for, at least in part, by reactions with trace impurities, such as copper. Cold-worked siloer exhibits an enhanced permeability to oxygen. THIS investigation was undertaken because of our interest in interactions between solute and lattice defects in metals. The reason for the choice of the O-Ag system for study is the anomaly in the solubility of oxygen in silver reported by Steacie and Johnson1 in 1926, specifically that isobars between 100 and 800 Torr show solubility minima at 400°C. It was also claimed that the copper impurity in the silver was not responsible for the minima. Recently, Eichenauer and Müller2 proposed that surface adsorption might have been responsible for this solubility anomaly, but no adsorption isotherms are available to check the pressure and temperature dependence observed at the low temperatures. Surface-tension measurements on silver made by Buttner, Funk, and udin3 suggest that a considerable fraction of a monolayer of oxygen exists on silver at 922°C and at an oxygen pressure of 150 Torr. To account for the pressure sensitivity reported by Steacie and Johnson1 at 300°C, the oxygen bound to the surface at 922°C cannot be considered relevant; the existence of a surface layer at 300°C characterized by a lower energy of binding would be required to explain the effect. All of our attempts to detect an interaction of this type with surface sites have failed. In this investigation we have not been able to associate the dissolution of oxygen in silver with a dislocation interaction. Evidently, severely cold-worked silver does not contain a sufficient number of trapping -sites for oxygen. Indeed, our work shows that oxidation of trace impurities in the silver was probably responsible for Steacie and Johnson&apos;s results. In addition, we have been able to establish the nature of the reaction with copper impurity in silver by thermodynamic considerations. EXPERIMENTAL PROCEDURE Measurements of both the internal oxidation rates and the solubility were made volumetric ally in a system equipped with a gas burette, a mercury manometer, a McLeod Gage, and a mass spectrometer. The silver and the silver-alloy samples were protected from contamination by mercury vapor from our pressure gages by suitably placed refrigerated (-78°C) traps. Silica reaction vessels were employed for experiments performed above 500°C. Spectroscopically pure gases were used in this investigation. The highest-purity silver used in the solubility measurements was 99.999 pet. By chemical analysis 2 ppm of Cu were found in this silver. The two Ag-Cu alloys used in this research were made up to 0.14 and 0.15 wt pet Cu starting with 99.999 pet Ag. An unexpected source of error was discovered in the course of the work. At temperatures near 800°C silver distilled from the hot silica reaction tubes into cooler regions of our system. Although the weight loss of silver from the sample was not important in itself, oxygen was being consumed by the silver distilled into the cooler part of the system to form a stable oxide phase from which the oxygen was not recovered. In a separate experiment conducted to determine the necessary correction, losses between 0.2 and 0.3 cu cm (stp) of 0, in 24 hr were measured at 810°C. On the assumption that the flux of silver from the vessel to form Ag2O de-

