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(Members are urged to send in for this column any notes of interest concerning themselves or their fellow-members.) Members who registered at Institute headquarters during September: William B. Phillips, Austin, Texas. Will L. Clark, Jerome, Ariz. Olof Wenstrom, Boston, Mass. Robert H. Richards, Boston, Mass. Anthony F. Lucas, Washington, D. C. Hennen Jennings, Washington, D. C. P. H. Griffin, New York, N. Y. Waldemar Lindgren, Boston, Mass. H. N. Lawrie, Portland, Ore. C. Snelling Robinson, Youngstown, 0. Clarence P. Linville, Everett, Pa. E. P. Ross, Johnson, Tenn. Members of the XII International Geological Congress who made the Institute Office their headquarters while in New York A. Bigot, Caen, France. E. Gentil, Paris, France. J. Caillebotte, Paris, France. H. Saint-Clivier, Paris, France. John Ashworth, Manchester, England. Tadasu Hiki, Kioto, Japan. T. Dahlblom, Falun, Sweden. L. Milch, Greifswald, Germany.

Jan 10, 1913

By P. R. Sahm, T. V. Pruss, D. A. Gandolfo

Cb 3 Sn powders were hot-pressed to estahlish densi-fication conditions and assess the possible use of such compacts in multiphase structures with ceramic or magnetic inclusions. Superconducting transition temperatures, Tc, were measured as a function of powder particle size, pressing temperature, and second-phase admixtures. The highest transition temperatures were found in coarse powder (5 to 500 u) compacts and were only weakly depended on pressing temperatures up to 1200°C. Fine powders (below 5 p), however, both single-phase CbsSn and Cb3 Sn-composites, showed a stronger dependence on hot-pressing temperatures. While Al2 O3 phase admixtures did not affect T, much more, a Pronounced T, degradation was observed upon the addition of ferrimagnetic Fe2 Mno. 5 Zno.5 O4 powders. Likewise, critical current densities, Jc, at 4.2° K were lowered appreciably in the ferrite doped samples, whereas the neutral A1203 inclusiow gave slightly improved Jc values. MULTIPHASE structures in electronically active materials have recently gained in interest as pointed out by Galasso in a review article.' In superconductors the most widely reported group are eutectic alloys which have been mostly investigated with respect to the proximity effect.2"4 In addition, "artificial eu-tectics", for example in the form of mercury-vycor glass sponge5 or fine niobium wires imbedded in a copper matrix,= have been prepared to attain higher critical fields and critical current densities. The idea was to enhance flux pinning phenomena by second-phase inclusions or ho1es. 7 Flux pinning should also be possible by the introduction of insulator inclusions into a superconducting matrix. Ceramic-metal combinations are most easily produced by hot pressing (= pressure sintering), as it permits an almost unlimited selection of material combinations. Cb 3 Sn has been prepared by vapor phase growths or in the form of thin films. Large compacts must be sintered,1° however, and normally present difficulties. At temperatures below 1350°C insufficient densities are obtained and at higher temperatures the loss of tin due to dissociation upsets the stoichiometry. Hot pressing is a natural approach to circumvent the problems, since one can obtain satisfactory densities despite low sintering temperatures. This work was mostly concerned with establishing conditions to hot press Cb3 Sn and Cb3 Sn multiphase structures (= composites), because no data are presently available, and to evaluate these with respect to their transition temperatures. However, a few measurements on critical current densities were also carried out, in particular to assess the value of ferrimagnetic inclusions. I) EXPERIMENTAL PROCEDURE The experimental procedure involved the preparation of feed material and powders, hot pressing, and evaluation of the resulting ingots. As no attempts were made to synthesize Cb3Sn from the elements during hot pressing, feed material had to be produced. For this purpose, a stoichiometric Cb-Sn powder mixture was compacted into a 3-cm-diam ingot (80,000 psi) and then inserted into a quartz tube and evacuated. Rf heating was employed. Above the tin melting point (230°C) liquid tin was forced out of the ingot, but was pulled back by capillary action after approximately 0.5 hr of continuous pumping at 700°C. Thus, it appears that evacuation at this temperature caused desorption of impurity gas creating wetting surfaces which, later on, favored alloying at relatively low temperatures. After a presintering period of 17 hr at 800°C, the quartz tube was tipped off and sealed into another such larger-diameter tube which, in turn, was connected to a pump system. This configuration permitted a temperature treatment at 1400°C, carried out for 65 hr, without collapsing of the glass walls or unfavorable chemical interaction of Cb3Sn with the quartz. The resulting material was

Jan 1, 1969

By P. F. Ahner, Leonidas Petrakis, Fred E. Kiviat

High gradient magnetic separations (HGMS) is a recently developed technique in which magnetic micron-size particles contained in a flowing fluid medium can be extracted from that medium. HGMS is presently utilized on a commercial scale in the kaolin processing and steel industries, and is being investigated for solid-liquid separations in numerous other processes. The performance characteristics of HGMS in the kaolin and steel industries are briefly summarized. A discussion of HGMS performance with respect to solids removal in a coal liquefaction process is presented to illustrate its applicability to difficult solid-liquid separations. Here, the properties of the solids suspended in the liquified coal are such that the HGMS separator achieves over 90% solids removal. The optimization of the HGMS operating parameters to yield these results will be discussed. To further illustrate the potential of HGMS, its performance with respect to the removal of oil and solids (over 98% removal) from a petroleum refinery process water effluent stream will be briefly considered.

