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By David L. Holt, Walter A. Backofen, Anwar-ul Karirn

ANNEALING the binary alloy, Mg-0.5 wt pct Zr in a hydrogen atmosphere results in extensive precipitation of zirconium hydride. The morphology of these precipitates has already been investigated.lm3 Squires, Weiner, and phillips4 found, following creep testing of the hydrided alloy at 450" and 500°C, that hydride denuded zones formed along grain boundaries that were aligned approximately normal to the tensile axis. Zone formation was attributed to diffusional creep by magnesium transport down the prevailing stress gradients.4'5 Recently, electron microprobe analysis by Karim, Holt, and Backofen6 showed the zones indeed to be zirconium-free and furthermore it was shown that in compression, zones formed along boundaries approximately parallel to the stress axis. Thus it was confirmed that the zones are formed by diffusional creep. Typical zone-formation is shown in Fig. 1. Karim, Holt, and Backofen also measured the con- tribution to the total strain, from diffusional flow, slip creep, and grain boundary shear: results showed that diffusional strain tends to saturate with increasing imposed strain, while grain boundary shear eventually becomes the dominant deformation mode. It was suggested in explanation that the precipitate-denuded zones are soft and easily sheared, such deformation contributing to the grain boundary shear measurement. A point to be emphasized is that the Mg-0.5 Zr alloy was not superplastic for the index of strain rate sensitivity of flow stress, m = a log log 6 was only 0.25. It was recognized that diffusional creep in a poly crystalline aggregate must be accompanied by boundary shearing.' The suggestion was that the non Newtonian-viscous character of diffusional creep in the Mg-0.5 wt pct Zr alloy is due to resistance to sliding of boundaries lying approximately parallel to the tensile axis.= The work described in the present note should be seen against the above background for it represents initial efforts to observe in a similar way diffusional creep in a superplastic alloy. Previous work with the ternary-alloy Mg-6 Zn-0.5 Zr, pellet extruded, showed it to be superplastic however it does not contain resolvable precipitates to act as internal markers for diffusional creep. Also the fine grain size of the alloy (0.5 p) made conventional metallography rather difficult. Solubility relationships in the Mg-Zn-Zr system are not well known, however making reasonable guesses about the probable phase diagram of the ternary, Mg-Zn-Zr, various heat treatments were attempted to produce hydride precipitation. Finally a 16-hr anneal at 495' in hydrogen was found to pro-

Jan 1, 1970

By N. Standish

By contrast with viscometry and sedimentation, no actual measurements of the applicability of Stokes' law to steelmaking have ever been reported; instead, the proof for and against Stokes' law has been inferred from the frequently observed phenomenon that inclusions collect in the top scum and the bottom of ingots. As such proof is unsatisfactory, continued reference to Stokes' law in steelmaking, even as a model for calculating the terminal velocities of rising inclusions, should be avoided. In fact, all conclusions that have been reached regarding size, density, and viscosity on the flotation of inclusions could be made more rationally without Stokes' law. A worthwhile exception is that of Mitsche1 who in 1934 used Stokes' law to show not only the expected effects of size, density, and viscosity on flotation but also the beneficial effects of i) melting under a slag layer (now Electro Slag Refining); ii) treating melts with synthetic slags (Ref. 2 gives recent applications); iii) bubbling an inert gas through the melt (recent applications in Refs. 3 and 4); and iv) altering the solid/ liquid interface characteristics. At this point, it should again be emphasized that the boundary conditions for which Stokes' law was derived place enormous restrictions on the law but its application, as noted earlier, can yield useful results in some fields under certain conditions and within well-defined limits. When compared with inclusions, which may be solid or liquid, spherical or irregular, dispersed or coagulated, interacting between themselves and the melt, floating in a liquid which is in a state of motion of varying intensity5 and contained in vessels of various shapes and sizes that may change with time as in a solidifying ingot, the conclusion is obvious-Stokes'

Jan 1, 1969

Abramovits, Charles, Student, Case School of Applied Science Cleveland, Ohio. '37 Abreu, Ovidio M., Student, Stanford Univ Stanford University. Calif. '38 Adams, Albert, Student, New Mexico School of Mines Socorro, N. M. '38 Adams, Benjamin C., Jr Address Wanted. '36 Adams, George H 44S2 4th St., Riverside, Calif. '36 Adams, Robert E., Student, Univ. of Wisconsin Madison, Wis. '38 Adkisson, J. F., Jr., Student, Agricultural & Mech. College of Texas, Box 2044, College Station, Tex. '38 Afshar, Hodji khan A., Amiranian Oil Co 10 Khiahan-i-Nusrat, Teheran, Iran. '35 Agvas, Cesar G., Student, Adamson School of Industrial Chem. & EngrgManila, P. I. '38 Ahlen, Louis S., Student, Univ. of North Dakota Univ. Sta., Grand Forks, N. D. '37 Albers, Francis C., Student, Univ. of Wisconsin Madison, Wis. '38 Alden, George P., Student, Univ. of Tulsa Tulsa, Okla. '38 Alexander, Donald B... Student. Univ. of Oklahoma, Sigma Nu House Norman, Okla. '37Alexander, Joe M., Student, Agricultural & Mech. College of Texas. Box 954. College Station. Tex. '38