Jan 1, 1964

By R. J. Stoehr, F. Howell

Since the ore mined in the Grants-Ambrosia Lake uranium area is taken from a water-saturated sandstone formation, part of the milling operation includes a drying process. The authors discuss the merits and disadvantages of two methods: natural drying and mechanical gas-fired drying; operating and cost data are included. In the Grants-Ambrosia Lake uranium area, approximately 7000 tons of ore are mined each day from a water-saturated sandstone formation. This ore is sampled and treated at four separate milling operations. Prior to the design and initial construction of the milling facilities, very little was known as to the physical properties of the ore which would be handled in the sampling and crushing facilities at the mills. All four mills initially were constructed without ore-drying facilities. As mining proceeded, it was determined that, even after drying the ore in place by use of drainage development underground, the moisture content of the ore as hoisted averaged approximately 18 pct. Actually, some of the mines produced ores which resembled cream-of-wheat in character and had a moisture content in excess of 24 pct. Each of the four mills had a differently designed sampling and crushing plant; therefore, the moisture content of the ores had varying adverse effects upon the operation of the plants. It was determined after some experience that most of the sampling and crushing facilities could not economically handle ore if the moisture content was greater than 10 pct—one operator considers 12 pct moisture as a maximum. In all of the plants, moisture content over 8 pct increases costs to some extent. In addition to the added sampling and crushing costs, some of the mills were located 15 to 20 miles from the mines and the high moisture content of the ores increased the haulage costs from the mine to the mill. It was generally conceded that it was necessary to employ some method to reduce the moisture content of the ore prior to sampling. Several methods were considered: 1) natural drying at the mine site, 2) mechanical gas-fired drying at the mill site, 3) mechanical gas-fired drying at the mine site, 4) radiant heating on concrete clabs, and 5) infrared shed drying. After preliminary investigation all possibilities but natural drying and mechanical gas-fired drying at the mill site were eliminated. NATURAL DRYING A cooperative feasibility study was conducted on the natural drying method. This study indicated the following: 1) During summer months May through September, ore piled 4-ft deep would dry from +20 to +10 pct HzO in 30 days with no mechanical turning. 2) In the first four days after piling a reduction of 5 to 7 pct total moisture content was accomplished by drainage and absorption of the water into the underlying ground on which the ore was piled, providing the ground surface absorption rates approached 10 gal per sq ft per day. After the first four days the decrease of moisture caused by drainage was practically nonexistent. 3) In the summer months drying by evaporation amounted to approximately 0.1 pct decrease in total moisture per day and mechanical turning of the ore could increase this as much as three times. 4) Climatic conditions of the area indicated that: a) Air temperatures averaged 64&apos; May, June, July, August, September; 41" October, November, December, March, April; 31" January, February. b) Average rain fall was 10 in. with 40 pct of this occurring during July and August. c) Only on very rare occasions were there over five consecutive days with below freezing temperatures. d) The annual evaporation rate was approximately 75 in. On the basis of this study, two of the mill operations decided to attempt to meet the moisture specifications by natural drying supplemented by mechanical turning. The other two operations decided upon rotary gas-fired drying at the mill site. The natural drying method has been quite successful in practice. The ore running approximately 18 pct moisture is trucked from the head frame and piled in rows about 6 ft wide, 3 1/2 ft deep, 5 feet apart during the months of May through September. In approximately 30 days this ore will

Jan 1, 1961

By O. N. Carlson, H. A. Wilhelm, J. M. Dickinson

On the basis of microscopic studies, melting-point observations, and X-ray analyses, a phase diagram is proposed for the Cb-V system. A complete series of solid solutions is formed with a minimum in the solidus at 1810°C near 35 wt pct Cb. No compounds or intermediate phases were found in the system above 650°C. THERE is an ever increasing need for better structural metals and alloys for use in nuclear reactors. In addition to the normal properties of engineering structural materials, such as high temperature strength, resistance to corrosion, and fabric-abil~ty, the nuclear properties of the material must be considered. In a nuclear reactor it is important to conserve neutrons, so a material which removes these neutrons from the reaction excessively is considered to have unfavorable nuclear properties. In nuclear-reactor design the engineer must have nuclear as well as other data available on alloys in order to make a wise selection of materials. Due to the fact that many of the common structural materials have undesirable nuclear properties, it is vital that new alloys of metals having more favorable nuclear properties be investigated. Columbium and vanadium are both high melting metals, both exhibit resistance to chemical attack, and no great difficulty is encountered in fabricating them into desired shapes. With proper treatment both metals can be cold rolled extensively without failure. In addition they have desirable nuclear properties for certain types of reactors. Therefore, the alloys of columbium and vanadium should be of interest in the atomic energy program. Since an alloy-development program is enhanced by a knowledge of the phase equilibria of the components, this investigation was undertaken to establish the phase diagram for the Cb-V system. According to the Hume-Rothery rules for alloying,&apos; the chemical similarity, crystal structure, and atomic-size factor are favorable for a complete series of solid solutions for this system. Both elements are in the same family of group V of the periodic table and thus are quite similar chemically. The crystal structures of columbium and vanadium are compatible for extensive solid solubility, since both have body-centered-cubic structures. The atomic diameters of columbium and vanadium are 2.85 and 2.62Å, respectively. This difference of slightly more than 8 pct is well within the 15 pct maximum difference allowed for extensive solid solubility. Experimental Procedures Source of Materials: Columbium powder and sheet trimmings were obtained from the Fansteel Metallurgical Corp. According to the manufacturer the metal contains less than 1 pct impurity. An analysis of the metal showed approximately 1800 ppm C in the powder while the sheet trimmings contained less than 500 ppm C. Spectrographic analysis showed minor amounts of Ca, Cr, Fe, Mn, Si, Ti, V, and Zr in both forms of the columbium. No commercial source of vanadium having the ductility and purity desired was available to the authors at the beginning of this investigation. As a result, all of the vanadium used in this study was prepared by the bomb reduction of vanadium pen-toxide with calcium employing the method reported by Long.&apos; Yields of massive vanadium normally were about 80 pct. Chemical analysis of the vanadium prepared in this manner showed the presence of 200 to 500 ppm N and 800 to 1000 ppm C. Minor amounts of Ca, Fe, Mn, Si, Zr, Cr, and Cb were detected by spectrographic analysis. This vanadium metal was ductile and was cold rolled into 5 mil sheet. Annealing was not necessary during this rolling and the metal retained its cold-rolling characteristic after are-melting. Preparation of Alloys: The Cb-V alloys were prepared by melting pieces of vanadium sheet togethel-with columbium in the form of sheet or pellets of powder. The melting was carried out under argon in conventional arc-melting equipment employing a tungsten electrode and a water-cooled copper crucible. Each alloy was remelted three or four times, inverting the alloy after each melting in order to assure complete mixing. Alloys normally were obtained as round flat disks, weighed approximately 70 grams, and had roughly the shape of a disk 1 1/2 in. in diameter and 1/4 in. thick. Half of each alloy

