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By N. A. Gokcen, J. Chipman

Aluminum and oxygen dissolved in liquid iron were brought into equilibrium with pure alumina crucibles and atmospheres of known H2O and H2 contents to study the reactions: 1—Al2O3(s) = 2 Al + 3 0; 2—Al2o3(s) + 3H2(g) = 2Al+ 3H2o(g); and 3—H2(9) +O = H2O(g). Aluminum strongly reduces the activity coefficient of oxygen and similarly oxygen reduces that of aluminum. Values of the product [% All" • [% O]3 are much smaller than those found in previous experimental studies and are of the order of magnitude of the calculated values. ALUMINUM is the strongest deoxidizer commonly A used in steelmaking, but the extent to which it removes dissolved oxygen has been debatable. The relationship between aluminum and oxygen has not been determined reliably not only on account of the usual experimental difficulties at high temperatures but also because of uncertainties in the analyses of very small concentrations of oxygen and aluminum. The earliest experimental attempt of Herty and coworkers' was followed by a more systematic study of Wentrup and Hieber.' These authors added aluminum to liquid iron of high oxygen content in an induction furnace and considered that 10 min was sufficient to remove the deoxidation products from the melt. Parts of the melts thus obtained were poured into a copper mold and analyzed for total aluminum and oxygen (soluble plus insoluble forms), assuming that the insoluble parts were in solution at the temperatures from which samples were taken. It is conceivable that the furnace atmosphere in their experiments, consisting of mainly air at 20 mm Hg pressure, was a serious source of continuous oxidation and therefore that their oxygen concentrations were correspondingly high. Scattering of their data was explained to be well within the maximum inaccuracy of 10°C in the temperature measurements and errors of ±0.002 pct each in the oxygen and total aluminum analyses. Maximum and minimum deoxidation values, i.e., values of the product [% All' . [% O] differed by factors of 10 to 15; mean values of 9x10-11 and 7.5x10-9 ere reported at 1600" and 1700°C, respectively. Hilty and Craftsv determined the solubility of oxygen in liquid iron containing aluminum, using a rotating induction furnace. Pure alumina crucibles used in their experiments contained the liquid iron which in turn acted as a container for slags of varying compositions consisting mainly of Al2O3, Fe2O3, and FeO. The furnace was continuously flushed with argon, and additions of aluminum and Fe2O3 were made in the course of each experimental heat. The inner surfaces of their alumina crucibles were covered with a substance other than pure Al2O3, containing both iron oxide and alumina. Although frequent slag additions can change the composition of slag in the liquid iron cup formed by rotation, the inner surface of the crucible must depend upon the transfer of oxygen or aluminum through the liquid iron for any adjustment in composition. It is not clear that their metal was in equilibrium with the crucible wall, but it is clear that it was not in equilibrium with Al2O3. Their deoxidation product, [% A].]" • [% O]3, varied by a factor of more than 50; the average values of 2.8x10- and 1.0x10-7 were selected for temperatures of 1600" and 1700°C, respectively. Aside from the experimental determinations, attempts have been made to calculate the deoxidation constant for aluminum indirectly from thermody-namic data. Schenck4 combined the thermodynamic data for Al2O3 and dissolved oxygen in liquid iron by assuming an ideal solution. His calculated values are 2.0x10-15 and 3.2x10-13 at 1600" and 1700°C, respectively. Later, Chipman5 attempted to correct for the deviation from ideality and derived an expression which led to deoxidation values of 2.0x10-14 and 1.1x10-12 at 1600" and 1700°C, respectively. The errors in these treatments originate mainly from inaccuracies of thermal data and uncertainties regarding the activity coefficients of dissolved oxygen and aluminum. The purpose of this investigation was to study the equilibria represented in the following reactions in the presence of pure alumina: Al2O3(s) = 2Al + 3O K = aAl2.ao3 [1] Al2O3(s) + 3H2(g) = 2Al + 3H2O(g) H2O K2 = aAl2(H2O/H2 ) [2] H2(g) +O = H2O(g) K3 = 1/ao (H2) [3] The experimental method consisted of melting pure electrolytic iron, usually with an initial charge of aluminum, in pure dense alumina crucibles under a controlled atmosphere of H,O and H2 and holding

Jan 1, 1954

By Paul A. Farrar, Harold Margolin

The Ti-Al-V system has been delineated from 50 to 100 wt pct Ti and front 600 to 1400°C by X-ray and ntetallographic techniques. Isothermal sections were delineated at 600, 700, 800, 900, 1000, 1100, 1200, 1300, and 1400°C. X-ray and metallographic evidence indicated the presence of two phases and in addition to those reported by previous investigators of the ternary system'-3. T HE titanium-rich region of the Ti-Al-V system as originally proposed by Rausch et UZ, and Jordan and Duwez3 was based on the binary Ti-Al diagram as proposed by Bumps et uZ. As additional phases have been shown to exist5-' in the region previously thought to be all a, the titanium-rich region of the Ti-Al-V system is open to serious question. Therefore, in order to determine the effect of the indicated changes in the Binary Ti-Al system on the ternary Ti-Al-V system the present investigation was initiated. EXPERIMENTAL PROCEDURE The experimental procedures in this investigation are the standard techniques used in the study of titanium-alloy systems which have been described in detail previously.1° Consequently, only details pertinent to the present work are discussed. The alloys employed for the delineation of this system were prepared from titanium obtained from the Bureau of Mines (73 BHN) 99.99 pct Al and 99.7 pct V, by the consumable arc-melting technique. The titanium was premelted into buttons of 25 g because spattering during the first melt caused uncontrollable loss of alloying constituents. Typical analyses of the material used have been reported elsewhere Attempts to hot roll the alloys prior to heat treatment were made on the first series of alloys prepared, but as the area of extensive hot workability was fairly restricted, see Table I, this practice was discontinued. The times of isothermal annealing are given in Table 11. A preliminary heat treatment of 96 hr at 1000°C was used prior to heat treatment at lower temperatures. The standard techniques used for polishing specimens involved belt-grinding, grinding on emery paper, polishing electrolytically, and etching with Remington "A" etch. In the low-alloy region it was found that the oRo etch, developed by Ence and Margolin,o gave better contrast between phases. Where this procedure did not differentiate between phases clearly, the method of stain etching developed by Ence and was used. It has been previously reported15 that filing of titanium alloys is not desirable for producing X-ray powders, since the powder becomes contaminated by the file material. Therefore, for those samples not brittle enough to Crush, the procedure described below was used. The specimen was wrapped in molybdenum sheet and placed in a vacuum system and evacuated to 0.03 p. Palladium-pur if ied hydrogen was then admitted to the system until a partial pressure of 5o Cm of H was reached. The specimen was heated in the system to the temperature of original heat treatment and was slowly cooled over a period of 3 hr to 400°C while the partial pressure of hydrogen was maintained. The sample was removed from the system after it had completely cooled and was crushed to a powder of 230 to 270 mesh. The powder, wrapped in molybdenum sheet, was dehydrogenated at the temperature of original heat treatment until a partial pressure of 0.03 p was maintained in the system for at least 1/2 hr. The tube containing the powders was quenched in iced brine while still connected to the vacuum system. After dehydrogenation, the powder, except for the