Jan 1, 1938

By D. H. Polonis, J. G. Parr

IN a recent paper on titanium-rich Ti-Fe alloys,' the hardness of quenched powder specimens was given, together with the amounts of martensitic a they contained. The values disagreed in two respects with previous determinations that Worner² made on lump specimens: 1—Some martensitic a' is formed in powder specimens of alloys containing up to 12 atomic pct Fe. On the other hand, Worner reports that, although his results on quenched lumps are not conclusive, alloys containing more than about 4 pct Fe "appear to consist of retained ß'' although they "may not be strictly unaltered solid solutions." 2—According to Worner, quenched lump specimens show a maximum hardness at a composition about 3.5 pct Fe. Earlier -experiments by the authors showed a maximum hardness of martensitic alloys in powder form at 10 pct Fe. Accordingly, more experiments have been conducted in order to confirm these differences and to determine with more accuracy the minimum composition of quenched lump specimens that consist entirely of retained ß. Lump samples?x?x1/20 in. were quenched from 1000°C in helium and, on the whole, showed similar hardness values to the specimens that Worner water-quenched from 950°C. The hardness curve the authors obtained shows a more pronounced peak, but in other respects it is very similar to Worner's curve (Fig. 1). While it was possible to make fairly reliable measurements of the percentage of martensitic a in quenched powders (by Geiger spectrometer), the solid samples allowed only approximate estimates to be made from visual observation of line intensities on film. Retained ? begins to appear in quenched lumps containing 0.2 pct Fe; at 2.2 pct Fe there are approximately equal quantities of ß and a'; in a 4.7 pct Fe alloy only a trace of a' appears. These results are represented as a curve on Fig. 1. The hardness and constitution of powder samples were verified by repeating both types of quenching experiment described in the earlier paper.' Identical results were obtained by the two methods, and they agreed with the figures published previously (Fig. 1). The appearance of the martensitic structure under the microscope is shown in Figs. 2 and 3. The mar-tensite needles are coarser in the lump; otherwise the structures are similar. Three conclusions are made at this juncture: 1— The properties of quenched lumps reported by Worner are confirmed, on the whole. The differences between them and the properties of powders are probably due to an effect of specimen size. Differences of behavior in various sizes of specimen have been observed previously in martensite reactions in Fe-Mn alloys3 and in Cu-A1 alloys. 2—The maximum hardness of both quenched lumps and quenched powders appears at a composition where martensite formation has all but ceased. In this respect, powders and lumps behave similarly. 3—Martensitic a' has essentially the same microscopic appearance both in powders and lumps. Acknowledgment This work is part of a program sponsored by the Defegce Research Board of Canada, Project No. 425.

Jan 1, 1956

By Lawrence H. Van Vlack

DISCUSSION, H. L. Burghoff presiding C. S. Smith (University of Chicago, Chicago)—The author is to be congratulated on his valuable contribution to the extremely meager absolute data on interface energies in metallic systems. It is interesting to note that the ratio between a and r iron boundary energies is approximately the same in the quaternary alloy with silicon, carbon, and sulphur as in the binary Cu-Fe alloy. A value for the copper grain boundary energy can be obtained from the author's data, Tables IV and VI, by utilizing the energy ratios of the following pairs of interfaces: copper-sulphide/copper, copper y iron, y iron/y iron, y iron/ copper, and copper/copper. This gives a value of 595 erg per sq cm for the free energy of the copper grain boundary. Despite this circuitous route and the assumption that the y iron boundary is identical in alloys with and without sulphur and does not change between 1105" and 100O°C, this figure is in remarkably close agreement with other values of the copper boundary energy, obtained by equilibration against lead and against the free copper surface. It therefore appears either that impurities do not greatly affect the grain boundary energies or that boundaries in even nominally pure metals are already sufficiently contaminated so that further change has little effect.