Jan 1, 1955

By F. W. Archibald

The American Smelting and Refining Co. has standardized its copper converting practice to attain a maximum unit blister production with a minimum of refractory consumption by careful location of the tuyeres and by applying magnetite coatings on the hard-burned magnesite brick linings. THE American Smelting and Refining Co. operates four primary copper smelters in the United States with a total of 17 Peirce-Smith type converters; 15 of them are 13 ft in diameter by 30 ft long, and two are smaller. Some details of operations vary with locale; however, fundamentals of design, operation, and maintenance are common to all plants. All converter shells are l-in. thick except for one new converter with riding rings on the ends which has a shell thickness of 1 1/2 in. More tuyeres can be installed with rings on the ends and hence more air can be used. Welded construction is replacing riveted. Minor shell repairs are made after each campaign. Principal causes of complete shell replacement are warpage and cracking resulting from localized over-heating. A 13x30 ft shell is being replaced after 34 years of operation and a total production of approximately 750,000 tons of blister. Converters are driven by 80 hp DC motors through company-designed worm gear reducers. Rotation is 0.38 rpm which permits the skimmer to spot the converter quite accurately for skimming slag. At three of the plants, protection against unscheduled air or power failures is afforded by the installation of emergency drives, consisting of auxiliary air motors or storage batteries. In order to avoid excessive splash out of the converters, the converter mouths are located as far back on the shells as existing flue facilities will permit. At one plant, the back of the mouth is only 13" to the rear of the vertical center line of the converter,. whereas it is 28" at another plant. Newly-lined mouth areas vary from 36 to 44 sq ft with effective operating areas about 25 pct less. It is important to keep the converter mouths as clean as possible. Dirty mouths create back pressures in the converters and as a consequence the tuyere air volumes are reduced. Mouths are generally cleaned by bumping with an empty ladle. Small mouth-cleaning rams have been used but unless extreme care is exercised the brickwork may be damaged. After considerable experimentation, the plants standardized tuyere elevations at 4 to 4% in. below horizontal center line at the shell with a downward pitch of about 13/16 in. per ft of length. Tuyeres are spaced at 6 in. centers except at the riding rings where there are no tuyeres. With this tuyere location, several of the plants now freeze magnetite slag in the bottoms, fronts, and backs up to the bottom of the tuyeres to control the internal shape of the converter. This has the effect of improving the agitation and mixing so that there is a marked increase in converting speed besides affording protection to the brick lining. Currently, the trend is to increase tuyere diameters from 1% to 2 in. to increase the air flow. At present, all converters are hand-punched but one converter is being equipped with a set of mechanical tuyere punchers. Rods for punching are % in. hexagonal smelter bar upset to ll/s in. and rods for cleaning tuyeres are upset to about 1% in., or larger, depending upon the tuyere diameter. Two of the plants use pneumatic reamers for cleaning the tuyeres between charges to minimize disturbing the coating on the inside of the refractory lining. Most of the puncher&apos;s platforms are pneumatically or hydraulically mounted so that a convenient punching position can be maintained regardless of tuyere position. For normal operations, tuyere air pressures vary from 15 to 13.5 psi, although some cycles such as magnetiting require pressures down to 8 psi. Air requirements vary from an average of 25,000 cfm on the newer installations down to 12,000 cfm on the older ones. Converter hoods are designed to protect the punchers from sparks, splash or hood accretions as well as to prevent the escape of objectionable

