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By R. G. Davies

Isothermal annealing of quenched Fe3Al reveals that the superlattice forms by the nucleation and growth of ordered domains. The activation energy for isothermal ordering and initial domain growth is -29 keal per mole, which is thought to he the activation energy for vacancy migration. Domain growth, in the absence of excess vacancies, has an activation energy of -100 keal per mole. A maximum in the hardness is observed when the structure consists Of ordered domains, -35-4 in size, growing into a short-range ordered matrix and, also as a function of degree of order, S, with an infinite domain size at S = 0.4 to 0.5. It is proposed that both maxima are due to a transition from deformation by unit dislocations to deformation by superdislocntions. FE-AL alloys containing more than -23 at. pet Al, when quenched from above the critical temperature for the Fe3A1 type of order, always contain the maximum possible amount of FeAl (B2) type of order' (see Fig. 1). Thus when the Fe3A1 superlattice forms, in a 25 at, pet A1 alloy, it does so from an imperfect FeAl structure. The formation of the Fe3Al superlattice has been the subject of several recent X-ray,1, 2 electrical-resistivity,3"6 and dilatometric3 investigations. McQueen and Kuczynski3 concluded that the isotherma1 ordering of Fe3Al, after quenching from 800°C (T,, the critical temperature for order, is 55O°C), takes place by the nucleation and growth of ordered domains. However, Feder and Cahn4 found that the isothermal ordering was complex but more like a homogeneous nucleation process than a nucleation and growth process. Lawley and Cahn2 reported that the degree of order is a continuous function of temperature which is a characteristic of a homogeneous nucleation process. Thus, the evidence is conflicting as to the nature of the Fe3Al ordering reaction. Lipson7 was the first to suggest that ordered domains should exist in the Fe3A1 structure since, as

Jan 1, 1964

Full details of the program for the New York meeting .were given in the January Bulletin and the announcement mailed to members. The following brief synopsis is offered simply as a reminder of the principal events of this meeting which Chairman Browne says is going to be "different." Time.-Monday, Feb. 14, to Thursday, Feb. 17, 1916, inclusive. Place.-Engineering-Societies' Building, 29 West 39th St., New York. Events.-Monday, Tuesday and Wednesday (Morning and Afternoon) Technical Sessions. Tuesday, 10 a. m.: Annual Business Meeting, Luncheon served in the-Engineering Societies Building between sessions each day.

Jan 2, 1916

By John D. Erickson

The concepts of long-range, open pit planning presented here are not new. The new idea is how the factors that control the location of the final pit limits can be considered altogether as one large, simultaneous equation. The X, or unknown quantity that is being solved for, is a long-range pit design that fulfills the desired objectives. The objectives of a final pit plan are determination of the extent and location of ore reserves, the allowable stripping necessary to uncover the ore, the final pit surface intercept, final pit topography and the contained metal values in both ore and waste.

Jan 4, 1968

By D. G. Swarthout, J. W. Savage, R. W. Guard

The Nb-N constitution diagvam has been examined from 1100° to 1450°C at 1 atm N2 pressure. On the basis of- our work plus that of Brauer and coworkers the stable nitvide Phases are: a) a solid solutlion, basic niobium (bcc) with nitrogen contained interstitially; NbN,, x = <0.40; b) ß NhNx,x = 0.39 to 0.45; hexagonal (W2C type) (NbzN); a =3.058A,c = 4.961 + 0.001A c) y NbNx x = 0.75 to 0.80; tetragonal (distorted NaCl type) (Nb4N3); a = 4.386A,c = 4.332 + 0.001A; d) 6 NbN,, x = 0.88 to 0.91; cubic (NaCl type); a = 4.3885 + 0.0007A; e) E NbNx x = 0.92 to 1.00; hexagonal (Tip type); a = 2.958A, c = 11.272 * 0.0015A. In addition, a metastuble transition phase 6&apos; has been observed and occurs only during the transition 6 — E. 6&apos; NbN has these distinguishing parameters: NbNx,x = 0.96; hexagonal (anti-NiAs type); a = 2.94A, c = 5 .46A. The high-temperature phase fields in the system are governed by two reactions: a) 6 + Y + E, a entectoid reaction occurring at 1270" ± 20° C; b) o + N2 = E, a 2 012 1 reaction occurring at 1370" ± 20° C. As part of a study of the electronic low-temperature. properties of transition metal nitrides, a number of specimens of niobium nitride were prepared. Preliminary observations on these specimens showed several points of disagreement with data already in the literature. It was, therefore, decided to determine the constitution diagram for the Nb-N system between 1100" and 1450°C. The objective of this research was to define the temperature and composition regions of stability of the various phases and to explain the differences among previously proposed diagrams. Review of Previous Literature. The research on the Nb-N system to 1964 has been summarized by storms.&apos; Although he does not present a phase diagram, his review covers the research on the system thoroughly. One of the earliest works on NbN is that of Ascher-mann and coworkers2 who studied the superconductivity of the cubic 6 NbN. Subsequent work by Horn and ziegler,3 Cook, Zemansky, and Boorse,1 Rögener:5 Schroder6 and sellmaier7 has qualitatively confirmed the first results on superconductivity, although their methods of preparation differed considerably. The constitution diagrams proposed for the system by Brauer and Esselborn8 and Elliott9 are shown in Fig. l(a)and (b). The diagram in l(a) has been modified by more recent results in Brauer&apos;s laboratory (private communication). Although 6&apos; is shown on this diagram, Brauer has indicated that he now believes 6&apos; to be a metastable phase in the temperature range shown. The diagram of Elliott, Fig. l(b), is based primarily on the results of Elliott and Komjathy10 on ni-trided niobium strip or rod specimens examined by X-ray diffraction and metallography. It will be noted that these proposed diagrams are in conflict with respect to the temperature and composition ranges of

