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By A. R. Marder

Fe-C alloys have been investigated at temperatures below the eutectoid transformation to determine whether the superplasticity phenomenon exists for these materials. As a result of void formation at the carbide-ferrite interfaces that led to void coalescence and premature failure, the large elongations characteristic of superplustic matem'als did not develop. It had generally been assumed that the conditions for obtaining superplasticity include: a) T/Tm =0.5, b) a strain rate sensitivity of the flow stress that results in m s 0.3, and c) fine-grain structure. Although superphsticity was not observed in the alloys studied, stress-strain rate measurements indicated that an additional condition for the existence of superplasticitv in Fe-C allovs is that the fine-wain structure be equiaxed in agveement with previous results on Zn-Al al1oys. MANY investigations have been made on the hot deformation of iron-base alloys in order to obtain an understanding of the variables that may affect high-temperature strength and ductility.'-6 However, only a few papers have studied the effect of carbon content on the high-temperature properties of iron-base alloys. Peretayat'ko and Zaikov and Wallquist and carlen4 studied the effect of carbon on strength in the y temperature range, and recently Robbins, Shepard, and sherby6 studied the influence of carbon on the torsional behavior of iron in the a and cu-plus-cemen-tite temperature ranges. In addition to studying hot deformation behavior, current investigations have also been concerned with the phenomenon of superplasticity in nonferrous materials, and, in particular, ferrous alloys. Superplasticity is identified by the increased resistance to necking that results from a high strain rate sensitivity of the flow tress. ackofenr proposed the strain rate hardening coefficient m = d log a/d log e as a measure of strain rate sensitivity and showed that m must be "-0.3 for superplasticity to exist. In addition, a fine-grained microstrcture,- a eutectic or eutectoid ornosition,- and a test temperature in the creep range (TITm > 0.5)" have also been advanced as necessary conditions for superplasticity. Although many workers""23 investigated superplasticity associated with the allotropic transformation in iron and the eutectoid transformation in steel, and recently chadler " and orrison have studied superplasticity in the a - y phase region, no one has measured the superplastic effect in the a + cementite temperature region. Therefore, the purpose of this investiga- tion was to determine the effect of carbon content, microstructure, and test temperature on the strain rate sensitivity of Fe-C alloys at temperatures below the eulectoid transformation. EXPERIMENTAL PROCEDURE Fe-C alloys with carbon contents of 0.2, 0.4, 0.6, 0.8, and 1.0 wt pct were prepared as 100 lb 54 in. ingots by vacuum-induction melting. The compositions are listed in Table I. All the ingots were soaked at 2100°F for 2 hr prior to rolling to 4 in. plate. The minimum finishing temperature for rolling was 1800°F. The plate was sheared in 10-in. lengths and tensile bars were rough-machined prior to heat treatment. The samples were heat-treated according to the schedule in Table I1 in order to obtain a fine-grained microstructure consisting of finely dispersed carbides in a matrix of small ferrite grains. All the tensile specimens were heat-treated prior to finish machining, and tensile testing was accomplished on an Instron machine in conjunction with a clamshell furnace. The buttonhead specimen configuration is shown in Fig. 1. An inert atmosphere, helium, was maintained during testing with a steel tube shroud in a bell jar arrangement. Predried helium was continually purged over the buttonhead specimens and exhausted through ceramic fiber at the bottom of the tube to keep oxidation at a minimum and to prevent

Jan 1, 1970

By R. F. Yap, N. Arbiter, C. C. Harris

A fully-instrumented driving mechanism has been constructed to study the power, aerating and solid suspension characteristics of several laboratory flotation machines. Machines operating over normal flotation speed ranges give substantially constant power numbers in liquid systems indicating that they operate close to fully baffled turbulent flow conditions. Owing to lack of geometrical scaling, power numbers for different sizes of cells of the same make are different. Even larger differences occur between cells of different manufacturers. For a given impeller, varying the tank size did not significantly affect the power consumption. Tank geometry and baffling has a slight effect on power consumption with shrouded impellers. For a given impeller speed, suspension height decreases while power increases with increasing mass of solids in suspension. While power consumption is hardly affected by particle size and clearance between the impeller and tank bottom, the suspension height significantly increases as clearance increases and as particle size decreases. Aeration was found to affect not only power but also the suspension characteristics of the machine. For closely sized particles, drastic sedimentation occurred when a critical value of the air flow number was exceeded. The critical air flow number decreased with increasing particle size. This effect was less marked with a wide particle size range. The self-aerating capacity of flotation machines increases with cell size at a given speed for the same impeller-shroud system, but decreases in the presence of frother and solids. Equations have been established for correlating all of the studied variables. Data taken from the literature were also correlated with these equations. The modern mechanical flotation machine has, over a period of 60 years, displaced other designs to become the dominant type. Once the correct chemical environmental conditions have been established, the major requirement, selective mineral-air adherence, is essentially a function of two operating variables — impeller speed and air rate. The special nature of this dependence is such that in the same apparatus, but under different operating conditions, adherence can be either prevented or destroyed. Another aspect of this observation is that flotation machines though differing in detail to a considerable extent can produce similar metallurgical results. It must be concluded that having determined the correct reagent levels, establishing the correct operating conditions are at least as important as machine design. Since about 1960, after 30 years of fairly stable circumstances, manufacturers have been developing new types of machines whose major feature is emphasis on pulp flow characteristics by virtue of new im-peller-stator designs. As far as can be ascertained, these new designs necessitate radical departures from accepted impeller speed-air rate values. Performance data are still awaited. Flotation machine development has been empirical, as is selection of operating conditions. The purpose of the work at Columbia University has been to elucidate the fundamentals of the subject so that the purely mechanical aspects of the flotation process can be understood in hydrodynamic terms.1-6 Hopefully, this might lead to intelligent selection of a particular machine and its operating conditions for a specific ore based on a minimum of testing, or even to new types of flotation machines. The present investigation is the most comprehensive to date.7,8 Earlier work, which was largely exploratory, has been repeated and extended and employes the most accurate measuring techniques available. Additionally, sufficient hydrodynamic data were obtained under flotation conditions to delineate the relevant operating ranges. Attention has been given to scaling in that several geometrically scaled tanks and impellers have been constructed and tested; it has been pointed out elsewhere that undistorted geometrical scaling rarely pertains in flotation machine design.5 This paper is divided into sections dealing with relevant combinations of the three phases: liquid, air and solid. Summaries of the literature have been given in previous publications2-4 and will not be repeated here. Recently, several reviews of mixing have appeared, which can be usefully referenced.9-11 It is of interest to note that flotation is not included