Jan 1, 1952

By J. P. Zannaras

Referring to the article by R. J. Charles and P. L. de Bruyn, let us assume that W = weight of glass bar; P = weight of hammer; e = total deformation; K = unit of deformation; K = potential stress energy; E = modulus of elasticity; L = length of the bar; 7 = coefficient of inertia; h = height, ft; V = velocity, fps; and a = cross section area. The portion of kinetic energy which is effective in producing stress energy in a fixed bar struck horizontally is given by the formula" 1 + 1/3 W/P K =1+1/3w/p/(1+1/2w/p)2 P.V2/2g = Ph where 1 + 1/3 W/P ? (1 + 3W/P/(1=1/2w/p)2 [8] Putting e e = W/P =------------ From the above equation it can be seen that the maximum transfer of kinetic energy to stress energy is when e = 0 or W/P = 0 which indicates that the weight of the hammer must be very large as compared with the weight of the impacted rod. Eq. 8 diametrically opposes the conclusions reached by the authors of this article. In fact, if their suggestions were followed to the extreme when e = co when P = 0, there would be no transfer of kinetic energy to stress energy at all, as 7, becomes zero. Eq. 8 presumes that the velocity with which the stress is propagated through the bar is infinite, whereas the authors claim that the compression waves reflected are reaching the struck end of the bar prior to the complete transfer of the kinetic energy to cause such modification of the conditions there as to make them reach the reverse conclusions demonstrated by the above formula. That such interference exists is unquestionably demonstrated by the authors and others. However, if my observations are correct, such interference for this specific experiment and also for practical comminu- tion is insignificant, and the conclusions of the authors are in error and must be reversed to comply with Eq. 8. Eq. 5, w. = AE 2/2, given by the authors on page 51, is derived from the following equation (Eq. 9): K = 1/2Pe, where P = Sa, S = ?E, e = EL, and L = 1. The above formula, Eq. 9, cannot be applied in this case. This formula is applicable for static loads where the load increases from zero up to its final value, P, in such a way that the deformation at different instants is proportional to the loads acting at those instants and actually represents the area of a right triangle in the strain load diagram of base e and height P. The typical photographs shown in Figs. 3 and 4 represent the familiar strain load diagrams, and since the line of the wave marks the existence and intensity of the strain with the unquestionable conclusion that such strain has been caused by the action of a load acting continuously all along the wave until it reached the horizontal axis, the work stored at this point is represented by the area under the wave line and the horizontal axis and not by the area of the fictitious triangle given by the authors. Then if this is correct, even visual estimation of these areas at gage stations given in the typical photographs of Figs. 3 and 4 suffice to contradict the authors' calculation given in Figs. 6a and 6b and Figs. 7a and 7b. The typical photographs presented by Charles and de Bruyn show a considerable variation of the intensity of the strain at different stations but very small variation of areas which actually represent the stress energy at the corresponding stations. And, apparently, by squaring the small quantities, the authors magnified their error tenfold. J. M. Frankland's paperV iscusses the relative strain intensity and not the total energy for different types of impact loading. He states in his paper, "The reader is explicitly warned not to confuse the results in this report with those obtained when the load is applied by a blow as from a hammer. In this case the peak load rises to very large but mostly unknown values. The accompanying large deflections and stresses are the result of high values of P, not of the dynamic load factor n. According to Frankland "the dynamic load factor" is the numerical maximum of the response factor. It therefore appears that the authors followed the same procedure in obtaining the relative strain energy ab-

Jan 1, 1957

By Thomas O. Ziebold, Robert E. Ogilvie

Experimental analysis of diffusion samples covering the entire Cu-Ag-Au system at 725°C has been carried out. Experimental coefficients are computed at twenty-eight points in the lermary system. It is found that the direct coefficienl for copper is nearly independent of the silver, content for 1ow-silver alloys and that the cross coefficient for silver correlates with the ther-nzodynatnic properlies. This lalter observalion leads to the conclusion that the mobililies of copper and gold are nearly equal. Thermodynarmic actiuilies for Cu-Ag-Au were computed from binary data and adjusted to be consistent with neasured tie lines across the two-phase region which extends into the ternary diagram at 725°C. Application of these therknodynarnic calculaliorls to the dijftcsion data terifies the Onsager reciprocal relations to the extent that experimental uncertainlies will allow. It is jound that diffusion neav the critical point of the 1200-phase field in the isothermal section at 725°C causes the composition gradients for copper and silver to approach infinity while the direct and cross coeficients become equal. From the standpoint of experimental Procedures, it has been shown that microanalysis of ternary samples may be carried out using a simple expression for the conversion of X-ray zntensily to composilion. The accuracy of' this method has been demonstrated by the analysis of standard alloys and by the reproducibility of composition determinations using differed X-ray lines. DIFFUSION in a ternary alloy system is basically and significantly different from diffusion in binary alloys. When only two elements are present there is just one independent composition variable, and the diffusion profile of composition vs distance across a sample must change monotonically through all solid solutions bracketed by the terminal alloys. If two-phase regions are stable for compositions between the terminal values, these must appear as planar interfaces in the diffused sample, with the composition profile exhibiting sharp discontinuities. Adding a third element may alter this picture completely. Because of the additional degree of freedom in the composition, we may see composition profiles which do not change monotonically through solid-solution ranges, and stable, nonplanar phase interfaces. Even though two-phase regions may lie between the terminal compositions of a diffusion sample on the constitution diagram, they will not necessarily appear in the diffusion profiles. In binary diffusion only one coefficient is needed to describe the interdiffusion process. From Fick's law this coefficient relates the diffusive flux of one component to its own composition gradient. In ternary diffusion we must allow for cross interaction between the two independent species. By a linear extension of Fick's law we introduce "cross" or "off-diagonal" diffusion coefficients which relate the flux of one component to the composition gradient of the other independent constituent. Thus, we require four interdiffusion coefficients to describe the transport process in a ternary system. In its most general form,1,2 the rate equation for mass transport by isothermal diffusion in a system of s components states that the diffusive flux of component i, Ji, is a linear summation of all driving forces, Xk, each force being multiplied by an appropriate compliance coefficient:

Jan 1, 1968

By M. C. Flemings, R. Mehrabian, G. E. Nereo

Analytic expressions were developed and applied in two preuious papers to predict effects of solidification variables on macrosegvegation. In this paper, experiments are reported to test the analyses quantitatively and qualitatively. Good quantitative agreement is obtained between predicted and measured compositions for the following types of segregation (all experiments were on laboratory-size ingots of Al-4.5 pct Cu alloy): (a) inuerse segregation, (b) composition distribution in ingots solidified with unidirectional heat and fluid flow, (c) simulated centerline segregation, and (d) segregation resulting fro, change in solidification cross section. In addition, qualitative agreement of prediction of theory with experiment is obtained for (a) segregation resulting from sudden change in heat extraction during solidification, (b) centerline segregation in a bidirectionally solidified ingot, and (c) under-riser positive segregation. IN two previous papers,''2 it was shown that a variety of apparently different types of macrosegregation result from the same basic mechanism. This mechanism is the flow of solute -rich liquid to feed solidification and thermal contractions. Analytic expressions were derived and examples given of their application to predict macrosegregation. Calculations were for Al-4.5 pct Cu alloy. In this paper, we describe a series of experiments designed to test results of the foregoing calculations. All experiments were on laboratory-size ingots (weighing 10 lb or less), of A1-4.5 pct Cu alloy, nominal composition. APPARATUS AND PROCEDURE A typical mold and chilling arrangement used for unidirectional solidification is shown in Fig. 1. Before casting, this assembly was placed in an electrically heated air furnace, and heated to 680" to 700°C. When the plaster reached the desired temperature (usually 1 to 2 hr after the furnace temperature), the mold was filled through a funnel extending from the outside of the furnace into the cavity. The assembly, with liquid, was held for several minutes to allow convection to subside. The coolant, either air or water (depending on the desired solidification rate), was turned on and the ingot solidified. With the four sides and top insulated, heat removal for solidification was through the bottom surface, resulting in a fully columnar structure. Design of the casting cavity was varied, as discussed below, to produce desired macrosegregation effects. However, for castings solidified with unidirectional heat flow, the top was always open to the furnace atmosphere and temperature. The bottom sur- face rested on a chill through which the heat was extracted, and all other surfaces were insulated with plaster. In one case to be discussed, the mold design was changed to obtain bidirectional solidification. In this case, two vertical chills were used (on two faces of the mold) and the bottom was insulated. In a last case discussed, that dealing with under-riser segregation, experiments were conducted in bottom-chilled sand molds at ambient temperature. Except for this one case, all experiments were conducted in the electrically heated air furnace as described above. In many ingots, thermocouples were used to obtain continuous records of solidification. In all cases, these were chromel-alumel couples with output continuously monitored on a twelve-channel recorder. Where thermocouples were employed, these were embedded in the refractory mold walls during mold-making so their heads extended to the vertical centerline. Six to nine couples were employed, placed at intervals from the chill. METAL-MEL TING PROCEDURE Aim analysis was A1-4.5 pct Cu; alloy was prepared from high-purity aluminum (99.9 pct Al) and A1-Cu master alloy (50 pctOFHC Copper, 50 pct high-purity Al). The A1-4.5 pct Cu alloy was cast in 5- to 7-lb ingots for subsequent remelting for the macrosegregation experiments. Melting for ingot casting was as follows. After preheating the melting crucible for an hour or more, the charge was melted and superheated to a temperature of 690" to 720°C. After stirring, the melt was degassed with chlorine for 10 min. After stirring again for 1 min, a gas sample, under reduced pressure, was taken. When the sample was gas-free, the crucible

Jan 1, 1969

By Henry D. Hibbard

THIS article aims to clarify the use of some terms often occurring in writings on iron and -steel, and also to suggest several new short abbreviated names for some of the things related to the subject which now are indicated by phrases. Crude or Crude Iron The name crude or crude iron is recommended for the iron produced from ore by the blast furnace. When necessary it can be specified as pig iron when in the form of pigs (usually cold), or as molten iron, commonly called hot metal to distinguish it from cold iron in pigs. The product of a direct process of making wrought or malleable iron from ore is not included.

Jan 1, 1924

By Robert Ammon

The purpose of this paper is to present a brief sketch of the development of this hindered-settling classifier, but primarily to show the actual results obtained in practice with the classifier working on the Butte copper ores at the Boston & Montana Concentrator at Great Falls and the Washoe Concentrator at Anaconda. The writer has endeavored to be as brief as possible, giving only the essential figures showing the efficiency of the classifier and leaving considerable interesting information to be derived by the reader from the statistics given. The paper deals with the practical work of the classifier under the following heads: 1. Desliming 4.0-mm. primary feed, overflow from 0.07 to 0.0 mm. 2. Desliming 1.25-mm. secondary feed, overflow from 0.07 to 0.0 mm. 3. Desliming Huntington Mill discharge, overflow from 0.07 to 0.0 mm. 4. Classifying 2.5-mm. deslimed secondary feed, overflow from 1.00 to 0.07 mm. 5. Classifying 2.5-mm. deslimed primary feed, overflow from 0.75 to 0.07 mm. 6. Classifying 4.0-mm. deslimed primary feed, overflow from 0.75 to 0.07 mm. 7. Classifying middling-tailing product from roughing tables, overflow coarse tailing. 8. Classifying 2.5mm. trommel undersize, not deslimed, overflow from 0.75 to 0.0 mm. 9. Remarks on the inner feed cone. In the following pages all overflows are given in quartz dimensions; i. e., by a 0.75-mm. overflow is meant an overflow whose maximum particle is a 0.75-mm. quartz grain. By the term "slime" is meant all material which will pass through a 200-mesh screen whose average opening is 0.07 mm. The development of the classifier was commenced in the fall of 1910,