Jan 1, 1968

By Dr. O’Neil Thomas J., Donald W. Gentry

"October. This is one of the peculiarly dangerous months to speculate in stocks. The others are July, January, September, April, November, May, March, June, December, August and February." -Mark Twain Although the preceding chapters describe the type of preliminary evaluation frequently performed in the mining industry, it has been assumed that input data are known with certainty-clearly an enormous simplification. In reality, estimates of ore grade, mining cost, metal price, etc., are subject to varying degrees of uncertainty due to the inability to predict the future with much precision. Management would probably not, for example, be equally attracted to two projects each offering a 20% DCFROI-one for a new gold venture in Nevada and the other for an undersea manganese nodule operation in the Pacific. The high risk with the latter investment tells the obvious; there is more to investment decision-making than a deterministic analysis using "best estimates" of input parameters. The science of decision theory normally distinguishes between decisions under uncertainty, where probabilities of various outcomes are unknown; and decisions under risk, where such probabilities can be estimated. The first case is rarely en- countered in practical capital investment analysis in mining, so the remainder of this chapter focuses on that class of problems where rational probability distributions for the future states of input variables can be estimated. RISK IN MINING Risk in the context of this chapter is the unforeseen deviation of individual cash flows from expected values for a capital project. For a mining venture, the source of this uncertainty could be any number of factors relating to such items as ore grade, ore reserve tonnage, operating costs, product prices, etc. With conventional *Adapted from O&apos;Neil, T.J., 1979, "Procedures for the Preliminary Financial Evaluation of Metal Mining Ventures," Computer Methods for the 80s, A. Weiss, ed., AIME, New York, pp. 556-573.

Jan 1, 1984

By Chairman Hill H. H.

The Round Table Discussion on Petroleum Engineering Research was planned for two purposes. First, to review the work that has been done by petroleum engineers during the past few years on problems that have confronted the industry and, second, to suggest new studies that should be made either on the problems that have been the subject of research during the past or on problems that have arisen only recently. It was believed that by holding an informal discussion of the problems in which the petroleum engineer is vitally interested it would be possible to bring out new ideas as to the relative importance of the different problems that should be studied and to suggest new methods for attacking some of the problems that have resisted the energies of research workers throughout the industry. It was also believed that by arranging for an exchange of experiences between the different engineers it would be possible to prevent a large amount of duplication of effort in conducting research work. Those engaged in petroleum engineering research realize that there is a broad field for study and by coordinating the work it will be possible to obtain the maximum amount of information in the shortest possible time. In planning the program on Petroleum Engineering Research, nine subjects were selected for discussion and each was assigned to a man who is doing work on that subject. These men were asked to lead the discussion on the problem assigned to each one. An effort was made to choose problems of a practical nature and upon which considerable work is being conducted at the present time. Some of the men who were asked to lead the discussion were unable to attend the meeting but most of them submitted a written discussion- of the subject assigned, with the result that there was a discussion of eight of the nine problems. Professor Uren, who was asked to lead the discussion on the subject of Recent Developments in Methods for Reducing Gas-oil Ratios, was unable to attend the meeting and advised that he did not have any new ideas to contribute.

