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By Ottokar Hofmann

In tank-lixiviation, the extraction of the silver from chloridized ore by solutions of hyposulphite salts is performed by filtration. The ore-particles are kept stationary, while the solvent moves down through the mass of ore. The quickness of extraction, other conditions alike, is in direct proportion to the rapidity of the movement of the solvent through the ore. The solution, if left in contact with the ore without moving, displays but very little solving energy. If the filtration is interrupted for ten or twelve hours, and thus solution and ore are left in complete contact for that length of time, it mill be found that, when filtration is started again, the outflowing solution is but very little more saturated with silver than it was at the time of interruption, and that the ten or twelve hours were almost a perfect loss in the total time of extraction. Notwithstanding the long contact, the solution had not become saturated with silver chloride to its full solving capacity. A rapid movement of the solvent through the ore is essential to a quick extraction. This fact is well known ; and the endeavor of leachers has been to hasten extraction by increasing the rate of filtration. Siphons, vacuumpumps and other devices have been used with more or less success; but none of them have given full satisfaction. I have found that, if chloridized ore is brought into rapid contact with a proper volume of moving sodium hyposulphite solution, the silver chloride contained in the ore dissolves almost instantly, and that it is rather the volume of the solvent than its concentrated state which produces this effect. Such favorable conditions cannot be attained in tanks. The rapidity with which a certain volume of the solvent can be brought into contact with the ore-particles is limited by the speed with which the solution descends through the ore; and thus the leaching-time in tanks cannot be shortened beyond the limit set by the filteringcapacity of the ore. In a trough, however, these favorable conditions can be attained by gradually introducing the ore into the moving stream of the solvent. The ore can thus be brought into

Jan 1, 1888

By Richard S. McCaffery, Clarence H. Lorig, Joseph F. Oesterle, Ira N. Goff, Oscar O. Fritsche

This paper is the report of an investigation to determine the relations between the viscosity, temperature and composition of that portion of the silica-alumina-lime-magnesia system(l) within which blast-furnace slags are found. The work reported in this paper was made possible by a grant of funds from Engineering Foundation. It is a part of a project, the determination of properties and structure of iron blast-furnace slags, upon which the laboratories of the University of Wisconsin have been at work for some years. Three papers have been published, which have marked the progress of the study of iron blast-furnace slags at Wisconsin. The first,(2) Desul-furizing Power of Iron Blast-furnace Slags, discussed the solubility of sulfur in slags that contained no magnesia and discussed the viscosity. The second, (3) The Constitution of Iron Blast-furnace Slags, developed the constitutions of slags in the four-component system, silica-alumina-lime-magnesia, and pointed out an explanation of slag structure. Composition of Iron Blast-furnace Slags,(l) which was the third paper, gave a method of calculating the percentages of the various compounds present in slags when the amount of oxide was given and also an interpretation of quaternary equilibrium. Illustrations were given of the mineral composition of blast-furnace slags and some applications of composition results were shown. Because of their relative ease of measurement, the viscosity-temperature relation for slags was selected as a factor which would be useful in correlating the physical properties of molten slags to furnace operation. With this in mind the viscosities of slags whose compositions are in that portion of the quaternary system, silica-alumina-lime-magnesia, in which blast-furnace slags occur, were thoroughly investigated.

Jan 1, 1932

By Irwin Roman

IT has been known for many years that when a wire is moved in a magnetic field, an electromotive force is developed which is proportional to the rate at which the wire is moved in a direction perpendicular to the magnetic intensity and to the magnetic intensity itself. If the ends of the wire are connected to a current indicator, the current behaves exactly as though the source of electromotive force were a galvanic cell. If a wire is wound circularly and rotated uniformly around a diameter not parallel to the intensity, the current flowing through the wire will be alternating of sinusoidal shape. If, instead of a single loop, there are several turns wound in a short helix, the electromotive force developed will be approximately proportional to the number of turns, the area of the helix and to the speed of rotation. The maxi-mum electromotive force will be devel-oped when the axis of rotation is perpen- dicular to the field; if this is not so, only the component of the intensity that is perpendicular to the axis of rotation is effective. This is the principle of the usual earth inductor. With a few turns of wire and comparatively slow rotation, FIG. 1.-MAGNETIC GRADIOMETER. it is possible to produce a considerable deflection of a galvanometer. If two similar coils are rotated together at two different points for which the magnetic field intensities perpendicular to the axes of rotation are different, the electromotive forces developed will each depend on its corresponding intensity and be a measure of that intensity. If the inten-sities are sufficiently different so that the electromotive forces may be compared, the two intensities may be compared. This is the principle

Jan 1, 1934

By P. V. Martin

Shortly after the middle of the nineteenth century, the invention of the regenerative open-hearth furnace and the development of the Bessemer process stimulated a rate of steel production whose magnitude created an unprecedented demand for pig iron. Revolutionary improvements in furnace design, auxiliary equipment, and operating technique enabled the blast furnace to keep abreast of the constantly increasing requirements of the steel mill. The development of improved blowing equipment, more adequate facilities for material handling, and particularly the general use of anthracitc and coke to replace the physically weaker charcoal, led to a rapid expansion in both the height and volume of furnaces. Beyond a doubt, however, the greatest single factor in blast-furnace improvement at this time was the use of the higher blast temperatures made available by the replacement of the recuperative iron pipe stoves by the regenerative Cowper brick stove. The increased blast temperature not only resulted in greater output but also reduced fuel consumption . far more than could be explained by the additional heat content of the blast. The progress of iron manufacturing was thus largely mechanical and thermal. Little attention was given the chemistry of the process until the publication of the investigations of Sir Lowthian Bell.1 Bell's painstaking experiments and investigations gave to the blast-furnace man his first real scientific conception of blast-furnace chemistry. Bell's recognition of the dual function of the carbon, the technical importance of reversible reactions and chemical equilibria, and his conception of the degree of oxidation of the carbon as shown by the analysis of the outgoing gases as a criterion of furnace efficiency, laid the foundation upon which has since been erected the structure of blast-furnace theory. Prominent among those availing themselves of Bell's data as .a basis of further study was M. L. Gruner,2 professor of the French School of Mines. His "Blast Furnace Phenomena," published in 1870, presented

