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By C. H. Samans, G. F. Tisinai, J. K. Stanley

The addition of nitrogen (0.10 to 0.20 pct) to Fe-Cr-Ni alloys of simulated commercial purity results in a real displacement of the u phase boundaries to higher chromium contents. The effect is small for the (Y + s)? boundary, but is pronounced for the (y + a +s)/(y + a) boundary. Although there is an indication of an exceptionally large shift of the n boundaries to higher chromium contents, especially in steels with nitrogen over 0.2 pct, the major portion of this apparent shift results from the fact that carbide and nitride precipitation cause "chromium impoverishment" of the matrices. The effect of combined additions of nitrogen and silicon to the Fe-Cr-Ni phase diagram is demonstrated also. Nitrogen can nullify the effect of about 1 pct Si in shifting the (y + o)/? phase boundary to lower values of chromium at all nickel levels from 8 to 20 pct. NItrogen can nullify this U-forming effect of about 2 pct Si at the 8 pct Ni level, but not at the 20 pct Ni level. The alloys studied were in both the cast and the wrought conditions. There are indications that the u phase forms more slowly in the cast alloys than in the wrought alloys if both are in the completely austenitic state. The presence of 6 ferrite in the cast alloys accelerates the formation of U. Cold working increases the rate of o formation in both cast and wrought alloys. THE major improvement in Fe-Cr-Ni austenitic alloys in recent years has been in the addition or removal of minor alloying elements to facilitate better control of corrosion resistance, sensitization, and heat resistance. One shortcoming of the austenitic Fe-Cr-Ni alloys, which never has been completely circumvented, is their propensity toward u formation. In the AISI-type 310 (25 pct Cr-20 pct Ni) and type 309 (25 pct Cr-12 pct Ni) steels, sufficient amounts of u phase can form, if service or treatment is in a suitable temperature range, to cause severe embrittlement. Also, there is a growing conviction that this phase may be contributory to some unexpected decreases in the corrosion resistance of certain of the 18 pct Cr-8 pct Ni-type steels. The present paper discusses the effect of nitrogen additions on the location of the (r+u)/d and the (y+a+u)/(y+a) phase boundaries in the ternary Fe-Cr-Ni system, for cast and wrought alloys of simulated commercial purity, and in similar alloys containing up to about 2.5 pct Si. The objective is to define compositional limits for alloys which will not be susceptible to u formation when used near 1200°F (650°C). An excellent review of the early studies of the u phase in the Fe-Cr-Ni system has been compiled by Foley.1 Rees, Burns, and Cook2 have determined a high purity phase diagram for the ternary system, whereas Nicholson, Samans, and Shortsleeve3 are- stricted themselves to a portion of the simulated commercial-purity phase diagram. Both groups of investigators show almost an identical position for the commercially significant (y+u)/y phase boundary. Further comparison of the work of the two groups indicates that, below the 8 pct Ni level, the commercial alloys have a decidedly greater propensity toward u formation than the high purity alloys. The two groups of workers agreed that both the AISI-type 310 (25 pct Cr-20 pct Ni) and the type 309 (25 pct Cr-12 pct Ni) steels are well within the (y+~) region and that the 18 pct Cr-8 pct Ni-type alloys straddle the U-forming phase boundaries. Nicholson et al.3 showed, in addition, that these boundaries shift toward lower chromium contents if greater than nominal amounts of silicon or molybdenum are added. The effect of nitrogen on the location of the s phase boundaries in the Fe-Cr-Ni system has not been known with any certainty. In 1942, an approach to this problem was made by Krainer and Leoville-Nowak,' but at that time they apparently were unaware of the slow rate of s formation in strain-free samples and aged their samples for insufficient times, e.g., 100 hr at 650°C (1200°F) and 800°C (1470°F). For this reason, it would be expected that their (y+ u) /y boundary would be shifted toward lower chromium contents (restricted ?-field) when equilibrium conditions were approximated more closely. Procedure for Studying the Alloys The alloys used were prepared in the following way: Heats of 200 lb each were melted in an induction furnace. A 5 lb portion of each heat was poured into a ladle containing an aluminum slug for de-