Jan 1, 1970

By Joseph R. Stephens

Deformation substructures of' polycrystalline tungsten, W-2, 9, and 24 pct Re, and W-3 pct Ta were studied by tra?zsrnission electron microscopy. The stress-strain curve for unalloyed tungsten showed gradual yielding followed by work-hardening. Electron nzicrographs indicated a gradual increase in dislocation density with increase in strain up to 5.0 pct. Dislocations, although frequently jogged, were straight over moderate distances and were in a randorn array. Stress-strain curves for alloy specimens of W-2 and 9 pct Re and W-3 pct Ta exhibited a drop in stress at yielding followed by only slight work-hardening. Electron micrographs of these specimens after strains of 0.05, 0.1, and 0.5 pct revealed no change in dislocation substructure from the unstrained specimens. After 2.0 pct strain, the three alloys exhibited dense networks. W-3 pct Ta was characterized by straight, frequently jogged dislocations comparable with the dislocation structure in unalloyed tungsten after a similar amount of strain. In contrast, W-2 pct Re exhibited dislocations that contained widely spaced jogs, while W-9 pct Re had developed a cell structure after the relatively srnall strain of 2.0 pct. The W-24 pct Re alloy contained a few dislocations after 0.1 pct strain, while after 0.5 pct strain twins were evident. Dislocation slip bands apparently preceded the twins. The stress-strain curve for the alloy indicated that twinning commenced after approximately 0.25 pet strain. These results indicate that the primary effect of low rhenium concentrations (2 and 9 pct) in tungsten is to increase dislocation multiplication after macroyielding by reducing the Peierls-Nabarro force (lattice resistance to dislocation motion). The dislocation bands that precede twins in W-24 pct Re may be caused by localized internal stresses resulting fro a metastable structure, for example, clustering of rhenium atoms. The effect of high rhenium additions (22 a 65 pet* Rproperties of tungsten. Klopp, Witzke, and Raffo5 reported bend transition temperatures as low as -100°F (200°K) for dilute electron-beam-melted W-Re alloys tested in the worked condition. Recrystalliza-tion increased the bend transition temperature, but alloys with 2 to 4 pct Re were still markedly superior to unalloyed tungsten. Fractographic examinations of tungsten and W- 3 pct Re and W- 5 pct Re alloys by Gilbert 6 revealed that these low rhenium alloys showed a greater tendency toward cleavage failure than did tungsten. Garfinkle 7 showed that rhenium additions, up to 9 pct, to (100) oriented tungsten single crystals increased the proportional limit stress and decreased the flow stress and the rate of work-hardening. In addition, while deformation in unalloyed (100) oriented crystals apparently involved both (110)100) and (112) slip, crystals with rhenium contents of 5 andpct or more deformed primarily by (112) slip. The mechanism by which high and low rhenium additions affect the mechanical properties of tungsten is still not well-established. The present investigation was undertaken to determine by transmission electron microscopy the effects of low rhenium additions, 2 and 9 pct, and a high rhenium addition, 24 pct, on dislocation substructure in the early stages of deformation of polycrystalline electron-beam-melted tungsten. Unalloyed tungsten and a W- 3 pct Ta alloy were included for comparison. EXPERIMENTAL PROCEDURES Materials. Triple electron-beam-melted tungsten, W-2, 9, and 24 pct Re, and W- pet Ta were used for this investigation. Chemical analyses of the cast ingots are given in Table I. A description of the starting metal powders and melting and fabrication pro-cedures for unalloyed tungsten and the W-Re alloys is reported.5 The W-3 pet Ta alloy was processed in a similar manner. Compression specimens measuring 0.300 in. (7.6 mm) in length by 0.130 in. (3.3 mm) in diam were machined from swaged rods. All alloy specimens were annealed in a vacuum of 8 x 10- 6 Torr (10'2iVper sq m) for 1 hr at 3600°F (2255°K). The recrystallized grain size ranged from 0.06 to 0.08 mm diam for the alloy specimens. Unalloyed tungsten was annealed at 2400° F (1589°K) for 1 hr to produce a recrystallized grain diameter of approximately 0.12 mm. Specimens were electropolished in a 2 pet NaOH solution to a diameter of 0.125 in. (3.18 mm) to remove surface notches resulting from grinding and to improve reproducibility of t data. The ends o the compression specimens were ground flat, parallel to each other, and perpendicular to the longitudinal axes with 4/0 emery paper. Compression Tests. The compressive stress-strain apparatus used for compression tests is described in detail by Stearns and Gotsky.9 Room-temperature compression tests were conducted at a crosshead speed of 0.01 in. per min (0.25 mm per min).