Jan 1, 1940

By P. V. Martin

Shortly after the middle of the nineteenth century, the invention of the regenerative open-hearth furnace and the development of the Bessemer process stimulated a rate of steel production whose magnitude created an unprecedented demand for pig iron. Revolutionary improvements in furnace design, auxiliary equipment, and operating technique enabled the blast furnace to keep abreast of the constantly increasing requirements of the steel mill. The development of improved blowing equipment, more adequate facilities for material handling, and particularly the general use of anthracitc and coke to replace the physically weaker charcoal, led to a rapid expansion in both the height and volume of furnaces. Beyond a doubt, however, the greatest single factor in blast-furnace improvement at this time was the use of the higher blast temperatures made available by the replacement of the recuperative iron pipe stoves by the regenerative Cowper brick stove. The increased blast temperature not only resulted in greater output but also reduced fuel consumption . far more than could be explained by the additional heat content of the blast. The progress of iron manufacturing was thus largely mechanical and thermal. Little attention was given the chemistry of the process until the publication of the investigations of Sir Lowthian Bell.1 Bell's painstaking experiments and investigations gave to the blast-furnace man his first real scientific conception of blast-furnace chemistry. Bell's recognition of the dual function of the carbon, the technical importance of reversible reactions and chemical equilibria, and his conception of the degree of oxidation of the carbon as shown by the analysis of the outgoing gases as a criterion of furnace efficiency, laid the foundation upon which has since been erected the structure of blast-furnace theory. Prominent among those availing themselves of Bell's data as .a basis of further study was M. L. Gruner,2 professor of the French School of Mines. His "Blast Furnace Phenomena," published in 1870, presented

Jan 1, 1940

CHARGING of an open-hearth heat is begun as soon as possible after the previous heat has been tapped. Ordinarily, about 40 min is required to drain and dress the furnace hearth, make up the taphole, and perform routine repairs of the banks. Repairs to the upper back wall and to the front wall and jambs can ordinarily be done while the first scrap charge is being heated and are therefore allowed to wait in order to save furnace time. Any major repairs required on the hearth or other parts of the furnace necessitate longer delays because such repairs must be made while the furnace is empty. Although charging practice varies considerably in detail with the type and proportions of scrap and iron used, many factors are common to all varieties of charge. Such general considerations are dealt with first. CHARGING EQUIPMENT AND STOCK HANDLING The physical equipment available for loading and weighing the stock, transporting it to the furnace, and charging it, may have a major effect on charging practice. Limestone, scrap, and all other solid melting stock are charged mechanically by means of equipment illustrated in Fig 2-1. The charging of hot metal is shown in- the background of this figure. Mechanical equipment and plant layout for stock handling are described in Chapter 1. It is generally desirable to charge the scrap about as fast as the physical volume can be put into the furnace. Charging scrap that is too heavy or adding the charge too rapidly (less than 45 min) may be detrimental to melting rate because of excessive chilling of the entire furnace system and also because of poor conduction of heat to the lower layers in the furnace; but the more common problem is the reverse one of charging light scrap fast enough to maintain a normal heat time. A limiting factor may be the height to which the scrap may he piled in the fur-

