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By O. W. Simmons, L. W. Eastwood, C. M. Craighead

H. Schwartzbart and W. F. Brown, Jr.—The authors have divided the effects of recovery on the true stress-true strain curve into two types; metarecovery, which effects only the first part of the curve or the yield strength, and orthorecovery, which effects the flow stress at any strain. Both of these are said to be true recovery effects, involve no recrystallization, and are explained by the removal of two different types of imperfections caused by work hardening. However, there seems to be some question as to whether the data are sufficiently conclusive to exclude, as an explanation of the authors' results, a mechanism based on the relief of residual stresses between the grains or slip bands and recrystallization. It appears that metarecovery could be interpreted in the same fashion as a customary interpretation of the Bausch-inger effect. The balanced system of internal stresses which exists between grains in a strained specimen due to varying orientation and, hence, yield strengths, of the different grains is responsible for a reduced yield strength in compression following pre-tension, and, similarly, for an elevated yield strength in tension following pre-tension. If the specimen is now heated so that the internal stresses are relieved by creep, then the yield strength in tension following tension will have been reduced and in compression following tension will have been raised. There seems to be a very strong case for the lack of recrystallization in the aluminum investigated by the authors, if one defines recrystallization as the presence of visually detectable new grains or accepts the X-ray evidence as conclusive. One must remember, however, that the appearance of spots on the back-reflection X-ray patterns cannot be taken as the time when recrystallization first started. The areas of recrystallized strain-free material must first have grown to a size large enough to give distinct spots on the patterns and this may take some time. Averbachl7 in an investigation of brass has shown that recrystallization can be detected by extinction measurements at temperatures lower than those based on hardness or X-ray line width determinations. It can be seen from fig. 10 that the rate of recrystallization is extremely low over a considerable time period at the onset of the process. Observations on the rate at which small amounts of recrystallization effect the flow stress would have given further insight as to whether undetectably small amounts of recrystallization might have been responsible for orthorecovery. Also, the question arises as to whether the effects observed in fig. 6 for various times and temperatures could not have been obtained if the time at 212°F were sufficiently long. In addition, the argument that the curve in fig. 10 is not sigmoidal seems weak in view of the scattering of the points. It is conceivable that an accurate determination of the curve for the first 100 hr would exhibit a relationship other than the one drawn. There is one point we would like to raise about the condition of the starting material. The authors annealed their material at 750°F for 15 min to remove the effects of any previous work hardening or machining strains. Reference to the work by Anderson and Mehl shows that this treatment may not have completely recrystallized the aluminum, so that the starting material may have had some strained areas. Higher temperatures or longer times may have been required to remove the effects of any small strains. We would like to mention some results of tests being conducted at the Lewis laboratory of the National Advisory Committee for Aeronautics in an investigation of the Bauschinger effect in relation to fatigue. Tests were performed on annealed electrolytic copper and several annealed brasses. Specimens were pre-strained 1 pct in tension and then tested in compression or tension with and without intermediate stress-relieving annealing treatments at 500°F for various times. Specimens heated at 500°F for 10 1/2 hr showed an elevation of the flow curve in compression and an approximately equal lowering of the flow curve in tension when compared with the curves for the un-heat treated specimens. After approximately 0.8 pct strain, all flow stresses coincided and were equal to the flow stress of the virgin material at this strain. This behavior is consistent with the metarecovery observed for aluminum by the authors and for which a residual stress model can be used. On the other hand, increas-

Jan 1, 1951

By E. C. W. Perryman

The wide grooves formed at the grain boundaries when high purity aluminum is attacked by hydrochloric acid or sodium hydroxide have been attributed by earlier workers to the high energy of the grain boundary material. The effect has been investigated for high-purity AI-Fe alloys with up to 0.055 pct Fe as a function of iron content and heat treatment. It is shown that the explanation given above is untenable, but that the results can be explained on the assumption that iron segregates to the grain boundary in solid solution. IN 1934, Rohrmann¹ showed that aluminum of 99.95 pct purity suffered intercrystalline corrosion when immersed in 10 to 20 pct hydrochloric acid, and that the susceptibility to intercrystalline corrosion depended upon the heat treatment given. The greatest susceptibility was found for specimens quenched from a high temperature (600°C) and the lowest susceptibility for specimens cooled slowly from that temperature. Lacombe and Yannaquis2 have shown that super-pure aluminum (99.9986 pct) annealed at 600°C suffers intercrystalline attack in 10 pct hydrochloric acid and that this attack is intensified by anodic dissolution in the same solution at a current density of 10 milliamperes per sq dm. No difference in extent of intercrystalline attack was found between the 99.993 and 99.986 pct Al, which led the authors to suggest that impurities played only a secondary role in the mechanism of intercrystalline corrosion. It was found, however, that the attack at the grain boundaries depended upon the relative orientation of the grains, large differences in orientation favoring rapid attack. Boundaries where the two neighboring grains were similarly orientated showed high resistance to attack as did boundaries between grains which were in twin relationship. These observations led Lacombe and Yannaquis to suggest that the intercrystalline attack was due to lattice discontinuities present at grain boundaries. Assuming that the grain boundary is a layer three to five atoms thick and has a crystal structure which is a compromise between the two neighboring grains it is clear that the discontinuities will increase with increasing difference in orientation between the neighboring grains and hence the increasing tendency to intercrystalline attack. Roald and Streicher³ investigated the effect of heat treatment of aluminum alloys ranging in purity from 99.2 to 99.998 pct on the corrosion resistance in 20 pct hydrochloric acid and 0.30N sodium hydroxide. They found that in hydrochloric acid the intercrystalline attack appeared to be determined by the type and quantity of impurities present and by the relative orientation of the grains. No difference in the susceptibility to intercrystalline attack was observed between specimens quenched and those furnace cooled, from 575°C. In 0.30N sodium hydroxide some materials exhibited intercrystalline attack, this taking the form of V-notches. Rohrmann¹ offered no explanation for the greater susceptibility to corrosion of material quenched from 600°C. It seems possible that this difference is connected in some way with a different distribution of impurity elements in the quenched and slowly cooled specimens. The fact that Roald and Streicher8 observed no difference between quenched and slowly cooled specimens may possibly be due to differences in either rate of cooling or silicon content or possibly both. Both these would be expected to have an effect on the distribution of impurity elements. Although the rate of cooling used by Rohrmann was slightly more rapid than that used by Roald and Streicher the position cannot be clarified because Rohrmann does not give the silicon content and Roald and Streicher give the silicon contents of only a few of their alloys. That Lacombe and Yannaquis2 found no difference in corrosion behavior attributable to impurities between the two materials they used may be because both were of high purity compared with the aluminum used by Rohrmann.¹ Although they found no difference in the corrosion behavior of their two materials it is possible that the results obtained by Lacombe and Yannaquis may, nevertheless, have been influenced by impurity distribution, since, on the transition lattice theory of grain boundary structure, it would be expected that sparingly soluble impurities would tend to segregate to boundaries where the orientation difference is such that there is a greater density of atomic sites of suitable size to contain them. It was considered worth while, therefore, to examine the corrosion properties of a series of materials of differing impurity content with the objects of confirming the experimental observations made