Jan 1, 1962

By G. V. Smith, C. O. Tarr, R. F. Miller

Metallurgists have long recognized that the Fe3C type of carbide is not a stable phase in steel and that, given sufficient time, it will decompose with formation of graphite, at least at temperatures below about 205o°F.l This slow graphitization has never been of much concern to users of hypoeutectoid steels, for in such steels it occurs only during prolonged exposure at elevated temperature. In I935, Kinzel and Moore2 reported complete graphitization of the carbide of a 0. I5 per cent carbon steel that had been in service for about 26,000 hr. at a temperature of about IIOO° to 12oo°F. In discussing this paper, both Mathews and Sauveur reported graphitization of slightly hypereutectoid steels. In an extensive investigation of graphitization of high-purity iron-carbon alloys, Wells1 produced graphite at 7oo°C. (I29o°F.) in an alloy containing as little as 0.13 per cent carbon, and showed that graphite may precipitate either directly from austenite or from decomposition of Fe3C, also that the rate of graphitization tends to increase with increase of temperature, at least up to a certain point, and with the presence of graphite nuclei. Jenkins, Mellor and Jenkinson,3 studying the structural changes in carbon steels ranging from 0.17 to 1.14 per cent carbon during stress-rupture tests at elevated temperature, observed graphite in an aluminum-killed 0.40 per cent carbon steel after fracture in 4340 hr. under a stress of 17,900 lb. per sq. in. at 840°F. and in many of the steels after stress-rupture test at 12oo°F. Unstressed samples of 0.60, 0.86 and I.I4 per cent carbon steels graphitized almost completely in 360 hr. at I200°F., while samples of 0.90 and 1.10 per cent carbon steels were somewhat graphitized, and steels containing 0.24, 0.40 and 0.57 per cent carbon showed no sign of graphite. Their results indicate that graphite precipitates the more readily the higher the carbon content, though this tendency is influenced to some extent by deoxidation practice, and that graphitization is accelerated by plastic deformation. In a recent study of the graphitization of 0.4 and 0.6 per cent carbon steel strip, made by M. A. Hughes,4 it was found that at I2oo°F. graphite appeared after 72 hr. in aluminum-killed but not in silicon-killed steel, and that the rate of graphitization was increased by slow cooling from I6oo°F. and by cold-working prior to annealing. In creep tests of normalized 0.15 per cent carbon steel killed with 2.5 Ib. of aluminum per ton, at the U. S. Steel Corporation laboratory at Kearny, spheroidization and some graphitization were found after 3000 hr. at Iooo°F. under a stress as low as 1500 lb. per sq. inch. .

Jan 1, 1944

By C. O. Tarr, G. V. Smith, R. F. Miller

Metallurgists have long recognized that the Fe3C type of carbide is not a stable phase in steel and that, given sufficient time, it will decompose with formation of graphite, at least at temperatures below about 205o°F.l This slow graphitization has never been of much concern to users of hypoeutectoid steels, for in such steels it occurs only during prolonged exposure at elevated temperature. In I935, Kinzel and Moore2 reported complete graphitization of the carbide of a 0. I5 per cent carbon steel that had been in service for about 26,000 hr. at a temperature of about IIOO° to 12oo°F. In discussing this paper, both Mathews and Sauveur reported graphitization of slightly hypereutectoid steels. In an extensive investigation of graphitization of high-purity iron-carbon alloys, Wells1 produced graphite at 7oo°C. (I29o°F.) in an alloy containing as little as 0.13 per cent carbon, and showed that graphite may precipitate either directly from austenite or from decomposition of Fe3C, also that the rate of graphitization tends to increase with increase of temperature, at least up to a certain point, and with the presence of graphite nuclei. Jenkins, Mellor and Jenkinson,3 studying the structural changes in carbon steels ranging from 0.17 to 1.14 per cent carbon during stress-rupture tests at elevated temperature, observed graphite in an aluminum-killed 0.40 per cent carbon steel after fracture in 4340 hr. under a stress of 17,900 lb. per sq. in. at 840°F. and in many of the steels after stress-rupture test at 12oo°F. Unstressed samples of 0.60, 0.86 and I.I4 per cent carbon steels graphitized almost completely in 360 hr. at I200°F., while samples of 0.90 and 1.10 per cent carbon steels were somewhat graphitized, and steels containing 0.24, 0.40 and 0.57 per cent carbon showed no sign of graphite. Their results indicate that graphite precipitates the more readily the higher the carbon content, though this tendency is influenced to some extent by deoxidation practice, and that graphitization is accelerated by plastic deformation. In a recent study of the graphitization of 0.4 and 0.6 per cent carbon steel strip, made by M. A. Hughes,4 it was found that at I2oo°F. graphite appeared after 72 hr. in aluminum-killed but not in silicon-killed steel, and that the rate of graphitization was increased by slow cooling from I6oo°F. and by cold-working prior to annealing. In creep tests of normalized 0.15 per cent carbon steel killed with 2.5 Ib. of aluminum per ton, at the U. S. Steel Corporation laboratory at Kearny, spheroidization and some graphitization were found after 3000 hr. at Iooo°F. under a stress as low as 1500 lb. per sq. inch. .