Jan 1, 1964

By Peter R. Drummond

THE final casting of refined copper has been re-J- stricted for generations by the following sequence of operations: Filling the reverberatory furnace, melting, skimming, blowing or flapping, and poling. The hoped-for 24 hr cycle, producing 300 tons or more, has been taken up largely with the necessary bat time-consuming tasks of cleaning the bath, sulphur elimination, and in turn removal of excess oxygen to produce tough-pitch copper. Incidental to comparatively slow melting under combustion gases, copper oxides react with the furnace lining, and the slag so-formed must be completely recycled. The three-phase arc furnace has eliminated some of the cycle stages, and telescoped the remainder into a continuous operation. Electrical energy, supplied to graphite electrodes enclosed in high grade refractories, rapidly melts copper cathodes and sustains a stream of metal, containing approximately 0.01 pct oxygen, without contamination from fuel. The arc was struck on the first large electric furnace for melting copper in the United States on April 13, 1949. The earliest use of this type of furnace was at Copper Cliff, Ont., in 1936, and an admirable description of their installation has been published? Copper, melted in the Baltimore furnace, is used to cast billets, and the installation differs somewhat from the Canadian, as will be described. The arc furnace is a heavy-duty, three-phase furnace, holding 50 tons, the general outline of which appears on Fig. 1. The steel shell is 15 ft ID with a bottom radius of 14 ft 2 in. The roof, separate and distinct from the body, consists of a 15-ft water-cooled, cast-steel ring of the same outside diameter as the furnace. The center line of the furnace lies 9 ft 6 in. from that of the trunnions, permitting a 5" backward tilt for skimming, and a 40" maximum nose tilt forward for complete draining. Normally, the furnace overflows by displacement, and the use of the forward tilt arrangement is restricted to covering charging delays. The charging slot, 3 ft 8 in. x 5 in., lies on the north center line, the tap hole on the south, and the 30x30 in. skim door 45" to the west of the slot. The original 20-in. graphite electrodes were replaced with 14 in. in December 1949. Three conventional direct current winch drives, governed by electrical controls, position each electrode which has individual mast supports and counterweights. An independent circulation supplies cooling water for the electrode glands, the roof ring, charge slot, and the skim door frame. Arc Furnace Refractories Hearth: Fused-in monolithic bottoms had been used in copper arc furnaces, installed prior to April 1949. These consisted of thin layers of periclase, successively fused in place over preliminary brick courses. Heat was obtained from the arc, using a T-like arrangement of broken electrodes resting directly on the periclase to be fused. The operation, taking weeks to perform, was very expensive. The chemically-bonded magnesite-brick bottom, installed at Baltimore, was the first of its kind and a radical departure from previous practice. It consists of a 1 to 6-in. layer of castable refractory laid on the steel shell, modifying it to a 12 ft 2 in. bottom radius. Two courses of 9x2 % -in. fireclay straights and keys follow. The third course is made of 9-in. magnesite blocks of special shape to form circles of an inverted arch. It was started by a four piece keystone with skew-backs forming the outer course. The fourth course also started on a central keystone, or button, of four 90" segments, 12 in. diam x 13 Vz in. deep, and continued with 13%-in. blocks. Skewbacks at the shell completed the course to produce a horizontal surface for the side walls with a single course of No. 2 arch fireclay against the steel. Dry chrome-magnesite cement was brushed over each course after laying, and a 1-in. expansion space between the brick and the shell was filled with the same mixture. The total bottom thickness, excluding the castable material, was 5 in. of clay plus 22% in. of chemically-bonded magnesite. Tap Hole: A 5-in. OD and 3-in. ID silicon-carbide tube constitutes the tap hole and is set tangential to the upper course of the furnace bottom. Molten metal fills the tube when the furnace is level and filled to capacity. Side Walls: The lining, against the shell, consists of a 9x4Y2x3 in. soldier course of fireclay, using straights and No. 1 arches to turn the circles. A second soldier course of 9x4'/2x2'/2-in. fireclay was laid in a somewhat similar fashion. Three courses of 13Y2x6x3 in. and 9x6~3 in. of final magnesite, laid flat, completed the lining, using Nos. 1 and 2 keys to turn the circles. Cardboard spacers were placed between every two bricks in horizontal courses, and a thin coat of chrome-magnesite cement filled the joints between the firebrick and magnesite. A sprung-arch spanned the skim door with jambs of suitable magnesite shapes. Charge Slot: The slot is 3 ft 8 in. wide x 5 in. high. A silicon-carbide sill of special shapes has a 30" slope to allow cathodes to slide easily into the bath. The original arch was flat, and composed of Nos. 1 and 2 wedge magnesite with a 6-ft radius. It projected 5 in. over the sill, and, being a flat arch, gave an 18 15/16-in. opening between the inner edge and the metal line. The whole assembly was later raised 9 in., and the flat arch replaced with an arch, the lower edge of which maintained the 5-in, width from the outer to inner edges as shown in Fig. 2. A water-cooled, cast-copper jacket protects the steel shell behind the slot.

Jan 1, 1952

By L. Sturkey

Quantitative X-ray diffraction studies of the precipitation of thorium in a Mg + 3.3 Th + 0.51 Zr alloy (HK31A) in both the as-cast and cold-worked states show that the precipitation may be described by f = 1 - e(t/T)1/2 over most of the aging period. The effects of cold work and temperature changes are determined. The precipitation of tile equilibrium Mg-Th compotund is Preceded by the formation of a transition phase of higher thorium content, with the cowposition Mg2Th and with a Laves-phase structure. Thorium-containing alloys of magnesium have assumed considerable commercial importance, especially for applications requiring exposures to temperatures greater than 500°F for appreciable lengths of time. The principal commercial wrought alloys in this system are HK31A, containing nominally 3 pet by weight thorium and about 0.7 pet by weight zirconium: and HM21A. containing 2 pet by weight thorium and 1 pet by weight manganese. These two alloys retain a high level of properties even after exposures of 100 hr at 600°F, while the more common alloys containing aluminum and zinc are badly overaged by these conditions. Therefore, in order to obtain some insight into the mechanism by which thorium produces high temperature stability in these wrought alloys, it seems desirable to investi- gate the effect of temperature and mechanical working on the precipitation processes involving thorium. Since the scattering power of a thorium atom for CuKa X-ray radiation is about ten times that for a magnesium atom, the magnesium alloys containing about 3 pet by weight thorium are ideal for quantitative measurements by X-ray diffraction of the amount of Mg-Th compound present at any time. Such quantitative measurements should be of theoretical importance in the study of the effects of time, temperature, and deformation on precipitation and precipitation hardening. Of the two commercial alloys, HM21A and HK31A, the latter is certainly the most suitable for study: In the Mg-Th-Zr2 system, only Mg-Th compounds need be considered, since no ternary compounds or binary Th-Zr3 and Mg-Zr4 compounds exist. The Zr is added to this alloy as a grain refining agent and probably plays a secondary role in precipitation. On the other hand. thorium and manganese form many

Jan 1, 1961

By J. R. Long, E. L. Anderson, R. S. Dean, F. S. Wartman

TITANIUM has been estimated to comprise about 0.65 per cent of the earth's crust and ranks fourth in abundance among the metallic elements suitable for engineering uses. In spite of this, applications of titanium for structural purposes have not as yet been developed. Its commercial exploitation has been largely confined to the chemical industries and the production of paints and pigments; no doubt because of the difficulties that lie in the way of preparing the pure ductile metal. Published reports, however, indicate that ductile titanium can be prepared and possesses some unique properties. This, together with the occurrence in this country of several important titaniferous deposits, has led the Bureau of Mines to investigate processes for the production of metallic titanium, 'methods required for its fabrication, and the properties that may be produced in this interesting metal. METHODS OF PREPARING METALLIC TITANIUM Many early investigators claimed to have isolated metallic titanium; but in the light of present knowledge, it seems that they really had produced various mixtures of carbides and nitrides that have a metallic appearance but are brittle, hard, and unworkable. The first metallic titanium pure enough to be malleable when cold was produced by Hunter1 in the early part of the twentieth century. Kroll,2 van Arkel,3 and Fast,4 subsequently confirmed the facts that pure titanium is a ductile metal and that small amounts of impurities were responsible for previous reports of brittleness. When heated, titanium reacts chemically with various gases, such as hydrogen, oxygen and nitrogen; if the volume of gas thus absorbed is too great, the metal becomes brittle and unworkable. The principal difficulty in making ductile titanium is prevention of such contamination, and the success of a given process may be judged by the completeness with which it solves this difficulty in a practical manner. Of the gases mentioned, hydrogen is the least objectionable, since its absorption by the metal is completely reversible and, although dimensional changes occur and absorption of hydrogen embrittles the metal, the gas may be removed readily by heating in a high vacuum for a short time. On the other hand, the absorption of oxygen or nitrogen results in the formation of permanent 'solid solutions of the oxide and nitride, which cannot be decomposed by the strongest reducing agents2 with sufficient completeness to restore the inherent ductility of the metal. Because of