Jan 1, 1954

By P. Duwez, C. B. Jordan

A. J. SHALER*—I should like to congratulate the authors for having carried out such a precise set of experiments. It has been found useful, in sintering experimental compacts in vacuo, to make certain that the residual gas is not one which reacts with the metal. Since traces of oxygen can be kept away only with great difficulty, the technique is often adopted of using a "getter " of powder in the vicinity of the compacts, and, in addition, of permitting a small hydrogen leak to flow into the vacuum chamber. Did the authors use similar devices? This paper brings up a question concerning the definition of the word ' sintering.' The authors restrict its use to the adhesion between particles. Kuczynski, in a paper presented at this meeting, applies the word to the growth of areas of contact between particles. I have used it to mean both these phenomena and also the dimensional changes which continue to take place after the first two have run their course. May I suggest that we should come to an agreement on the use of these words ? Fig 1 and 2 show an interesting feature: extrapolation of the curves to zero time does not give a densification parameter of zero. The higher the temperature, the higher is the intercept on that axis. These observations agree with the concept of a practically instantaneous densification taking place while the compact is being brought to heat. Such a change may be brought about by plastic deformation and primary creep. The stress pattern causing this first rapid flow is, to my mind, due to the force of attraction between the surfaces of opposite particles in the regions immediately flanking their common areas of contact. The stress is not temperature-sensitive, but at room temperature plastic deformation only proceeds until the metal in the area of contact can support it elastically. As the metal is heated, the elastic limit falls, and further plastic flow occurs. At the higher temperatures, this is followed by primary creep, and finally by the steady-state rate-reaction which the authors are seeking. If they were to recalculate their densification-parameter values, using, not the initial density of the cold compact, but the density after the compacts have been brought to temperature, the systematic deviations from linearity in Fig 3 and 4 might be eliminated. Such initial densities might be obtained by extrapolating the curves of Fig 1 and 2 to zero time. I am naturally pleased to see that such a very well done series of experiments leads to a heat of activation (for the densification process in hydrogen) that is much higher than that for self-diffusion, in confirmation of the less elaborate results reported by Wulff and myself (Ind. and Eng. Chem., (1948) 40, 838). J. T. KEMP*—I would like to comment on Dr. Shaler's remarks. There are apparently different interpretations of the word "sintering." It seems to me that an accurate definition of our word is essential in all metallurgy. May I point out, in this connection, that in practical metallurgy the word "sintering" has been applied to a bonding process in the preparation of ores and flue dust for fur-nacing. It would be unfortunate if in the area of powdered metallurgy we should establish a definition that is essentially different in meaning. F. N. RHINES*—I think that I can answer the question by saying that I see no essential difference between the use of the term "sintering" in extractive metallurgy and in powder metallurgy; physically the same things are going on. I admit sintering is used for different end purposes in the two cases. When we resort to the sintering of lead ore mixture we are doing so to obtain a chemically reactive, loose texture of some rigidity. This is only a difference in use. After all, in powder metallurgy we sometimes deliberately produce a very porous material which has just a little strength, just as in the case of sinter cake. P. DUWEZ (authors' reply)—We agree that it would be helpful to have well-established definitions of such terms as "sintering." Since the question has now been raised, the time might be appropriate for its consideration by some suitable committee of one or more of the metallurgical societies. In answer to Dr. Shaler's first question, no getter nor hydrogen leak was used in our vacuum experiments, except insofar as the guard disks (used to reduce friction between specimens and trays) may have acted as getters. Dr. Shaler's statement that extrapolation of the curves of Fig 1 and 2 does not lead to zero densification at zero time apparently overlooks the logarithmic

Jan 1, 1950

By L. H. Van Vlack, T. Y. Tien, R. J. Martin

The phase relationships of the MnTe-MnS system were studied by DTA procedures. There is an eutectic at 810°C with about 10 mole pct MnS-90 mole pct MnTe. An eutectoid occurs at about 710°C with approximately 7 mole pct MnS where the MnTe(NaCl) solid solution dissociates on cooling to MnTe(NiAs) and MnS. There is very little solid solubility of MnTe in MnS. ALTHOUGH MnS may exist in three different crystal forms,' only the NaC1-type phase is stable.2 Above 1040°C, MnTe also has the cubic NaC1-type structure. Below that temperature, MnTe changes to the NiAs-type structure.3 This phase transition is rapid for both heating and cooling. As a result the high-temperature crystal form of MnTe cannot be retained at room temperature. Because MnO, MnS, and MnSe are all stable with the NaC1-type structure, and MnTe has this structure at high temperatures,4 solid solution formation could be expected among these compounds. It is interesting to note, however, that a complete series of solid solutions exist only in the MnS-MnSe system,' and that the solid solution is quite limited in the MnO-MnS system.' The MnSe-MnTe system possesses a complete series of solid solutions at high temperatures with separation at lower temperatures.7 Although ion size may be critical in the miscibility of MnO-MnS, it is quite possible that the bond type plays a more important role with the miscibility of MnSe-MnTe. This would permit us to speculate that the miscibility gap would be extensive in the MnTe-MnS system. EXPERIMENTAL Preparation. The samples were prepared by mixing and compacting MnTe and MnS powders. The MnS was previously prepared through the sulfur reduction of Mnso4.8 The MnTe had been prepared by mixing and compacting double vacuum distilled metallic manganese and high-purity tellurium in stoichiometric ratio modified with 1 wt pct excess tellurium. The compacted powders were put in a graphite crucible which was sealed in an evacuated vycor tube. The free space in the vycor tube was made minimal to reduce the loss of tellurium. The sealed assembly was then heated slowly to about 500° C where the free manganese and tellurium reacted vigorously, melting the MnTe which formed. Only one phase, MnTe, was detected by X-ray powder patterns and metallographic techniques. Each compact of MnTe-MnS was placed in a graphite crucible and then sealed in an evacuated vycor tube. The samples were heated at 1250°C for 4 hr and furnace-cooled. Microscopic examination revealed no third phase beyond MnS and MnTe. A typical microstructure is presented in Fig. 1. Identification. X-ray powder patterns were obtained using 114.6 mm Debye-Scherrer camera and Fe-Ka radiation. Mixtures of cubic MnS and hexagonal MnTe were observed in all of the compositions prepared. No lattice parameter change was noticed among different compositions, indicating no solid solution could be retained at room temperatures between these two end-members. A lattice parameter of 5.244Å for MnS was obtained by the Nelson and Riley9 extrapolation method using the diffraction lines of (h2 + k2 + 12) equal 12, 16, 20, and 24. The values, a = 4.145Å and c = 6.708Å, for hexagonal MnTe were obtained from the (006) and (220) lines in the back-reflection region. These values agree well with the values reported by Taylor and Kag1e.10 Differential Thermal Analysis. A differential thermal analysis procedure was used to determine phase relationships since the high-temperature equilibrium conditions could not be retained for examination at room temperature, even when the sealed samples (~0.5 g) were quenched in water. The samples were sealed in an evacuated 4 mm vycor tube with a recess in the bottom to accept a thermocouple. An Al2O3 reference was similarly prepared and the two placed within a piece of insulating fire brick to dampen spurious temperature changes within the furnace. The furnace was controlled by a mechanically driven rheostat which increased the temperature at a rate of about 15°C per min. Known phase changes in the Pb-Sn system1' and the a-to-ß quartz inversion12 were used for calibration