Jan 1, 1955

By W. L. Dumoulin

A RATHER detailed description of the entire plant and leaching process was given in a paper recently presented to the Institute,1 so this paper will cover briefly only the crushing practice of the New Cornelia Copper Co. for the year, 1918. The ore, which is mined by steam shovels and loaded in side-dump cars, passes through two crushing plants. These crushing plants were constructed to crush the ore required by a leaching plant of 5000 tons daily, capacity, during a crushing period of 16 hours. The primary plant reduces to less than 3 in. (76 mm.) and the secondary plant, to ¼ in. (6.35 mm.), which is sufficiently fine for satisfactory percolation of solution through the ore in the leaching tanks, and to give good extraction. The ore is then conveyed to the leaching tanks by a system of 28-in.. (71-cm.) belt conveyors. On the way to the leaching tanks, it passes through an automatic sampling plant, where a 1-per cent. sample is taken. From the primary crushing plant, the ore is conveyed by two 36-in. (91-cm.) belt conveyors to a 10,000-ton steel storage bin, with a reinforced-concrete flat bottom. The ore discharges from the bottom of this bin onto four 20-in. belt conveyors, which deliver it to the four units of crushers in the secondary crushing plant. There is no storage bin between the mine and the primary crushing plant. The ore breaks very coarse, is very hard, and contains a great many boulders, some as large as 4 ½ by 4 1/2 by 10 ft. (1.3 by 1.3 by 3 m.). Jams form in the bowl of the coarse crusher in the primary crushing plant, as a consequence, and are freed by means of an immense steel hook operated from a 40-ton electric traveling crane.

Jan 8, 1919

By Henry N. McCarl

Mineral aggregates are those natural and manufactured industrial mineral and rock materials that provide bulk and strength in port- land cement concrete, bituminous concrete mixes, and plaster or stucco surfaces. They may also provide special characteristics such as thermal and acoustical insulation, weight, surface textures, abrasion resistance, and impermeability to various concrete products and mixes. Most aggregates are relatively inexpensive materials, but despite their low unit value, the tonnages consumed each year in the United States and throughout the world make mineral aggregates by far the leading industrial mineral and rock materials both in terms of volume (tonnage) produced and value consumed. The major use of mineral aggregates is in concrete. The low cost, high bulk aggregates tend to keep concrete construction costs at levels competitive with other building materials. Other major uses of aggregates include highway and railroad base or ballast materials, graded fill, and various industrial uses such as metallurgical and chemical fluxes and raw materials-these latter uses are often considered nonaggregate uses of mineral aggre- gate materials. The value of mineral aggregate materials at the production plant varies from less than a dollar per ton to several dollars per ton for most crushed stone and sand and gravel, $2 to $15 per ton for slag and expanded clay or shale aggregate, to as much as $60 to $100 per ton for expanded perlite or vermicu- lite. The more specialized the use, the higher the price-although there are naturally some exceptions to this rule of thumb. The first four parts of "Construction Materials" deal with the utilization of four major mineral aggregate types: Crushed Stone, Light- weight Aggregates, Sand and Gravel, and Slag. While the focus will be on the use of these industrial minerals and rocks there will also be some coverage of sources, mining and production, especially in those cases where they are