Jan 1, 1946

By M. Kaufman, A. E. Palty

THE mechanical properties of an alloy under given test conditions are in the main determined by the grain size and the phases present, their distribution and temperature-time dependence. Alloys intended for use at elevated temperatures, especially multiphase alloys depending upon precipitation reactions for strengthening, are particularly susceptible to changes in the phases present. Unfortunately, for modern superalloys, very little published information exists which relates the mechanical properties to the basic structural variables. The present work is an effort to do this for an alloy widely used in aircraft gas turbines and related applications, Udimet 500.* The mechanical properties inves- two heats is within the normal specifications for the alloy. However, the elements which most strongly influence the phase reactions, C, Ti, Al, and Mo, are comparable in the two heats. There was no intentional addition of boron, but normal pickup is about 0.002 pct. The vacuum-melted material was received in the form of centerless ground wrought bar stock % in. diam. PHASE STUDIES A) Methods—Two methods were used to determine the phases present: X-ray diffraction measurements on electrolytically digested residues, and microscopic examination. No single commonly used electrolyte will reveal all the phases that may be present; therefore, two complementary acids were selected: HCl and H3PO4. The necessity for this will be shown below in the results. The acid concentrations used were 10 pct HC1 in alcohol and 10 pct H3 PO4 in water. The current density was maintained below 0.1 amp per sq in. Residues were rinsed and separated by centrifuga-tion. X-ray diffraction patterns were obtained in a 114 mm Debye-Scherrer type camera using cobalt radiation. A temperature of 2250°F was selected to obtain as complete solutioning of phases as possible without encountering incipient melting. Specimens were

Jan 1, 1961

By C. A. Stickels, C. M. Yen

Secondary recrystallization was investigated in vacuum-melted electrolytic iron to which 70 pm N was vacuum-meltedadded. The secondary texture is "near {554}<225>" for material cold-rolled 75 to 90 pct, the sharpness of the texture increasing with increased rolling reduction and with decreased annealing temperature. At reductions of 95 and 97.5 pct the secondary texture is &apos;"near {322)(296)". Both secondary orientations also exist as major components of the primary re-crystallization texture. Development of a strong "near {554) (225)" secondary texture appears to depend on the evolution of the Primary texture to a transition texture depleted in orientations near the secondary orientation before the onset of secondary growth. A variety of qualitative experinzents have been used to show that nitrogen is important in limiling primary grain growth. The presence of nitrogen does not seem essential for the establishment of a transition texture, but a loss of nitrogen during annealing may facilitate growth of grains in the secondary orientation. Secondary grains we shown to form initially at the specimen surface. This is not thought to indicate that surface energies are important in the growth process. It is proposed that the quasi-two-dimensional character of surface grains permits discontinuous growth parallel to the surface before secondary growth of interior grains is possible. An earlier study of recrystallization textures in 90 pct cold-rolled electrolytic iron showed that secondary recrystallization occurred after annealing for several days at 700C1 This type of secondary recrystallization, which had not been reported previously, results in the formation of a strong texture, best described by the indices "near {554}(225)". The purpose of the present work was to investigate the effect of various processing variables on secondary recrystallization in this material and determine the mechanism of secondary grain growth. LITERATURE REVIEW An understanding of the mechanism of a secondary recrystallization process depends on knowing: 1) how grains in the secondary orientation come to be in the primary recrystallization texture; 2) why normal grain growth does not occur; and 3) what factors determine the strength of the secondary texture. For secondary growth of grains of a particular orientation, a certain minimum fraction of the grains must be in that orientation after primary recrystallization. This requirement is apparently satisfied "naturally" in certain systems, i.e., when the primary texture obtained by rolling and recrystallizing material initially randomly oriented contains a sufficient fraction of primaries in the secondary orientation. However, in other cases, e.g., {110}<001> and {100}<001> secondary growth in silicon iron,2 it is necessary to enhance the fraction of primary grains in the secondary orientation by rolling and recrystallizing textured material. In the present case, the "near {554}<225>" orientation is contained within the spread of orientations found in the primary recrystallization texture of iron or bbc iron-base alloys. In systems where the main driving force for secondary growth is the reduction in total grain boundary energy, secondary growth is observed only when normal grain growth is minimized. Four ways in which normal grain growth can be limited are: 1) Limitation by a strong primary texture. When a very strong primary texture consisting of a single component or twin-related components develop, most primary grains are separated from one another by relatively immobile small-angle grain boundaries. The classic instance of this is secondary growth into the primary cube texture in some fcc metals. 2) Limitation by precipitates. Precipitates present in the proper volume fraction with a suitable dispersion will limit primary grain growth. The role of MnS inclusions in impeding normal grain growth in Si-Fe is well-documented.5 3) Limitation by sheet thickness. Normal grain growth slows drastically when the mean grain diameter is of the same order as the sheet thickness. This effect has been used to obtain secondary recrystallization in thin sheets of high-purity silicon iron.&apos; 4) Limitation by solute impurities. It is well-established that certain impurity elements in solution can have a large effect on grain boundary mobility.&apos; However, there does not seem to be any secondary recrystallization process in which primary grain size stabilization has been shown to be due to the drag exerted on grain boundaries by dissolved impurities. In certain systems, e.g., secondary recrystallization in silver,&apos; the means by which normal grain growth is limited has not been identified, and solute-impurity limitation might be suspected. In order to understand the factors which determine secondary texture strength in three-dimensional growth, it is necessary to examine in more detail the current picture of general secondary recrystallization processes. Following Cahn,9 it is assumed that the primary grains have a range of sizes and that secondary growth of one of the large grains in this distribution is possible when it exceeds a critical size with respect to its neighboring grains. The critical size depends on the ratio ?S/?p, where ?s is some sort of average grain boundary energy between the potential secondary and the primary grains and ?p is some sort of average grain boundary energy between primary grains. For a constant primary grain size, the critical size for secondary growth increases as ?$/?p increases. May and Turnbull5 have incorporated the