Jan 1, 1980

By C. C. Henning

IN any line of human activity logical handling requires an amount of attention to each phase that is in keeping with the importance of that phase. In the complex and rapidly changing field of modern ferrous metallurgy it is difficult to allot the proper degree of attention to the various subjects, and it is not surprising that one or more subjects of major importance should for a time be neglected. Bessemer steel is a good example of a neglected subject. It is largely taken for granted, inferior physical quality is assumed for it, and there is a lack of interest in its technical development and commercial application. Availability of supply is limited. In spite of this technical indifference, limited source and lack of commercial exploitation, bessemer steel constitutes a large volume of tonnage in daily use. There is a fairly prevalent idea even among technical people that bessemer steel survives mostly because of low cost and by reason of a dwindling residue of applications sufficiently common to permit of its use. The difference in cost between bessemer and open-hearth steel is governed largely by the relative cost of manufacture of pig iron and the market price of steel scrap. During the past four years there have been long periods of time during which the price of steel scrap was low enough so that open-hearth steel was considerably lower in manufacturing cost than bessemer. As to application, it may be pointed out that thousands of tons of this material in the form of bessemer screw steel are applied for moving parts in automotive and other equipment, Some forging work and considerable heat treating, particularly carburizing, is performed on this grade. From an economic standpoint the most effective contribution of bessemer steel lies in its free machining qualities. It is by far the freest

Jan 1, 1935

By R. I. Jaffee, J. G. Dunleavy, H. R. Ogden, Webster Hodge

INTRODUCTION EARLY in 1946, at the instigation of the U. S. Army Signal Corps, the authors made an extensive survey of the available literature covering high-strength, high-conductivity alloys. For the purposes of the Signal Corps, the minimum tensile strength to be considered was 150,000 psi, together with at least 50 pct IACS conductivity. No single alloy was found to be in production, or mentioned in the literature, which had this combination of properties. The constitution diagram for the iron-copper alloys is not well established in the respect to the two-liquid region. Fig I, taken from the 1939 Ed. of the ASM Metals Handbook, shows a two-liquid system existing above the liquidus temperature, but disconnected from it. Although Ruer and Fick1 and Ruer and Goerens2 definitely established the fact that two liquid layers could occur in the system, Ruer3 was not convinced that the line off immiscibility was closed in the center. Maddocks and Claussen4 found no liquid immiscibility at temperatures up to 1570°C in alloys containing almost no carbon. Although the diagram would indicate age hardening of the copper-iron alloys is possible, Hanson and Fords found that little hardening effect could be obtained. However, Gordon and Cohen6 reported that the response is better if the quenching is very efficient. Although iron has a potent effect on the conductivity of copper, Hanson and Ford5 have shown that, when hot-rolled copper iron alloys are aged at 1200°F, the electrical conductivity decreases sharply with iron contents up to 0.2 pct Fe, and then levels off at conductivities between 60 to 70 pct IACS with compositions to 2.1 pct Fe. They also found that aging did not increase the tensile strength, although increasing iron content gave a moderate improvement in strength. High-strength copper-iron alloys, usually containing about 60 pct iron, had been reported by Smith and Palmer.7 They found that an alloy containing 56.4 pct copper, 0.21 pct magnesium, balance iron, would give a maximum tensile strength of 170,000 psi and an electrical conductivity of slightly over 35 pct IACS. The alloy could also be processed to have 150,000 psi tensile strength, and about 41 pct IACS conductivity. Decreasing the iron content increased the electrical conductivity, but the long mushy range during freezing of the alloy containing 10 to 20 pct iron made the production of good castings difficult, In spite of the difficulty encountered by Smith and Palmer with alloys of around 20 pct iron, it was decided to make a rather complete investigation 0f this type of alloy. As a result 0f this investigation, a method of alloying and fabricating copper alloys containing 10 to 15 pct iron was developed

Jan 1, 1948

By Robert A. Rapp

The standard molar free energy of formation of MOO,was determined between 750o to 1050o C in galvanic cell measurements involving the solid electrolyte Zr0.85 Ca0.15 O1.85. Use of the reference electrodes Fe + FexO and Ni + NiO Provided two independent determinations of FoMoo2. The data may be represented by the equation FMoO2 = -137,500 +40.0T(°K) which is in good agreement with previous CO/CO, gas equilibrium measurements of Gleiser and Chipman. The standard molar free energy of formation of molybdenum dioxide has recently been reported over a large range of high temperatures from the calorimetric experiments of King, Weller, and christensen.' Gleiser and chipman2 have also recently determined ?FmoO2 from 926o to 1068oC in CO/CO2 gas equilibrium measurements. Earlier gas equilibrium studies have been made and reviewed by Gokcen.' Over the temperature range studied by Gleiser and Chipman their values for ?FoMoO2 differed from the calorimetric data by about 700 cal. This investigation was undertaken to resolve this difference and thereby establish the two phase mixture of Mo + MOO, as an electrode of well-known oxygen potential for galvanic cell applications at very high temperatures. The galvanic cell technique involving the Zr,.,, Ca,0.15 01.85 solid electrolyte was introduced for the