Jan 1, 1969

By NIM REPORTS

Notices CONFRRENCE ON COMPUTRRS IN ENGINEERING The above Conference, which is being organized by the Electrical and Electronics Board of the Institution of Engineer!!, Australia, ill to be held in Sydney on 16th and 17th May, 1974. It is being planned in conjunction with the Sixth Australian Computer Conference and will be held during the week immediately preceding the latter. NIM REPORTS Report No. 1495 A method for the determination of ekmental sulphur on the surface of sulphide minerals. Report No. 1500 The chromatographic separation and determination of nobk metals in matte-leach residues. Report No. 1505 The production of the refractory aluminium saicates, with special refererence to South Africa. Report No. 1517 Computer-controlled flotation plants in Canada and Finland. Report No. 1526 Heat of solution as a control of aqueous-ammonia concentration. Report No. 1529 South African costs of equipment for the metallurgical industry (197).

Jan 9, 1973

By D. Gaboury

Most hydrothermal gold deposits contain quartz with fluid inclusions that record composition and temperature of the mineralizing event. Quartz has strong chemical and mechanical resistances to weathering and abrasion. Hence, quartz fragments occurring on weathered surfaces or in till cover have the potential to be used for gold exploration as far as we know what is the signification of the volatile content of the fluid inclusions for gold mineralization. In this perspective, fluid inclusions were analyzed in pure quartz by solid-probe mass spectrometry from gold deposits hosted in the Archean Abitibi (Canada) and Paleoproterozoic Birimian (Western Africa) belts. This approach provides the relative abundance of volatiles from families of fluid inclusions decrepitating at specific temperatures (Td). Fluids in gold deposits are aqueous-dominated with CO2 and minor H2S in accordance with their fundamental role as gold-ligands (H2S) and for pH-buffering (CO2) of the gold-bisulfide solubility. However, there is no correlation between CO2 and H2S contents of fluid inclusions with the gold grade of the samples. This suggests that trapped fluids represent a marginal portion of the fluid history recorded by the veins. However, in barren quartz veins away from known mineralization CO2 or H2S are commonly absent. Wide ranges in H2O Td from 250 to 450°C and CO2-H2O Td offsets are common, indicating phase separation. The content of other volatiles seems to be more dependent on the geological setting. Helium content may suggest proximity to a crustal-rooted fault, a critical parameter for gold mineralization. Methane and especially C2H6 appear to be indicative of fluid buffering by organic-rich material such as graphitic-shale. These rocks may provide an important primary gold reservoir from sedimentary metal-rich pyrite. Nitrogen-bearing fluids may be related to a magmatic component but can also be derived from organic-rich sedimentary rocks. For gold exploration, quartz containing CO2- H2S-bearing fluid inclusions and having H2O-CO2-Td offsets within 250 - 450°C has strong vectoring potential.

Jan 1, 2011

By J. A. Ruppert, P. M. Sullivan

A process was developed on a laboratory scale in which iron contained in galvanizers ' zinc dross and aluminum contained in scrap zinc die - cast alloy were removed simultaneously . Melting both types of scrap together in the proper proportions resulted in formation of solid intermetallic Fe₂ Alg , which was readily removed from the melt by filtration and centrifugation . Recovery was 87 to 92 percent of the available zinc . Refined filtrate contained 0.02 percent iron and 0.20 percent aluminum. Aluminum content was further reduced to less than 0.01 percent by fluxing with zinc chloride . Low temperatures employed and inexpensive equipment requirements make the process commercially attractive .

Jan 1, 1964

By S. K. Palm

The commercial life of a mineral product can be long, but like hitech electronic and software products that soar in sales one day to become obsolete and forgotten a few months or years, mineral products have a life cycle. They are invented, commercialized and eventually decline and disappear from commerce. Diatomite and perlite are produced from mined minerals, but almost none of the products in these industries spring from the earth in finished form. Straight calcined diatomite was invented over 100 years ago and flux calcined diatomite just a few years later. The expansion properties of perlite were discovered and first exploited about 75 years ago. Some applications for diatomite have disappeared or currently face competition from newer products. Dry cleaning was once the second-largest application for diatomite but disappeared over 30 years ago. Both the beer and wine filtration markets for diatomite face increased competition from membrane filters. A pipeline of new, proprietary products is critical to the future of the industry.

Mar 2, 2022

By D Connelly

High purity alumina (HPA) is a processed, premium non-metallurgical alumina product characterised by its purity level ie 99.99 per cent (4N) or 99.999 per cent (5N). The market price, application, and performance of HPA varies widely according to its degree of purity. High purity alumina is a form of aluminium oxide (Al2O3) with growing use in LED lighting and coating on the ceramic cathode separator sheets in lithium ion batteries. This paper describes a study on producing high purity alumina from feldspar to produce HPA and potash for use as a fertiliser. The project was positive for a number of reasons. High potassium extraction (>99 per cent) was achieved in the alkali leach using sodium hydroxide as the leaching agent and full dissolution of the alkali leach residue was achieved using hydrochloric and sulfuric acid. The potential to produce a high purity alumina product via a complex but promising emerging process was assessed. The potential to extract potassium as potassium sulfate to add value to the project was also demonstrated with the test work.