Jan 1, 1914

By W. M. McKewan

Magnetite pellets were reduced in flowing hydrogen at pressures up to 40 atm over a temperature range of 350° to 500°C. The rate of weight loss of oxygen per unit area of the reaction surface was found to be constant with time at each temperature and pressure. The reaction rate was found to be directly proportional to hydrogen pressure up to 1 atm and to approach a maximum rate at high pressures. The results can be explained by considering the reaction surface to be sparsely occupied by adsorbed hydrogen at low pressures and saturated at high pressures. PREVIOUS investigation1,2 have shown that the reduction of iron oxides in hydrogen is controlled at the reaction interface. Under fixed conditions of temperature, hydrogen pressure, and gas composition, the reduction rate is constant with time, per unit surface area of residual oxide, and is directly proportional to the hydrogen pressure up to one atmosphere. The reduction rate of a sphere of iron oxide can be described3 by the following equation which takes into account the changing reaction surface area: where ro and do are the initial radius and density of the sphere; t is time; R is the fractional reduction; and R, is the reduction rate constant with units mass per area per time. The quantityis actually the fractional thickness of the reduced layer in terms of fractional reduction R. It was found in a previous investigation2 of the reduction of magnetite pellets in H2-H,O-N, mixtures, that the reaction rate was directly proportional to the hydrogen partial pressure up to 1 atm at a constant ratio of water vapor to hydrogen. Water vapor poisoned the oxide surface by an oxidizing reaction and markedly slowed the reduction. The enthalpy of activation was found to be + 13,600 cal per mole. It was also found that the magnetite reduced to meta-stable wüstite before proceeding to iron metal. The following equation was derived from absolute reaction-rate theory4,8 to expfain the experimental data: where Ro is the reduction rate in mg cm-2 min-'; KO contains the conversion units; Ph2 and PH2O are the hydrogen and water vapor partial pressures in atmospheres; Ke is the equilibrium constant for the Fe,O,/FeO equilibrium; Kp is the equilibrium constant for the poisoning reaction of water vapor; L is the total number of active sites; k and h are Boltzmann's and Planck's constants; and AF is the free energy of activation. Tenenbaum zind Joseph5 studied the reduction of iron ore by hydrogen at pressures over 1 atm. They showed that increasing the hydrogen pressure materially increased the rate of reduction. This is in accordance with the work of Diepschlag,6 who found that the rate of reduction of iron ores by either carbon monoxide or hydrogen was much greater at higher pressures. He used pressures as high as 7 atm. In order to further understand the mechanism of the reduction of iron oxide by hydrogen it was decided to study the effect of increasing the hydrogen pressure on rebduction rates of magnetite pellets. EXPERIMENTAL PROCEDURE The dense magnetite pellets used in these experiments were made in the following manner. Reagent-grade ferric oxide was moistened with water and hand-rolled into spherical pellets. The pellets were heated slowly to 550°C in an atmosphere of 10 pct H2-90 pct CO, and held for 1 hr. They were then heated slowly to 1370°C in an atmosphere of 2 pct H2-98 pct CO, then cooled slowly in the same atmosphere. The sintered pellets were crystalline magnetite with an apparent density of about 4.9 gm per cm3. They were about 0.9 cm in diam. The porosity of the pellets, which was discontinuous in nature, was akrout 6 pct. The pellets were suspended from a quartz spring balance in a vertical tube furnace. The equipment is shown in Fig. 1. Essentially the furnace consists of a 12-in. OD stainless steel outer shell and a 3-in. ID inconel inner shell. The kanthal wound 22 in. long, 1 1/2, in. ID alumina reaction tube is inside the inconel inner shell. Prepurified hydrogen sweeps the reaction tube to remove the water vapor formed during the reaction. The hydrogen is static in the rest of the furnace. The sample is placed at the bottom of the furnace in a nickel wire mesh basket suspended by nickel wire from the quartz spring. The furnace is then sealed, evacuated, and refilled with argon several times to remove all traces of oxygen. It is then evacuated, filled with

Jan 1, 1962

Butte Meeting. It is expected that the technical program of the Butte meeting will be a sufficient attraction to draw to that point every member who can possibly get away long enough to attend. Attention is particularly called to the papers already sent in by the Committee on Precious and Base Metals, and the papers which they are undertaking to secure for the meeting, as published on pp. ix and x. In addition to these papers, and in addition to those published in this Bulletin, the office of the Institute already has in hand several papers, which will be published at an early date, and presented in abstract and discussed at the meeting. Of the discussion of these papers a special feature is being made. Therefore, let every member come to Butte who can do so. A meeting of the Board of Directors of the Institute will also be held at Butte. In the May Bulletin, we stated for the information of members the railway excursion rates for two suggested tours in connection with the meeting, which included Glacier Park before the meeting and Yellowstone after it. The itineraries there given and rates quoted were planned and offered by the railway agent bureaus in New York. Circulars have been sent to the members of the Institute by A. E. Vaughan & Co., a tourist trip agency of New York City, containing details of tours nearly identical with those outlined in the Bulletin, and quoting rates covering, the expenses of the whole trip. Members can easily compare these rates with the rates offered by the railroads, and published in the last Bulletin. In order that no misapprehension may exist, the statement is here made that the Institute and its officers have no connection whatever with Vaughan & Co.'s proposals and that any such tours will be strictly personal matters between said company and the members of the Institute who take them. Nevertheless, the Board of Directors calls to the attention of members the circumstance that tours such as these will afford an opportunity for members to travel together to the meeting.