Jan 1, 1955

By R. E. Miner, J. K. Jackson, T. L. Wilson

A study has been made of the transformation kintetics and morphology in gas-quenched Fe-5 pct Ni alloys. Massive ferrite or nlurtensite was found to form at temperatures in agreement with earlier results. Surface Preparation appeared to have the greatest effect in determining the type of transformation Product formed. Grain size change, subgrain morphology, and dislocation densities of massive ferrite were essen-tinlly invariant functions of cooling rate. The mar-tens~tic structures were typical of transformation in austenite grains of the order of the sample thickness. BY introducing the idea of a massive transformation in dilute iron-base alloys, Gilbert and Owen in 1962 served to stimulate a closer inspection of phase transformations in ferrous systems.&apos; A wealth of literature has followed concerned with the morphology,2-5 transformation mode,2&apos;3&apos;6-8 and thermodynamics of transformations.2,5,9 interesting controversies have resulted from the experimental observations; closer attention has been paid to the definition of massive transformations and still to be resolved is the Ms temperature of pure iron. The purpose of this note is to report some observations made in a detailed kinetic and metallographic examination of the transformations that occur in gas quenched Fe-5 pct Ni. The results are pertinent to the Ms extrapolations that have been made for pure iron and the observed substructure of massively transformed alloys. PROCEDURE The material used in this investigation was cast as a 5 lb vacuum induction melted, carbon deoxidized heat. The ingot was hot worked to a strip 0.10 in. thick, decarburized, and cold rolled to final thickness. Specimens were punched to sample shape and decarburized in wet hydrogen at 600°C. The chemistry of test specimens is indicated in Table I. Samples 1 by $ in., by thickness were resistance heated and rapidly cooled in an apparatus as illustrated in Fig. 1. Similar helium quenching apparatuses have been described previously;2,6&apos;10 the distinguishing feature of this apparatus is the nozzle geometry. Twin nozzles with three convergent-divergent, sonic ports were mounted perpendicular to the specimen for quenching purposes. The specimen temperature was monitored with an intrinsic 3 mil chromel-alumel thermocouple spot welded to the specimen and con- nected to an oscilloscope. A single switch stopped the current flow through the specimen, opened a gas flow solenoid valve, and initiated an oscilloscope sweep. Cooling rates were read directly from a photograph of the oscilloscope trace and the quenching rates quoted later were made at a point immediately prior to transformation on the trace. All quenching studies were conducted either under a vacuum of 10"-5 torr or a helium atmosphere. The sample geometry was chosen large enough to permit optical and electron microscopy to be performed on samples. The samples were electropolished for both types of examination in a perchloric-acetic solution. Electron microscopy was conducted on a HU 11A microscope operated at 100 kv. Dislocation densities were determined at the same general operating reflection (110) on foils thick enough to produce Kikuchi patterns. RESULTS AND DISCUSSION Transformation Temperature Observations and OD-tical Microscopy. The quenching results of this study are illustrated in Fig. -2(a). Samples which were treated in the as-decarburized condition transformed in a manner similar to that reported previously—a decrease in transformation temperature occurred as the quenching rate increased until a plateau region was reached. Within the scatter of the results obtained here, there was no break in the transformation temperature-quenching rate plot. This suggests no change in transformation mode. Further metallographic examination confirmed this observation. A martensitic transformation was obtained for similarly treated samples except that the surface was first electropolished. The Ms temperature was approximately 85°C below the plateau determined for the unpolished samples and cooling rates as slow as 2000°C per sec were adequate for producing marten-site. The transformation temperatures are in gen- / STRIP SAMPLE / COPPER I SAX/-/}&apos; / CONNECTOR-----A f, SON|C poRTS AC LEADS 3j ]j"I rofe[ \ Mf/yinK/iKj&apos;&apos;^ THERMOCOUPLE X-(^—iie"s*c? &apos; LEADS 1 INCH Fig. 1-—Schematic scction of gas quenching apparatus

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.&apos; 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 D. L. McDanels, J. W. Weeton, R. W. Jech