Jan 1, 1975

Jan 1, 1968

By M. Cohen, P. E. Beaubien, D. Caplan

Addition of 1 pct Mn to Fe-26 CY ca/(ses a12 increase in scaling rate at 870° and 1090°C. Whereas only the rhombohedral oxide, formrs on tire manganese-free alloy, with manganese present major amounts of the spinel .Mn0 . Pro-duced through which diffusion is more rapid. At 1090°C the scale formed OH Mn is highly, irregular due to blistering and sintering, This is a second cause of faster oxidation Since much of the blistered oxide does not participate as part of the protective layer. The oxide phases that form at different times depend on the degree to which the surface metal has become depleted, in manganese by preferential oxidation. PREVIOUS work has demonstrated that the oxide scales formed on chromium steels at high temperature consist of both spinel and rhombohedra1 phases and may contain considerably more manganese, especially in the spinel phase, than is present initially in the alloy.' It has further been observed that a binary Fe-26 Cr alloy scaled less rapidly than an equivalent commercial steel and formed no spinel.' However, it is not established if manganese in chromium alloys actually promotes formation of spinel nor if it affects oxidation resistance. It seemed expedient to try to evaluate the specific effect of manganese by oxidizing an Fe-26 Cr alloy containing a small manganese addition under the same experimental conditions as used for the manganese-free alloy.' EXPERIMENTAL Table I shows the composition in weight percent of the Fe-26 Cr-1 Mn alloy and of the Fe-26 Cr alloy used as reference. The alloy with manganese was prepared from the vacuum-melted Fe-26 Cr alloy by remelting in vacuum with the proper manganese addition. The ingot was rolled to strip and specimens 1.1 by 0.035 by 5 cm and 1.1 by 0.5 by 4 cm cut out or machined. As previously described2 the specimens were smoothed by abrading through 600-grit Sic paper, electropolished in acetic-perchloric acid (20 to 1) to remove 10 p of metal, annealed at 1100°C in 20 torr of purified argon, and electropolished to remove a further 10 p of metal. They were then either rinsed in water and methanol and blotted dry or immediately etched for 45 sec in deoxy-genated 4 N HC1 under cathodic polarization at 10 pa per sq cm. Oxidation runs were carried out in a vertical tube furnace in a flow of purified oxygen at 1 atm while the weight increase was recorded continuously on a* automatic balance.3 Alter 98 hr at 870.C or 20 hr at 1090"C the specimens were removed to flowing argon and observed microscopically for evidence of spalling as they cooled. Subsequently they were examined by X-ray diffraction, the oxides analyzed chemically, and metallographic cross settions prepared. RESULTS Figs. 1 and 2 show the oxidation curves of Fe-26 Cr-1 Mn at 870' and 1090C. The Fe-26 Cr reference runs are shown as dashed curves. Experimental details of the six runs are included in Table 11. The first 20 hr of run 1 is repeated in Fig. 2 to show the effect of temperature. At 8'70°C the weight gain for Fe-Cr-Mn is about twice that of Fe-Cr (curve 1 vs curve 2, Fig. 1). Both alloys show regular oxidation curves. On parabolic coordinates (right-hand scale of Fig. 1) the

Jan 1, 1965

By Nicholas J. Grant, Bill C. Giessen

The phase diagram Ta-Ni has been treated repeatedly; investigations up to 1958 are summed up in Ref. 1. Since then, an equilibrium diagram has been presented by Kornilov and Pylaeva.2 They found the following intermediate phases: Ta2N; (CuA12 type),3 "TaNi" (p phase, Mo8Co7 type), TaNiz, and TaNi3 (Ticus type).5 Of these, the crystal structure of TaNin remained unknown, and no further literature reference was found. According to Kornilov and pylaeva,&apos; TaNil melts peri-tectically at 1420°C. A 5-g alloy of Ta + 66.7 at. pct Ni was prepared by inert-gas arc melting using 99.95 pct Ta rod (NRC-MG) and 99.9 pct Ni shot (Fisher). The weight loss after three melts was <2 mg; a microscopically homogeneous ingot resulted. As shown in Ref. 6, no oxygen or impurity pickup needed to be considered under these conditions. Specimens were annealed for 15 hr at 1350°C, 15 hr at 1205oC, and 15 hr at 875°C with subsequent powdering and stress relief for 10 min at 1350°C, 15 min at 1205oC, and 1 hr at 875oC, respectively. Micrographs and X-ray analysis (114.6 mm Norelco camera, filtered CuK, radiation) showed homogeneous one-phase alloys in the as-cast state and at all three temperatures, all having the same powder pattern except for broad lines in the as-cast state. From this it is concluded that either TaNi, melts congruently or its melting behavior is a transition between the eutectic and peritectic cases. Since it was found that very slight surface contamination on stress relieving can considerably alter the structure of similar phases (e.g.,TaPt3),7 patterns before and after stress relief were also compared; they were identical except for line broadening in the non-stressrelieved condition, and two faint impurity lines in the stress-relieved samples. The powder pattern is presented in Table I. All lines were accounted for by indexing with a tetragonal lattice, 0Yh- I4/mmm , with a, = 3.154 ± 0.001A, co = 7.905 * 0.002i and c/a = 2.506. The atomic positions are: 1000; 1/2, 1/2, 1/2] + 2Ta(a): 000; + 4Ni(e): *(002); z = 0.333 * 0.002 This proves TaNi2 to be a representative of the MoSiz (Cllb) type. The variable parameter was determined by grouping the visually estimated intensi- ties according to their 2 indices, plotting their F values, and varying z for best fit, which was obtained at z = 0.333= 1/3. Since absorption was not considered, only the relative intensities with similar 28 can be expected to agree. The interatomic distances are presented in Table 11; the standard error is *o.o2A, resulting chiefly from the uncertainty in z. These distances indicate a coordination number 10 for both kinds of atoms and show that the pseudohexagonal layers along (110) are rather un-distorted. The mean atomic volume of 13.1A3 is close to the Vegard&apos;s law value of 13.3A3, showing