Jan 1, 1969

By AIME AIME

THE outstanding event of the past month has been the conference in Washington of the representatives of about seventy-five of the leading national, regional and local engineering and allied technical organizations." Editorial-July Journal-The American Society of Mechanical Engineers. ?Prominent engineers from all parts of the country, representing more than 75 engineering societies, with an aggregate membership of more than 125,000, participated in the organization.? Editorial-July Journal-American Institute of Mining and Metallurgical Engineers. ?Civil, mechanical, electrical, mining, heating, sanitary, chemical, agricultural, marine, refrigerating, radio, fire protection, automotive, industrial and military engineers, as well as some architects and geologists assembled at Washington, D. C. to consider the advisability of organizing a national federation capable of bringing the united efforts of the entire membership to a focus on any problem of mutual interest." American Architect, June 23, 1920. ?"the object sought places it above any engineering gathering ever held." Editorial-American Machinist, June 17, 1920. "The youngest. profession has voluntarily accepted its responsibility in human affairs and is now reaching forward to a career of essential service." "This is .unmistakably the meaning of the Organizing Conference " Editorial-Industrial Management, July; 1920.

Jan 1, 1920

By A. English, M. J. Greaves

This paper is a discussion of pelletizing systems developed over the last 15 to 20 years for the agglomeration of fine iron ore concentrates. Five different pelletizing systems are now in commercial use in the U.S., and for the purposes of comparison, estimated operating costs and capital costs for the construction of four of them are included. Capacities of existing commercial plants are examined, together with the anticipated capabilities of plants presently being built or projected. In just 21 years, the art of pelletizing iron ore concentrates has reached maturity with mastery of the separation of finely divided concentrates from low-grade iron ores and their agglomeration into pellets containing about 65% iron. This accomplishment has firmly established an improved feed for the blast furnace, with resulting benefits in faster reduction of burdens and greater economy of operation. The present production capacity of pelletizing facilities in the U.S. is about 20 million tons annually, and it is expected to double by 1970. The rapid expansion of pelletizing in the U. S. has been closely observed in European countries, notably in Sweden and Germany, where pioneer agglomerating techniques were developed early in this century. PELLETIZING PLANTS AND PROCESSES Twelve commercial pelletizing plants are presently in operation in the U. S. and Canada. An equal number are in various stages of engineering and construction, and many others are being considered. The pelletizing systems currently employed include shaft furnaces, horizontal grates using three different airflow arrangements, and grate kilns. The capacity of pelletizing operations in North America is shown in Table I. All of these plants are integrated with the mine, the ore-crushing facilities and the concentrating plant and are equipped with storage and loading-out facilities for the product. Normally, a pelletizing plant is composed of a shuttle conveyor above the receiving bins; balling equipment; a heat-hardening system including the necessary auxiliaries such as furnaces, burners and controls, ducts, fans, meters and instrumentation; fines recovery equipment; and product handling equipment. Also included are the necessary auxiliaries and utilities such as boilers or power plant; oil unloading and storage facilities; coal unloading storage, grinding and distribution facilities; and bentonite or other additive storage and distribution equipment. Shaft Furnaces: Vertical shaft furnaces used in this country are limited to the treatment of natural magnetite concentrates containing minor quantities of residual fluxing elements. Fig. 1 shows a perspective view of the shaft furnace with balling, feeding and pellet-handling equipment. An elevation of a shaft furnace is shown in Fig. 2. A commercial shaft furnace unit is about 45 ft in height with a 6 x 14-ft rectangular cross section and has a nominal production rate of approximately 50 gross tons of pellets per hour. The upper portion of the furnace shaft is constructed of fire brick as protection from the high process temperatures involved. The lower portion of the shaft is lined with abrasion-resistant brick to withstand erosion from the fired pellets. Rectangular-shaped, refractory-lined and insulated combustion chambers are situated on each side of the shaft furnace. Oil or gas is burned to produce high temperature gases which are forced into the furnace through ports located near the top of the combustion chambers. The shaft furnaces operate at a pressure of approximately 6 psi. Pellets are fed into the top of the furnace by a pantographic feeder designed to produce the desired pellet stockline. The furnace is kept filled at all times and pellets are fed and discharged from the furnace at the same rate. In early development, 10 to 12 lb of solid fuel was incorporated in the pellets. However, the main source of process heat is now derived from the combustion chamber gases. In some instances no solid fuel is used. The pellet movement in the shaft furnace is counter-current to the flow of gases. Ambient air is introduced through louvers at the lower end of the furnace, at about 6 psi, cooling the fired pellets and supplementing the heat of the high temperature gases from the combustion chamber ports. The zone of maximum temperature, located above the ports of the combus-