Jan 1, 1963

By D. B. Masson

THE martensitic transformations occurring in ,B AgCd as a result of cold working at room temp.. a-ture and cooling below room ten1erature have been reported by Masson and Barrett.1 These authors found that simple cooling to 127°K caused body-cen-tered-cubic AgCd containing 53.0 at. pct Ag to transform to an orthorhombic structure, while mechanical deformation at room temperature was found to induce a hexagonal structure. Furthermore, they observed that the temperature at which martensite first appeared on cooling, Ms, and that at which the transformation was complete, MI, as well as the corresponding temperatures for the reversion on heating, As and Af, could be ascertained to ± 1°K, although they report these temperatures only for the alloy containing 53.0 at. pct Ag (51.96 wt pct). In the work reported here the temperature range of the transformation on cooling was studied as a function of composition, and further consideration was given to the relationship between the original bcc structure and the transformation-induced hexagonal and orthorhomic structures. EXPERIMENTAL PROCEDURE All samples studied were in the form of annealed 300-mesh powder and were prepared from the constituent metals in the manner described previously.1 The compositions in weight and atomic pct are listed in Columns 1 and 2 of Table I. Method of determination of transformation temperature was the same as utilized previously1—measurement of the intensity of X-rays diffracted by the transformation product. A diagram of the low-temperature sample mount is shown in Fig. 1. This mount was made to fit a GE model XRD-3 X-ray diffraction unit. The powder sample was varnished to a flat surface on a copper sample holder at A. This sample holder was soldered to a brass rod, B, which was in turn soldered to a second copper rod, C. Rod C could be cooled to 78°K by inserting it into a reservoir of liquid nitrogen, D. A heater coil of nichrome wire was wound at E on the sample holder, and by varying the current through the heater a thermal gradient could be maintained along brass rod B. Sample temperatures from 78°K up could be attained and held as long as needed. The sample holder and nitrogen reservoir were insulated by Styrofoam, F, and the sample itself was protected by two radiation shields, G. The X-ray beam entered and left the sample chamber through windoys in the supporting container at H and I. Sample temperatures were measured by a calibrated copper-constantan thermocouple varnished to the sample holder just out of the beam at A. The low-temperature structure had been found to possess orthorhombic symmetry1 and the expected positions of the diffraction maxima were known. Therefore the sample was aligned in the X-ray beam and the Geiger counter of the spectrometer set to receive the most intense orthorhombic reflection without an adjacent interfering cubic reflection. This

Jan 1, 1961

By Alexander R. Troiano, Eugene P. Klier

Since the very early work on quenched structures, where the products of the martensite transformation had been recognizedl this transformation has provoked much interest and study. Theoretically it was desirable to account for the extreme changes in physical properties resulting from the formation of martensite, as that information might lead to results of practical importance. Desirable information, therefore, was on the kinetics of transformation from austenite to martensite as well as on the physical constitution of the martensite. It has been known for years that certain steels quenched to room temperature will contain persistent austenite as well as martensite. Other important and now well-established features of the martensite reaction are the Progress of the reaction on cooling only and the independence of Ar" with cooling velocities exceeding the critical velocity. These facts have been experimentally determined by various methods and investigators. Notable among the early investigators were Tammann and Scheill and Wever and Engel.2 Unfor- tunately, all investigators did not fully interpret their results, which, coupled with several investigations that yielde. conflicting resultsl caused a state of confusion that was not clarified until quite recently. X-ray investigations of martensite in steels3-6 indicate it to be a supersaturated solid solution of carbon in alpha iron. This supersaturation is evident from the appearance of a body-centered tetragonal structure with the addition of sufficient carbon. The dimensions of this tetragonal structure are a linear function of the carbon content within the limits of accuracy of the determinations. Slight tempering of tetragonal martensite leads to the destructi,, of the tetragonal structure, which is converted to a body-centered cubic strutture. This phenomenon has precipitated a controversy concerning the nature and mechanism of the tempering of martensite The facts in the case have been outlined by Epstein7 and more recently by Antia, Fletcher, and Cohen.8 The significance of Ar" on theoretical grounds is not clear, however, although considerable study has been devoted to its determination. The precise position of Art' aS a function of carbon content was first stated by Greninger and Troiano9 for steels covering a range of approximately 0.7 to 1.8 per cent carbon. The effect of alloying elements on Ar" has recently been the subject matter of investigation by various methods.

Jan 1, 1945

By A. H. Daane, W. J. Wunderlin, B. J. Beaudry

In an earlier study at this laboratory,' the authors investigated the solid solubility of holmium in copper, silver, and gold, and also the melting point of the first eutectic on the noble-metal-rich side in these systems. In a more recent study of the solubility of rare-earth metals in gold at this laboratory. Rider, Gschneidner, and McMasters2 noted that the reported eutectic temperature in the Ho-Au system was lower than the corresponding eutectics in systems of dysprosium and erbium with gold. A redetermination of the eutectic temperature in the Ho-Au system with the same alloy used in the original study showed the eutectic temperature to be 810°C instead of 778°C as originally reported. (The eutectic temperatures in the Ho-Ag and Ho-Cu systems were also redetermined and were found to be 795" and 864°C compared to 787° and 868°C, respectively, which were reported originally.) Since this large change in eutectic temperature made the solubility data appear questionable, new alloys were prepared, heat-treated, and examined metallographically. The procedure of sample preparation and heat treatment used in the original work' was followed except longer heat treatments were used ranging from 7 days at 600°C to 6 hr at 900°C. The results obtained are shown in Fig. 1. There is good agreement between the two studies in the maximum solubility of holmium in gold. Since the slope of the solvus is less pronounced in the present study, the extrapolation to 550°C shows a solubility of 1.6 wt pct Ho in gold compared with 1.0 wt pct in the original work. Since additional data were obtained and longer heat treatments were used in the present study, the larger value is believed to be more accurate. The authors wish to express their appreciation to P. Rider and K. A. Gschneidner. Jr., for bringing this discrepancy to their attention, and also to D. Thompson for assisting in the present study.