Sep 1, 2024

T. L. Joseph (University of Minnesota, Minneapolis, Minn.)—Mr. Killian is to be commended for his inquiry as to why a decrease of 15.3 pct in coke consumption was accompanied by a decrease of only 1.9 pct in the CO/CO, ratio of the top gas. Although I encouraged him in his efforts to find a more significant and meaningful carbon-gas ratio, the method, calculations, and conclusions originated entirely with the author. Whether we agree entirely with the assumptions underlying his corrected ratio, I think we can all agree that a better ratio is needed, and that he deserves credit for recognizing the need and for doing something about it. To be most useful, the CO/CO, ratio for the top gas should reflect the way in which the carbon is gasified in the furnace, particularly the degree of oxidation and the extent of heat release. The effect of CO, from the decomposition of the limestone and the effect of the reaction of CO and water vapor in the upper part of the furnace on the quality of top gas should be isolated from the effect of more basic reactions such as ore reduction and solution loss. The reduction of sufficient iron oxide to produce a ton of pig iron and the accompanying changes in the proportions of the carbon gases can be expressed on a mol basis as follows: 16.5FeA + 16.5CO - 33FeO + 16.5CO, [1] 33FeO + 33.0CO = 33Fe + 33CO. [2] The CO, from eq 1 normally escapes from the furnace because it does not react with coke at the temperature of its release. However, as much as 50 pct of the CO, from eq 2 may react with coke because it is released at temperatures above 1000°C: 16.5CO. + 16.5CO2 = 33CO [3] Thus we see that the gasification of 16.5 mols or about 200 lb of C in solution loss will regenerate as much CO as is required for the last stage of reduction. The pronounced effect of reaction 3 on the final proportions of CO and CO, in the gas stream is obvious. It also follows that reaction 3 materially alters the heat release but it cannot be entirely eliminated in Lake ore practice. A previous study indicates that 100 lb of solution loss may be necessary to maintain the reducing power of the gas. A reduction of 100 lb from the normal level of 200 lb would mean that over 1,-000,000 additional Btu are released in the process. The 100 lb of additional carbon which will be burned at the tuyeres will release 437,000 Btu and the absorption of 583,000 Btu by eq 3 would be prevented. One of the chief objectives of sized ore burdens and more uniform distribution of the gas stream is to minimize solution loss by completing reduction higher in the furnace and in turn to release the maximum amount of heat from the carbon charged. It was odd, indeed, to find that the use of sized ore and sinter in the careful tests reported by Dobscha had changed the CO/CO2 ratio very little. The author has stressed the close correlation between his corrected CO/CO, ratio and production. A correlation of fuel consumption and the CO/CO, might be more logical. Fuel consumption on the prepared ore burden decreased 15.3 pct whereas the corrected CO/CO2 ratio decreased 19.0 pct. Although this agreement between the reduction in fuel and the reduction in corrected CO/CO2 ratio is not as close as the correlation with tonnage, it indicates the corrected ratio is superior to the normal ratio which decreased less than 4 pct. Some adjustment should be made for the reaction: CO + H2O = CO2 + H2 —700Btu [41 because there is evidence to indicate that about one half of the hydrogen is formed in the upper 21 ft of

Jan 1, 1953

Frother-Depressant Interactions in Two and Three Phase Systems

Sep 13, 2010

By Gilbert Santoro

In this study tantalum carbide filaments of various compositions in the fcc region were prepared by heating a tantalum wire in a measured amount of hydrocarbon vapor. Such properties as tensile strength, transverse-rupture strength, microhard-ness, lattice parameter, electrical resistivity, and magnetic susceptibility were measured as a function of carbon content. Some of these properties showed a maximum or a minimum between a carbon-to-tantalum molar ratio of 0.80 and 0.85, the composition range at which the color changes from metallic gray to gold. The recent literature has shown that in the large homogeneity range for the fcc phase of the Ta-C system there is a marked variation for such properties as microhardness,' electrical resistivity, ' thermodynamic properties,z'3 lattice parameter and superconductivity. In addition to the technological interest of these variations within a single phase, important theoretical benefits may ensue from a continuous investigation of properties vs stoichiometry in tantalum carbide and other refractory defect compounds, for example the effect of the metal-metal distance and crystal structure upon bond formation and electronic structure. Whether there is a similar variation of the tensile strength is of practical importance. Of theoretical interest would be the knowledge of how the magnetic susceptibility varies with carbon composition. Both of these properties, as well as the transverse-rupture strength, were determined for tantalum carbide filaments in the fcc phase region. The lattice parameter, the electrical resistivity, and the microhardness were also investigated and compared with values appearing in the literature. EXPERIMENTAL PROCEDURE Material Preparation. Tantalum carbide specimens of various carbon content were prepared by statically carburizing high-purity tantalum filaments in measured quantities of pure hydrocarbon gas. The filaments, which were 10 mil in diameter and 20 or 40 cm long, were heated by their own resistance at temperatures from 1800" to 2200°C using alternating current in some cases and direct current in others. Tantalum wire from two sources was used. One had a reported purity of 99.96 wt pct Ta. The ingot from which this wire was drawn had been electron beam-melted as the final step in its purification. The reported purity of the other wire, which was prepared by a powder metallurgical process, was 99.88 wt pct Ta. The impurities of both tantalum sources are listed in Table I. Spec-trographic-grade toluene or research-grade propane was used as the source of carbon. The apparatus for the carburization of the filaments is shown schematically in Fig. 1. The system was evacuated to 1 x l0-' mm of Hg when using the electron beam-melted tantalum and to less than 1 p when using the less-pure tantalum. Neither the difference in the base vacuum, nor in the hydrocarbon gas, nor in the tantalum filaments resulted in any detectable difference in the properties studied in this work. However, only the carbides prepared from the higher-purity tantalum were used in the magnetic-susceptibility determinations. Once the base vacuum had been reached, the system was flushed with the hydrocarbon and re-evacuated. For any desired composition of the carbide, hydrocarbon gas was allowed to leak into a constant-volume flask to a predetermined pressure. These pressures were measured by a Cartesian diver-type mercury gage with a range from 0 to 20 mm and with a scale expanded to nine times. The hydrocarbon gas was allowed into the reaction tube and the filament was quickly raised to the reaction temperature. The slack in the filament due to thermal expansion was taken up by drawing out the movable electrode with a screw mechanism. The filaments were kept at temperature from a few minutes to 200 hr. In the majority of cases the heating time was 6 hr. The reaction tube was cooled by cascading water on its outer surface. By means of an ammeter and a voltmeter the electrical resistance at the reaction