Jan 1, 1928

By R. I. Jaffee, C. M. Craighead, G. A. Lenning

The a-p type alloys are subject to a loss of tensile ductility with increasing hydrogen content. No hydride phase is visible in embrittled a-B type alloys. The embrittlement encountered appeared to be of the strain-aging type. Both compositional and structural factors are shown to influence the hydrogen tolerance of a-B type alloys. IN previous papers by the authors, the effects of hydrogen on the structure and mechanical properties of high-purity and commercial-purity titanium were described. It was shown that the presence of hydrogen in these materials results in a hydride phase having low solubility at room temperature. This insoluble phase had little effect on the room temperature tensile properties, but did decrease the notch-bar toughness to a large degree. The binary Ti-H systemz indicates that hydrogen is much more soluble in £ than in a-titanium. Hence, it is not expected that the embrittling effects of hydrogen in a-fi alloys will be similar to that found in a alloys. Materials of this type studied in the present investigation included two commercial a-/3 alloys, Ti-8Mn and Ti-4A1-4Mn; three high-purity Ti-Mn alloys containing 3, 6, and 9 pct Mn; and four high-purity Ti-Mo alloys, Ti-SMo, Ti-10.9M0, Ti-13.1 M0, and Ti-20M0, of the a-B and B types. Materials and Fabrication—The high-purity base alloys, composed of iodide titanium and high-purity metals, were prepared by double arc melting under argon in a water-cooled copper crucible, using a water-cooled tungsten electrode. Radiographic examination showed homogeneous Ti-Mo alloys. The analyses of the alloys used in this investigation and their as-fabricated hydrogen content, where determined, are listed in Table I. The two commercial alloys were 1/2 and % in. diam bar stock. No analyses other than hydrogen content were made. In order to maintain the a-B relationship existing in the two commercial alloys, vacuum annealing and hydrogenation were carried out in the B field prior to final fabrication. Fabrication then was done by The three high-purity Ti-Mn alloys were forged to 3/4 in. diam rods at 1600°F, and hot swaged to 1/2 in. diam at 1380°F. Machined and degreased bar stock was vacuum annealed or hydrogenated, fabricated to 1/8 in. diam rod by hot swaging, annealed in argon for 1 hr at 1380°F !75O°C), and water quenched. The six yz lb Ti-Mo ingots were forged to Wi in. diam rods at 1600°F and then hot swaged to % in. diam. Machined and degreased specimens were vacuum annealed or hydrogenated in the /3 field, prior to fabrication, by swaging to Y4 in. rod at 1400°F. After fabrication, the 5 and 10.9 pct Mo alloys were annealed in argon for 16 hr at 1290°F and water quenched to give equiaxed a-B structures. The higher molybdenum content alloys were argon annealed 4 hr at 1470°F and water quenched to give all £ structures.

Jan 1, 1957

By AIME AIME

A COMMITTEE of the American Railroad Association, consisting of Samuel Rea, president of the Pennsylvania System, Chairman; F. D. Underwood, president of the Erie Railroad; A. T. Dice, president of the Philadelphia & Reading Railroad; E. J. Pearson, president of the New York, New Haven & Hartford Railroad; and C. H. Markham, president of the Illinois Railroad (who was represented at the meeting by J. F. Porterfield of his system), met at the headquarters of the American Institute of Mining and Metallurgical Engineers, New York, on April 6 to assist in the Coal and Coke Committee&apos;s attempt to solve the bituminous coal problem. There was a full discussion of the questions of seasonal freight rates and the purchase by the railroads during the slack months of their winter&apos;s supply of coal. As the question of fuel supply for the railroads is in the hands of the Railroad Association, as well as the coal operators, it was vital that it should be represented at this important conference. The situation from the point of view of the mine operators was presented very thoroughly. No seasonal rates were considered necessary for the year 1920, and, as this matter requires considerable thought, it will be given thorough study and worked out in detail at a subsequent meeting.

Jan 1, 1920

The following is an incomplete list of members and guests who called at Institute headquarters during the period May 10, 1919, to June 10, 1919. Victor C. Alderson, Golden, Colo. E. Fleming L&apos;Engle, Joplin, Mo. H. Foster Bain, Washington, D. C. W. D. Mainwaring, Philadelphia, Pa. J. J. Beeson, Master Engr. Sr. Gr. W. G. Matteson, Fort Worth, Tex. Gunnar Bergstrom, Stockholm, Sweden. C. E. MacDonald, Toronto, Can. Corry T. Brown, Socorro, N. Met. Francis Nicholson, Charlotte, Va. Geo. W. Brymer, Scranton, Pa. Richard A. Parker, Denver, Colo. W. Burns, Santa Fe, N. Mex. S. B. Patteson, Jr., Allentown, Pa. H. C. Carlisle, The Pas, Manitoba, Can. G. W. Ponton, Ottawa, Canada. J. Carson Adkerson, Va. W. H. Reber, Ensign, Dallas, Tex. C. A. Cheney, Madison, Wis. T. C. Roberts, New York City. J. M. Clark, Toronto, Canada W. A. Rukeyser, Ensign, N. F. Corps. John J. Doyle, St. Louis, Mo. W. H. Staver, New York City. H. S. Emlaw, Grand Haven, Mich. D. B. Sterrett, Washington, D. C. E. Francis, Birmingham, Ala. T. V. Smirnoff, Russia. Rudolf Gahl, Denver, Colo. E. C. Templeton, New York City. Emit Gthmann, Baltimore, Md. H. A. Toblemann, Ajo, Ariz. A. B. Hardie, Philadelphia, Pa. Priestley Toulmin, Birmingham, Ala. Clarence M. Haight, Franklin, N. J. M. R. Walker, Baltimore, Md. N. V. Hansell, New York City. H. A. Wheeler, St. Louis, Mo. C. A. Heberlein, Nacosari, Mexico. D. B. Whitaker, Santiago, Cuba. Pierre Hines, Milwaukee, Wis. A. W. G. Wilson, Ottawa, Can. J. S. Hollings, Brymoba, Wales. Oba Wiser, Hanover, N. Mex C. Vey Holman, Maine. Harry J. Wolf, Denver, Colo. H. H. Howry, Denver, Colo. F. H. Wood, Mt. Vernon, N. Y. W. Spencer Hutchinson, Boston, Mass. J. P. Wood, New York, N. Y. W. B. Kunhardt, Reading, Pa. Lawrence Addicks sailed on May 17 for London; lie expects to be absent from this country for several months. Roy H. Allen has accepted a position in the research engineering department of the National Aniline & Chemical Co., Buffalo, N. Y. H. Foster Bain has gone to Shanghai, China, where he expects to remain for some time. Percy J. Bandy has accepted a position with the Cia. Minera del Mirasol S. A., at Cusihuiriachic, Chihuahua, Mexico. J. H. Batcheller is now with the North Midas Copper Co., St. Helena, Alaska.