Jan 1, 1944

By Aloys IL. Wollrab

The conditions that govern the selection of a suitable type of detachable bit for the small, isolated mine, for rock work and tunnel contracting and for the large mine are quite dissimilar, therefore intelligent recommendation of a particular make of detachable bit must involve consideration of many factors. Small, isolated mines may find advantageous a detachable bit that a larger mine cannot afford to use, because of more efficient conventional steel sharpening and steel-distribution practice, but a small mine may not be able to adopt the practice of large mines because of necessary capital expenditures for equipment for detachable-bit sharpening and . shanking. The perfect detachable bit, suitable for all conditions, has not yet been developed. Conventional Drill Bit in Use That drilling speed increases as the bit gauge decreases is well known, but since small-hole drilling is the exception rather than the rule the rate of such increase is probably not so generally recognized. Tests carried out by Holman Brothers Limited1 with 7/8-in. steel show that when bit sizes are reduced from I½ in., I3/8 in., and I3/16 in-; to I3/8 in., I1/4 in.., and I1/8 in., respectively, the volume of the hole is I 5 per cent less and the drilling speed 49 per cent higher. Recent tests by the United States Bureau of Mines2,3 in hard, uniform basalt showed that a reduction in average gauge from 2.077 to I,826 in., a difference of o.25I in., gave an increased drilling speed of 65 per cent. The limits to which a drill hole may be extended depend upon the loss in gauge. Since the conventional bit with 5° and 14° clearances did not possess sufficient reaming edge to permit the smaller gauge changes that were required for small-hole drilling, experiments were made in bit construction with a view to correcting this deficiency. This resulted in the adoption of the full reaming cross bit (having full Carr charac1erist.i~~) by some mines4 In order to drill holes having a minimum taper from the collar to the bottom of the hole, full reaming center-hole cross bits with 90° cutting angles, 5° and 14° wing tapers, forged on 7/8-in. quarter-octagon drill rods, are used in the conventional drill-steel practice at Hollinger. In the steel shop close attention is given to the condition of the dies and dollies, which, combined with a rigid inspection of the finished bit, ensures as nearly a perfect bit as call be forged economically. Shop practices and costs have been discussed in earlier paper.5,6 The method by which the radii for the special dies used for forging bits of different gauge diameter are obtained so that the points and reaming edges cut circles of the same diameter has been discussed by Hibbert. Requiremnts for Detachable. Bit for Hollinger Under these it it was appar- ent that it would be difficult to introduce a detachable bit that could successfully compete with the conventional bit already in

Jan 1, 1943

By Aloys IL. Wollrab

The conditions that govern the selection of a suitable type of detachable bit for the small, isolated mine, for rock work and tunnel contracting and for the large mine are quite dissimilar, therefore intelligent recommendation of a particular make of detachable bit must involve consideration of many factors. Small, isolated mines may find advantageous a detachable bit that a larger mine cannot afford to use, because of more efficient conventional steel sharpening and steel-distribution practice, but a small mine may not be able to adopt the practice of large mines because of necessary capital expenditures for equipment for detachable-bit sharpening and . shanking. The perfect detachable bit, suitable for all conditions, has not yet been developed. Conventional Drill Bit in Use That drilling speed increases as the bit gauge decreases is well known, but since small-hole drilling is the exception rather than the rule the rate of such increase is probably not so generally recognized. Tests carried out by Holman Brothers Limited1 with 7/8-in. steel show that when bit sizes are reduced from I½ in., I3/8 in., and I3/16 in-; to I3/8 in., I1/4 in.., and I1/8 in., respectively, the volume of the hole is I 5 per cent less and the drilling speed 49 per cent higher. Recent tests by the United States Bureau of Mines2,3 in hard, uniform basalt showed that a reduction in average gauge from 2.077 to I,826 in., a difference of o.25I in., gave an increased drilling speed of 65 per cent. The limits to which a drill hole may be extended depend upon the loss in gauge. Since the conventional bit with 5° and 14° clearances did not possess sufficient reaming edge to permit the smaller gauge changes that were required for small-hole drilling, experiments were made in bit construction with a view to correcting this deficiency. This resulted in the adoption of the full reaming cross bit (having full Carr charac1erist.i~~) by some mines4 In order to drill holes having a minimum taper from the collar to the bottom of the hole, full reaming center-hole cross bits with 90° cutting angles, 5° and 14° wing tapers, forged on 7/8-in. quarter-octagon drill rods, are used in the conventional drill-steel practice at Hollinger. In the steel shop close attention is given to the condition of the dies and dollies, which, combined with a rigid inspection of the finished bit, ensures as nearly a perfect bit as call be forged economically. Shop practices and costs have been discussed in earlier paper.5,6 The method by which the radii for the special dies used for forging bits of different gauge diameter are obtained so that the points and reaming edges cut circles of the same diameter has been discussed by Hibbert. Requiremnts for Detachable. Bit for Hollinger Under these it it was appar- ent that it would be difficult to introduce a detachable bit that could successfully compete with the conventional bit already in