Jan 1, 1963

By H. S. Choi, I. Iwasaki

The effect of temperature as a parameter in ore flotation has not been systematically studied, although for some ores it has been known for many years that selectivity and grade of concentrates can be improved by conditioning or flotation at moderately or substantially elevated temperatures. In 1934 Coghill and Clemmer' stated that "elevated temperature seems to make hard water more tolerable" in soap flotation, and a cursory examination of the literature indicates that some earlier investigators were aware of beneficial results obtained by floating above room or mill-water temperatures. Tartaron,2 Mitchell et aL.,3 and particularly Hamilton et al.' have shown that high pulp temperatures give high recovery and selectivity in soap flotation of fluorite from a variety of gangue minerals. As a result of the work of these and of other investigators, high-temperature flotation has been successfully applied to industrial flotation of fluorite. Falconer" stated: "The use of high temperature conditioning (above 35°C and preferably above 60°C) in connection with the soap flotation of nonsulphide ores, as covered in B. Kalinowski's French Patent No. 847,-215, Dec. 7, 1938, is claimed to provide more intense activation, better separation of values from gangue, and reduction in quantity of reagents." The various investigations quoted above clearly demonstrate the improvements to be realized by increasing the temperature at which fluorite is floated, and it is surprising that little or no information has been published on systematic investigation of the same method as applied to other ores. The present article describes comparative results obtained at room temperature and at elevated temperatures by otherwise conventional flotation of iron ores, by single-mineral flotation in the Hallimond tube, and by contact-angle measurements. It is not the purpose of this article to evaluate the commercial feasibility of high-temperature flotation of iron ores, although preliminary calculations show that the cost should not be excessive and that the gains to be realized may well outweigh the extra cost of heating the pulp. LABORATORY FLOTATION TESTS Experimental: All flotation tests reported in this article that use a laboratory Fagergren cell were made on ore A (either containing the primary slimes or deslimed at 20µ as specified) following the preparatory procedures described elsewhere- or this particular ore. Ore A contained hematite, goethite (limonite), quartz, chert, and minor quantities of magnetite. Confirmatory tests made on other ores show that the same general principles apply to ores of similar mineralogy and structure. Conditioning was performed by adding the collector, as a soap, to the vigorously stirred and heated pulp. Conditioning time was 5 min, the pulp contained 60 pct solids, and the equivalent of 0.5 Ib of fatty acid was added per ton of original (unde-slimed) ore. The pH of the pulps ranged from 6.8 to 7.1 (25°C measurement) during conditioning. The pH was adjusted in the flotation cell by addition of either NaOH or H2SO4, the values reported being those of the pulp cooled to 25°C. During conditioning and flotation the pulp temperatures were controlled within l° to 2° by immersion of Variac-controlled Calrod elements in the conditioning beaker and in the cell, and by the use of water, heated to the appropriate temperature, for transferring pulp, adjusting pulp volume, and washing. Selectivity indexes have been calculated on the same basis as those reported in the article by Cooke, Iwasaki, and Choi." Results: Table I gives the results obtained by floating deslimed ore A with pure oleic acid at 25" and 50°C, and Fig. 1 shows the selectivity indexes plotted against pH of flotation. Inspection of Table I

Jan 1, 1961

By J. F. Martin, L. M. Melnick, R. Rapp, R. C. Takacs

Recent studies on the determination of hydrogen in steel have shown that the hot-extraction method for removing hydrogen from a solid sample is preferable to its removal from a molten sample by vacuum fusion or by fusion in vacuum with tin. A number of techniques are available, however, for determining the hydrogen so extracted. They include: thermal conductivity, gas chromatography, pressure measurement before and after catalytic oxidation of the hydrogen to water and removal of the water, and pressure measurement before and after diffusion of the hydrogen through a palladium membrane. These techniques have been evaluated on the basis of initial cost, maintenance, speed and accuracy of analysis, and applicable concentration range. The results of this study showed that the palladium-membrane technique is best suited for routine use. FOR some time investigators have been concerned with the origin, form, and effect of hydrogen in steel. In such stdies', the analysis for hydrogen constitutes one of the most important phases. It is quite apparent that the results for hydrogen concentrations in a given steel are dependent on the method of obtaining the sample, storage of the sample until analysis, preparation of the sample, and analysis of the sample, including all the facets inherent in the calibration and operation of an apparatus for gas analysis. There are a number of means available for determining hydrogen. This is a critical study of some of the more common techniques in use today. In most conventional melting and casting methods, hydrogen concentrations of 4 to 6 parts per million (ppm) in steel are quite common. Because of the undesirable effects of hydrogen on steel there has been increased use of techniques such as vacuum melting,' vacuum casting, and ladle-to-ladle stream degassing, which lower the hydrogen content to levels on the order of 1 to 2 ppm. Therefore, the method used for determining hydrogen in steel must be sensitive and precise. In any analytical procedure for gases in metals there are two distinct operations—the extraction of the gas from the metal and the analysis of the extracted gas. To extract the gas from the steel, three methods have been employed: 1) fusion of the sample with graphite at high temperature; 2) fusion with a flux, such as tin, at a lower temperature; and 3) extraction of the hydrogen from the solid sample at a temperature below the melting point of the steel. Fusion with graphite is the least-acceptable method. The blank in this method is higher and more variable than in either of the other two methods. The hydrogen fraction of the total gas composition usually is between 10 and 50 pct; thus, a larger analytical error is possible. The vacuum-tin fusion4 extraction of hydrogen is probably the most rapid method in use today; the extraction time is usually about 10 min. However, with this system a bake-out of the freshly charged tin for 2 hr is necessary and a change of crucible and a charge of fresh tin are required after each day of operation whether one or thirty samples have been analyzed. In addition, frequent checks of blank rates are required since CO and Na are continually being given up by the steel samples dissolved in the tin bath. The composition of the gas in this method lends itself readily to analysis; although the hydroge concentration may fall to as low as 50 pct, more often it is above 90 pct, thus allowing a more precise analysis (because of less interference from other gases). In 1940 ewell' published the hot-extraction method for extracting hydrogen from the solid sample, comparing analysis for hydrogen extracted at 600°C with similar analysis for the gas extracted at 1700°C by fusion with graphite. Good agreement for hydrogen was obtained between these two methods, provided sufficient time was allowed for extraction at the lower temperature. carsone obtained good results in his comparison of this hot-extraction method with vacuum-tin fusion. Subsequent work by Geller and sun7 and Hill and ohnson' has shown that steel samples should be heated to at least 800°C to effect the release not only of the diffusible hydrogen but also of the "residual" hydrogen that may be present as methane. Since the rate of evolution of hydrogene9l0 depends on such factors as sample size and composition, thermal history, and extent of cold work, a fixed extraction time is not possible. Extraction times of 30 min are normal, but 2 hr are not unusual. Induction or resistance heating may be used in the hot-extraction method. With resistance heating the