Jan 7, 1919

THE iron and steel industry spent $1,170,000,000 in 1952 for improvements and expansion, with an accompanying rise of 9 million net tons in capacity to 117.5 million tons of ingots and steel for castings. More than $4,600,000,000 has been invested in the seven years following World War 11. An average of $1260 has been laid out every minute during those seven years for an increase in capacity of 25.6 million tons.

Jan 3, 1953

By C. H. Herty

During the past few years frequent reference has been made to a certain type of iron known generally as "bad iron," "dirty iron," etc., both by steelmakers and producers of iron castings.&apos; Open-hearth operators have claimed that such iron was entirely within the specified limits as to ordinary analysis but it was difficult to make good steel from it. It has been suggested that unreduced material in the iron is responsible, also that most trouble is experienced with the iron when the blast furnace is delivering a product whose analysis is more or less variable. Other reasons suggested are the use of scrap or borings, unsintered flue dust in the burden, a cold hearth, too high a proportion of unsuitable grade of ore, or the form of the carbon in the iron. Accordingly, a subcommittee was appoint at the fall meeting of the Open-hearth Committee of the Americ m Institute of Mining and Metallurgical Engineers held in Pittsburgh in October, 1926, to investigate the reasons for the occurrence of this bad iron and suggest methods for preventing its appearance. This committee enlisted the cooperation of the U. S. Bureau of Mines Experiment Station at Pittsburgh in undertaking a study of pig-iron quality as affecting the open hearth. The following is a report of the findings to date: Method of Testing Inasmuch as a "spell" of dirty iron occurs without warning and may be of short duration, it was deemed advisable to attempt to simulate the conditions existing before and during such a period and compare the quality of the steel made with the normal and with the dirty iron. The Bourne-Fuller Co., of Cleveland, which has a single furnace that supplies

Jan 1, 1929

By C. M. Marquardt

The minerals industry in general should bow in homage to the uranium milling industry. Those in the uranium milling industry have "spark- plugged" more progress in the application of instrumentation and control in the last 5 years than has been made in 50 years by the industry in general. Part of this real progress in the application of instrumentation and control has been born of necessity because milling processes are more complex. A lot of the impetus for instrumentation and control has come from the fact that the industry is new, the processes are new, and there are no hide-bound conventions to overcome. Part of it comes from the fact that much of processing is more chemical than metallurgical and the thinking of the process chemist has carried over into the plants. One need not elaborate on what instrumentation has done for the process chemical industries. Each new mill that is constructed, if of appreciable size, seems to employ more instrumentation than the previously completed plant. Plant A makes a break-through in the application of instrumentation to a particular part of the process and Plant B accepts this and goes one step farther, while Plant C, anticipating Plant B, makes two or three more steps before Plant B is entirely leveled out.