Jan 1, 1964

By L. F. Stutzman, George Thodos

A new graphical method, which is a modification of that proposed by Buckley and Leverett&apos;, is presented for the determination of the displacing fluid saturation at breakthrough for frontal drive mechanisms. This method is applicable to all types of fluids, which may be either soluble or insoluble in the oil at reservoir and injection conditions. This adaptation has unique advantages in handling problems in which either liquids or gases are the displacing fluids. INTRODUCTION In the displacement of oil from a reservoir with water or gas injected into the formation or by the expansion of a gas cap, the determination of breakthrough or flood advance is important to define the initial and subordinate phases of the operation. Thus, the determination of the displacing fluid saturation at breakthrough is fundamental to the establishment of over-all reservoir production. Methods for analyzing this saturation and for calculating production rates have been previously reported1,&apos;*".". All of these methods utilize the rate equation for linear flow developed by Buckley and Leverett1. xs = QT - Qn/A? (?fu/?Su) .....(1) Sp where f, represents the volumetric fraction of displacing fluid in the total flowing fluid within the reservoir as is defined by: fo = 1/1+ ku/kd µn/µn .......(2) Since the time, ?, necessary for the flood front to travel distance ?X is unknown, Eq. 1 cannot be solved directly for the determination of the breakthrough point. Therefore, it becomes necessary to solve Eq. 1 simultaneously with an equation representing a fluid bal-ance on the reservoir behind the flood front. To develop such a fluid balance equation, it has been assumed that the fluid composition is uniform in the reservoir behind the flood front. Based upon this assumption, the follow-ing equation can be derived from a material balance on the displacing fluid: x = (?X) s Dx ? x=0 [ Snx - Sn1 ] dx = QT - Qn/A? It is readily observed that the displacing fluid satura-tion behind the flood front, Snx, may be determined through the simultaneous application of Eqs. 1 and 3. Although the approach may vary, this objective re-mains the same for all of the previously mentioned graphical or analytical methods. After Sox is determined, it is then possible to calculate the gas-oil and/or water-oil ratios and the rates during the initial and subordinate phases of production, and thereby define the production history of the reservoir. In this presentation only the de-termination of Snx is considered, and the reader is re-ferred to the several texts in reservoir engineering2,4,5 for a complete discussion on the methods for establish-ing the production history of a reservoir. PROPOSED METHOD FOR DETERMINING Snx Since at any time the quantity, QT - Qn/A? is a constant, Eq. 1 may be expressed in differential form as follows: dxsn = QT - Qn/A? d (? fn/? Sn) sn .....(4) Eq. 4 may be substituted directly into Eq. 3 to obtain: ? fn/? Sn ? [ Snz - Sn] QT - Qn/A? d ( ? fn/? Sn) = QT - Qn/A? ...............(5) Eq. 5 simplifies to the expression: ? fn/? Sn ? [ Snz - Sn] QT - Qn/A? d ( ? fn/? Sn) = 1 ....(6)

Jan 1, 1958

Jan 1, 1961

Executive SCOTT TURNER, Chairman FREDERICK M. BECKET JOHN A. MATHEWS H. A. GUESS WILLIAM WRAITH Finance HENRY KRUMB, Chairman PAUL D. MERICA ROBERT E. TALLY Admissions JOHN M. LovEJOY, Chairman GEORGE D. BARRON ERLE V. DAVELER WILLIAM H. BASSETT W. D. B. MOTTER, JR. Alternates HENRY B. FERNALD F. JULIUS Polls AMOR F. KEENE