Jan 1, 1964

By J. A. Fellows, Earnshaw Cook, Howard S. Avery

HEAT-RESISTANT alloys of the higher nickel and chromium ranges have been empirically developed through the practical experience of the past two decades to a position of significant industrial importance. Few quantitative metallurgical data have been published, however, that could be related directly to engineering specifications and design. Inadequate information led to the inauguration, in 1934, of a research directed toward a better understanding of the properties and performance of these steels at temperatures between 1400° and 2000°F. Investigation was concentrated upon the 26 per cent Cr:12 per cent Ni type, which owes its extensive application to an excellent combination of economy, strength and surface stability at elevated temperatures. The requirements of the oil-refining industry have been prominent in influencing the scope and evolution of the program. Creep strength at the maximum allowable deformation ranges has been employed to evaluate the merits of these alloys. Permissible elongation rates vary from 1.0 per cent in one million hours for turbine blades to 1.o per cent in ten thousand hours for tube supports and heating-furnace equipment. Throughout this paper, the limiting creep stress (L.C.S.) for a given temperature has been defined as the stress that will produce a uniform elongation rate (stage II) of one per cent in ten thousand hours as extrapolated from rates per hour or tests of less than two thousand hours duration. Certain current specifications require a minimum L.C.S. of 1600 lb. per sq. in. at 1800°F., of 3600 lb. per sq. in. at 1400°F., and include a tensile acceptance test at room temperature, after aging for 24 hr. at 1400°F., as a measure of embrittlement in service. An elongation of 4 per cent is considered satisfactory for general use, while a minimum of 9 per cent is prescribed on occasion. Many differences of opinion concerning the properties and performance of these alloys appear to result from their sensitivity to unappreciated variations of chemical analyses which, for example, may change their L.C.S. from 3500 to 350 lb. per sq. in. at 1800°F. Within broad chemical limits, it is possible to balance the six or more important elements present to attain desirable properties with a number of different combinations. One method for adjusting the composition has been described under the sponsorship of the Alloy Casting Institute,1 in which, by maintaining a substantially austenitic* alloy, optimum creep strength and minimum tendency to embrittle after aging are expected. Unless, however, the limits are undesirably dose, it has not been possible to ensure a moderate range of strength and of ductility by any chemical specification alone. There is a definite need

Jan 1, 1942

By Allan T. Gwathmey, J. Bruce Wagner, Kenneth R. Lawless

Between 250°and 550°C in oxygen pressures of 10 to 760 mm Hg, the relative oxide thicknesses formed per unit time on the (100), (111), (110), and (320), decreased in this order. The predominant oxide orientation was (001)Fe3O4 and [110] Fe304 A small amount of a-Fe2O3 which exhibited a fiber pattern was observed in all oxides except those formed at 250°c in low oxygen pressures. IN order to understand the properties and behavior of thin films, particularly their role in oxidation reactions, there is needed more experimental information on their structure, composition, and kinetics of formation on clean surfaces of known geometry. The present study was undertaken to obtain such information for thin oxide films formed on single crystals of a body centered cubic metal, iron. Crystals in the form of spheres were used in order to expose all possible crystal planes to the reactant gases simultaneously. Using this procedure the relative rates of oxidation of all planes could be estimated, and those exhibiting, maximum or minimum oxidation rates could be selected for a more intensive study of composition and expitaxy. Since the three types of information were obtained from the same experiment using a single crystal substrate, any differences between rate, composition, and expitaxy on the several planes may be attributed to the influence of the substrate structure. EXPERIMENTAL The starting material for this study was Armco iron rods, 6 in. long and 3/8 or 1/2 in. in diam. The analysis was reported to be 99.8 pct Fe, with major impurities listed as carbon 0.018 pct, manganese 0.027 pct, phosphorous 0.005 pct, sulfur 0.029 pct, silicon 0.005 pct, and copper 0.11 pct. These rods were cleaned with ether and heated in an atmosphere of moist hydrogen at 950°C for three days. The purpose of this treatment was to decarburize the iron and to provide a consistent, uniform grain size for the growth of single crystals. Crystals in the form of rods, up to 1/2 in. in diam and 2 in. long, were prepared by a method similar to that reported by Leidheiser and Buck.2 Spherical crystals, 3/8 and 1/2 in. in diam and having a handle about 1/4 in. in diam by 112 in. long, were machined from the single-crystal rods. After machining, each crystal was etched in aqueous nitric acid to remove the cold-worked metal at the surface. Before each experiment, the crystal was polished mechanically with metallographic papers through number 4/0, and electropolished in a perchloric acid-acetic anhydride mixture by the method of Jacquet and Rocquet.3 The crystal was washed in running distilled water for at least 5 min. The excess water was blotted with a soft paper tissue, and the crystal was placed in an all-glass reaction vessel (See Fig. 1). The crystal was heated at 550°C in hydrogen for at least 8 hr prior to the oxidation in order to reduce any surface oxides and to help relieve any stresses in the metal. For the structure and composition studies, flat faces were cut parallel to (loo), (111), and (110) planes on the iron crystal spheres. The orientation of these faces was checked by a back-reflection X-ray diffraction technique. The faces were within 2 deg of the desired orientation. The flat faces were treated as described above except that prior to electropolishing they were given an additional polish on metallographic felt saturated with twenty-minute levigated alumina. After the crystal was heated in hydrogen for at least 8 hr, the temperature was adjusted to the desired value, and the hydrogen was evacuated. Commercial oxygen which had been purified by successively passing through tubes containing magnesium perchlorate, Ascarite and magnesium perchlorate was admitted to the desired pressure as measured by a mercury manometer. A resistance-wound furnace surrounded the oxidation chamber, and a temperature controller maintained the temperature within ±3°C. The furnace was provided with a small window for observing the crystal. The changes in thickness of the oxide films were followed by observing the interference colors. The oxidation was allowed to proceed until the desired oxide thickness was attained on a given crystal plane. The experiment was terminated by evacuating the oxidation