An investigation was conducted to determine the stress-strain behavior and tensile properties of metallic composites and to relate them to the properties of the base materials. Tungsten fibers in a copper matrix were chosen as the model sys-tem to investigate these properties. Room-temp-atzcre tensile and dynamic -modulus tests were used To meet the demand for improved structural materials, considerable effort is being directed toward the creation of composite materials. One such material under investigation is the fiber-reinforced composite, in which a ductile, relatively weak matrix is reinforced with high-strength fibers. Some work has been done on the reinforcement of metallic matrices with metallic fibers1"5 as well as with ceramic-whisker reinforcement. Summaries of much of this work have been published by Mach-lin7 and Baske~.&apos; Work done at the Lewis Research Center of NASA-" has involved the making and testing of composites composed of a fiber, tungsten, and a matrix, copper, which were insoluble in each other. An equation was presented that showed the relation between the ultimate tensile strength of the composites and the strengths and relative volume percents of the components. This relation was verified experimentally for composites containing fibers extending the full length of the specimen as well as for composites reinforced with short-length or discontinuous fibers. Since the publication of the preliminary results,9&apos;10 additional data have been obtained. This paper presents these data as well as analyses of the stress-strain behavior and the mechanics of deformation of uni-axially oriented, fiber-reinforced metallic composites, in which both components are mutually insoluble. EXPERIMENTAL PROCEDURE Composites were made with high-purity copper as the matrix and 3-, 5-, and 7-mil-diam tungsten wire as the reinforcing fiber. Some specimens were made in which the reinforcing fiber extended completely through the full length of the composite test section. The fibers were placed in ceramic to determine these properties. The composites were reinforced with either continuous or discontinuos tungsten fibers. The tensile properties and stress-strain behavior of composites reinforced with either type of reinforcement were similar, and in both cases the full strength of the fiber was utilized. tubes to insure axial orientation and infiltrated with copper at 2200°F for 1 hr. Composites were also made with short-length (3/8 in. long) 5-mil-diam tungsten fibers as discontinuous reinforcement and infiltrated under the same conditions. The reinforcing fibers in both cases were longitudinally oriented in the direction parallel to the tensile axis. Room-temperature tensile tests were conducted on the composites, stress-strain curves were obtained, and ultimate tensile strength and elongation were measured. Dynamic modulus of elasticity, based upon the resonant frequency induced in the flexural mode of vibration, was also determined for the composites and the components. A more detailed discussion of the experimental procedure employed in this investigation may be found in Ref. 11. RESULTS Room-temperature tensile tests were conducted on 3-, 5-, and 7-mil-diam tungsten wires which had been annealed for 1 hr at 2200°F and show the tensile strength of the wires to be 331,000, 327,000, and 290,000 psi, respectively. Since all the composites were infiltrated under these conditions, the strengths were considered to represent the strength of the fibers in the composites throughout this study. The average tensile strength of the annealed copper used as the matrix in this investigation was 27,800 psi. Results of room-temperature tensile tests on composites reinforced with 3-, 5-, or 7-mil-diam tungsten fibers are plotted as functions of composition in Fig. 1. These data represent previously reported datag,l0 as well as data more recently obtained. The line shown on each curve represents a prediction of tensile strength as a function of composition; the calculation of this line was explained in Refs. 9 and 10. A similar plot is presented in Fig. l(d) for composites reinforced with short-length, discontinuous 5-mil-diam tungsten fibers. Again, both the previously reported and the recent data are presented, and the line shown is identical to that shown for