Jan 9, 1958

By M. E. Fine, A. A. Hendrickson

The critical resolved shear stress and the strain rate dependence of the .flow stress are reported for Ag base A1 single crystals up to 6 at. pct A1 over a temperature ralzge of 4.1° to 470°K. At room temperature the solid solution strengthening is attributed brinciballv to an increase in the dislocation density on alloying. At 42°K the alloys are relatively stronger because the dislocations widen on alloying and are more difficult to cut. In a previous paper, the authors reported the characteristics of strain aging in Ag base A1 alloys.&apos; The variation of the magnitude of the strain-aging yield drops on aging time, aluminum content, and testing temperature strongly indicated that the phenomenon is due to the Suzuki mechanism. In the present paper, the critical resolved shear stress (CRSS) and the strain rate dependence of the flow stress are reported for Ag base A1 single crystals as functions of testing temperature (4.2" to 470") and A1 content (0 to 6 at. pct). Observations on the substructure and measurements of the effect of crystal growth rate are also given. The experimental re- sults appear explainable considering that aluminum additions increase the dislocation density and dislocation width. EXPERIMENTAL METHOD The methods of specimen preparation and tensile testing were described previously.&apos; Briefly, single crystals were grown from the melt in a vacuum (1 x lom5 mm Hg), shaped into tensile samples by a combination of abrading and acid etching, annealed at 850°C in a vacuum (1 x 10" mm &Z) for several days, and furnace cooled to 200&apos;. The testing was done on a model TM Instron using friction grips and a strain rate of about 5 x 106 per sec unless otherwise noted. The various testing temperatures were obtained by immersing the specimens in the following media: 1) 4.2&apos;: liquid helium using the apparatus of G. Byrne;&apos; 2) 77°K: liquid nitrogen; 3) 200"K: dry ice and acetone; 4) 415&apos; K: ethylene glycol; 5) 460&apos; \ to 480&apos; : molten Stanalax. In the determination of the CRSS above 296&apos; K, a ratio method was used; the crystals were strained initially at the higher temperature, 0.1 pct shear strain, and then tested at 296&apos; . The ratio of the flow stress between the two temperatures was then multiplied by the average CRSS at 296&apos; K to obtain the value of the CRSS at the high temperature. (The test at 296&apos; showed a yield point resulting from strain aging; the flow stress at 296&apos; K was taken as the lower yield stress.) The ratio method, we believe, is more accurate than using different crystals since

Jan 1, 1962

By J. W. Hickman

INTRODUCTION MOST of the electron diffraction studies of oxide films which form on metals and alloys have been carried out by oxidizing the specimens in an auxiliary furnace, cooling down to room temperature and then inserting the sample in the camera and taking a diffraction photograph. Prior to the investigations at the Westinghouse Research Laboratories, no systematic study of the oxidation of any metal or alloy had been made over a wide temperature range with photographs taken at elevated temperatures except the investigation of iron by Jackson and Quarrell.1 In order to interpret the electron diffraction patterns of the oxides which may form on alloys, it seemed advisable to make a comprehensive study of the oxidation of the purest metals available in the bulk form. Thus far studies have been completed on iron,2 cobalt,2 nickel,2 chromium ,2 Copper2, tungsten,3 molybdenum,3 beryllium ,4 uranium,4 thorium4 titanium5 and zirconium5 while investigations are in progress on tantalum and columbium, The next step in the study is concerned with the investigation of the oxides which form on binary alloys. Some of the alloys consist of a single, phase while others contain more than one phase. The identification of the oxides on the surface of a binary alloy may be more difficult than on the metals since oxides containing both metals as well as those of each metal separately may form. Furthermore, it is possible that solid solutions of the various oxides may occur. In these cases the electron diffraction method may not be able to identify the oxidation products unequivocally, Examples of solid solutions which may present difficulties are MoO2 and W02, Fe3O4 and FeO-Cr2O3 since the diffraction patterns are so similar. Since electron diffraction can determine only the structure of the oxide, the chemical composition must be inferred from the diffraction data based upon X ray diffraction investigations of specimens of known composition. It is possible in some cases to decide whether a given solid phase reaction can occur on the surface of an alloy by a study of available thermodynamic data. Table I lists the equations for the formation of most of the oxides from the elements together with values of log10 KR for the several temperatures. [OF°R = -RT In KR 0 and-~75T = logi°KR] [?F°R] and KR are the free energy and equilibrium constant of the reaction. Calculations may be made using the values given in Table I to determine whether a given solid phase reaction is possible thermodynamically. For example, a solid phase reaction should occur between aluminum and chromium&apos; oxide to form