Jan 1, 1943

By D. C. Hilty

C. E. SIMS*—This is a most interesting and important paper. It is important from two standpoints. First, it has as-spects of being highly accurate and therefore extremely useful to the operating man in showing how far he can go in oxidizing carbon from a steel bath while retaining chromium and the effect of temperature on this limit. Fig 5 is undoubtedly the highlight of the paper, and in checking it against available data, it seems capable of predicting actual results. The other important aspect is the apparently significant fact that reactions within a steel bath, such as between carbon and an oxide, are homogeneous reactions which do not require equilibrium with an external phase. In this case, it seems obvious that the reaction and the equilibrium established were between dissolved carbon (or carbide) and dissolved CrO. The classic reaction which requires equilibrium with solid Cr2O3, did not fit the data. It was recognition of this principle that made Fig 5 possible. G. R. FITTERER*—It might be in assuming the CrO is dissolved in the metal that you are actually approaching the correct factor for activity. If that is true and if the percentages that were used represent an equilibrium constant, this would mean, in effect, that the minimum temperature at which chromium would be reduced by an equivalent percentage of carbon would be about 2400°F from your equation. I believe that is in line with my experience particularly if you deal only with low percentages. J. CHIPMAN†—The author in discussing the equilibrium of the reaction offers two choices of explanation. One postulates an effect of chromium on the activity coefficient of carbon; it would require perhaps more data than are available to us at the present time to test this thoroughly. The other postulate which he offers is a reaction with CrO. This explanation is consistent with physical and chemical principles provided that a solid phase CrO is present, and that is essentially what he is postulating. If that oxide were present one would expect and, I think, almost require the kind of results that he obtained. The weakness of that postulate lies in the lack of confirmatory evidence regarding the presence of solid CrO as such. This should offer an interesting subject for further study. It seems a pity that the theoretical arguments revolving around CrO should have obscured for a moment the very real excellence of this paper. E. CARTER*—I think there is a problem that presents itself in practical application where the metal temperature is from 3200 to 3300 and the temperature desired in pouring is considerably less. B. M. LARSEN†—I get the impression from some of the work that was done earlier that there is probably a localized zone of high temperature at the slag metal interface which effects the removal of the carbon. Without that I wonder if the effect is possible to have a localized superheat and an equilibrium there favorable to low carbon content and

Jan 1, 1950

By P. Greenfield, C. C. Smith, A. M. Taylor

J. E. Harris (Berkeley Nucclear Laboratories, England)—Greenfield et al.11 attribute abrupt changes in slope of their log o/log i curves for heat-treated Mg/0.5 pet Zr alloy (zA) to &apos;atmosphere&apos; locking. It is proposed here that a more reasonable explanation of the apparent strengthening at low rates of strain can be based on precipitation either during the preanneal or during the creep tests. All the tests were carried out above 0.5 Tm where solute atmospheres are likely to be largely evaporated2 and can migrate sufficiently rapidly so as not to impose any &apos;drag&apos; on the moving dislocations. McLean3 has derived an expression for determining the temperature Tc above which, due to the high-migration rate of the atmospheres, Cottrell or Suzuki locking can play no part in determining creep strength. This expression, which holds for an applied shear stress of not greater than 5 X 107 dynes per sq cm is: Tc/Tm= 7/6.8 - log10? where i = secondary creep rate The values for T, corresponding to the maximum and minimum reported creep rates at each temperature have been calculated from the data of Greenfield et al. These are given in Table VII. All the test temperatures were above T,, the margin being greater for the higher temperatures and for the lower strain rates where the breaks in the log s/log ? curves occurred. Dorn and his collaborators14, 17 have studied systematically the effect of solute hardening on the creep properties of an A1/3.2 at. pet Mg alloy. In the temperature range where strain aging occurred in tensile tests, abnormally high-activation energies for secondary creep were obtained but at temperatures above 0.43 Tm, solute alloying did not have any effect on the creep parameters. Moreover, there have been no reports of any strain aging phenomenon during elevated temperature tensile tests with ZA material.18 Instead of the observed strengthening being due to atmosphere locking, it is now proposed that precipitates play an important role in enhancing the creep strength of the material. There are two possibilities—precipitation of zirconium hydride during the high-temperature preanneal and/or precipitation of the hydride or a-zirconium during creep. On the basis of the former the results can be interpreted in terms of a critical stress being necessary to force the dislocations through or over preexisting precipitates. From the latter, if the strengthening is due to pre- cipitation during the test then hardening should be associated with a critical strain rate. At low rates of strain, time is available during the tests for precipitation to occur either directly onto dislocations (thus pinning them) or generally throughout the matrix (which would impede dislocation movements). Examination of the data of Greenfield et al. suggests that both mechanisms may be operative since they observed precipitation during creep and also found that their alloys exhibited high-creep strength in the early stages of the low-stress tests, i.e., before creep-induced precipitation had time to occur. It is not easy to understand why they considered that precipitation of zirconium hydride is unlikely to occur at 600°C while it can take place in tests in air at as low a temperature as 200°C. Precipitation of the hydride during the preanneal cannot be ruled out merely on the basis of metallographic examination. Hydride precipitates in ZA type alloys are very small and can only be accurately resolved in the electron microscope.9 For example, in this laboratory20 hydride platelets with major dimensions <(1/10) µ have been observed by electron transmission through thin film specimens of hy-drogenated ZA material. Complex interactions between dislocations and such particles are illustrated in Fig. 12. Additional evidence for precipitation during pre-annealing is provided by the data presented in Greenfield&apos;s Fig. 1 and Table IT. These show that the creep strength at 200o and 400°C increases with the time of preanneal at 600°C. Such increases cannot be explained on the basis of increases in grain size alone for further improvements in strength were observed when the material was annealed for longer times than that required to stabilize the grains. Although the main discussion is confined to ZA material, similar arguments can be used against the strain aging hypothesis proposed to explain the binary Mg/Mn alloy data. In this case no precipitation is possible during the preanneal, but precipitation-hardening during creep can occur.