Jan 1, 1965

By Pol Duwez, Paul Pietrokowsky

Titanium and vanadium form a complete series of solid solutions at temperatures above 885°C. Below 885°C, vanadium is slightly soluble in a titanium (about 1.5 pct V at 650°C) and a two-phase a plus ß region extends to 18.5 pct V at 650°C. The alloys up to 15.0 pct V transform into a supersaturated a solid solution upon rapid cooling from the ß phase. Above 15 pct V the ß phase may be retained. The temperature at which the ß to a' martensite-type transformation takes place decreases steadily from 850°C for 0.5 pct V to 325°C for 12.7 pct V. IN the transition series of the periodic table, the six elements that have a body-centered cubic structure at all temperatures, namely chromium, vanadium, molybdenum, columbium, tungsten, and tantalum, fulfill the conditions required for the existence of complete series of solid solutions with the high temperature (ß) form of titanium. Up to the present time, complete solubility at temperatures above the a to ß transformation of titanium has been observed in the systems Ti-MO,1,2 Ti-Cb,2 Ti-W,³ and Ti-Ta.³ The two remaining elements, vanadium and chromium, have atomic diameters which are not as near that of titanium as the atomic diameters of the four other elements. Both are smaller than titanium and the difference, expressed in percentage of the titanium atomic diameter, is 12.3 pct for chromium and 8.2 pct for vanadium. On this basis, the least favorable conditions for solid-solution formation are therefore encountered in the Ti-Cr system. The early work on the Ti-Cr system4 revealed that up to about 1100°C the solubility of chromium in ß titanium is limited, and an intermetallic phase exists near the stoichiometric composition TiCr2. It was found later' that a continuous series of solid solutions extends across the diagram at higher temperatures, and TiCr2 forms by solid state reaction. It became apparent therefore that complete solubility would also be present in the Ti-V system. The present investigation confirms this possibility. The iodide-refined titanium used in this investigation was purchased from the New Jersey Zinc Co., New York. According to the manufacturer, a typical analysis of this product is 0.0065 pct Mn, 0.0022 pct Fe, 0.0015 pct Cu, and 0.0042 pct Pb. The Vickers hardness number (10 kg load) of the crystalline bar varies between 55 and 80. The vanadium metal was purchased from the Westinghouse Electric Corp., Lamp Div., Bloomfield, N. J. The purity of this metal is probably above 99 pct V. The oxygen content must have been quite low, since a bar about 1/4 in. square could be cold rolled into a ribbon 0.020 in. thick. To prepare the various alloys on the titanium-rich side of the diagram, the titanium rod was machined into cups 1 cm in diam and about 1.5 cm high. Pieces of vanadium, of the proper weight, were cut from the square bars and placed within the cups prior to melting. On the vanadium-rich end of the diagram, single pieces of the metals, in the proper weight pro-

Jan 1, 1953

By S. A. Levy, J. D. Wood, J. F. Libsch

A study of the preparation and characteristics of a sevies of Fe-C-N alloys has been conducted. X-ray, microhardness, and metallographic data from a series of single-phase alloys produced by controlled nitriding of carburized foils are reported. The existence of a, y', and E carbonitride has been confirmed and it has been found that y' is the hardest of these three phases. These results have been used to interpret the structures observed in nitrided steels. ThIS paper represents the second phase of a three-part study concerning combined nitriding and induction hardening. The first part' reports the engineering properties (depth hardness, fatigue, corrosion, and tempering resistance) obtained when seven different steels were subjected to the combined surface hardening treatment. The present paper concerns the nature of a series of single-phase alloys prepared by controlled nitriding of carburized foils. The objective was to simulate the various layers present in the as-nitrided steels. A third paper will report data concerning induction hardening of the foils and will attempt to set forth a mechanism for the high hardness obtained with the combined treatment. Studies of nitriding generally fall into three categories. The first were performed with bars using conventional nitriding procedures, thus yielding a nitrogen gradient and a series of phase layers near the surface.2-6 Such studies were complicated by the presence of the nitrogen gradient, for it was difficult to correlate microhardness and X-ray data with chemical composition. There were always complications due to penetration of the X-ray beam which yielded data influenced by underlying layers. Similar difficulty was experienced in taking hardness readings on a tapered specimen or in isolating various layers by machining. A second group of investigations was primarily concerned with establishment of regions in the Fe-N7,9,10 or the Fe-C-N phase diagram.8 These were performed with powder specimens which were nitrided in am-monia-hydrogen mixtures to produce a series of single-phase specimens. Unfortunately, the powder specimens were not amenable to either microstruc-ture or microhardness study. The third was essentially a compromise of the first two. Bose and Hawkes," using 0.005-in.-thick iron foils, were able to obtain microhardness, microstruc-ture, composition, and X-ray data on Fe-N alloys of eutectoid composition by equilibration in ammonia-hydrogen mixtures at 1290°F. To obtain a better understanding of the nitriding process and the characteristics of the phases in ni-trided plain carbon steels, the present investigation was undertaken. The investigation was patterned after that of Bose and Hawkes with the exception that nitriding was performed at 930°F (below the eutectoid temperature), as in the case of conventional nitriding. In addition to gaining some understanding of the source of high hardness of nitrided and induction-hardened steels, found in a previous study,' this investigation was concerned with the possibility of controlling the phases present in a nitrided case through the use of ammonia-hydrogen mixtures. Knowledge of the properties of each phase and the conditions necessary for its production could make it possible to avoid undesirable surface layers in nitriding. EXPERIMENTAL PROCEDURE Specimen Preparation. A specimen thickness of 0.005 in. was selected to insure that a uniform nitrogen concentration could be obtained in a reasonable length of time by equilibration in an ammonia-hydro-gen atmosphere. The starting material was a vacuum-cast ingot of pure iron, the major impurities of which were: C, 0.003 pct; 0, 0.014 pct; S, 0.004 pct; N, 0.003 pct.* 'Compositions are quoted as weight percent except as noted. The ingot was heated to 1500°F in a nitrogen atmosphere and forged to a slab 0.25 in. thick. The material was then reduced to 0.005 in. by alternately cold rolling and annealing in a nitrogen atmosphere at 1600°F. This material was divided into five groups of specimens which, after carburizing, yielded the following carbon contents: 0.003, 0.28, 0.36, 0.46, and 0.74 pct C. The carburizing was performed by the Leeds and Northrup Research Laboratories, using a controlled "carburizing potential" which yielded a uniform carbon concentration throughout the cross section.