Jan 1, 1965

By J. T. Norton, A. L. Mowry

The monocarbides of the A subgroup elements in the fourth and fifth group of the periodic table in addition to being hard and refractory are of special interest in that they are isomorphous in crystalline structure. They are cubic with a sodium chloride type structure in which the metal atoms are essentially close packed in a face-centered cubic arrangement with the carbon atoms placed in the interstices between. Interstitial structures of this close packed type were first investigated systematically by Egg1 and he gave the rule for their formation, stating that the radius ratio of the nonmetal to the metal atom should not exceed the value of 0.59. The carbides of interest are those of titanium and zirconium of the fourth group and vanadium, columbium and tantalum of the fifth group. Table 1 shows the radius ratio using the Goldschmidt radii for 12 coordination for the metal atoms and the diamond radius for the carbon atom. It will be noted that while there is considerable variation in the size of the metal atom, in all cases the ratio is smaller than the limit of 0.59 placed by Hägg. It has been known for some time that these cubic carbides are soluble in one another, at least to some extent or, in other words, the metal atoms can be replaced, one by another without destroying the stability of the structure. Since the stability of these close packed interstitial substances appears to depend more upon geometry than upon the exact chemical nature of the atoms involved, it is of interest to examine the possibilities of replacement in these carbides in some detail. Hume-Rothery2 has pointed out the importance of the difference in size of solute and solvent atom as a factor in limiting the solubility in simple binary solid solutions. Largely on an empirical basis, he states that if the difference in size between solvent and solute atom is more than 14-15 pct of the solvent atom, the range of solubility is very restricted. The atom size was based on the distance of closest approach in the elements involved. While there is some question as to how one should calculate the size of the metal atom in the carbide structures, reference to Table 1 will show that zirconium is the largest and vanadium the smallest of the group and that the difference is about 15 pct. The Ti-Zr difference is about 9 pct and the others are smaller. Thus one would predict that if the size factor controls the solubility, all of the pairs except VC-ZrC would have wide or complete solubility whereas this latter pair is on the border line and might have restricted solubility. The purpose of the present investi- gation was to examine the solubility of the several pairs of carbides by heating them together until equilibrium was established and then examining the product by X rays. Previous Work Agte3 and his associates prepared various transition metal carbides and determined the melting points of binary mixtures. He concluded from the shapes of the melting point curves that there was extensive solubility in the case of the cubic carbides. Umanskii and his colleagues made an investigation of a number of pairs of the cubic carbides, using X rays and plotted lattice parameter vs. composition curves for the systems TaC-Tic, CbC-Tic, TaC-ZrC and CbC-ZrC. All pairs showed a continuous series of solid solutions. The first two pairs gave a linear relation while the latter two showed a negative deviation from Vegard&apos;s law. Kiefer and Nowotny, in a paper which became available after the present work was well advanced, investigated the binary pairs of the five cubic carbides by means of X rays. Relatively few points were obtained and results indicated that in some cases, at least, equilibrium was not reached at the temperatures used. The results indicated that solubility in the VC-ZrC system was not complete. All of the results of previous investigations indicated the desirability of a more detailed study. Materials The raw materials used were mono-carbides of titanium, zirconium, vanadium, columbium and tantalum and were the purest which could readily be obtained commercially. Spectrographic qualitative analysis showed that the CbC and TaC contained less than 1 pct

Jan 1, 1950

By Appendix by A. S. Goldoni, R. E. Tate, E. M. Cramer

The diffision coefficient for self-diffision of plutonium in the temperature range 350" to 440°C has been measured by using puZ3 as the tracer isotope. Autoradiopaphic techniques were used to inzlestigate the possibility of grain boundary diffusion, hut only evidence for volume diffusion was found. The least-squares fit of the data gives the .following equation for the diffision coefficient: The computer least-squares technique for fitting the nonlinear equations is outlined. DETERMINATION of the self-diffusion coefficient of the fcc 6 phase of plutonium is in principle a straightforward experimental task. However, the chemical reactivity, the intense @ activity, and the toxicity of plutonium put limitations on the experimental techniques. The technique selected included roll bonding for preparation of the diffusion couples and pulse-height analysis of the @-particle activity to determine the distribution of the PU tracer in the diffusion couples. EXPERIMENTAL PROCEDURE Couple Preparation. Cylinders about 0.5 in. in diam were cast from two special stocks of plutonium, one of which had been enriched in puZ3&apos;, as shown in the isotopic analyses listed in Table I. For each roll-bonded composite sheet, a cylinder 0.437 in. in diam and 0.190 in. thick was turned from each kind of plutonium on a lathe in a 98 pct He atmosphere. The two freshly machined cylinders were positioned face to face in a tube of commercially pure aluminum which had been sealed at one end by welding and the assembly was evacuated on a vacuum manifold overnight. The elapsed time between machining the faces of the plutonium cylinders and evacuating the loaded tube was about 15 min. After overnight evacuation of the assembly the indicated vacuum was 1 X 10"5 torr or better. The aluminum tube was then warmed and pinched off with a cold-welding tool. The pinched-off weld was also fusion-welded as an additional precaution against leakage of air into the evacuated assembly. The assembly was immediately heated for 30 min in a 250°C furnace and reduced in thickness by being passed through a rolling mill with rolls heated to 200°C. The rolling schedule (four passes of 100 mils each with a 10-min reheat after two passes) reduced the thickness of the assembly to 100 mils. The rolled assembly was allowed to cool normally in air. The rolled assembly was sheared at the edges and the aluminum peeled from the composite plutonium sheet. The elliptical sheet was about 0.060 in. thick and usually three 0.383-in.-diam disks could be punched from its central portion. The sheet was heated on a hot plate to the ductile low 0 range (140°C as measured by temperature-indicating pellets) before each disk was quickly punched. The quality of the bond in the disks was evaluated by metallographic examination of the scrap sheet adjoining the hole left by the punch. Only well-bonded specimens without oxide in the interface (as indicated by metallography) were considered satisfactory for further use. More than half of the specimens so examined contained sufficient oxide in the interface to be rejected. Diffusion Anneal. Each disk to be diffusion-annealed was wrapped in 1-mil-thick tantalum foil and sealed within a Pyrex capsule evacuated to 1 x 10~5 torr or better. This capsule was then sealed within another Pyrex capsule at a similar pressure. The diffusion anneals were carried out at temperatures between 350" and 440°C in Marshall furnaces adjusted to have a temperature gradient of not more than *1/2"C over a 5-in. length. This gradient was then further smoothed by using a nickel tube as a liner in the furnace. The liner was divided into three longitudinal cavities by a septum of nickel sheet to which two calibrated Chromel-Alumel thermocouples were attached. One thermocouple was used for controlling the furnace temperature by means of a Brown Pyrovane controller equipped with a Capaciline anticipation circuit; the second thermocouple was monitored twice daily with a Leeds and Northrup K-2 precision potentiometer.