Jan 1, 1948

By D. R. Shreve, L. W. Welch

An analysis for predicting the behavior of reservoirs exploited by a combined gas drive and gravity draintrge mechanism is presented. The method allows the prediction of gas-oil ratio, oil production rate, and cumlulative oil recovery for high relief pools in which pressure is maintained either by gas injection in the crest of the structure or by a combined gas and water injection program. The reservoir is represented by a model of communicating vertical prisms, with production or injection at the upper or lower boundaries of each prism. The theory of macroscopic gravitational segregation of two fluids within a reservoir is discussed in detail for all ranges of saturation. A description of an application of the analysis to a typical field problem is also presented. INTRODUCTION The importance of gravity drainage as a producing mechanism has become more evident during recent years. However, many of the reservoir engineering methods which have been used in the past for predicting the behavior of reservoirs producing by this mechanism have required such drastic, simplifying assumptions that the calculated performance in some cases does not resemble the actual pool behavior. The mechanism of fluid displacement in porous media has been described by Buckley and Leverett,1 and a simplified "one zone" method based on this development has been used by engineers to evaluate the performance of gravity drainage pools. Welge2 has modi- fied this approach so that the calculation time has been markedly decreased. In addition, many other authors3-9 have analyzed various phases of the gravity drainage producing mechanism during the past several years. In a high relief reservoir in which gravity may affect the recovery mechanism, there are two main types of exploitation programs which can be employed: 1. Reservoir pressure is maintained by injecting gas into the upper portions of the reservoir. In this case gas is not released from solution and, if dynamic pressure drops are small compared to the total pressure of the system, the gas and oil may be considered to be immiscible and to have constant fluid characteristics. 2. Reservoir pressure is not maintained and a "natural depletion" mechanism prevails. The term "natural depletion" can be used to describe the production history of a reservoir in which a gradual decline of average pressure and the associated release of solution gas continue throughout the producing life. In this mechanism gas is released throughout the reservoir and there is a continuous change in the characteristics of both the gas and the oil phases. It is the purpose of this paper to present the results of mathematical developments by which a method has been devised for evaluating the performance of reservoirs which are depleted by the first type of exploitation program discussed above. It should be mentioned that, although throughout the analysis gas and oil represent the two fluids existing in the reservoir, the method is equally applicable to gravity drainage in water-oil systems. STATEMENT OF PROBLEM In reservoirs of pronounced structural relief, where the closure of the oil section is great or the angle of

Jan 1, 1957

By George A. Roberts, Arthur H. Grobe

Tensile, hardness and density properties are presented for a new 18-8 stainless steel powder for the —50, —100, and —140 mesh cuts and also for a prepared blend containing 62 pct —325 mesh powder. The data were obtained for sintering temperatures of 2100° to 2350°F and for compacting pressures of 30 to 50 tons per sq in. A NEED has existed for a stainless steel powder of uniform composition from particle to particle, which could be readily molded and sintered, and which would yield moderate to high strength and ductility when processed into parts by standard powder metallurgy techniques. The literature reveals that previously available powders have been lacking in one or several of these desirable characteristics. Originally attempts to produce stainless steel powder by blending and diffusing elemental metal powders were unsuccessfu1.l A truly alloyed stainless steel powder reported by Dale in 1947 yielded reasonable mechanical properties but required an exceptionally high sintering temperature of 2400°F in pure dry hydrogen.&apos; A later development reported in 1950 by Stern provided a stainless type powder of excellent molding properties&apos; but this powder likewise required sintering at temperatures over 2350 °F and did not yield properties as high as those previously reported. The powder was not truly alloyed prior to sintering. A prealloyed stainless steel powder has been developed that is readily moldable at 30 to 50 tons per sq in., can be sintered at low temperatures (2100" to 2300°F), and which yields compacts with both high strength and ductility. The powder is prepared by the liquid disintegration method," direct from molten metal. Early attempts to make 18 pct Cr, 8 pct Ni stainless steel powder by this method did not produce a usable powder for pressing and sintering because of a preponderance of spherical particles with relatively high, surface oxide content. It was learned that silicon additions to iron alloys, powdered by this process, gave bright powder of low-oxide content and such additions were made to 18-8 stainless with similar results. An attendant change in particle shape occurred reducing the number of spheroids materially. The optimum composition which has been studied and is reported here contains approximately 2.5 pct Si. Silicon additions to stainless steel of the 18-8 type improve the resistance to high-temperature oxidation and promote the formation of ferrite. Corrosion resistance is, however, controlled by chromium content and even though ferrite may exist the corrosion resistance is not generally impaired.4 Experimental Procedure A detailed study of the 18-8, 2.5 pct Si powder has been made in accordance with the following outline: 1. Four blends a. — 50 mesh b. —100 mesh