Jan 1, 1970

By E. G. Banks

This paper is presented in the belief that metallurgists and chemists will be interested in the practice of grinding in tube-mills in connection with stamps, especially since the records of working here given extend over a lengthy period of time (since May, 1905). The ore from the Waihi mine—more especially that produced in the upper levels—contains a large proportion of hard, chalce-donic quartz, and the gold exists in an exceedingly fine state, conditions which necessitate very fine crushing in order to obtain a high extraction of the precious metal. Before the introduction of tube-mills at the 90-stamp Waihi mill, it was found necessary to stamp through 40-mesh (1,600 holes per sq. in.) woven wire-screens, having a fairly high discharge. The pulp then graded : Per Cent. Per Cent. On 50-mes11, .- 0.1 On l20-mesh ..8.77 On 60-mesh, . 8-74 On 150-mesh,.... 7.48 On 80-mesh,... 16.06 Passed 150-mesh,... 55.19 On 100-mesh,.... 3.66 The stamp-dnty was 2.89 short tons per stamp per day, or a total of 260 tons daily. Although the extraction on this pulp was: gold, from 88 to 90 per cent., and silver, from 74 to 78 per cent., it was recognized that finer grinding of the sands would prove beneficial, provided a machine could be found to do this work economically. Various grinding-mills and pans were tried, but without satisfactory results, the particles of sand being so hard that the capacity of any of the machines was too small to be economical. The results of grinding in tube-mills in other countries were so satisfactory that three tube-mills mere erected at the Waihi

Jan 1, 1908

By E. T. Turkdogan, P. Grieveson

Experimental results are presented for the rate of solution of nitrogen in a iron in the temperature range 750° to 873°C and in 6 iron in the temperature range 1410° to 1470°C. It is shown that the rate controlling process is diffusion of nitrogen into the iron. The diffusiting of nitrogen in a and 6 iron is derived from the results, and the temperature dependence of the diffusivity is represented by the equation D = 7.8 x e- 18,900/RT sq cm per sec. The solubility of nitrogen in a and 6 iron, in equilibrium with 1 atm pressure of nitrogen, has been measured. Using these results and other available data, it is found that the variation of the logarithm of nitrogen solubility with the reciprocal of absolute temperature is nonlinear. In an Appendix, some results of Darken and Smith are presented for the rate of solution of nitrogen in iron using ammonia + hyidrogen mixtures. These data also support the view that diffudsion of nitrogen in iron is the rate-controlling process when ammonia + hydrogen mixtures are used. A considerable effort has been made to obtain data on the solubility1-5 and diffusivity of nitrogen in a iron6-l2 because an understanding of the effect of nitrogen on the properties of steel must be based on an accurate knowledge of the properties of nitrogen in pure iron. However, no information is available concerning the kinetics of solution of nitrogen in a and 6 iron. Recently the authors13 have investigated the rate-controlling mechanism operating in the kinetics of solution of nitrogen in y iron. This study was directed to determine the rate-controlling processes for similar reactions with a and 6 iron as well as to establish values for the solubility of nitrogen in equilibrium with nitrogen gas in a and a iron. EXPERIMENTAL The procedure used in experiments to determine the rate of solution in cylindrical iron rods was the same as that described in a previous communication.13 Ferrovac E grade iron used in all experiments contained the following impurities in weight percent: C, 0.005; Mn, 0.001; P, 0.002; S, 0.006; Si, 0.006; Ni, 0.025; Cr, 0.002; V, 0.004; W, 0.02; Mo, 0.01; Cu, 0.001; Co, 0.01; 0, 0.007. After cleaning the surface, the iron rods were treated in an atmosphere of purified hydrogen for 17 hr before the reacting gas was introduced for known experimental times. After quenching, the samples were sectioned radially and analyzed for nitrogen. In addition to experiments using rods, iron foils were used in the measurements of solution rates of nitrogen in a iron. The foils of two different thicknesses were prepared by cold rolling Ferrovac E grade iron cylindrical rod to 0.051 and 0.152 cm. Foil samples were used in a rectangular form 5 cm long and 1.25 cm wide. The specimens were thoroughly cleaned of surface oxide with fine emery cloth and degreased with carbon tetrachloride immediately before entry into the furnace. The experimental procedure was the same as that used in the study with rods. At the completion of an experiment, the foil samples of the nitrogenized iron were analyzed for nitrogen after discarding 0.3 cm from the perimeter of the specimen. Iron foils were nitrogenized and denitrogenized in the a and 6 range with a gas mixture of 95 pct N and 5 pct H for times varying from 5 min to 2 hr. Results obtained for the average composition of nitrogen in iron for these experiments are presented in Fig. 1. Prior to the denitrogenization experiments, the samples were saturated with nitrogen at 1000°C and 0.67 atm N, giving a uniform nitrogen concentration of 0.0204 pct. According to the known a-y phase boundary in the Fe-N system,14 this composition lies within the ferrite region at temperatures 750" to 850°C. Use of this initial nitrogen content insured that reaction occurred between the gas and a single solid phase, a iron. Examples of the results for the mean concentrations of nitrogen in cylindrical iron rods, 0.356 cm radius for both the a and 6 ranges are given in Fig. 2. Typical examples of the results obtained for the radial distributions of nitrogen in rods are presented in Fig. 3. It appears that the results for radial distributions can be extrapolated to constant surface compositions which agree with the equi-