Jan 1, 1962

By G. H. Wigton

THE oxidized ores of the Chief Consolidated Mining Co., in the Tintic mining district, have never yielded to metallurgical treatment by any standard method except smelting. These ores occur in comparatively small, irregular, siliceous replacements in limestone, and carry silver, lead, and gold as the principal valuable metals. But smelting methods. are unsatisfactory, as the lead smelters heavily penalize the large amount of silica in the ore, and the copper smelters do not pay for the lead in the ore; therefore, an alternative treatment method was sought. In 1916, a metallurgical research department was organized under; the direction of the writer. Many standard metallurgical methods were investigated but, on oxidized ore, none yielded extraction great enough to make its use feasible. During some chloridizing-roasting experiments in 1916, made in an assay muffle, the writer noted that large proportions of the silver, lead, and gold were volatilized from certain oxidized ores. The possibility of volatilizing the metals and collecting them in the form of a concentrated fume was then studied. But, as it was shown that satisfactory chloride-volatilization would be limited to a furnace feed with a high fusing temperature of the non-volatilizable constituents, and as the results obtained in a rotary kiln with the admixture of common salt with the ground ore were inferior to muffle tests, further volatilization experiments were discontinued until 1920. During this interval, an experimental sulfidizing and flotation plant was operated. After the cessation of attempts to float the oxidized minerals, chloride-volatilization experiments were resumed. The highly infusible flotation. tailing from the operation of the flotation mill was used for furnace feed. The first tests made on flotation tailing mixed with 10 per cent. common salt showed -nearly complete lead and gold volatilization but considerably less volatilization of silver than in muffle tests. In attempts to increase the silver extraction, the percentage salt in the furnace feed was varied, and it soon became apparent that the use of smaller percentages of salt raised the fusion temperature of the furnace charge and thus,

Jan 8, 1925

Aamot, Olav Crone, (M&apos;29) Chem. Engr., Elektrokemisk, Raadhusgt. 23, Oslo, Norway. Abadilla, Quirico A., (M&apos;3S) Min. Engr., Dir., Bu¬ reau of Mines. Manila, P. I. Abbott, Agatin Townsend, (J&apos;40) Jr. Engr., Phelps Dodge Corp., Morenci Branch, Morenci. Ariz. Abbott, Clarence E., (M&apos;04) Civil Engr., V.P., Ten¬nessee Coal Iron & R.R.. Co., 1542 Brown -Marx Bldg.. Birmingham, Ala. Abel, Raymond L., (M&apos;28) Chem. Engr., Command¬ing, Officer, Pittsburgh Chem. Warfare Procurement Dist., American Bank Bldg., Pittsburgh. Pa. Abel, Walter D., (M&apos;32) Alin. Engr.. Div. of Cor¬porations, State of Calif., 403 California State Bldg., Los Angeles, Calif. Abele, Louis T., (M&apos;31) 423 3rd Ave., Decatur, Ga. Abell, Leo M., (M&apos;31) Mill Supt. & Met., Emperor Gold Min. Co., Led., Vatnkoula, Fiji Islands. Abernathy, G. E., (M&apos;38) Geol., Kansas State Geol. Survey, Pittsburg, Bans. Abernathy, Jack H., (M&apos;40) Peer. Engr., c/o Superior Oil Corp., National Bank of Tulsa Bldg., Tulsa, Okla. Aborn, Robert H., (M&apos;10) Asst. Met., Research Lab., U. S. Steel Corp., Lincoln Highway, Kearny, N. J. Abraham, A. W., (M&apos;39) Peer. Engr., Richfield Oil Corp., 555 S. Flower St., Los Angeles, Calif. Abramowitz, William L., (J&apos;40) Chem., Industrial Inventions, Inc., 103 Foster Se., Peabody, Mass.

Jan 1, 1942

By D. A. Dukelow, G. Derge

THE Fe0-Mn0-Si0, system has many properties of fundamental interest besides its occurrence in steel-making. The system is the simplest ternary complication of the FeO-SiO, binary whose electrochemical properties have been explored. Therefore, knowledge of the conformance of Fe0-MnO-SiO, melts to Faraday&apos;s laws, as investigated by current efficiency measurements, would be of particular value in interpreting the observed variations of electrical conductivity, while at the same time furthering understanding of this type of molten oxides. Pertinent to the present study are conductivity data of FeO-SiO, and Fe0-Mn0-SiO2, and current efficiency data of FeO-SiO,. Concerning conductivity measurements, Inouye, Tomlinson, and chipman&apos; studied the specific conductivity, a, of molten wüstite as a function of temperature and oxide addition. Liquid FeO in equilibrium with an iron crucible was found to have a o whose magnitude and positive temperature coefficient were undeniably semiconductor in nature (o1400°C = 269 ohm-1cm-&apos;). Additions at 1400°C of 5 mole pct MnO decreased o to 227 while a similar addition of SiO2, dropped o to 125. Additional SiO2 produced an exponential decline of o. It was reasoned that because wüstite is known to be oxygen excess and since there is no pronounced change in oFeO upon melting, that UFeO(L)is p-type semiconduction. Björling2 gives o data for the ternary FeO-MnO-SiO, using fayalite, 2 FeO-SiO,, as the base substance. Upon addition of tephroite, 2 MnO-SiO, he noted a small decrease in o. Taking his data for 1400°C with one-half of the fayalite replaced by tephroite, o shows a decrease of some 40 pct relative to o for fayalite, i.e., 5 ohm-&apos; cm-&apos; to 3. In line with this observation, Mori and Matsushita3 reported that substituting MnO for FeO in the silicate de- creased o slightly except at high SiO, contents where the decrease was somewhat larger. Investigating the ionic nature of FeO-SiO2 melts, Simnad, Derge, and George4, found the anode current efficiency, C,, as calculated from iron anode weight loss, was temperature invariant and in general increased with SiO2 concentration. The authors suggested the possible presence of electronic conduction to explain why C+ < 100 pct. Fig. 4 includes their results as curve (b). EXPERIMENTAL METHOD Apparatus—Fig. 1 is a schematic section of the electrolysis cell used in all the current efficiency runs. The melt was contained in an iron crucible that rested on an alundum ring. A centered silica tube came within 1 in. of the surface of the melt and initially served as a thermocouple well and then as the guide for the anode.