Jan 1, 1969

By R. P. Laforce, ZJ. R. Low, A. M. Turkalo, D. F. Stein

The interaction of the crlloying eletnenls, nickel and chromium, with the impurity elements, antimony, pIzosphorus, tin, and arsenic, to producse reversible temper brittleness in a series of high-purity steels containing 0.40 wt pct C has been investigated. The alloyed steels contained approximately 3.5 pcl Ni, 1.7 pct Cr, and 0.05 to 0.08 pct of the particular irnpurity to be investigated. Susceptibility to teirlper embrittlement was measured by comparing the notched-bar transition temperature of each steel after quenching from the final temper and after very slow cooling (step cooling;) following the final temper. A plain carbon steel without alloying elements, bu/ ud/h 0.08 pel Sh, does not embrittle when step-cooled through the emzbrittling range of temperatures. The same embrittling treatment, applied to a steel with about the same antinzony content but with nickel and chvonziunz added, causes a 700°C increase in transition temperature. If chromium or nickel is the only alloying element, the increase in transition temperature is only 50%, again with antimony present. A carbon-free iron containing nickel, chromium, and antimony shou~s a 200°C shift in transition temperature for the same thermal treatment. Specific alloy-impurily interactions are also observed for the other impurity elements, phosphorus, tin, and arsenic. Additional investigations involving electron microscopy, trzicrohard-ness tests of vain boundaries, minor additions of zirconiutn and the rare earth and noble metals, nzainly with negative results, are also described. HE particular type of embrittlement investigated is that which is encountered in alloy steels tempered in the temperature range from about 350" to 525&apos;C or slowly cooled through this range of temperatures when tempered above this range. This type of embrittlement is sometimes called reversible temper brittleness to distinguish it from the embrittlement indicated by a minimum in the room-temperature V -notch Charpy energy vs tempering-temperature curve encountered in the range 28 0" to 350°C. Temper brittle-ness seriously restricts the use of many alloy steels since it precludes tempering or use in the embrittling range of temperatures and may significantly raise the ductile-brittle transition temperature of heavy-section forgings and castings tempered above the embrittling range, since such sections cannot be sufficiently rapidly cooled after tempering to avoid embrittlement. The very voluminous literature of temper brittle-ness up to about 1960 has been reviewed by woodfine&apos; and LOW.&apos; Of particular significance to the present investigation was the demonstration by Balajiva, Cook, and worn3 that high-purity Ni-Cr steel does not exhibit temper brittleness and the subsequent detailed and systematic study by Steven and Balajiva~ of the effect of impurity additions on the susceptibility to embrittlement of Ni-Cr steels. Steven and Balajiva showed that, of the impurities which may be found in commercial steels, Sb, As, P, Sn, Mn, and Si could all produce temper brittleness in a high-purity Ni-Cr steel. The principal purpose of the present investigation was to study the effects of particular alloy-impurity combinations on susceptibility to temper embrittlement. The steels used were high-purity 0.30 to 0.40 wt pct C steels containing 3.5 wt pct Ni and 1.7 wt pct Cr, separately or in combination. The susceptibility of these steels was then determined when approximately 500 ppm by weight of antimony, arsenic, phosphorus, or tin were added as an impurity. The melting, casting, and forging practices used in the preparation of the materials investigated are described in Appendix A. Table A-I in this appendix shows the analysis of all steels to be discussed. The steels were produced as 20- or 2-lb heats. The smaller heats were used after it had been demonstrated (see Appendix B) that a small, round, notched test specimen could be used to measure the shift in the ductile-brittle transition temperature caused by temper brittleness with about the same result as that obtained by Charpy testing. HEAT TREATMENT Unless otherwise noted, all steels were tested for embrittlement in the tempered martensitic condition. A typical heat treatment for a 0.40 C, 3.5 Ni, 1.7 Cr steel was: 1 hr at 870"C, in argon, quench into oil at 100"C, quench into liquid nitrogen, temper 1 hr at 625"C, and water-quench. The warm oil quench was used where quench-cracking was encountered; otherwise the initial quench was into room-temperature oil or water. For other compositions austenitizing temperatures were 50°C above Acs with the remainder of the thermal cycle the same. Steels in this condition, with no further heat treatment, are designated as non-embrittled. The above quenching and tempering cycle for the 0.40 pct C steels resulted in as-quenched hardnesses of 48 to 53 RC and as-tempered hardnesses of 24 to 31 Rc except in the case of the plain nickel or plain carbon steels. In these, the as-tempered hardness was as low as 80 to 90 Rg. No attempt was made to adjust the tempering temperature to obtain the same hardness in ali steels since it was felt that a uniform thermal cycle was more important than exactly equivalent hardness values. Pro- the standard quench and temper described above, the standard embrittling treatment was "step-cooling". For this the thermal cycle was: 593"C, 1 hr; furnace-cool to 538"C, hold 15 hr; cool to 524"C, hold 24 hr; cool to 496"C, hold 48 hr; cool to 468&apos;C, hold 72