Jan 1, 1964

By J. G. Mravec

The SO-called gas-tube method as developed by Hare, Peterson and Soler for determining the type and content of gases in molten steel is particularly adapted for determining the hydrogen content in molten steel. The development of this method has been quite completely discussed in a paper published by these co-workers in 1937.' The purpose of this paper, therefore, is to familiarize with this method those who arc not well acquainted with it. The principle of the gas-tube method is as fol1ows: One end of an evacuated gas tube is sealed with a very thin copper disk. when this end of the tube is immersed in a test spoonful of molten steel, the copper disk melts almost immediately, allowing a portion of the steel to be drawn up into the evacuated space. Any dissolved gases, and those produced by reactions between the components of the steel during freezing, are liberated into the unfilled portion of the sampling tube. A blob of steel that forms on the immersed end of the tube after the sample is drawn up seals up the end, making it airtight again. The evolved gases are subsequently drawn out of the gas tube through a suitable valve and analyzed by means of an Orsat apparatus. The volume of the sampling tube is sufficiently large so that a pressure considerably below atmospheric is produced after the gas evolution. The low pressure is necessary to effect a nearly complete removal of gases from the steel. Construction oF Sampling Tube The sampling gas-tube apparatus is illustrated in Fig. I. Seamless S.A.E. tubing used for section A is first cut into 7-ft. lengths. Through cxperiencc it was found necessary to use two sizes of tubing for this part, "is-in. and 3/4-in., respectively. The smaller tubing is generally used on the hotter or lower carbon steels. If the inside wall of the tube A is coated with grease or oil, it should be burned out before being machined as follows: One end is machined on the inside diameter to a depth of 34 in., so as to form a seat for a 7/8-in copper disk, which is stamped out from 0.005 to 0.007-in. thick copper sheet. The other end is machined on the outside diameter to a diameter of 1 5/16 in., to take a small reservoir or trap member I?. About 4 in. below the same end, a 1/2-in. hole is drilled to receive the 1/2-in. copper tubing leading to the large reservoir C. The inside of tube A is then thoroughly cleaned with a round wire brush, such as is generally used to clean out shotgun barrels. It is very important that the surface be absolutely free of rust or any nonmetallic coating. Next, the thin copper disk is soldered in place with low-melting silver solder.