Jan 1, 1964

By Richard G. Warren, Michael M. Minor, Gayle J. Thomas

INTRODUCTION Los Alamos Scientific Laboratory (LASL) initiated its effort in the Hydrogeochemical and Stream Sediment Reconnaissance (HSSR), a part of the Department of Energy (DOE) National Uranium Resource Evaluation (NURE) program, in late 1975. Since that time, LASL has completed sampling of the Rocky Mountain states of New Mexico, Colorado, Wyoming, and Montana at a density of about one water or waterborne sediment sample per 10 km2 and has sampled about 85% of Alaska at a density of at least one per 25 km2. Analytical results for these samples are reported by National Topographic Map Series (NTMS) quadrangle. All collection and analytical procedures have been standardized from the outset (Sharp, 1977; Sharp and Aamodt, 1978). Until early 1979, only uranium results were reported for these samples; thereafter, results include analyses for at least 42 additional elements in the sediment samples. Each LASL HSSR report also includes a discussion of the relationship of these analytical data to known or possible uranium resources; more recently, these data have been re- examined for their relationship to resources for other metals (Beyth et al., 1980; 1980a). Analytical results for LASL HSSR sediment samples are both comprehensive, including 43 elements, and precise, allowing close inter-comparison of data between quadrangles. A recent HSSR study shows that elemental concentrations in sediment samples from the Dixon Entrance quadrangle in Alaska are insensitive to the choice of sieve size fraction and do not vary significantly between stream locations a few meters apart, except for extremely high elemental concentrations (Warren, et al., 1980). Elements that are particularly reproducible include thorium, hafnium, and the .ram earth elements. Fortunately, sensitivities for these elements, which are often associated with uranium, are excellent by the nondestructive analytical technique of neutron activation analysis that LASL employs. Wet chemical techniques may not always give reliable results for these elements in sediment samples due to the difficulty of dissolving the resistate mineral phases that normally contain these elements. LASL has recently open-filed results for a large number of NTK3 quadrangles; to date analyses have been reported for about 605 of the quadrangles within the states of the Rocky Mountain region. As a result, uranium analyses am now complete for all sediment samples collected from the state of Wyoming. We have chosen to sumnarize the analytical results for Wyoming and portions of adjoining states hereafter termed "the ~ Wyoming regionn (Fig. 1 ) ; nearly 24 000 uranium analyses and nearly 17 000 analyses for 42 additional elements are available for sediment samples collected from this area. The DOE open-file report numbers are shown in Fig. 1 for reports open filed before September 1, 1980. The Wyoming region provides an ideal area for an examination of regional geochemical patterns ex- hibited by sediment samples. It is endowed with a variety of exposed geologic units and with widely distributed uranium districts associated with host units of several ages. However, discussion will focus on rock units of Eocene and Precambrian ages because they are widely exposed and host the majority of the uranium resources within the region (Fig. 2). Miocene units are also widespread and host mineralization in the Brown&apos;s Park Formation (Fig. 2). Epigenetic uranium mineralization occurs in sand- stones of lower Eocene age such as the Wasatch, Battle Springs, and Wind River Formations whereas vein type mineralization occurs in Precambrian crystalline rocks. Precambrian rocks within the region consist of a variety of granitic and meta- morphic rocks with l .4-2.8 billion year ages, except within the Uinta Arch, where they consist of a sequence of very low grade metamorphosed or unmetamorphosed 1.0 billion year old sedimentary rocks. Precambrian rocks have served as sources for much of the clastic material comprising the lower Eocene units, and may also have provided the source for the uranium in mineralized Eocene or Miocene sandstones (Stuckless, 1979). The remainder of this paper describes the relationship of elemental concentrations in HSSR sediment samples to exposures of Eocene and Precambrian rocks. The results are used primarily to determine where these Precambrian rocks might provide the most suitable uranium source areas and to infer the direction of transport toward adjacent depositional basins during the Eocene.