Jan 1, 1969

By R. Silverman, H. Blumenthal

It was found that despite the similarity of chemical analyses of different titanium carbides used as base materials for cermets, the physical properties, especially transverse-rupture strengths, of test bars were different. Hence this metallographic study attempts to link physical properties to micro-structures. It is shown that microstructure, grain shape, and grain growth are functions of three interrelated factors: 1—powder production procedure, 2—surface conditioning of the particles, and 3—impurities either contained in the original powder or acquired during ball milling. An explanation is offered for the "coring effect," long observed, but heretofore of unknown origin. The explanation is based on assumption of an oxide film and on chemical analyses which substantiate these findings. TITANIUM carbide has become in recent years a material of great interest in the high temperature field. Consequently, many manufacturers in the United States and Europe are producing titanium carbide for cermet applications as well as for additions to the well known tungsten carbide tools. All present commercial processes of titanium carbide production utilize the chemical reaction of titanium dioxide and carbon to form as nearly as possible stoichiometric Tic. This reaction is carried out in three ways: 1—in a menstruum of molten metal,&apos; 2—in the solid state, either in a protective atmosphere2 or in vacuum;" or 3—in an are-melting operation. In spite of the fact that the pure carbides obtained in these operations are almost identical chemically, the physical properties vary considerably when they are combined with a binder (Ni, Co) to form cermets. This fact led the authors to examine metal-lographically nickel-bonded titanium carbide in order to find the possible reasons for this behavior. Materials and Methods Five different titanium carbides were used in this investigation. They are identified in Table I. The first four materials were used in the as-received condition. Material E, received in lumps, was crushed to —100 mesh and carried through a flotation process in order to bring its graphite content in line with the other products. A Galagher flotation cell was used with pine oil as frothing agent. The chemical analyses of the investigated materials are given in Table 11. The binder used was carbonyl nickel of 9 to 14 microns particle size, supplied by A. D. Mackay. The materials were ball milled at a ball to charge ratio of 6:1 using procedures described under "Experiments and Results." All particle sizes mentioned are averages determined with a Fisher Sub-sieve Sizer. Test bars (lx0.40x0.16 in.) were prepared by 1—hot pressing to 85 to 95 pct of theoretical density at pressures between 1 and 1½ tsi and temperatures from 1600" to 1800°C, 2-—-cold presssing after 3 pct camphor had been added, or 3—wet pressing, both 2 and 3 at pressures between 5 and 10 tsi. All pressed bars were sintered in a vacuum of 105 to 10-6 mm Hg for 2 hr at 1350 °C. Transverse-rupture strengths were determined by breaking on a Baldwin Universal Testing Machine over a 9/16 in. span. Densities were measured by water displacement. The preparation of the specimens for micrographs was done according to Silverman and Doshna Luscz." All magnifications are at X1000. A sodium picrate electrolytic etch was used. Experiments and Results The influence of ball-milling procedure, ball-milling medium, pressing procedure, and sintering procedure on the microstructure of 80/20 — TiC/Ni were investigated. Ball Milling of Materials A, B, and C in a Steel Mill: Figs. 1 and 2 show microstructures of hot-pressed and vacuum-sintered test bars of materials A and B after the respective materials had been ball milled to 2.1 microns particle size in a steel mill and mixed with 20 pct Ni binder. Material A (Fig. 1) shows considerable grain growth. Also evident is a tendency of the carbide grains to coalesce. The density is 98 pct and the low transverse-rupture strength of 111,000 psi is probably caused by many large grains and an unfavorable packing factor. Almost all grains show a slight indication of "coring." Material B (Fig. 2), although showing grain growth, still has many small particles and a better distribution of binder and carbide due to the relative absence of the coalescing tendency. "Coring" can be observed in almost all grains. The high transverse-rupture strength of 179,000 psi and the density of 100 pct are believed to be due to the many small grains completely surrounded by the binder phase. There is also a preference to form spherical grains with material A, while most grains of material B preserve their angular shapes. Material C, of which no picture is given, stays between A and B in every respect. Rounding of some grains can be observed as well as coring, but the latter to a lesser degree than with material B. Its densification is good and the transverse-rupture strength obtained is 142,000 psi. Ball Milling of Materials A, B, C, and E in a WC Mill: When the Tic powders were ball milled to 2 microns particle size in a we mill, then ball-mill mixed with 20 pct Ni binder, hot pressed, and vacuum

Jan 1, 1956

By Paul L. Burkhart

THOUGH at an earlier period brief studies had been made by the Lehigh Navigation Coal Company Inc., it was not until 1924 that J. B. Warriner, then general manager, called for a comprehensive study of timber treatment, treating solutions, and related operations. Because of the impossibility of procuring accurate costs and other pertinent treating data, the American Wood Preserving Co. agreed to treat a car of mine timber at the Port Reading plant of the Reading Railroad. Fig. I shows the treating data collected at the treating plant on this timber, which was installed in two of the Lehigh mines in 1925. [ ] In the latter part of 1925, a survey of the tunnels and gangways of the company&apos;s Panther Valley mines was made to determine the proportion of gangways or tunnels timbered, cause of timber failures, approximate life of the average wood set and the approximate percentage of treated timber that it would be profitable for the company to use. Results of this survey appear in Table I. Because of the high percentage of timber that gave evidence of failure by rot, as noted in Table I, and a further study of the life of the sections where timbering was necessary, it was deemed advisable to fill 25 per cent of the yearly gangway timber requirement with treated timber. QUANTITIES AND TREATMENTS OF TREATED TIMBER PURCHASED In the early days of the program, none of the large treating companies would accept orders for small quantities of treated timber and the company was compelled to state the quantity of timber it would require at each operation in each month of the year. Table 2 is an example of the tabulation given to the various treating companies to assist them in making a yearly bid on the minimum quantity of timber required by the company during the year 1931. A bill of timber was calculated for each car, the weight of which would be between the minimum and maximum weight permitted for railroad-car loading. The treating company was paid for timber upon delivery at the operation specified. As extra cars of timber were required, they were ordered from the treating company, and payment was required on their receipt, on the basis of the piece or tonnage