Jan 6, 1913

By Howard L. Hartman, Eugene P. Pfleider

The paper relates to the laboratory and pilot plant studies that have been carried out by Sherritt Gordon Mines Ltd., Metallurgical Research Div., in developing the ammonia pressure leach process for extracting Cu, Ni, Co, and S from high grade Ni concentrate produced from Lynn Lake ores and describes in some detail the chemistry of the process. In the leaching stage air, ammonia, and sulphide minerals react in a sequence of inter-related steps to produce soluble amines of Ni, Cu, and Co in a solution containing sulphate, thiosulphate, polythionate, and sulphamate ions and a controlled amount of free ammonia, the iron and other components of the concentrate remaining as an insoluble residue which is separated from the leach solution by filtration. The pregnant solution produced in leaching is further treated by boiling to precipitate and remove Cu as cuprous sulphide, leaving a copper-free solution with a composition suitable fir subsequent precipitation of Ni and Cu by H,. The progress of individual reactions in leaching and copper removal is discussed. The effect of variables such as temperature, pressure, ammonia concentration, particle size, and agitation is outlined and the technique of thionate control in leach solutions is described. Data is given for the physical and chemical reactions involved in precipitating copper from leach solutions as Cu sulphide by boiling.

Jan 1, 1956

By Charles S. Barrett

Metallurgists are making use of the stereographic projection to a steadily increasing extent. In the last five years no less than 20 papers in American metallurgical journals alone have employed the stereographic projection; recent books on physical metallurgy, plastic deformation, X-rays and their applications, and crystal structure use it. Mineralogists have used the stereographic projection for many years in the description of symmetry classes and crystal planes, for it presents an accurate, easily understood plot of any angular relations in crystals with all unessential features eliminated (such as the accidentally determined size and shape of crystal faces). In metallography and physical metallurgy the projection is much used for the analysis of markings appearing on polished grains: slip lines, twins, cracks, structures formed by precipitation, magnetic-powder patterns, etch pits, etc. Data from certain types of X-ray photograms are most conveniently analyzed by its use, particularly those for determining the orientation of single crystals or the preferred orientation of grains in an aggregate. Calculations of how to tilt or cut a crystal parallel to a certain crystallographic plane or to reflect X-rays from a certain plane are rapidly carried out. It has been adopted almost universally to the exclusion of other methods by those studying the deformation of metallic crystals, save where the accuracy required is greater than a few tenths of a degree. Any directional property in a crystal or polycrystalline material can be shown on a stereographic projection; for example, the modulus of elasticity, yield point, etc. It has been repeatedly called to our attention that there is need of an elementary explanation of the principles and methods of stereographic projection, covering all the common applications in the field of metallography and physical metallurgy. The following contribution is an attempt to fill this need. The description of each method is supplemented with references to publications in which the method has been used. For convenience to the reader, these illustrative references are chiefly in American journals, particularly the publications of the American Institute of Mining and

Jan 1, 1937

By Lawrence Wright

ON the steep slopes of the Kootenays, 15 to 20 miles south of Nelson, B. C., are a number of ore occurrences (Fig. 1). One that has many features of mining and geological interest is the Ymir Yankee Girl Gold Mines Limited. Discovery of this property was made in 1899 at an outcrop on the hillside, some 2500 ft. above the village of Ymir. This proved to be a branch vein 20 ft. long leading to two main veins, the Yankee Girl and the Spur (Fig. 5). The ore pattern as now delimited shows a remarkable increase in ore area with depth below the outcrop. This recent development is of interest particularly in view of the fact that the property was nearly con-demned, professionally. Another interesting feature, which has a bearing upon its history, is the persistent excess of tonnage produced from period to period over visible reserves. This is not ascribed to lagging mine development, but to detailed vein characteristics, which we will endeavor to make clear by giving a description of the veins and their surroundings.

Jan 1, 1938

By Bernard Kopelman

The particle size of tungsten metal powder used to make tungsten wire for use in radio tubes and incandescent lamps must be closely controlled if the highly desirable feature of nonsagging is to be achieved. Our knowledge of factors affecting the particle size of tungsten powder has been limited by lack of any extensive literature on the subject. There are no publications in the English language discussing this problem in any but the most general terms. In this paper the published experimental information on this subject is reviewed. Based on the published data and some experimental data obtained by the author, a general theory of factors affecting the particle size of tungsten metal powder obtained by reduction of the oxides in hydrogen is offered. The various forms of the tungsten oxides and tungsten metal used in the experiments described were those prepared by commercial processes, so that the experimental results obtained on a laboratory scale would be more directly applicable to actual commercial reduction conditions. The effects of time of reduction, temperature. and concentration of water vapor in the hydrogen were observed by counting microscopically the particle distribution obtained. Such a method is very time-consuming, but the data obtained by this method are more reliable and more fundamental than those obtained by any of the several instruments available for measurement of particle size. Review of the Literature Until the year 1912, tungsten powder was prepared commercially by reducing tungstic oxide in the presence of carbon or one of its compounds (such as glycerin or ethylene glycol), usually by firing the mixture in a closed fire-clay crucible. The resultant powder contained about 95 per cent tungsten.' When reduction of tungsten oxides by H² was introduced Commercially,²,³ several investigators became concerned with the equilibria between tungsten metal and its several oxides in an atmosphere of hydrogen and water vapor.4-l3 While these many investigations were in fair agreement on the equilibria WO³ + H2 = WO2 + H2O and WO2 + 2H2 = W + 2H2O there arose considerable controversy as to the chemical composition of the oxide between WO³ and WO², because of very poor agreement on the equilibrium constant of this intermediate oxide reduction reaction. The controversy on this intermediate oxide has been in progress for more than 100 years. Thus Malguti,14 in 1835, decided it was WO2.56 In 1857, Riche15 proposed the formula W³O8. Usler,16 in 1865, confirmed the existence of WsO8 and further proposed W1O1l. In 1897, Desil7 agreed upon W²O8 but also found W5O14.