Jan 1, 1963

By T. V. Falkie, Robert Stefanko

Coal Mine Health and Safety Act in December 1969, the most comprehensive law in the history of the coal mining industry went into effect. Hardly anyone will quarrel with the philosophy that all miners in the United States should be afforded equal protection under the law irrespective of the state in which they reside. Furthermore, it appears contrary to any sound judgment that a competitive advantage should appear to accrue to a mine operator because his operations are located in a state with less stringent (and costly) health and safety laws Yet it is recognized that the quality of state mine laws and inspection agencies has been extremely variable. Why, then, should there be such resentments over the passage of the Act? The Act was written In the highly emotionally charged atmosphere following the Consol No. 9 disaster: the technical expertise within the industry was discredited and largely ignored; and those directly involved in the writing of the Act were non- technical and largely nonmining people. The result was that the law incorporated certain requirements that were impossible to meet (automatic brakes are the classical example) and certain parts that were too ambiguous and consequently open to interpretation.

Jan 1, 1971

By F. Lynwood Garrison

As the specific gravity of low-carbon steel seems to be attracting considerable attention as a means of détermining the quality and value of the metal, I give here a few of the results of a long series of experiments recently made by me. The object of the experiments not requiring more than approximate correctness in the third place of decimals, the extreme precautions to obtain accuracy which are sometimes used in determining specific gravity were not observed. But each piece of metal was carefully cleaned to remove scale, rust, dirt, etc., thoroughly dried, and then weighed in air, all weighings being carefully repeated. The balance used was one of Troemner's finest, and sensitive to one-tenth of a milligram. Before weighing in distilled water, care was taken to remove, as far as possible, the air enclosed in the pores of the metal. The pieces of metal were suspended by means of fine silk thread, a correction of course being made for the weight of the silk. Care was taken that the water in which they were weighed should he always at the same level, and, as far as possible, at the same temperature.

Jan 1, 1887

It is too late to gloat over the fact that minerals have furnished 67 per cent of the primary wealth of Pennsylvania. It is high time that Pennsylvanians do some constructive thinking in terms of the present and future with regard to the remaining mineral heritage.1 The discovery of a deposit of pyrophyllite in the South Mountain region of the Cumberland Valley indeed emphasizes the increased importance of a more exact inventory of the metal and mineral resources of Pennsylvania. Active geological work and research have been in progress for many years, but the surface has been little more than scratched. The results obtained thus far in survey and research have been extremely beneficial to the Commonwealth and a fundamental hub of its industrial development. But these activities are wholly inadequate in the face of rising demands for the future. The need for an accelerated program to continue a study of Pennsylvania's mineral resources on a more intensive scale was never more acute.2 In part, this has been caused by the insatiable demands of industry. It has been caused also by wasteful use and ignorance.

Jan 1, 1950

By F. Keller, A. H. Geisler

Some of the important changes that take place in the structure of aluminum alloys are largely submicroscopic in character. This is especially true of the changes that accompany age-hardening and recrystallization. Although improved metallographic practices have been helpful in indicating some of these changes indirectly, results are limited by the resolving power of optical microscope lenses. The electron microscope, however, provides a new and important means for investigating the fine structure of metals in a range not possible heretofore. It is anticipated that this microscope will yield new and useful information to metallurgists when suitable techniques are developed to utilize its very high resolving power. Several unique methods have been devised to permit the examination of the structure of opaque samples by the electron microscope. These are all based on the principle of producing a very thin film that represents the prepared surface of the opaque metal sample, which can be examined in the transmission microscope; however, the available methods fall into three categories, depending on the manner in which the film is obtained. With the oxide-film method developed by Mahl3-17 a very thin surface layer of the metallographic sample is converted into oxide by a thermal, chemical or anodic treatment. After removal from the metal sample by one of the special techniques described by Evans,19 this oxide film is examined in the electron microscope, and is found to portray the structure of the original metal surface. Films formed on aluminum,8,17 nickel,7 iron 3,8,9 and a nickel-beryllium alloy" by heating; films formed on iron 8.9 by chemical action; and films formed on aluminum3-16 by anodic oxidation, have been examined. The other two are more recent methods developed principally in this country; they consist in making the thin film in the form of a mold or replica of the surface contour of an etched sample. With the negative replica method the thin film is made directly on the sample; with the positive replica method it is made on a negative reproduction of the sample. With both methods, variations in the contour of the metal sample, caused mainly by etching, produce variations in the thickness of the surface replica. Because of the direct dependence of the amount of scattering of the illuminating electrons upon the thickness of the film penetrated, only the variations in thickness of the replica are responsible for tones in the electron image. Negative replicas are made by forming a thin film of some suitable material directly on the polished and etched metal sample, and derive their name from the fact that after removal from the sample they have the high and low points in the sample reversed, as shown in Fig. I. Nega-