Jan 1, 1942

By E. J. Whittenberger, F. N. Rhines

The formation of casting porosity is viewed as a nucleation and growth process with solidification shrinkage and gas precipitation as cooperative driving forces. Experimental evidence evaluating the individual contribution of each force confirms the major premise that microporosity is nucleated by gas precipitation. Additional data disclose the effects of the identity and quantity of gas, the temperature range of freezing, and the freezing rate upon the amount and distribution of porosity in light metal castings. POROSITY, as it commonly occurs in metal castings, has been ascribed to the contraction that accompanies the freezing of the metal, to the evolution of dissolved gases from the liquid during cooling and freezing, and, most frequently, to a combination of the two.&apos;" "Pipe" formation is universally associated with solidification shrinkage; spherical "gas holes" are understood to be retained bubbles of gas; but the origin of microporosity, which usually exhibits an irregular or scalloped outline when viewed in cross-section, remains a subject of continuing debate.", &apos; The present research constitutes an endeavor to show what parts are played respectively by shrinkage and gas evolution in the initiation and growth of micropores and what factors determine their number, size, shape, and distribution. A preliminary outline of the mechanism of cavity formation in castings will serve to coordinate the experimental findings that will be reported and will at the same time provide a general theory of the origin of casting porosity. Basic to the mechanism of cavity formation that will be proposed are the following established facts: 1—that liquid metals and alloys contract gradually during cooling down to the freezing range, where a relatively large contraction, known as the "solidification shrinkage," takes place (exceptions in bismuth and gallium being noted);&apos; 2—that the capacity of liquid metals and alloys for dissolving gases diminishes with cooling until, in the freezing interval, a large decrease takes place (various exceptions being noted and excluded from present consideration) ;.: and 3—that freezing proceeds counter to the direction of heat flow, the freezing front most commonly being composed of a system of tapering dendrite stalks and side arms which are coarser the lower the freezing rate and which protrude into the liquid for a distance that is greater the longer the freezing temperature interval of the alloy and the smaller the thermal gradient.&apos; The mechanism by which voids are generated

Jan 1, 1953

By Howard Eavenson

IN a discussion in the Transactions of the Institute (vol. xl, page 835 et seq.) the writer gave some data about explosions of gas and dust in the coal mines of the United States, Canada, and Mexico, among which was a table showing all explosions in the countries mentioned, in which five or more fatalities occurred, which had been reported from Oct. 2, 1871, to Mar. 28, 1908. Since the elate of publication of this discussion, as opportunity has presented, this table has been extended by the addition of explosions occurring since that time, by including several explosions which occurred prior to 1871, and by adding to the list all explosions of gas or dust in which less than five fatalities occurred, this having been clone by a search through the various mine reports, partly by the author and partly by the Librarian of the United Engineering Societies, and by correspondence with mine inspection officials of the various States, to all of whom acknowledgment of assistance is made: The accompanying list, Table I, is believed to be as nearly correct as such a statement can be made, and, while only extending to Dec. 20, 1912, contains all of the data now available in reports. An analysis of Table I, showing the occurrence by months of all explosions of gas or dust, including those in which less as well as those in which more than five fatalities occurred, is given in Table II. The statements are frequently made, in technical papers, that the number of explosions, particularly the larger ones, is increasing more rapidly than the number of mines, and that the number of fatalities caused by mine explosions, particularly the more disastrous ones, is increasing faster, proportionally, than the production of coal. To investigate the correctness of these statements Table III was prepared, in which the number of accidents and fatalities is taken from Table I, using those in the United States only, on account of lack of other data for Canada and Mexico; the production is from the records of the U.S. Geological Survey, and the number of mines, from 1909, from the same source; previous to 1909, the number of mines, for the decennial periods, was furnished by the U. S. Census records.

Jan 10, 1914

By C. K. Stoddard, R. G. Knickerbocker, E. T. Leidigh, R. R. Lloyd, K. L. Mattingly

DURING July 1941, a study was initiated at the Boulder City Experiment Station of the Bureau of Mines on proposed methods for the production of magnesium metal. The major emphasis was placed upon development of a process that required minimum reliance upon the sale or utilization of by-products to meet economic competition from processes already in operation. It was felt that only through the development of such a process could the nation&apos;s vast resources of dolomite and low-grade magnesite that are contiguous to large blocks of firm power be properly utilized. During the course of exploratory tests on electrolytic and electrothermic processes, it was noted that fused magnesium chloride would dissolve appreciable quantities of magnesium oxide in the presence of small amounts of boric oxide. The use of such a solution in an electrolytic cell employing a graphite anode and a steel cathode caused the deposition of magnesium at the cathode and the evolution of carbon oxides at the anode with minimal amounts of chlorine The original results were sufficiently encouraging to justify further investigation on an enlarged scale, since the process seemed to present advantages over other proposed processes 1.2 for the direct use of magnesium oxide in a fused chloride electrolyte. When work on a larger scale was initiated it was found that the addition of boric oxide to the electrolyte resulted in dispersion of metal and constantly decreasing current efficiency. However, by this time solubility studies had shown that magnesium oxide was sufficiently soluble in magnesium chloride melts containing no boric oxide to effect the same anode reactions, and development work was continued. Over two years of development work have been carried out, and considerable work remains to be done. It is desired to emphasize the fact that many of the difficulties met require a long period of continuous operation before they become serious enough to affect operation of the cell. ACKNOWLEDGMENTS This paper is one of many reporting on various aspects of the Bureau of Mines&apos; program initiated in August 1939 by passage of the Strategic Minerals Act the scope of which was greatly expanded by subsequent legislation. Some of these papers are published as war minerals reports. others as bulletins, technical papers, reports of investigations