Jan 1, 1961

By N. A. Gokcen, S. Fujishiro

THE pressures of Mn(g) in equilibrium with Mn7C3 and graphite have been measured by McCabe and Hudson&apos; and Butler, McCabe, and paxton2 by means of graphite, zirconia, and Ta-Mo Knudsen cells. The details of corrections in their measurements were somewhat different from those of the authors;3&apos;4 hence, further independent results at higher temperatures seemed to be warranted. The stoichiom-etry of Mn7C3 in equilibrium with graphite and its stability up to the vicinity of 1373 °K have been satisfactorily established by Kuo and persson5 (See also Hansen and Anderko).&apos; The method of investigation has been described elsewhere in detail.3&apos;4 The carbide was prepared in the manner described by McCabe and Hudson. In order to avoid possible reactions with refractory oxides, only graphite cells were used for the effusion experiments. The authors&apos; results are the runs M-1, M-2 and M-3 listed in Table I, together with those of McCabe and co-workers. The heat capacity, c;, of Mn7C3 is not known; therefore McCable et al. plotted their data as log P vs 1/T and the results were only expressed first&apos; as, and later2 as logP=- -14220+6.06 without further thermodynamic calculations. However, it is possible to estimate closely the necessary heat capacity as follows: When H; - &98 of CrC, or MnO, is plotted vs x, corresponding to various compounds, with temperature as the parameter, it is seen that the relationship is represented by a line of very small curvature. This relationship is also followed by other compounds which do not exhibit solid state phase transformations. In view of the fact that chromium and manganese have the atomic numbers of 24 and 25 respectively, it is reasonable to assume that H; - H;ae vs x for MnC, follows the same type of lines as those for CrC, for which adequate data are available.7 Similar correlations appear to be valid for oxides of two neighboring metals for which reliable data are available. The procedure is thus to move the lines for CrC, vertically (parallel to the axis of HT - H298) until they coincide with the points for MnC1/3, the only carbide of manganese whose H; - HO298 has been

Jan 1, 1963

[A Aamodt, H. C., A55 C Sales, Great Lakes Carbon Corp., 18 E. 48th St., New York, N.Y. Aamot, Olav Crone, M29 2126 M4, Galveston, Texas Aaronson, Hubert I., M47 E Prin Res Engr, Ford Motor Co., Met. Dept., Sci. Lab., Engrg. Staff, P. O. Box 2053, Dearborn, Mich. Aase, Glenn D., M54 D Prin Engr, Air Pollution Control Dist., 1480 Mission St., San Francisco 3, Calif. Abbett, W. A., Jr., M56 8629 Beverly St., Duluth 8, Minn. Abbey, Robert G., M56 138 Dartmouth Rd., Bala Cynwyd, Pa. Abbott, A. G., Jr., J52 E Test, Dev Engr, Chrysler Corp. Missile Div., P. O. Box 2628, Detroit 31, Mich. Abbott, Mark L., A58 Rep, General Refractories Co., 208 S. LaSalle St., Chicago, Ill. Abbott, W. K., J58 E Dev Met, Int&apos;l Nickel Co., Inc., 67 Wall St., New York 5, N.Y.]

Jan 1, 1961

By A. Goldberg, T. E. Tietz, J. E. Dorn

INTRODUCTION THE concept that the flow stress for plastic deformation of metals in the work hardening range is a function of the instantaneous values of the strain, strain rate and test temperature was promulgated by Ludwick1 many years ago. This thought has been more or less a guiding principle throughout the immediate period of development and growth of mechanical metallurgy. It has been tacitly assumed or knowingly applied to many investigations on the plastic behavior of metals. Spurred by the theoretical interest and practical importance of the plastic behavior of metals, many investigators have attempted to determine empirical, semi-empirical, and theoretical equations2-18 relating the flow stress with the strain, strain rate, and temperature. Although the applications of such equations have been extensive, the experimental verification of the existence of a mechanical equation of state is weak. Furthermore, some of the available evidence on plastic flow appears to suggest that no simple functional relation between stress, strain-rate, strain and temperature can exist. The mechanical equation of state demands that [a = a(e, E, T)[I]] Where o = true stress = load/instantaneous area e = true strain = In Instantaneous gauge lengths initial gauge length e = true strain rate T = temperature] This formulation implies that [da] is an exact differential. Therefore the flow stress (o) is a function of the instantaneous values of the strain (e) the strain rate (e) and the temperature (T), independent of the previous mechanical or thermal history. Having reached a certain strain, strain rate, and temperature, by any path whatsoever, a definite and fixed value of the flow stress should be obtained. One check on the validity of the mechanical equation of state is easily obtained. Consider the following three tests on the stress-strain curve at constant temperature: [A. Strain at rate [ei] B. Strain at rate i2 C. Strain at rate [Ei] to [ei] and continue test at i2] A graphical representation of possible results of such tests are shown in Fig I. If, for strains greater than [el], test (C) data coincide exactly with the data from test (B), as shown in Fig Ia, a limited verification of the mechanical equation of state is obtained for the range of strain rates from [ii] to [e2]. But if the data from test (C) differ from those of (B) beyond strain [el] where the strain rate is changed from [e,] to [€2i] the mechanical equation of state is not

Jan 1, 1948

The income of the Institute is derived mainly from dues, advertising in MINING AND METALLURGY and sale of publications. These sources are fortunately supplemented by the interest from invested funds now amounting to over $500,000, a considerable portion constituting endowments for especial purposes. Aside from the medal funds the principal endowments are the James Douglas Fund for support of the Library, the Rocky Mountain Fund and the Seeley W. Mudd Memorial Fund. The income from the two last named is available for the support of research and a variety of other purposes. It is allocated by the Board of Directors upon recommendation of standing committees in each case. An active effort is being made to increase the endowments of the Institute, it being felt that unusual opportunities exist for research and public and professional service through the organization.