Jan 1, 1947

By Bernard Kopelman

The particle size of tungsten metal powder used to make tungsten wire for use in radio tubes and incandescent lamps must be closely controlled if the highly desirable feature of nonsagging is to be achieved. Our knowledge of factors affecting the particle size of tungsten powder has been limited by lack of any extensive literature on the subject. There are no publications in the English language discussing this problem in any but the most general terms. In this paper the published experimental information on this subject is reviewed. Based on the published data and some experimental data obtained by the author, a general theory of factors affecting the particle size of tungsten metal powder obtained by reduction of the oxides in hydrogen is offered. The various forms of the tungsten oxides and tungsten metal used in the experiments described were those prepared by commercial processes, so that the experimental results obtained on a laboratory scale would be more directly applicable to actual commercial reduction conditions. The effects of time of reduction, temperature. and concentration of water vapor in the hydrogen were observed by counting microscopically the particle distribution obtained. Such a method is very time-consuming, but the data obtained by this method are more reliable and more fundamental than those obtained by any of the several instruments available for measurement of particle size. Review of the Literature Until the year 1912, tungsten powder was prepared commercially by reducing tungstic oxide in the presence of carbon or one of its compounds (such as glycerin or ethylene glycol), usually by firing the mixture in a closed fire-clay crucible. The resultant powder contained about 95 per cent tungsten.' When reduction of tungsten oxides by H² was introduced Commercially,²,³ several investigators became concerned with the equilibria between tungsten metal and its several oxides in an atmosphere of hydrogen and water vapor.4-l3 While these many investigations were in fair agreement on the equilibria WO³ + H2 = WO2 + H2O and WO2 + 2H2 = W + 2H2O there arose considerable controversy as to the chemical composition of the oxide between WO³ and WO², because of very poor agreement on the equilibrium constant of this intermediate oxide reduction reaction. The controversy on this intermediate oxide has been in progress for more than 100 years. Thus Malguti,14 in 1835, decided it was WO2.56 In 1857, Riche15 proposed the formula W³O8. Usler,16 in 1865, confirmed the existence of WsO8 and further proposed W1O1l. In 1897, Desil7 agreed upon W²O8 but also found W5O14.

Jan 1, 1947

By K. Huessener

INTRODUCTION THIS paper attempts a survey of the principles involved in the com¬bustion of blast-furnace gas in boilers and stoves. I do not expect to be able to give much information which is actually new, since the laws and methods of calculation pertaining to this problem are contained in a number of handbooks; but, so far as I am aware, they have nowhere been treated exclusively with reference to this one problem of gas economy. The road which I follow is less that of strictly scientific research than that of the practical engineer, faced by operation problems which require a. quick solution by easily accessible means without a prolonged search in scientific handbooks. For this purpose I give a number of tables and diagrams, as well as the results of many years' practical working experience as combustion engineer, which I trust will prove useful to my readers. HISTORICAL Considering the tremendous efforts made during the last 15 years to increase the efficiency of all kinds of power plants, it is highly surprising that until a comparatively very recent period, progress in the economical combustion of blast-furnace gas has been so slow. The explanation generally put forth, that fuel has been so cheap that it was hardly worth while to bother about this economy, will not hold good, for the greatest trouble was taken to improve the combustion of coal and to reduce the consumption of steam. The most plausible explanation appears to be the general impression, held as an accepted fact, that gas, especially hot, uncleaned, blast-furnace gas, was not a good boiler fuel. Many managers of high repute had tried unsuccessfully to improve the conditions of its use, and the dictum went forth that efficiencies of over 60 per cent., particularly with high loads, were impossible. This was the situation not only in the United States, with very cheap fuel, but also in Europe, where fuel prices are twice and three times as high.