Jan 1, 1932

By E. R. Macgregor, J. Halpern

The kinetics of the reduction of cupric salts in aqueous solution by molecular hydrogen to metallic copper are described. The rate of reduction appears to be homogeneously determined and shows a marked inverse dependence on the hydrogen-ion concentration. Precipitation of copper is more rapid and complete in sulfate than in per-chlorate solution. A mechanism is proposed which accounts for this behavior. IN recent years a number of important hydro-metallurgical processes have been developed in which metallic salts in aqueous solution are reduced by gaseous hydrogen, usually at elevated temperatures and pressures, and thereby precipitated either in metallic form (e.g. Ni, Co, Cu) or as an oxide or salt of lower valence (e.g. VO, and VA).1 ; In connection with applications such as these, most of which involve heterogeneously catalyzed reactions, recent demonstrations"-" that salts of certain metals, e.g., Cu", Ag&apos; , Hg" , and Hg,+ can activate molecular hydrogen homogeneously and thus catalyze hydrogenation reactions in aqueous solution, are of considerable interest. For example, cupric salts catalyze the otherwise slow homogeneous reduction of dichromate and ceric salts by hydrogen, i.e. Cr,O; + 3HL. + 8H1 ? 2Cr&apos;- -- 7H,O [1] 2Ce &apos; ! H.. ? 2Ce" | 2H . [2] In perchloric-acid solution, where Cu" is not appreciably complexed, the kinetics have been found" " to be of the form -d [H,] _- k,[HJ [CU"]&apos; dt -LCu+&apos; l+ (k ,/k,) [H] [3] where, K = 1x10"&apos; exp [—26,600/RT] 1, mole-&apos; min-&apos; ... [4] and (k ,/k,) - 0.25 at 110°C. This kinetic behavior suggests that the reaction proceeds by a mechanism involving the following sequence of steps kt Cu" +H1CuH&apos; 4-HA [51 k., CUH + CU" ? + 2Cu f H&apos; [GI 6Cu&apos;+ Cr,07 + 14H&apos;+ 2Cr- (¦ 7H,O + 6Cu" [&apos;i] fast The catalytic activities of the undissociated cupric-sulphate and cupric-acetate complexes are somewhat greater (by factors of 6.5 and 120, respectively) than that of the simple hydrated Cu" ion.&apos;" These considerations suggest that, in contrast to the other hydrometallurgical hydrogenation reactions, referred to previously, reduction of cupric salts in aqueous solutions by molecular hydrogen, should proceed homogeneously, and in the absence of an extraneous catalyst. This has already been demonstrated" for acetate-buffered solutions where the product of the reaction is a precipitate of cuprous oxide (Cu,O). In the present paper, an investigation of the reduction of cupric salts in perchloric-acid and sulphuric-acid solutions, leading to the precipitation of metallic copper, is described. Some results for the sulphate system, which is of considerable hydrometallurgical interest, have been reported previously by Ipatieff",&apos; " and by Schaufel-berger&apos; but these do not give a quantitative account, of the kinetics and mechanism of the reaction. Experimental The reduction experiments were conducted in the 1-gal, titanium-lined autoclave described in connection with earlier related investigations," " and using a similar procedure. The solutions were made up by dissolving known amounts of CP grade Cu(ClO,), and HC10, (or CuSO, and HISO,) in dis-tilled water. While in the autoclave, the solution was stirred and maintained at a constant temperature (±loc) and under a constant partial pressure of hydrogen gas. It was ascertained that the agitation (provided by two 3-in., four-bladed stirrers, rotating at 900 rpm) was adequate so that at no time was the rate of reaction limited by solution of hydrogen. During each experiment, samples of the solution were withdrawn periodically, allowed to cool, and the concentration of copper remaining in solution was then determined electrolytically. Some Cu&apos; is formed during the reaction, but undergoes disproportionation when the solution is cooled to room temperature. For purposes of kinetic calcula-

Jan 1, 1959

By Harry C. Wiess

PETROLEUM has come to be one of the most important and essential of the mineral re- sources of the nation. It is the most advantageous source of mineral fuels and of lubricants, and as such it has provided in large measure for the phenomenal development in our modern means of travel and in the facilities for transportation by land, sea, and air. Without it the remark- ably rapid industrial and social progress in our country during the past thirty years could not have occurred. In view of its important place and the fact that petroleum is irreplaceable and existent in limited quantities in nature, its conservation is essential.