Jan 1, 1944

By A. Phillips, D. H. Killpatrick, V. Kerlins

The tensile failure of two dispersion-strengthened Ni-20 Cr alloys was studied and compared to the fracture of a similar alloy with no dispersoid. The fracture characteristics were studied using electron fruc-tography and transmission electron microscopy. In all three cases, the mode of failure was found to be microvoid coalescence. The failure in the dispersion-strengthened alloys was found to have initiated at the particles. The size of the dimples in the fracto-graphs was found to be related to the spacing of the particles, but not to the total elongation before failure. The elongation before failure was found to be related only to the amount of dispersed phase. These results are compared to those predicted by a theoretical model of ductile failure. THE continually increasing strength requirements for creep-resistant materials capable of long life at service temperatures above 1600"F (875"C) have developed considerable interest in dispersion-strengthened nickel-base alloys. The high yield strength of these alloys results from the presence of a dispersion of fine particles of thoria which act to impede the normal motion of dislocations. Many theories have been presented to explain the role of the dispersion in the strengthening of these alloys. A good review of these theories is presented by Ansell.' The ultimate strength of such alloys is not only related to mechanisms raising the yield strength, but also to the amount of work-hardening which occurs. This work-hardening is determined by the work-hardening rate and the amount of plastic strain before failure. Since fracture limits the amount of plastic strain, a study was undertaken to gain a better understanding of the fracture mechanisms in these alloys. Electron fractography and transmission electron microscopy were used to study the fracture characteristics in two dispersion-strengthened alloys and one similar alloy containing no dispersoid. These results are related to the tensile properties of the dispersion-strengthened alloys at room temperature and at 2000°F (1095°C). The results are also related to a theory for ductile fracture. 1) PROCEDURE Standard tensile specimens were made from 0.020-in. (0.051-cm) sheets of three different Ni-Cr alloys. The alloys had nominal compositions of Ni-20Cr, Ni-20Cr-2Th0,, and Ni-2OCr-4Th0,. All alloys were supplied in an annealed condition. The specimens were fractured in tension at room temperature to determine the effect of the dispersed thoria on the fracture appearance. The tensile properties were determined from the average of a minimum of three ten- sile specimens for each condition tested. The fracture surface of all of the alloys was examined by electron microscopy using standard two-stage plastic-carbon replica techniques. Thin foils of the dispersion-strengthened alloys were used to investigate the size and spacings of the dispersion particles, and to investigate sections of the 2 pct thoria alloy taken from areas of the specimen adjacent to the fracture surface. The thin foils were mechanically ground to 0.010 in. (0.025 cm) and then chemically polished to approximately 0.001 in. (0.003 cm) in a solution of 29 g of ferric chloride and 10 ml hydrochloric acid in a liter of water. The polishing solution was maintained at 150"F (65°C) during the thinning operation. The final thinning of the foil was done electrolytically using a solution of 700 ml ethanol, 100 ml 2-butoxy eth-anol, 120 ml distilled water, and 78 ml perchloric acid (70 pct). The potential was maintained at 15 v and the bath temperature at -20°~ (-29°C) during the thinning operation. The final thickness of the foil was approximately l000A as determined from the width of twins boundaries observed in many of the foils. 2) RESULTS AND DISCUSSION The electron fractographs of the fracture surfaces of the alloys are shown in Figs. 1, 2, and 3. The normal interpretation of this type of fractographz is that, as a result of the difference between elastic and plastic properties of the matrix and the particles or other inhomogeneities in the alloy, microvoids are formed

Jan 1, 1969

By F. Bonaccorso, G. Venturello, C. Antonione

A study of the effect of small amounts of interstitial impurities on recovery and re crystallization in high-purity iron (99.995 pct) has been undertaken. This paper gives results on the effect of carbon, introduced in small-dosed amounts @om 0.0005 to 0.0086 wt pct) by heating iron specimens of high purity m a static atmosphere of CO + ,. The materials prepared in this way, cold-worked 80 pct and subjected to a series of isochronal and isothermal annealings, were submitted to examinations by X-rays, micrographs, and hardness tests. It was observed that effect of carbon is remarkable in the sense of blocking recovery of mechanical properties up to the temperature at which re crystallization begins. On the contrary, carbon has a negligible effect on primary re crystallization temperature, when compared with the known effect of substitutional impurities. This is in agreement with the high mobility of interstitials. In effect, only a slight decrease, -20 pct, of the grain-boundary motion rate was noted, due to the interaction between the grain boundary and the carbon atoms. On the other hand, in the samples in which the carbon content is above the solubility limit at the temperature at which re crystallization occurs, a slight increase of nucleation frequency is noted due to the presence of precipitated carbides. ThE effect of purity on recovery and recrystalli-zation phenomena has been known for a long time; however, recently, new attention has been given to the problem since new methods for obtaining metals with extremely low impurity contents have become available. Most of this research work essentially concerns the effects of impurities which cause precipitates or give rise to substitutional solid solutions. The works of Bolling and winegard1 and Aust and utter' on lead, and Vandermeer and orddon' on aluminum, 01sen4 on nickel, and Abrahamson and Blakeney on iron should in particular be referred to. On the effect of this type of impurities a quantitative theory has been formulated by Detert and Lucke6 and has lately been discussed critically by Cahn7 and Gordon and vandermeer.' Interstitial solid solutions in iron have not as yet been studied. As the study of the effect of interstitials is of great interest both from a theoretical and practical standpoint, it was deemed useful to examine the effects of carbon, nitrogen, and boron on recovery and primary recrystallization of iron. There already is some work by Chaudron et al.,1-" on the effect of interstitials on iron; their work, however, mainly concerns secondary recrystallization. The present paper refers in particular to the effect of carbon. EXPERIMENTAL PART Preparation of Materials. The pure iron used for this research was obtained from FeCIS recrystal-lized and purified by extraction with isopropylic ether. From the ferric chloride purified in this manner, hydroxide was precipitated with a solution of very pure ammonia, and, by calcination in pure sintered alumina crucibles, oxide was obtained. Reduction of the oxide to iron sponge was performed in sintered alumina tubes with very pure hydrogen at 650°C; at the end of the operation temperature was increased to 900°C. Specimens for the experiments were obtained by sintering the sponge at 1480°C in pure Hz after a pre sinter ing treatment at 900° C. It is important to note that the treatment at 1480°C in Ha produces a further purification from more volatile elements such as zinc, cadmium, arsenic, lead, and tin. Details on the preparation and characteristics of this type of very pure iron are given in a previous work.'' Only the complete analysis performed by neutron-activation methodsz3 is given here, Table I. Some of the specimens prepared in this way were carburized in the 0 region with very low amounts of carbon by treating them at 700°C in a static atmosphere of Ha containing a definite amount of CO. The set-up used is described in Fig. 1. A gas-tight quartz tube containing the specimen to be carburized and an internal friction control specimen, after being evacuated, was filled with Hz