Jan 2, 1916

By Leonid V. Azároff, H. N. Murty

In the course of investigating the soft X-ray absorption spectra of Ni-A1 alloys, Das and ~zfiroff' observed an anomalous brittleness in solid solutions containing 3 to 5 at. pct Al. The brittle alloys also had anomalously larger lattice constants' and the fine structure at the nickel K edge of an alloy containing 3.96 at. pct A1 unexpectedly resembled that of NiA1 more closely than that of a solid solution containing 5.51 at. pct Al. This led them to suggest the possibility that short-range ordering was produced during the homogenization anneal to which each specimen was subjected, similar to that observed in dilute solutions of vanadium and of chromium in nickel: as the principal cause of embrittlement. It was decided to repeat the above measurements on several alloys in this composition range in order to test the above hypothesis and to establish the effect of heat treatment on the alloys' properties. Three solid solutions in this range were prepared containing, respectively, 3.16, 4.00, and 4.54 at. pct Al, in addition to alloys containing 6 and 9 pct. All samples were annealed in vacuum at 1800°F for 100 hr and then furnace-cooled to room temperature, as in the original study. The lattice constants of these alloys are shown by open circles in Fig. 1 and can be seen to fit quite closely the curve established by the previously measured values, indicated by solid circles. (The experimental errors are of the order of the size of the circles, +o.oo~R.) The three alloys containing 3 to 5 at. pct A1 were found to be quite brittle, as previously reported. In order to test the possibility of ordering, these three solid solutions were again heated to 1800°F and, following a few hours at that temperature, were quenched directly to room temperature. All three alloys lost their brittleness. In order to verify the proposed ordering hypothesis, it was decided to subject each sample to several cycles of quenching and annealing (disordering and ordering). After some experimentation, it was established that reproducible results could be obtained by anneal-

Jan 1, 1968

By O. D. Gonzalez

Existing measurements for the steady-state permeation of hydrogen in a iron above 100°C have been examined for contribution of determinate errors. The analysis leads to a recommended equation for the permeability of hydrogen in a iron: o= (2.9 ±0.5) x 10-3 exp - (8400 ± 400)/RT cu cm (ntp H2) cm-1 sec-1 atm-1/2 THE permeability of a iron to hydrogen has been the subject of numerous investigations over the past 40 years, and at present there are thirteen sets of published results for the rate of steady-state permeation of hydrogen in a iron above 100°C. The numerical values in each set of results are entirely self-consis-tent, but the spread among the sets is too large to be attributed solely to experimental error, i.e., to error other than in the specimen itself. Several reasons have been advanced to explain the disparities, but to date the relative importance of experimental inaccuracy to the spread remains uncertain. The purpose of this report is to examine in detail the sources of determinate errors inherent in the experiments and to assess as far as possible the contribution of the errors to the results. The ultimate goal is the selection of values for the permeability and heat of permeation most nearly representative of hydrogen in a iron. The analysis is limited to those experiments in which the permeation rate was observed at steady state—a condition in which traps for hydrogen within the metal are filled to a fixed level15 so that the trapping mechanism is not reflected in the rate of passage of the gas. Furthermore only data are examined in which surface processes are judged to have little or no influence on the flow. It is hoped with these restrictions to obtain values of the permeability and the heat of permeation which will be as closely related as possible to the mechanism of lattice diffusion. I) DEFINITION OF TERMS; UNITS In this report the data for permeation are given in terms of a coefficient oj permeability, ?, which is defined by the equation: jt=?A/?x{p1/2-po1/2) [1] where jt is the total flow of gas normal to the surface of a membrane of planar geometry, e.g., a disc, of area A and thickness ?x; pi and po are the pressures in the input and output sides, respectively. For flow radial to the walls of a membrane of cylindrical geometry, e.g., a tube, the corresponding equation is: where 1 is the length of the cylinder, and ri and ro are the inner and outer radii, respectively. The flux normal to the surface is given by Fick's law: j= -D(dc/dx) [3] At steady state the concentration gradient will be constant, and integration of Eq. [3] gives for the total flow through a disc of area A and thickness Ax: h =-DA(co - ci) [4] where c, and ci are the concentrations of solute at the output and input surfaces, respectively. When surface control is absent, co and ci are given by Sievert's law c = Kp1/2, and substitution therewith into Eq. [4] gives directly Eq. [I] where ? = DK. Integration of Fick's Eq. [3] in cylindrical coordinates will give Eq. [2] where again ? = DK and is thus shown to be independent of geometry (provided that surface control is negligible). The coefficient of permeability, or simply the permeability,* must be expressed in proper units. In *The term permeability will refer in this report always to the coefficient defined above; permeation will be used to specify the general phenomenon of gas passage through a membrane. this report ? will be expressed in the units of cu cm (ntp H2) cm-1 sec-1 atm-1/2. The variations of D and K with temperature are given by D = Do exp(-Ea/RT) and K = KO exp(-?Hs/RT) where E, is the activation energy for diffusion and AH, the heat of solution, each usually expressed in calories per mole of solute. The variation of permeability with temperature will thus be given (for conditions where surface control is negligible) by ? = ?o exp(-?Hp/RT) where ?0 = DoKo and ?Hp = Ea + ?Hs. The units of ?0 are the same as those of 6, and??Hp will be expressed in calories per mole H. 11) SUMMARY OF PERMEABILITY RESULTS Table I gives the values reported to date for the permeability of H2 in a iron in terms of ?o and ?Hp. Except where noted the parameters listed were taken directly from the numbers reported by the various investigators with only a change in units. The temperature limits within which the listed ?o and ?Hp hold are given in column 7; the limits marked in parentheses in this column indicate the entire temperature range covered in each investigation. The listed values of ?o and ?Hp are those giving a linear plot of ln? against T-1 at the higher temperatures in each set of measurements, and thus presumably represent the case for which surface control was negligible. Column 6 gives values of 9 at a representative tem-

Jan 1, 1970