Jan 1, 1939

By S. Jana, C. M. Wayman

The reverse transformation of bcc martensite to the fcc phase was studied in an Fe-33.95 wl pct Ni alloy by nzeans oj dilatometry, melallography, and electron microscopy. Upon "slozc" heating (-1°C per min) length cJmnge us temperature plots showed u gradual contracLion over the temperature range 200" to 280"C ,followed by a more abrupt contraction beginning a1 -280°C. Howet,ev, zchen the heating rate was increased -4°C per tnin, no gradual contraction was observed and only the abrupt contraction starting at -2BO"C was found. Thus on slower heating- the AS "temperature" for the subject alloy, unlike the MS temperature, is better defined as a range of temperatures. Both optical and transmissiorl electron microscope observations showed that some of the martensite plates exizibited a partial loss of transformation twins during reversal. The midvib region of the martensite plates disappeaved relatively early duirng the reversal. Metallographic observations slowed that the earliest detectable stage of the rezlerse tvansforrvration begins (axd Moues inulardly) at The Martensens i te - parent interface. At higher temperatirres, the. formation of martensitically reversed jcc plates within the bcc martensite plales was observed. It is concluded that the reverse transformation consists of a diffusion less process (martensitic); but this is ps-obably aided by a prior or simultaneous dijjusiorz-comltvolled process, at leasl in the case of slower heat-ing&apos; experiments. ALTHOUGH numerous investigations have dealt with the parent-to-martensite ("forward") transformation (fcc — bcc) in Fe-Ni alloys, comparatively little is reported on the ("reverse7&apos;) martensite-to-parent transformation.&apos;-4 Even though such reverse transformations have been studied in detail in some nonferrous systems, one of the difficulties of studying the reverse transformation in most ferrous mar-tensites is that the martensite decomposes by tempering during heating. However, carbonless Fe-Ni alloys do not exhibit this difficulty since the transformation in these alloys is completely reversible. The present investigation represents an attempt to shed more light on the nature and mechanism of the martensite-to-parent transformation. 1) EXPERIMENTAL PROCEDURE 1.1) Alloy Prepatation. Fe-Ni alloys of compositions near 34 wt pct Ni were prepared from zone-refined iron (99.994 wt pct Fe) and high-purity nickel (99.999 wt pct Ni) by induction melting in recrystallized alumina crucibles in an argon atmosphere, with prior vacuum evacuation to 10"3 mm Hg. The alloys were homogenized by induction stirring in the molten state for 5 min. After solidification, the alloys were further homogenized in evacuated quartz capsules for 96 hr at 1230°C. 1.2) Dilatometry. Slices of the ingot were hot-forged (750°C in air) into approximate rod form and these specimens were then hot-swaged (750°C in air) into long cylindrical rods 0.55 mm diam. From the rods, specimens about 1 in. long were cut. These were then vacuum-annealed for 24 hr at 1200°C, cooled to room temperature, and subsequently transformed to martensite in liquid nitrogen (whereby about 40 pct transformation was obtained). Dilatation measurements were made by observing length changes in a vacuum dilatometer with an externally mounted LVDT sensing element. 1. 3) Preparation of Electron Microscope Specimens. Slices of the ingots were cold-rolled (with intermediate vacuum anneals) to -0.020 in. Out of these rolled sheets, specimens (about 1 by 1 in.) were cut. These were then vacuum-annealed, transformed to martensite by cooling in liquid nitrogen, and subsequently heated from room temperature to various temperatures to effect either partial or complete reverse transformation. These specimens were then chemically polished to 0.002 in. in l:l HsOz (30 pct) and &PO4 (85 pct) solution, and thinned to electron transparency in an electrolyte consisting of 150 g CraOs, 750 ml glacial acetic acid, and 30 ml ~~0.~ Observations were made with a 100-kv Hitachi HU-11 electron microscope equipped with an HK-2A tilting device. 1.4) Optical Microscopy. Metallographic observations were made with a Leitz MM5 metallograph on the same 0.020-in. sheet specimens as were used for electron microscopy and on bulk specimens which were 0.2 in. or more on a side. The chemical thinning solution when cooled below 20°C also served as an etchant for this alloy. Observations of surface relief were made with a Zeiss interference microscope employing a Thallium light source of wavelength 0.54 p. Specimens for interference studies were prepared by two-stage polishing on Buehler vibromet polishers using 0.3 and 0.05 p alumina abrasives. 2) EXPERIMENTAL RESULTS 2.1) Comparison of the MS,AS, and Af Tempera-tures wTth Previous Re sults. The AS aLd Af tempera -tures of several Fe-Ni alloys were determined dila-tometrically. The MS temperatures of the same alloys were determined by continuously lowering the temperature using a mixture of isopentane and liquid nitrogen and observing the highest temperature at which a prepolished specimen showed surface upheavals. For the present the As temperature is defined as the temperature at which an abrupt decrease in length occurs in the dilatation plot. The Ms,As7 and A determinations in the present investigation and those of Kaufman