Jan 1, 1939

By D. D. Lambert, E. C. Simmons

Rare earth element (REE) analyses of cumulus plagioclase from the PGE ore zone of the Stillwater Complex indicate that pressure fluctuations in the magma chamber acting on a single magma cannot be responsible for features such as the reappearance of Mg-rich olivine on the liquidus. The data are more consistent with a model involving multiple injection and mixing of geochemically distinct magmas. The PGE-bearing sulfides of the J-M Reef are associated with such a reappearance of olivine, indicating that the unusual location of the ore zone may be related to the addition of new magma, compositionally distinct from magmas that produced cumulates in the Ultra- mafic zone of the complex.

Jan 1, 1984

By S. T. Wlodek

The surface and subscale oxidation reactions were followed by means of continuous weight-gain and metallographic techniques over the range 1600" to 2200°F (871° to 1204 °C) for up to 400 hr. Full identification of all scale and subscale reaction products was obtained by electron and X-ray diffraction. At or below 1800°F (982°C) a linear rate of reaction (QL = 46.0 kcal per mole) governed the oxidation process, extending for up to 100 hr at 1600°F (871 "C). During linear oxidation the surface scale consisted of an amorphous SiO2 film overgrown with Cr 2O 3 and NiCr204. This initial linear process was followed, and above 1800°F completely replaced, by two successive parabolic rate laws (Qp = 60 and 57 kcal per mole). This parabolic reaction involved the formation of a complex scale consisting of Cr2 O3 and smaller amounts of NiCr2O4. Parabolic oxidation appeared to coincide with the disruplion of the silica film present during linear oxidation and was followed by subscale (internal) oxidation of crystobalite and NiCr2O4. The balance between the subscale and surface oxidation reactions controls the oxidation of this commercial alloy. The amorphous silica film appears to result in the linear rate and diffusion through Cr2O3 is the more likely rate-limiting step during parabolic oxidation. THE oxidation of a multicomponent composition is a complex phenomenon not presently amenable to a rigorous classical interpretation. Nevertheless, even a qualitative understanding of the scaling and subscale reactions that occur in a commercial composition can illuminate the reactions that limit its high-temperature stability in an oxidizing environment. This study of the oxidation of Hastelloy Alloy X presents the first of a series of studies with the above approach in mind. Hastelloy X exhibits one of the best combinations of strength and oxidation resistance available in a wrought, solution-strengthened, nickel-base alloy. Although during long time exposure some precipitation of M6C and M23C8 carbides as well as a complex Laves phase occurs, the amounts are probably small enough to have no appreciable effect on the chemistry of the matrix. Radavich has identified the oxidation products on Hastelloy X oxidized for 5 min to 10 hr at 1115°F as NiO and the NiCr2O4 spinel. Oxidation for 5 to 15 min at 1500°F produced a scale of spinel, NiO, and a rhombohedra1 phase, probably Cr2Os. Sannier et 2. have reported continuous weight-gain data for Hastelloy X at 1650" and 2010°F and internal-oxidation measurements after 150 hr at 2010°F. In addition, much of the data on binary Ni-Cr alloys recently reviewed by Kubaschewski and okins&apos; and Ignatov and Shamgunova4 as well as studies of binary Ni-Mo alloys5 are also pertinent to the oxidation of this composition. EXPERIMENTAL Continuous weight-gain measurements and metallographic measurements of subscale reactions were the main experimental techniques used in this study. X-ray and electron diffraction backed up by a limited amount of electron-microprobe analysis served to characterize the nature of the scale- and subscale-reaction products. Two heats of commercial sheet of the composition given in Table I and identified as A and B were used in the bulk of this study. Internal-oxidation measurements were made on a third heat of material in the form of a 0.5-in.-diam bar. In order to assure homogeneity, all heats were reannealed 4 hr at 2175°F prior to sample preparation. weight-Gain Measurement. All specimens (1.5 by 0.4 by 0.03 in.) were abraded through 600 paper, electropolished, and lightly etched in an alcohol-10 pct HCl solution. An electrolyte of 150 cu cm H,O, 500 cu cm HsPO4 (85 pct conc), and 3 g CrO3 at a current density of 0.9 amp per sq cm or a solution of 10 pct HaW4 in alcohol used at 4 v and 0.3 amp per sq cm was used for electropolishing. The resultant surface exhibited a finish of 3 ± 1 p rms. Continuous weight-gain tests were made at 1600°, 1700°, 1800°, 1900°, 2000", and 2200°F on auer&apos; type balances capable of recording a total weight change of 110 mg with an accuracy of k0.1 mg. All tests were made in air dried to a dew point of -70°F and metered into the 2-in.-diam reaction