Jan 1, 1945

By A. R. Rosenfield, G. T. Hahn

The effects of adding 0.025 volume fraction of uniformly dispersed Tho2, particles (particle size -380A) to iron have been examined. The two-phase alloy is shown to have a lower transition temperature and higher cleavage stress than either the base metal or a solid-solution alloy (Fe-3Si). Enhanced cleavage resistance is thought to be a consequence of particles limiting the effective slip distance of dislocations and thus the ability of Pile-ups to act as stress concentrations. Further calculations predict that the highest cleavage strengths are attained with uniform dispersions, where the particle size is kept small (-250A) and the volume fraction is relatively large (-2 to 3 pct). ALTHOUGH a fine dispersion of second-phase particles is known to impart strength to an alloy, the possibility that it might also impart ductility has only recently been explored. Outstanding progress has been made in the production of ductile tungsten alloys containing Tho21,2or (Hf, Zr)C,= but the contribution of the dispersion to ductility has neither been optimized nor fully analyzed. The basis for such an analysis has now been developed1,5 and is discussed briefly below. Experiments on iron are also described that attempt to isolate one of the possible ductilizing mechanisms and test its effectiveness. In analyzing these experiments, it is important to remember that cleavage obeys a stress criterion, and that either a lower ductile-to-brittle transition temperature or a higher cleavage stress is evidence of enhanced resistance to cleavage. DUCTILIZING MECHANISMS Slip Blockage. Cleavage fracture of polycrystals may be initiated at the intersection of a slip band (or twin) and a neighboring grain or a large (-1 to 5 µ) particle. The slip band generates a stress concentration at its head, and either the grain boundary fractures, the particle fractures, or cleavage is initiated in a neighboring grain. Experimental evidence for this viewpoint is found in metallographic studies8-8 and from the observation that cleavage is usually preceded by yielding.8-11 The potency of a slip band as a stress concentrator

Jan 1, 1968

By R. J. Allio, C. H. Randall

An automatic system for the quantitative analysis of microstrulcture has been developed. The method is based upon a statistical model of microstructure and employs phase intercept fractions as fundamental random variables. Structure-sensitive properties can be related explicitly to parameters derived from the distribution of phase intercept fractions in a matevial. THE properties of mechanical dispersions and of multiphase alloys art: intimately related to the amount, distribution, and intrinsic properties of the dispersed material. In many instances, the amount and distribution of the dispersed substances have greater influence on the over-all system properties than do the properties of the dispersed substance-Such properties as average particle size, particle size distribution, homogeneity of dispersion, volume per cent of dispersed phase, and mean free path between dispersed particles thus become critical in any evaluation of the properties of a dispersion. Unfortunately, techniques for evaluating quantitatively the microstructural characteristics of dis~ersions rapidly and accurately have not been available. If necessary, qualitative or semiquantitative devices can be utilized. Some such devices are based merely upon comparison of the specimens in question with arbitrary standards. Other techniques involve superposition of grids and tabulation of particles observed in different portions of the grid. Still other methods based upon lineal analysis entail counting lines or points at random upon the specimen to be examined; the number of intersections of the line with a given phase, the proportion of line or number of points which are covered by given phase are variously recorded. The mathematical basis for most of the techniques of quantitative metallography have been amply reviewed in the literature.&apos;l2 without exctption, however, the standard techniques for quantitative metallography are either imprecise or are exceedingly tedious to apply. For example, tranverse integrating stage for rapid manual areal fraction measurement has been designed by Smith.3 Howerer, as in the case of any device based upon visual eslimation, operator bias may cause appreciable error. Various flying spot devices, on the other hand register particle size distributions in 2-dimensions, from which it is difficult to derive other properties of the dispersion. Thus, e.g., the 3-dimensional distribution cannot be inferred unless certain rather broad assumptions are made concerning the randomness of distribution and geometrical configuration of dispersed phase. Many limitations to the application of quantitative structural analysis can be overcome by the use of the microdensitometer, which is essentially an automated areal analysis device. Data obtained in the form of a distribution of phase intercept values are

Jan 1, 1962

By C. F. Smart

Since their comparatively recent development, the alloys of cadmium with silver and copper or nickel, and of cadmium with nickel alone, have been used somewhat extensively as liners for connecting rod bearings and main bearings of internal-combustion engines. In applications where, owing to high bearing loads and elevated operating temperatures, the life of bearings made from tin-base or lead-base alloys has been undesirably short, cadmium alloys have demonstrated considerable advantage. In applications where, because of deflection of the crankshaft or bearing supports, copper-lead bearing metals have resulted in scored shafts, cadmium alloys have been operated successfully. The somewhat serious drawback to these alloys in bearing service has been that they, in common with copper-lead and certain lead alloys, are susceptible to corrosion in the presence of lubricants containing organic acids. These acids may be fatty acids (such as oleic acid) derived from animal or vegetable oils that sometimes have been compounded with engine lubricants, or they may be compounds of an acid nature developed by the oxidation of some petroleum oils under operating conditions. The result of such acid attack on bearings is illustrated by the photograph, Fig. 1. The petroleum oils that are most readily oxidized to a corrosive condition, and the operating conditions under which such oxidation most readily occurs, are now rather well known among the technical organizations of both the petroleum and automotive industries. While means have been developed for inhibiting petroleum oils against such acid formation, the knowledge that this condition may be encountered has been a factor in restricting the wider use of cadmium alloys as engine bearings. It has appeared, therefore, that their value and use would be enhanced if they could be made corrosion-resistant to a degree compatible with the most severe operating requirements. Investigation, undertaken with this in view, indicated that the addition of fractional percentages of the metal indium to the cadmium bearing

Jan 1, 1938