Jan 1, 1968

By C. F., Christopher

The purpose of any steel-making process is to convert the two raw materials iron and scrap into steel. The chemical analysis of the steel is set within certain limits which involve the physical proper- ties desired by the consumer. In addition to tensile properties a large number of other properties are at the present time desired, such as impact resistance, grain size regulation, specific degrees of hardenability, resistance to creep at high temperatures, resistance to expansion on heat treatment, the ability of the product to undergo extreme amounts of cold deformation without failure, and many other properties which are not directly related to the specified chemical analysis. Variations in many of the properties of plain-carbon steels are due to elements and compounds not included in the ordinary "five element"§ composition. Such variations result primarily from variations in oxidation of the metal in the furnace and from the deoxidation method employed. The steel-making problem therefore resolves itself into two distinct sections: (I) the elimination of certain elements from the charge (the iron and scrap entering the furnace) and the addition of materials to meet the chemical specifications for the heat; (2) the manipulation of the slag and metal and the process of deoxidation to give certain desired properties in the steel entirely apart from the chemical analysis specified. Each of &apos;these two ends must be attained

Jan 1, 1957

By Herbert A. Franke

ANY analysis of the aluminum situation, particularly of the factors involved in the current shortage of the metal, must consider the rapid march of events since the Munich fiasco of September 1938. At the end of 1938, the sole domestic producer of aluminum had more than a year&apos;s supply of metal on hand, having added 50,500 tons of output to stocks during the year. Not until war had been declared between the United Kingdom and Germany in September 1939 did the United States begin to make serious preparations for defense and a limited national emergency was proclaimed. Congress reflecting popular opinion, was not disposed to spend huge sums for armaments in 1939, and although 31,413 tons of aluminum was withdrawn from stocks. 107,500 tons was still on hand at the end of the year. When the President initiated a 50,000-airplane program in May 1940 the country was astounded, as previous Congressional authorizations of approximately 6000 planes for the Army by July 1911 and 3000 for the Navy by 1914 had been considered adequate.

Jan 1, 1941

By A. G. Allten

EXAMINATION by metallographic and X-ray diffraction means of an austenitic steel containing 0.06 pct C, 1.26 pct Mn, 0.38 pct Si, 21.15 pct Ni, 18.72 pct Cr, 3.07 pet W, and 9.14 pet Mo indicated that a - phase, which had not been reported in the literature, existed in this steel. Both the s and car-bides of the M6 C type were present in the as-forged steel. The carbides dissolved completely in the aus-tenite at approximately 2200°F (1205°C). The s, however, was present at all temperatures up to the solidus temperature of the steel at approximately 2375°F (1300°C). Ferrite was not found in any condition of the steel. The faintness of some of the s lines in X-ray patterns obtained from samples of the steel made it desirable to isolate the s prior to verification of its entire pattern. It was found that the s phase could be separated from samples of the steel which had been water quenched from 2250°F (1230°C) by dis-

Jan 1, 1955