Jan 1, 1963

Discussion of the paper of C. A. HEBERLEIN, presented at the New York meeting. February, 1915, and printed in Bulletin No. 98, February, 1915, pp. 497 to 504. H. D. PALLISTER, State College, Pa. (communication to the Secretary).-Mr. Heberlein gives some valuable information concerning a branch of mining and metallurgy which has been neglected during the last few years. His results with fine grinding and Deister tables suggest the possibility of the mining and concentration of a number of low-grade cinnabar deposits. Can Mr. Heberlein give us some idea as to the loss of quicksilver during this concentration? The advances in the metallurgy of quicksilver have been few and far between in recent years. Most miners prefer to go on in the old way with the Scott furnace and the large brick condensers. The large brick condensers are expensive to build and to keep in repair, and on account of the thick walls are extremely slow in cooling the gases; and further, the joints between the bricks offer passages for the quicksilver vapors unless great care is used in building the condensers.. The escape of vapor at the top of the furnace is another source of loss and I see the people at the Oceanic mine have overcome this in a very ingenious way. I recently returned from the Terlingua quicksilver district of Texas, where the only operating company, the Chisos Mining Co., is operating a 20-ton furnace of the Scott type. They had considerable difficulty with getting a good extraction, clue to too high heating in the furnace with its .resulting diffIculties. Recently they have been running with ashpit and fire doors closed and have been able to keep the heat lower, thus getting much better results and practically verifying the results obtained at the Oceanic mine.

Jan 5, 1915

By N. A. Gokcen, J. Chipman

A revised treatment of the authors' published data eliminates the complex relation previously proposed between concentration of silicon and activity coefficient of oxygen in liquid iron. Revised values of the thermodynamic properties of the liquid solution are presented. IN a recent experimental study of the reaction SiO2 (s) = Si+ 2O; Kf, = [% Si] [% O]² [1] the authors' found a substantially constant equilibrium product in liquid iron at 1600°C of 2.8x10-5 They also reported extensive data on the reactions: SiO2 (s) + 2H2 (8) = Si- + 2H2O (g); K'2= [% Si] (H2O/H2 [2] and H, (g) + 0 = H2O (g); K'3 =( H2 O [31 (H2) [%O] From the results on reaction 3 and earlier data of Dastur² on this same reaction in the absence of silicon, they determined the activity coefficient of oxygen, f0, on the basis of the definition K3 = (H2O)/ (H2)f0 [% 0] where K, is the equilibrium constant and f0, is taken as unity in the pure Fe-0 system. Similarly values of fsi were deduced from results on reaction 2. In a more recent study" of analogous reactions in the system Fe-A1-0, it was found impossible to reconcile the results on reaction 3 with Dastur's data; accordingly the latter were ignored and the equilibrium results were extrapolated to find a value of K, at zero concentration of aluminum. This procedure failed to locate the cause of the discrepancy but it did yield reasonable values of activity coefficients. It also avoided introduction of the complex empirical relation between the oxygen activity coefficient and the concentration of the added element. The same type of discrepancy exists for system Fe-Si-0.' In the earlier paper an attempt was made to fit both sets of data by a single curved line (Fig. 6 of ref. l), the form of which is contrary to the theoretical requirement of a finite slope at infinite dilution. In the light of experience on the Fe-A1-0 system the discrepancy must be recognized as one which can be resolved only by more refined measurements. Accordingly Figs. 6 and 10 are retracted. It is pointed out also that until the discrepancy is resolved Figs. 7, 8, and 11 are subject to some uncertainty. Qualitatively the following conclusions still appear valid: 1—The activity coefficient of oxygen is reduced by addition of silicon. 2—In dilute solutions the activity coefficient of silicon increases with its concentration. 3—With respect to equilibrium in reaction 1, the above effects are approximately compensating. The discussion of K'1 in the previous paper requires no revision. It was pointed out that the constancy of the product [% Si] [% 0]² ndicated a compensating effect of the activity coefficients of silicon and oxygen. Therefore, as a very good approximation, K1 = K'1 and the following average values are suggested both for K, and K', at the temperatures 1550°, 1600°, and 1650", respectively, 1.0x10-", 2.8~10-" and 5.5 ~lo-'. Revision of the thermodynamic treatment is necessitated by the recent appearance of new data, based on a combination of combustion and solution calorimetry,' which yields for the heat of formation of low-cristobalite from the elements, the value —209,330 ±250 cal per mol at 25°C. This is about 4000 cal larger than the value previously accepted. The new value for cristobalite is used, together with Kelley's tables of high-temperature heat contents" and entropies and with Korber and Oelsen's' heat of fusion of silicon to obtain the following equation for the standard free energy of cristobalite in the temperature range 1700" to 2000°K: Si (1) + 02 (g) = Si02 (crist.); ?F° = -217,700 + 47.OT [4] The free energy of solution of 0, in liquid iron is:8 O2 (g) = 20 (in Fe); AF° = -55,860 - 1.14T [5] and these two equations are combined to give: Si (1) + 20 = SiO, (crist.); AF° = -161,840 + 48.14T [6] ?F°1873 = -71,700 cal. From the experimental value of K, = 2.8x10-5, Si + 20 = SiO2 (crist.); ?F°1873 = -39,000 cal. [7] The combination of Eqs. 6 and 7 yields the free energy change when liquid silicon dissolves in iron to form the dilute solution of unit activity (1 pct). Si (1) = Si; ?F°1873 = -32,700 cal. [8] The heat effect in this process according to Korber and Oelsen' is an evolution of 28,500 cal per gram