Jan 1, 1962

By E. A. Steigerwald, G. L. Hanna

The influence of work-hardening exponent on the variation of fracture toughness with material thickness was studied for high-strength steel, aluminum, and titanium alloys. The results indicate that, when materials are compared at similar fracture toughness to yield strength ratios, the material with the lower work-hardening exponent undergoes the transition from flat to slant fracture at a larger thickness than material with a high work-hardening exponent. In the thickness range where complete slant fracture is obtained the reverse is true and a lower work-hardening exponent results in a lower fracture toughness. The influence of work-hardening exponent on fracture toughness is, therefore, dependent on the particular fracture mode. In the transition region a low work-hardening exponent is beneficial for fracture toughness while in the 100 pct slant region it is detrimental. THROUGH the use of fracture mechanics analyses, the influence of geometric variables on the crack propagation resistance of structures can be interpreted with reasonable consistency. However, in order to gain a more complete understanding of the fracture process, the influence of material parameters on crack propagation must be defined and coupled to the macroscopic fracture mechanics approach. The work-hardening exponent, which characterizes specific material behavior, may serve as an effective parameter to allow some degree of coupling to be accomplished. In the extension of a crack in a specimen of suitable dimensions the propagation process occurs in a stable manner when the crack extension force is balanced by the resistance to crack extension, which exists in the material at the crack tip. As the applied stress, and therefore the crack extension force, on the specimen increases, the resistance also increases primarily because the effective plastic zone at the crack tip, which is the main energy absorption process, becomes larger. Since the work-hardening rate of a material influences the stress-strain relationship, it will also influence the energy absorption process in the plastic enclave at the crack tip and hence should have an effect on crack propagation. A number of studies have been made correlating the strain-hardening exponent or the strain to tensile instability with the ability of a material to resist fracture. Gensamer1 concluded that a low-strain-hardening exponent would result in a steep strain gradient at the base of a notch. He reasoned that a large work-hardening coefficient would result in high-energy ab- sorption due to the increased area under the stress-strain curve. Larson and Nunes2 experimentally observed in Charpy tests on steels heat-treated to below 200,000 psi yield strength that the energy to failure in the fibrous mode, i.e., above the brittle-to-ductile transition temperature, was logarithmically related to the strain-hardening exponent. In order to avoid the complicating effects present in studying materials which undergo a brittle-to-ductile transition, Ripling evaluated the notch sensitivity of a variety of fcc metals with varying work-hardening exponents.3 The results indicated that the relative notch sensitivity, as determined from tests on specimens with a sharp notch, decreased with increasing values of strain hardening. Although the energy required for ductile or fibrous fracture increases with increasing work hardening, high-strength steels often exhibit improved crack propagation resistance when heat-treated to obtain low values of strain hardening.4,5 An analysis of whether low strain hardening is beneficial or detrimental to crack propagation resistance must depend on the particular fracture criterion involved. At temperatures where the material is relatively ductile and the development of a critical strain is required for fracture, high strain hardening increases the energy required to produce failure. In the transition region and below, however, a critical stress law appears to be valid6 and a low rate of work hardening may produce superior resistance to semibrittle crack propagation. The experimental program is aimed at studying these possibilities and determining the specific influence of strain hardening on the crack propagation resistance of several high-strength materials. MATERIALS AND PROCEDURE The alloys, chosen as representative of various classes of high-strength materials, are summarized in Table I. The heat treatments evaluated along with the smooth tensile properties are presented in Table 11. Pin-loaded sheet tensile specimens were employed to determine the smooth tensile properties. A strain gage extensometer (measuring range 0.200 in.) was used at a strain rate of 0.02 in. per in. per min. The work-hardening exponents were determined from the stress-strain curves generated in the smooth tensile tests and the assumption that the portion of the stress-strain curve beyond the yield point can be described by the power relationship: where a is the true stress, P is the true plastic strain,

Jan 1, 1969

By R. M. Hudson, P. E. Perry

Weight-gain data obtained by nitriding low-carbon sheet steel in an amrnonia CNH,) atmosphere indicated that the process obeyed a parabolic rate law. The calculated actization energy for nitriding in the range 964" to 1268°F agreed reasonably well with published data. At 1358"F, rate data indicated that the activation energy decreased. Weight-gain data obtained by uszng mixtures of NH3 -Nz at 1268°F containzng jrom 10 to 100 zol pct NH3 also obeyed a parabolic rate law. The rate of 'nitriding increased with an increase in the NH3 content of the gas Mixture. It is well-known that steel heated in gas mixtures containing ammonia (NH3) takes up much larger quantities of nitrogen than steel heated in nitrogen, both gases having a total pressure of 1 atm;' this phenomenon can presumably be attributed to the catalytic decomposition of NH3 on the steel surface to furnish nascent (monatomic) nitrogen. This process was studied bv Brunauer. Jefferson, Emmett, and Hend-ricks at furnace temperatures of 752" and 831°F2 using mixtures of NH3 in Hz. Englehardt and wagner3 reported that, at a furnace temperature of 914°F and under their experimental conditions, both nitriding and denitriding were controlled by the rate of gas-metal reactions at a steel surface rather than by the rate of diffusion of nitrogen in iron. The present study was undertaken to obtain information on the kinetics of nitriding low-carbon steel strip at higher temperatures so that practical rates for short-time strip-annealing treatments could be estimated. Variables studied included time: temperature, and NH, content in the annealing atmosphere. Mechanical and chemical characteristics of steel nitrided in this manner will not be considered in the present article. MATERIALS AND EXPERIMENTAL WORK The samples used were from a commercial low-carbon steel, 0.0244 cm thick, in the cold-reduced condition. The chemical composition of this steel is given in Table I. Panels were cut to 5.1 by 17.8 cm, degreased in toluene, and weighed just before treatment. Four specimens were nitrided under each of the experimental conditions. A study was made of the nitriding rate of steel in a 100 vol pct ammonia atmosphere, 740 mm pressure, at five specific temperatures within the range 964" to 1358°F. The nitriding rates of steel in ammonia-nitrogen gas mixtures containing 10, 18, 26, 50, and 100 vol pct ammonia, 740 mm total pressure, at 1268°F were also determined. All atmospheres used were dried by successively passing them through drying towers packed with soda lime and with Linde Molecular sieve Type 4A. Quoted gas compositions refer to those entering the furnace. Specimens were held in the constant-temperature zone of a vertical annealing tube furnace for times of 14, 3, 5, 10, or 15 min. Gas flow rates were maintained at 3.8 cu ft per hr, which was nineteen volume changes per hour for the system used. The rate of flow was selected to provide a high level of free NH3 for cracking on the steel surface where the ammonia gas is most effectively used as a nitriding agent. The vertical annealing tube furnace consisted of a Hevi-Duty tube furnace with a 2 1/2-in.-ID mullite ceramic high-temperature tube. The constant-temperature zone (controlled within 10°F) was about 10 in. long. After each specimen was degreased, a hole was punched in one end, for attaching the specimen by hook to a chain so that it could be lowered into or raised from the high-temperature portion of the tube by means of a power-driven winch. A stainless-steel access port with O-ring seals was connected by suitable glass-to-metal seals to the cool upper portion of the furnace tube. After the weighed specimen was placed in the access port, the furnace tube was evacuated to approximately 10"3 torr, and then the system was flushed thoroughly with the atmosphere under study. When the gas flow rate and constant-temperature zone of the furnace were established, the specimen was lowered into the constant-temperature zone. The atmosphere flowed from the top to the bottom of the vertical furnace tube and was then vented. For all these runs, during the first 3 min of the time the specimen was in the constant-temperature zone of the furnace the specimen was heating up to the tempera-