Jan 1, 1952

By R. L. McCarron, G. R. Belton

The diffusivities of oxygen and sulfur in liquid iron have heen determined hy a capillary technique in which the surface concentrations of the solutes were established by means of appropriate H2/H2 and H2S/H, gas mixtures. Total diffusate and concentration profile results are shown to be in good accord, yielding for- 1560 and Supporting results at 1660°C are also presented. The conditions necessary to avoid gas transport control in this type of experiment are discussed. IN spite of their importance in understanding the kinetics and mechanisms of refining reactions, the dif-fusivities of oxygen and sulfur in liquid iron are not well established. Accordingly, as a first step in studies of rates of solute absorption from the gas phase into liquid iron, new measurements of these diffusivities have been made and are presented in this paper. The only published results for the diffusion of sulfur in pure liquid iron are those of Kawai.&apos; He used a diffusion couple technique in which two cylindrical specimens, one containing sulfur and the other with negligible sulfur concentration, were joined together and held in a refractory capillary. After an experiment, the sample was quenched and the concentration distribution of the solute determined. Kawai recognized that significant changes in solute distribution occurred during melting and freezing and he attempted to correct the concentration profiles for these effects to give a sulfur diffusivity of 4.6 x 10-6 sq cm per sec at 1560°C. The method of correction, however, was not rigorous and the uncertainty in this result cannot be easily assessed. Koslov et a1.2 have reported the diffusivity of oxygen in iron as 7.8 x 10"3 sq cm per sec at 1660°C. This value appears to be unreasonably high but, unfortunately, details of their experiments are not available. Shurygin and Kryuk have used the rotating disc method for a study of oxygen diffusion in liquid iron. In their experiments a silica disc was rotated in liquid iron containing oxygen, and the rate of formation of liquid iron silicate was determined by measuring the decrease in weight of the disc. On the assumption that the rate of dissolution was controlled by the diffusion of oxygen in the iron, the diffusion coefficient was computed to be 5.2 x sq cm per sec at 1550°C. However, the Levich equation, which was used to interpret the rate data, was originally de- rived for the case of mass transfer between a solid disc and a single-phase liquid. The hydrodynamic and diffusion boundary layers in the iron stirred by a disc, via coupling of the silicate melt, may be appreciably different from those predicted by Levich&apos;s equations. Recently, Novokhatskiy and Ershov, using an identical experimental method to that of Shurygin and Kryuk, obtained a diffusivity for oxygen in liquid iron of 1.22 x 104 sq cm per sec at 1550°C: no reasons were offered for the disagreement. Schwerdtfeger5 has also recently studied the diffusivity of oxygen in liquid iron. He reacted shallow melts of liquid iron, 0.5 to 1.0 cm deep and contained in high-purity alumina crucibles, with appropriate H20-HZ-He mixtures. The total sample was analyzed, without sectioning, to obtain the average concentration of diffusate. A value at 1610°C of D = 12(3) x 10-5 sq cm per sec was obtained from the results of twenty experiments.= Oxygen profile measurements, which were carried out in three additional experiments using long capillaries and the semiinfinite boundary conditions, indicated a diffusivity about half that computed from the shallow bath experiments. Possible sources of error in Schwerdtfeger&apos;s study will be discussed later in this paper. EXPERIMENTAL TECHNIQUE The essential arrangement of the diffusion cell is shown in Fig. 1. The liquid iron was contained in an alumina capillary, 3 to 4 mm diam and 3 to 9 cm long, which was supported by a hollow alumina pedestal and this whole assembly was held within a movable alumina reaction tube. This tube, which was about 7 mm in bore

Jan 1, 1970

By B. J. Bayles, J. A. Ford, M. J. Salkind

Unidirecltonally sulidijied AL-A13Ni was found to exliibit excellent microstructural stability to 508°C. Above this temperature the transverse micvostructure coarsened in a manner analogous to "Ostwald Ripening", but the A13Ni whiskers did not spheroidize or foreshorlen. The coarsened whiskers reinforced the matrix as effectirely as the finer whiskevs and no reduction in room-temperature tensile strength mas noled after exposure for 90 hr at 96 pel 0-1 the absollite melling temperature of the eulectic. The slability was auribucted to the presence of low-index (and presurrlably lor-energy) cvystallographic interjaces between phases. The driving force for microstructural coursening was presumed to be the reduction of interfacial energy between the phases. Finer micvoslructures were found to coarsen more rapidly than those which weve imitally coarser. FIBER-reinforced metals have been shown to exhibit good short-time tensile properties at elevated temperatures.1-4 However, many elevated-temperature applications require long-time property retention which dictates the necessity for microstructural stability. Two types of instability are common to fiber composities. The first, chemical instability, is due to chemical reaction between the matrix and the relnforcing phase. Petrasek and weeton5 have shown that chemical reaction between tungsten wire and several copper alloy matrices resulted in degradation of the reinforcing phase and reduction in strength of the composite materials. Kreider and Leverant4 have noted similar behavior for boron-reinforced aluminum above 400°C. The second type of instability arises in systems in which the phases are essentially chemically stable with respect to each other and is characterized by spheroidization and/or agglomeration of the reinforcing phase. parratt6 has reported this type of degradation to occur relatively rapidly at moderate temperatures for nickel and cobalt alloys containing whiskers of silicon nitride, aluminum oxide, and silicon carbide and has identified it as "physicochemical" instability. Although both types of instability are the result of the differences in the chemical potential between the fibrous phase and the matrix, they may be differentiated by the fact that the second type, "physicochemi-call&apos; instability, results in little or no net change in the composition or amount of each phase present. It does, however, result in a morphological change which profoundly affects the strengthening mechanism. Chemical interaction. on the other hand, results in a change in the composition and amount of one or more of the phases present, or the appearance of a new phase. In the latter case, the appearance of the mi-crostructure, that of a fibrous phase aligned in a matrix, may remain essentially unchanged, but the chemical degradation of the fiber results in markedly reduced composite strength. This investigation was initiated in an attempt to evaluate the microstructural and related mechanical property changes of a whisker-reinforced eutectic composite after elevated-temperature exposure. The composite system chosen for this study was the uni-directionally solidified A1-A13Ni eutectic which Lemkey. Hertzberg, and Ford7,8 demonstrated to behave as a whisker-reinforced composite. The microstructure of the unidirectionally solidified material consists of aligned needles of A13Ni in a matrix of aluminum. EXPERIMENTAL PROCEDURE Eutectic ingots having a nominal composition of 6.2 wt pct Ni were produced by melting 99.99+ pct A1 and 99.99 pct Ni in recrystallized alumina crucibles in an argon atmosphere. The analyses of the starting materials are listed in Table I. These master heats were cut up? melted by induction heating, and unidirectionally solidified vertically in 1/2 -in.-ID. 6-in.-long graphite crucibles, in a dynamic argon atmosphere. Therma1 gradients in the liquid were determined by measuring the temperature in the melt as a function of apparatus travel distance, as discussed by Kraft and Albright.9 The rate of solidification was assumed to be equal to the uniform rate at which the crucible assembly was withdrawn from the furnace. Ingots representative of nominal growth rates of 2, 5, and 11 cm per hr were produced and the solidification parameters are summarized in Table 11. Ingots used for met allographic analysis were pre-