Jan 1, 1964

By Werner Engelmaier

Constant-strain rotating-bending fatigue tests were conducted on 2024-T3 aluminum alloy constant-strain McAdams-type specimens in ultrahigh vacuum, 10-lo Torr, and in atmospheric air. In the elastic strain range the ratio of vacuum-to-air fatigue life varied from about 10 at a maximum bending stress of 37 ksi to near unity at the endurance limit. After 10 pct of the fatigue life, intergvanular microcracks appeared on the surface which did not appreciably propagate further until the last 5 pct of the fatigue life, where trans-granular cracks propagating first in slip-plane cracking and then in the stress-dominated mode led to failure. No reduction in residual tensile strength resulted from the fatigue of the specimens short of fracture. The mechanism governing the fatigue life operates between the occurrence of the intergranular microcracks and the initiation of the trans granular cracks. MORE than forty years ago, Mc Adam&apos; and Haigh2 published the first systematic studies showing that the fatigue properties of metals were greatly affected by environments. Gough and sopwith3 found in a later study that the fatigue life of metals subject to corrosion can be substantially increased by reducing the atmospheric pressure. This and their later work4&apos;5 was, however, done in a vacuum of only 10~3 Torr. Subsequent studies6-&apos; at a vacuum of 10- Torr showed that water vapor is the primary agent affecting the fatigue properties of aluminum alloys. Only recently was ultrahigh vacuum used for fatigue studied.l Various theories have been put forward to explain this phenomenon. Wadsworth and Hutchings12 postulated that cracks form very early in the test and that the chemical attack by the corrosive components of the atmosphere at the root of the cracks enhances the speed of the crack propagation. They also noted that chemisorbed layers of gas on the fresh crack surfaces prevent cold welding of the crack and make the separation irreversible. Ham and Reichenbach, on the other hand, reported that their specimens showed no cracks after a run in vacuum, even though the runs were long enough to have caused the specimen to fracture in atmospheric air. They also reported that in argon at 0.14 Torr the fatigue life was about 5 times longer than at 7 X lo-&apos; Torr. Jacisinll reported that fatigue tests in dry nitrogen showed results close to those obtained in ultrahigh vacuum. The present experiment was designed to determine the extent to which cumulative fatigue damage, i.e., the accumulation of length of fatigue cracks, reduces the residual static tensile strength of an aluminum alloy in both ultrahigh vacuum and atmospheric air and to obtain fatigue-life data for this aluminum alloy at various strains. It was also hoped that information ex- tracted from these data might support one of the previously mentioned theories. 1) EXPERIMENT Test Specimens. The material tested was commer-cial 2024-T3 aluminum alloy. In addition to aluminum, this alloy contains approximately 4.4 pct Cu, 0.5 pct Fe, 0.5 pct Si, 0.6 pct Mn, 1.5 pct Mg, plus smaller amounts of other elements.13 Some typical mechanical properties of this aluminum alloy are listed in Table I.&apos;~ The material was used in the as-received condition, i.e., solution-treated and cold-worked. Mc Adam-type1&apos; constant-stress specimens and notched fatigue specimens were machined from 4 -in. -diam rods, Fig. 1. The critical section of the McAdam-type samples was polished with a 45-deg spiral motion to a No. 8 microfinish. The notch in the notched specimens was machined on a lathe with a forming

Jan 1, 1969

By John Fritz

MEMBERS of the American engineering societies who were in London and Paris during the last days of. June and early July were present at many interesting gatherings. The official delegates of the Founder societies were the specially honored guests of the foreign engineering societies, but other American engineers took part in the meetings. The members of the delegation were: Arthur S. Dwight, Charles F. Rand, William Kelly, representing A. I. M. E.; Ambrose Swasey (general chairman of deputation), Ira N. Hollis, Jesse M. Smith, representing A. S. M. E.; F. B. Jewett, A. E. Kennelly, G. O. Squier, representing A. I. E. E.; Charles T. Main, Robert A. Cummings, John R. Freeman, D. C. Henny, representing A. S. C. E.

Jan 1, 1921

By R. T. Hukki

THE accepted basis of comparisons between mills of different diameter is the percentage critical speed. If n = actual mill speed, rpm, nc = calculated critical speed, rpm, np = calculated percentage critical speed, and D == inside diameter of the mill in feet, then n, In the following analysis capacity, T, is expressed in short tons per hour, tph, and power consumption, P, in kilowatts, kw. Accordingly power consumption per unit of capacity, P will be expressed in kilowatt hours per short ton, or kw-hr per ton. In all equations D refers to the inside diameter of the mill in feet and v to the peripheral speed of the mill in feet per minute inside the liners. &apos; Comparison between separate mills must be based on equivalent grinding conditions, i.e., same feed, same size distribution of feed, same size distribution of product, and same percentage of solids. In addition, comparisons between separate rod mills must be based on the same rods, same type of liners, and same percentage rod load. Comparisons between separate ball mills presuppose the same balls, similar liners, and same relative ball load. The practical np-range through which the equations apply varies, being narrower for fine grinding in ball mills and wider for coarse crushing in rod mills. The Relationship between Capacity and Speed It is the general belief that the capacity, T, of a tumbling mill is directly proportional to the speed of the mill, other things remaining constant.&apos; Mathematically this is represented by the equation T - c¹ n tph [4] where c, is a factor related with the grinding characteristics of the ore, method of reduction, and the units chosen. It is proposed here that the general equation relating mill capacity and speed should be of the form T = c¹ nm tph [5] In other words, the capacity should be proportional to the mill speed raised to power m, the numerical value of the exponent being 1 5 m 5 1.5, depending on the circumstances. Eq. 5 can also be written in the following forms: T = c, (np)m tph, and [6] T=Ca vm tph, [71 where v = peripheral speed of the mill in feet per minute. If the observed capacity of a mill at speed n¹ is = T¹ tph, the capacity T² of the same mill at speed n² should be T² = T¹ (n²/n¹)tph [8] The Relationship between Power Consumption, Mill Diameter, and Speed The only well known theoretical deduction relating power consumption, P, and mill diameter appears to be the formula of duPont introduced by Gow, Guggenheim, Campbell, and Coghill.&apos; According to duPont, the power required to operate a mill is a function of the mass of the balls, of the lever arm of the ball mass, and of the speed of the mill. The ball mass per unit of mill length is proportional to the square of the diameter, the lever arm is directly proportional to the diameter, and the critical mill speed or any percentage thereof is inversely proportional to the square root of the mill diameter. Following this reasoning, the original duPont formula is of the form P = c4D² c D • c6D-0.5 = c7D2.5 [9] If the mill speed in the above equation is expressed in terms of Eq. 3, the duPont formula may be written as follows: P=f1(D2) f2(D) f³(—np) or [10] vD P = c np D2.5 kw [11] Eq. 11 may also be derived from the mechanical principle of force, which is equal to mass x acceleration. Power necessary to operate a mill may be considered to be an homogeneous linear function of the force developed. Ball or rod mass per unit of mill length is a function of D2. The acceleration factor of the ball or rod mass is a function of the peripheral speed of the mill. Thus P = f4(F) = /x(D2) f5(v) Indicating that v = nDn, and n = c9 np /vD, the above equation becomes P = f2 (D2) -fa (D ca np/vD

Jan 1, 1955