Jan 1, 1970

By H. H. Klepfer, K. J. Gill, P. Chiotti

SOME observations relative to the U-Zn system have been made by other investigators. Chipman1 and Carter2 have reported the preparation of several U-Zn alloys and point out that these alloys are generally difficult to prepare. Chipman1 reported evidence for a high melting compound at about 90 atomic pct Zn and the possible existence of a eutec-tic between the compound and uranium. Raynor," in a theoretical discussion of the alloying properties of uranium, included zinc among the elements predicted to have little or no solubility in a, p, or y uranium. In the present investigation, thermal analyses, X-ray, metallographic, and vapor-pressure data were obtained to determine the phase boundaries. The relatively high zinc pressure over most of the alloys at temperatures of 900 °C and above proved troublesome and special techniques had to be employed in preparing suitable alloys. Materials and Preparation of Alloys The metals employed in this investigation were Ames Laboratory biscuit uranium containing less than 500 ppm total impurities and Bunker Hill slab zinc or Baker Analyzed reagent granulated zinc, both with a purity of 99.99+ pct. Due to the high vapor pressure of zinc and the high reactivity of both uranium and zinc with oxygen at only moderately high temperatures, alloys were prepared in closed containers which had either been evacuated or evacuated and filled with helium. High purity magnesia, magnesia containing 10 pct calcium fluoride, and tantalum proved to be suitable crucible materials. Tw-o different procedures, described below, were used to prepare alloys, the latter being the most satisfactory. The metals, uranium turnings and granulated zinc, were cleaned with dilute nitric acid, rinsed, dried, and placed immediately in a helium-fill'ed dry box. The two metals were placed' in a 10 mil Ta crucible. The charge was enclosed in the tantalum crucible by welding on a preformed tantalum cap. This assembly was enclosed inside a stainless steel (AISI 309) bomb. The bomb was made by welding a piece of stainless steel plate on each end of a stainless steel pipe. All these operations were carried out in a helium atmosphere. These assemblies were heated in a muffle furnace at temperatures between 1100" and 1200°C for 10 to 15 min or held as long as 15 to 20 hr in the 950" to 1000°C temperature range before quenching. Spectrographic and chemical analyses showed no tantalum pickup by the alloys, indicating no reaction between the alloys and the crucibles. However, some of these crucibles failed, probably due to imperfections in the welds of the stainless steel or tantalum crucibles. The second and most satisfactory method was to prepare the alloys by powder metallurgy techniques. The procedure was to press degreased and acid-etched uranium turnings with granulated zinc into 20 g compacts under 20,000-psi pressure. The compacts were placed in MgO crucibles, and sealed in evacuated Vycor or fused silica tubes. The alloys were then heated as long as two weeks at about 550°C in a muffle furnace. The pressed compacts were observed to expand by several volume percent during heating and it was necessary to make allowances for this expansion in order to avoid breaking the crucible and Vycor tube. This method was found very satisfactory for preparing alloys which were suitable for thermal analysis or vapor pressure studies. Experimental Methods and Results The phase diagram for the U-Zn system at 1 atrn pressure, shown in Fig. 1, is based primarily on vapor pressure measurements and on thermal analysis taken at temperatures below 950°C. Fig. 2 shows the U-Zn diagram at 5 atrn pressure, constructed on the basis of thermal analysis of alloys in sealed containers up to 1150°C and on the basis of metallographic, X-ray, and analytical data. The alloys sealed under vacuum were actually under their own vapor pressure and those sealed in an atmosphere of helium were under an additional pressure due to the helium. At temperatures up to 1100°C the zinc pressure is 5 atrn or less for these alloys; consequently the maximum pressure over the alloys sealed under a helium atmosphere was 10 atrn or less at temperatures up to 1100°C. Changes in pressure of this order of magnitude do not appreciably alter the position of most solid-solid or liquid-solid phase boundaries. In constructing the phase diagram for a pressure of 5 atm, the effect of pressure on all phase boundaries except those for liquid-vapor or solid-vapor regions was considered negligible.

Jan 1, 1958

By R. A. Wood, R. I. Jaffee, H. R. Ogden, D. N. Williams

The tensile properties, creep resistance. and thermal stability of highly alloyed Ti-Al-Zr alloys were examined. On the basis of these studies, the Ti-7Al-1ZZr composition was selected for more complete evaluation. The alloy was found to be weldable and free from excessive directionality. In addition, it developed maximum properties without requiring heat treatment other than an annealing operation in the alpha field. The alloy was recommended for scale up and is presently being investigated on a production-level basis. One of the more attractive properties of titanium alloys is their ability to withstand stress at moderately high temperatures, and a considerable amount of effort has been devoted to increasing the maximum service temperature of titanium alloys. This work has suggested that the optimum alloys for high-temperature service will be single-phase a (close-packed hexagonal) alloys containing significant amounts of aluminum. However, the maximum amount of aluminum which can be alloyed with titanium is between 6 and 8 pct,l since at high-aluminum contents an embrittlement reaction occurs in the anticipated service temperature range, 800" to 1100°F. It has been shown that the embrittlement reaction involves decomposition of the high-aluminum a phase to one or more new phases.' Since this reaction does not occur at intermediate or low-aluminum contents, it was felt that intermediate Ti-A1 alloys might be strengthened by a-soluble ternary additions without inducing the embrittlement reaction. The first alloying addition considered was tin, which shows extensive solubility in a titanium and has moderate strengthening tendencies. Unfortunately, it was soon apparent that tin also promoted the embrittlement reaction, and that to obtain a stable alloy, the aluminum content had to be reduced as the tin content was increased. The second alloying addition considered was zirconium, which is similar to tin in its effects on titanium. This element did not contribute to the embrittlement reaction and, in fact, appeared to increase the maximum amount of aluminum which could be alloyed with titanium without inducing instability. This paper describes an investigation of the Ti-A1-Zr a alloy region. Alloys containing from 4 to 12 pct A1 and from 6 to 15 pct Zr were examined. The properties of these alloys are described and the bases for selecting an optimum composition is outlined. This composition, Ti-7A1-12Zr, is presently being scaled up in tonnage quantities, and is being evaluated extensively throughout the industry. In addition to presenting the basis for its selection, this paper presents a description of the properties developed in laboratory material as determined during the alloy investigation. These properties suggest that this alloy can fill an important position in applications requiring light weight, fabrica-bility, weldability, and strength to 1000oF or higher. EXPERIMENTAL PROCEDURES Titanium alloy ingots were prepared by inert electrode arc melting under an argon atmosphere. Alloying elements used were 110 Bhn titanium sponge, high-purity aluminum, and reactor-grade zirconium. Pancake-shaped ingots were prepared weighing approximately 300 g. The composition of the ingots was checked by weight measurements before and after melting. The pancake ingots were forged at 2000°F to approximately half their original thickness to give a flat plate roughly 1/2 in. thick. This plate was then rolled at 1800' to 1600°F to 0.250 in. thick. All of the alloys examined fabricated well. However, alloys containing 15 pct Zr tended to overheat due to exothermic oxidation, and scaling was excessive. As might be anticipated from its effect in decreasing the ß transus, increased zirconium appeared to improve fabricability somewhat, especially during rolling at lower temperatures. Except for a limited study of heat-treatment response, all alloys were examined in the a-annealed condition. Prior to heat treatment the a and ß tran-sus temperatures were determined by metallo-graphic examination of samples quenched after annealing at 50-deg intervals in the transformation region. These data are shown in Fig. 1. Recrystal-lization appeared to occur in about 1 hr in the range 1300º to 1500ºF. Therefore, alloys were annealed for 1 hr at 1550ºF (4 and 5 pct Al), 1600ºF (6 through 7-1/2 pct Al), or 1650°F (8 or more pct Al). This produced an equiaxed a grain structure. In most alloys, a "ghost" structure was visible after the a-annealing treatment, as shown in Fig. 2. This structure apparently resulted from the acicular

Jan 1, 1963