Jan 1, 1954

By R. M. Treco

THE thermal expansion of pure beryllium was first investigated by Hidnert and Sweeneyl in 1925 on a single cast specimen stated to be of 98.9 pct purity. A study of the coefficients of expansion by X-ray methods has recently been made by P. Gordon' on powder obtained from the Brush Beryllium Co. assaying 97 wt pct beryllium with about 2 wt pct oxygen as an impurity. The present work pertains to about the highest purity metal obtainable from the Brush Beryllium Co. This metal was in the form of lumps which were remelted and cast in a high-vacuum induction furnace. Expansion data in this paper deal with beryllium extruded into bars of rectangular cross-section 4x1/2 in. and the anisotropic expansivity of a large cast single crystal. Analysis of the cast metal after extrusion indicated a beryllium assay of 99.28 pct. Impurities were 0.170 pct Fe, 0.140 pct Al, 0.014 pct Mg, 0.086 pct Si, 0.020 pct Cu, 0.020 pct Mn, 0.007 pct Ni, 0.007 pct Ca, 0.080 pct C, 0.179 pct 0, (as BeO). Part I—Extruded Metal Thermal expansion measurements were obtained with a fused quartz-tube differential type dilatom-eter3 on specimens cut from the extruded sections in longitudinal and transverse directions. Specimens were machined to a diameter of 0.250 in. and a length of 3.000 in. One sample from each direction was vacuum annealed for 1 hr at 800°C. The dilatometer tube was placed in a vertical tube furnace having a uniform temperature zone which was twice the length of the sample. Average time for a complete thermal cycle to 500°C was 18 hr. Temperatures were observed with two chromel-alumel thermocouples attached to the sample about 1 M in. apart. The accuracy of temperature measure-ment was ± 0.25°C. An argon gas atmosphere sur-rounded the specimen to prevent oxidation of the beryllium. Samples were identified as shown in table I. Results - The original plan included measuring expansion up to 1000°C, but the first two runs on sample No. 5 showed that a permanent contraction occurred from