Jan 1, 1980

By L. W. Holm, G. G. Bernard

Laboratory experiments were conducted to determine the effect of foam on gas flow in porous media. Previous studies have indicated that foam may be applicable as a restrictive agent in influencing underground gas flow. Foam was found to be exceedingly effective in reducing the permeability of porous media to gas. Consolidated and unconsolidated sands with specific permeabilities of 100 to 146,000 md had, in the presence of foam, gas permeabilities that were less than 1 per cent of the specific permeability; in many cases the gas permeability was practically zero. Foam reduced the gas permeability of loose sand to a much greater degree than that of a tight sand. For example, the permeability of a 125,000-md sand was reduced to 3 md while the permeability of a 4,000-md sand was reduced to 7 md. This effect should cause, to some degree, selective plugging of high permeability channels in various oil displacement processes. The presence of oil in a porous medium decreased the effectiveness of foam in reducing gas permeability; apparently oil acts as a foam depressant. However, it was found that certain foaming agents were very effective in reducing permeability even. in the presence of oil. Also, continuous injection of other foaming agents increased their effectiveness, when oil was present. The effect of foam on permeability of porous media to gas was studied as a function of foaming agent concentration and injection rate, absolute permeability, total pressure, pressure gadient, length of porous system, brine concentration and time. INTRODUCTION The use of surface active agents in flood water to increase the recovery of oil has been studied in the laboratory and in the field for decades, with rather limited success. In recent years a new approach to the problem was proposed. Instead of using an aqueous solution of surfactant as an oil recovery agent, Bond and Holbrook2 proposed that the oil recovery agent be a mixture of surfactant solution and gas. In their method, a water-soluble surface active agent with foam-producing characteristics is injected into an underground formation as an aqueous slug. This slug is followed by gas to produce a foam within the rock. Foams are defined as "agglomerations of gas bubbles separated from each other by thin liquid films&apos;! A foam is fundamentally an unstable system. The foam process for oil recovery has since been studied by other investigators. Fried 1 has shown that foam can displace from porous structures oil that normally is unrecoverable by conventional water or gas drives. This superior oil displacing action is believed to be the result of several factors: (1) foam introduces into the reservoir many resilient interfaces of various sizes and curvatures, which increase the probability that a proper combination of forces for oil displacement will be created; (2) foam has appreciable viscosity which improves mobility ratio and contact efficiency; (3) foam accentuates the trapped gas effect because high gas saturations are possible without producing high gas/oil ratios. Deming4 studied the effect of various foam properties on the displacement of liquid. He found that (1) high foaming ability favors high displacement efficiency; (2) high foam stability is not necessary for high displacement efficiency; (3) displacement efficiency decreases with increase in plasticity of foam and (4) displacement efficiency is unaffected by surface tension of foaming agent solution. One of the important aspects of this oil recovery process is the effect of foam on gas permeability. The simultaneous flow of gas and liquid in porous media has been studied by numerous investigators and a large amount of literature exists on the subject. In this previous work, gas and liquid have generally been considered as essentially independent phases, whose flow characteristics are related through the saturation parameter. Foam, however, is a material with properties that are considerably different from those of its components; for example, the viscosity of a foam is greater than either of its components. Gas flow in a porous medium in the presence of foam should be considerably different

Jan 1, 1965