Jan 1, 1945

By F. X. Tartaron

The comparison of actual energy of comminution with theoretical surface energy presents a wide gap. On the other hand, Solid State Theory presents a viewpoint that places actual energy of breakage in understandable relationship with the theoretical strength of materials. Therefore the physical background of comminution theory favors Kick's Law. A11 discussions of comminution theory appear farfetched when comparison is made of the actual energy of crushing with the energy input called for by thermodynamic theory. Quoting Taggart:' "The work required to crush halite was of the general order of 0.10 kg cm per sq cm of new surface produced (based on air percolation) whereas the theoretical energy of this new surface, based on thermodynamic calculation by a number of investigators is 1 x 10'4 kg cm per sq cm or the work expended was 1000 times the new surface energy — an indicated crushing efficiency of 0.10 pct." It is frustrating to probe into comminution theory with the prospect that the energy one is measuring is 99.9 pct waste energy. Do the laws of comminution deal with energy consumed by mineral breakage or with waste heat energy? There seems to be something wrong. Solid State Theory may provide the answer. The following statement is particularly apt to the situation: "The fundamental fact about metals and indeed all materials, is that they are much weaker than they should be. By this, we mean that if the stress required to cause a given plane of atoms to slip past a neighboring plane is calculated from basic atomic and Solid State Theory, the resulting number is some 1000 times too large to be accounted for by errors in approximations or assumptions. We know now that this is due to dislocations, and we even have a fair understanding of the details of the process by which dislocations make a crystal weak. An intriguing possibility which suggests itself, results from the discovery of very minute metallic crystals which are apparently free of dislocations and which have a strength approaching the ideal value for a perfect crystal." Here we have an extraordinary reversal of the situation. The actual energy for breakage is only one-thousandth the theoretical instead of 1000 times as given by Taggart. But the theoretical energy presented by Solid State Theory is the binding energy between atoms (or ions) in a crystal lattice. This is the kind of energy envisaged by Kick's law. The surface energy or the other hand corresponds to Rittinger's law. Comparison of theoretical with actual, on the basis of surface or on the basis of Rittinger's law, seems to present an absurdity; yet Rittinger's law is confirmed by an impressive array of experimental evidence. Comparison of theoretical and actual on the basis of Solid State Theory or, in other words, on the basis of Kick's law seems reasonable; yet there is not a scrap of evidence to verify Kick's law. The situation appears paradoxical, but the developments in the next section of this paper show that the paradox is a hint to a strange turn of events. ENERGY-WEIGHT FRACTION FUNCTION charles3 has presented a final equation showing the relation of energy consumed in comminution to size of product n{-1) a-n + 1 E is energy, C is proportionality constant, a and n are constants and k is the coarsest size in the ground product (size modulus). In the course of study of a large amount of experimental evidence, Charles discovered that in the general case, a-n + 1 approximated zero. For some reason, a and n differed only by unity. But if a—n + 1 = 0 within experimental error, then it is obvious that Eq. 1 states that it requires an infinite amount of energy to grind any sample to any size. This is a case where nature gives an absurd answer to a man-made mathematical equation, hence something is wrong with the equation. An adjustment must be made to fit the equation with experience. TO reveal the nature of the adjustment, the differential equation preceding the second integration in Charles' group of equations is given When Charles integrated the quantity xa-n it was almost impossible for him to realize that he was dealing with a special case of integration and that

Jan 1, 1962

By Robert L. Braun, Arthur E. Lewis

A contained nuclear explosion is proposed to produce a chimney of broken ore well below the water table. After the chimney is filled with water and reaches hydrostatic equilibrium, oxygen, under pressure, is introduced near the bottom of the mass of broken rock. The increase in solubility of oxygen at high hydrostatic pressure is sufficient to initiate the oxidation of the primary sulfides, chalcopyrite and pyrite. The oxidation and dissolution of these sulfides produces enough heat to increase the temperature of the ore and water in the chimney as much as 100° to 150°C. The rate of dissolution of chalcopyrite exposed to oxygen-saturated solutions becomes so rapid under these conditions that the rate of copper extraction becomes limited by the accessibility of chalcopyrite to the solutions. Means of solution control and distribution of oxygen in the chimney have been examined in some detail. A calculational model utilizing experimentally determined solubility rates is utilized to estimate the temperature rise in the chimney and the rate of copper recovery as a function of initial temperature, sulfide content of ore, and size of ore particles. It is concluded that the high pressure and high temperature obtainable in a large ore mass broken by nuclear explosives promise the rapid and economic recovery of minerals not considered soluble by conventional in-situ leaching methods. The method should be broadly applicable to the kinds of deep deposits that must be depended on for future production and with minimum disruption of the environment.

Jan 1, 1974

By R. Kossowsky, W. C. Johnston

CONSIDERABLE experimental data has been published in recent years on a few directionally solidified aluminum base eutectic alloys. The mechanical properties of A1-Al3Ni eutectic have been studied by Lemkey, Hertzberg, and Ford1',2 and Salkind et a1.3 where the alloy has been demonstrated to behave as a whisker reinforced composite with superior properties over many aluminum alloys. It has also been shown4 that aluminum reinforced with Al3Ni whiskers R. KOSSOWSKY, Member AIME, is Senior Metallurgist, Westing-house Research Laboratories, Pittsburgh, Pa. W. C. JOHNSTON, formerly with Westinghouse Research Laboratories, is Head, Department of Physics, George Mason College, University of Virginia, Fairfax, Va. Manuscript submitted March 28. 1969. IMD__________________ possesses excellent microstructural stability to 508°C. The effect of lamellar structure in A1-CuA12 eutectic was compared3 to the rod-like structure in the A1-Al3Ni eutectic, and was found to be better in low cycle fatigue. This was rationalized3 by postulating that a lamellar reinforcement was more efficient at blocking lateral plastic flow than a rod-like rein-forcement. The purpose of this study is to explore eutectic strengthening in other aluminum base alloys. The Al-Pd system5 forms a eutectic with aluminum and the intermetallic compound A13Pd at 25 wt pct Pd, and has a melting point of 615°C. A13Pd has an orthorhom-bic structure and transforms to hexagonal Pd2A13 at 785°C. Growth characteristics, compression yield strength, creep rupture, and microstructural stability of directionally solidified and as-cast alloys are investigated. Eutectic ingots having a nominal composition of 25 wt pct Pd were prepared by first compacting 1/2 in. diam by a in. high pellets from 99.99+ pct A1 and 99.99 pct Pd powders. The pellets were then levitation cast into 1/4 in. diam by 2 in. long copper molds. Directionally solidified bars were prepared by grinding the cast ingots to 0.195 in. diam, inserting them into alumina