Jan 1, 1968

By R. V. Ramani, Y. J. Wang, R. Mishra

Abstract-The validity of the concept of mine resistance in ventilation systems is examined. Mine resistance is the proportionality constant relating the head generated by a fan and the quantity flowing through it when the fan is connected to a mine circuit. It is the total effect of all branches in the mine ventilation circuit. Its importance lies in the fact that for new mine projections, and in the extensions of existing mines, ventilation calculations are usually based upon adjustments of the mine

Jan 11, 1978

By G. Derge, Ling Yang, John Henderson

Self-diffusion coefficients of aluminum have been measured by the capillary reservoir technique in liquids of the CaO-SiO2-Al2O3 system containirg equimolar portions of CaO and SiO2, in the temperature range 1400o to 1520oC. For the melt containing 6.0 mole pct Al2O3 the results can be expressed by the equation D = [(4.3 * 0.3) x 104] exp [(-85,000 * 17,500)/RT]cm~sec-&apos;; for the melt containing 12.5 mole pct Al2O3, by the equation D = (5.4 i 0.2) exp [(-60,000 * 10,00O)/RT] cm2sec-: A concept of the constitution of these melts has been developed which proposes that the ratios O/(Si +Al) and Al2 O3/CaO determine the dominant species present. In recent years considerable work has been done on the system CaO-SiO2-A12O3 with a view to elucidating the structure of liquids of this system. Conductivity,&apos; electromotive force, 3 density,4 and activity5 measurements have all contributed to an understanding of the problem, but as yet there is no entirely satisfactory theory that can account for all the observed experimental results. In principle, measurements of self-diffusion coefficients should give further information which is not obtainable in other ways. Self-diffusion coefficients for calcium, silicon,6 and oxygen7 have been measured. The present investigation extends the series of self-diffusion coefficient measurements to include aluminum and develops a concept of constitution for liquids of this system which is compatible with all the observed experimental results. EXPERIMENTAL The technique used in the measurement of the self-diffusion coefficient of aluminum was the capillary-reservoir method,8 A126 being incorporated in the melt contained in the capillary and inactive aluminum in the reservoir melt. This method was chosen because it minimizes convection currents&apos; in the diffusion zone. It has the further advantage that the boundary conditions are such that a slight modifica- tion of Anderson and saddington&apos;s8 solution of Fick&apos;s Law yields an equation which requires measurements only of the length and the initial and final average specific activities of the solidified melt within the capillary. where t is the diffusion time, Co is the average specific activity of the melt contained in the capillary at t = O, C is the average specific activity of the melt contained in the capillary at t = t, 1is the length of the capillary, and D is the self-diffusion coefficient. The equation is of somewhat different form than those previously applied to high temperature diffusion studies because the specific activity of the melt in the reservoir was negligible both before and after diffusion. The diffusion apparatus is shown schematically in Fig. 1. It consisted of a silicon carbide resistance furnace containing the inactive melt in a graphite crucible, internal dimensions 7/8 by 4 in., enclosed in a 1 1/4- by 1 1/2- by 30-in. McDanel tube. A 3/8-in. McDanel thermocouple sheath to which graphite radiation shields and a graphite protection tube were affixed, was used to hold the capillaries. By moving this sheath through a rubber seal at the top of the McDanel tube the capillaries could be lowered into or raised out of the melt. Temperature was measured with a Pt-Pt 13 pct Rh thermocouple which was calibrated periodically against a standard couple. The furnace temperature was controlled to ± 5°C through the thermocouple immersed in the melt. Temperature variation over the depth of the melt was less than ± 3°C. The inactive melts were prepared from pure native quartz, assaying better than 99.9 pct SiO2, A.R. grade A12O3 and A.R. grade CaCO,. The same quartz and CaCO3 were used in preparing the active melts, the

Jan 1, 1962