Jan 1, 1970

By W. F. Holbrook, C. E. Wood, E. P. Barrett

THE high-frequency laboratory induction furnace with a rotating liquid crucible enables research workers to conduct certain investigations heretofore very difficult or impossible to realize because vessels are not obtainable that are capable of resisting the physical conditions or the chemical actions to which they are subjected. Applications for patents covering commercial uses of centrifugal liquid crucibles were filed in the United States in 1919,1 in France in 1925, in the United States and Germany in 1926,2 and in France in 1927.3 In 1928, Schuette and Maier,4 of the Bureau of Mines, reported the construction of a whirling table beneath a high-frequency induction furnace. SHAPE AND SIZE OF CENTRIFUGAL LIQUID CRUCIBLES When a vessel containing a liquid is rotated, the inner surface of the liquid takes the shape of a paraboloid under the simultaneous action of centrifugal force and gravity. The internal dimensions of the paraboloid rotating liquid crucible are governed by the diameter of the container and .its speed of rotation. Melted Wood's metal was poured into cylinders 2 1/2 in. in diameter, which were rotated at 180, 300, and 468 r.p.m. When the metal was solid, a slurry of plaster of Paris was added, and rotation continued until the plaster had set. Fig. 1, a photograph of sections of the cylinders, shows the effect of speed of rotation of the containers upon the shape of the inner surface of the rotating liquid crucibles. TYPES OF FURNACES Three types of high-frequency induction laboratory furnaces using liquid-metal crucibles were designed and constructed by the Blast

Jan 1, 1938

By Oscar A. Glaeser

The purpose of this paper is to present specific applications of the fundamental principles of mine ventilation and the means that are at our disposal to accomplish the task of maintaining a healthful atmosphere in small mines and prospect openings. Need for Ventilation The range between the maximum oxygen content of normal air and the minimum safe content is extremely narrow, approximately 5 per cent. Furthermore, normal air must not be contaminated with irrespirable gases beyond definite limits, and mineral dusts are equally dangerous, although health impairment is slower. Tables 1 and 2 present the more common respiratory hazards encountered in mines. Adequate ventilation is the most effective means available to reduce or eliminate these hazards. Mechanical ventilation has reached varying stages of perfection in coal mines and in many large metal mines. The small mines and prospects rarely have a definite plan of ventilation. As a consequence underground workers are frequently exposed to health and accident hazards. The annual reports of fatalities in mines in the Western States record yearly many deaths due to asphyxiation and deaths caused by men falling down shafts and raises after being overcome by gas. At times mining operations are slowed up because there is no means of driving out the gas after blasting. For that reason only one shift is worked, or men wait from one to several hours for the "smoke" to clear out. Methane gas in varying quantities is widely distributed throughout California. It has been encountered in explosive concentrations in the coast range. It has been found on the Mother Lode and in quicksilver mines near the coast. Methane explosions have not been uncommon in underground operations and in tunneling jobs. These conditions and records indicate that there is room for improvement in mine-air conditions.

Jan 1, 1937

By H. H. Hausner, H. S. Kalish

IN recent years zirconium and beryllium have become of great interest because of their special properties. Zirconium is known for its remarkable ability to absorb the gases oxygen, nitrogen and hydrogen, has extensive use as a getter in electron tubes and has indeed become the primary metal in the manufacture of photoflash bulbs. Beryllium is of interest because it is light metal and as such could supplement aluminum and magnesium in light metal applications. The late development of zirconium ductile at room temperature produced by both the iodide and calcium reduction methods has stimulated frequent articles on this metal in the recent literature. Beryllium is so closely akin to zirconium crystall-ographically that some investigators are of the opinion that beryllium should be ductile at room temperature if it were pure enough. The crystallographic relationship of beryllium and zirconium, the interest created in Be-Zr alloys where the zirconium is believed to function as a remover of solid solution gas in beryllium and thus increase the ductility of the beryllium rich alloy, and the amenability of these metals to powder metallurgy methods stimulated this approach to an investigation of the system Be-Zr. Preliminary work showed that a zirconium compact melted when pressed in contact with a beryllium compact and sintered at 1250°C, without any apparent final bond between the two. A search of the literature revealed nothing pertinent about the Zr-Be system or about any zirconium - beryllium alloys. In 1936 Misch³ vaguely identified the compound ZrBe2 and reported that it has a deformed cubic structure. Two patents," , were issued to Mach-lett Laboratories in 1943 and in 1946 regarding the addition of small amounts of zirconium to beryllium in order to improve the ductility of beryllium so that it may better be used for X ray windows. In considering the alloy formation of beryllium with other metals Raynor thought it unlikely that zirconium would have any solid solubility in beryllium. From this indefinite information nothing could be ascertained about the possible use of Zr-Be alloys or the use of these metals intimately. A brief preliminary study of the system Zr-Be by powder metallurgy methods was undertaken to provide regions for more intense studies of Zr-Be alloys and as a guide to a formal phase diagram determination. The investigation described in the following pages was completed in less than three weeks. The usefulness of this rapid method to facilitate phase diagram studies and new alloy investigations will be shown. Materials and Procedure Zirconium powder was produced from crystal bar stock of high purity supplied by Foote Mineral Co. The crystal bar was converted to hydride at 730°C, crushed, and the hydride decomposed in purified argon at 850°C. 60 pct of the powder passed through a 325 mesh screen, 40 pct of the powder was —140 mesh, +325 mesh and the average particle size of the total was 8 to 12 microns. The beryllium powder was used as re-

Jan 1, 1951