Jan 1, 1951

By H. J. Morel-Seytoux

Methods of predicting the influence of pattern geometry and mobility ratio on water flooding recovery predictions are discussed. Two methods of calculation are used separately or concurrently. The analytical method yields exact solutions in a convenient form for a unit mobility ratio piston-like displacement. A few typical pressure distributions, sweep efficiencies and oil recoveries are presented for various patterns. For non-unit mobility ratio, one may resort to a numerical method, such as that of Sheldon and Douherty. 1,2 Because the domains of applicability of the analytical and numerical techniques overlap, the exact solutions provide estimates of the errors in the numerical procedures. The advantages of the analytical and numerical methods can be combined. To develop a numerical technique as independent of geometry as possible, the physical space is transformed into a standard rectangle. The entire effect of geometry is rendered through one term, the "scale-factor", derived from mapping relations. The scale factor can be calcu-lated from the exact unit-mobility ratio solution for the particular pattern of interest. By this means recovery performances for arbitrary mobility ratio con be obtained for many patterns. A sample of tesults obtained in this manner is presented. INTRODUCTION Pattern geometry and mobility ratio are two major factors in making a waterflood recovery prediction. Because assisted recovery has become increasingly important to the oil industry, pattern configuration and mobility ratio also assume a greater significance in the assessment of the economic value of recovery projects. The influence of pattern geometry and mobility ratio in shaping a recovery curve and on the other quantities of interest to the reservoir engineer is the main subject of this paper. Much effort has already been spent on estimating quantitatively the influence of either pattern or mobility ratio or both on oil recovery. The literature reports many investigations of this nature. 3-9 However, many results or methods of recovery prediction presented in the literature cannot be considered fully satisfactory. Even for unit mobility ratio and piston-like displacement, where analytical solutions are available, the literature shows discrepancies. For non-unit mobility ratio, the divergence in the results is extreme. For infinite mobility ratio in a repeated five-spot, depending on the investigator, the sweep efficiency ranges from 0 per cent to 60 per cent. With respect to the influence of pattern on recovery, only the repeated five-spot has received much attention. Other confined patterns and pilot configurations have received very little attention. Two calculation methods are presented in this paper, either separately or concurrently: the analytical method of potential theory and the numerical method of finite-difference approximation. The analytical method is more restricted in scope than the finite-difference method, but it has the definite advantage of providing exact solutions within its range of applicability. If a unit-mobility ratio piston-like displacement is assumed, the analytical approach is possible. A few typical results are reported in this paper; the detailed description of the general method and of a great variety of results will be the subject of other articles. Fornon-unity mobility ratio, we must resort to a numerical scheme. The numerical technique is that which was described by Sheldon and Dougherty.l,2 It is not limited to piston-like displacement. However, mainly single interface results will be presented here. Because the respective domains of applicability of the analytical and the numerical method overlap, useful comparisons of exact and numerical solutions can be made for a variety of patterns. The advantages of the analytical and numerical approaches can be combined. The reason for the success of this analytical-numerical approach can be summarized in the following two points: 1. The numerical solution for arbitrary mobility ratio can be programmed most efficiently when the physical space in which the displacement actually takes place is transformed into a standard shape; and 2. This can be done with remarkable simplicity whenever an analytical expression for the pressure

Jan 1, 1966

By B. Foley

IN THE first report1 disks of steel of known composition and history were exposed, under carefully prescribed conditions, to impacts of explosion products resulting from the explosion of 50-gm. charges of explosives having well-determined rates of detonation, or to impacts of a weight of known mass, falling through measured distances. There were available at the Bureau of Mines Experiment Station five explosives, having rates of detonation of 5716, 4470, 3190, 2296, and 1523 meters per second. It being desired to continue this exposure downward, and no explosives with a lesser rate of detonation being at hand, resort was had to the impact of a falling weight, the differences in the velocity of impact being attained by varying the height from which the weight was allowed to fall. It was expected by this means to demonstrate experimentally that the development of Neumann bands was determined by the suddenness of the deformation of the metal in which they were produced. But while progressive deformations were obtained with the explosive arges, the falling-weight method failed because the tup was ogressively deformed, also the disks were so flattened as to be unfit for metallographic examination. A return, therefore, has been made to the use of explosives as the source of energy, an air-gap being interposed between the charge of explosive and the disk, as the effect diminishes as the function of the distance between the charge and the object on which the products of its explosion, or the shock wave from its detonation, impinge. The details of the arrangement of the tests by the air-gap method were the same as those described by Foley and Howell,2 except that after the system had been set up and the desired apparent specific gravity for the explosive

Jan 2, 1926

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

THERE has been considerable discussion among A metallurgists and others interested in the development of titanium alloys as to the nature of the fine line markings which appear in the microstruc-tures of unalloyed titanium and a titanium alloys. These markings have been variously called x phase, etching pits, and twins. Figs. 1 to 3 show typical examples of the line markings as they appear, respectively, in high purity iodide titanium, commercial purity titanium rod, and a high purity a alloy containing 0.2 pct N. The commercial purity titanium sample, shown in Fig. 2, contains both line markings and some spheroidal retained-ß phase. The ß phase seen in this micrograph should not be confused with the line marking, since it originates from the iron present as an impurity in the sponge metal. It will be noted that the line markings seem to have a definite orientation pattern with respect to the grains and in titanium noticeably resemble the precipitation occurring in artificially aged high purity Al-Cu alloys. This similarity led to the hypothesis that the line markings observed in titanium and a titanium alloys resulted from precipitation of some unknown phase. With the work of Hägg1 and McQuillan2 as background, there was little reason to suspect hydrogen of being responsible for the line markings. Hagg reported that up to 33 atomic pct H was soluble in a titanium, and McQuillan reported an a solubility of up to 8 atomic pct. Furthermore, McQuillan examined metallographically a 5.1 atomic pct H alloy quenched from 300°C and stated that the structure appeared to be a single-phase a solid solution. Line markings have been observed frequently and were first noted by Jaffee and Campbell3 in 1949. They did not observe them in vacuum-annealed titanium and a Ti-O and Ti-N alloys, but did observe them in Ti-H alloys containing 0.25, 0.5, and 1 atomic pct H. They stated that it was doubtful that the markings were a second phase based on hydrogen, because 1 atomic pct was far from the accepted solid solubility of hydrogen in titanium. Finlay, Resketo, and Vordahl,4 in 1950, concluded that the markings were a true second phase as a result of observations on their etching characteristics. Anderson5 Commenting on these markings in titanium was inclined to associate them with hydrogen, but was loath to call them titanium hydride, because of the reported high solubility of hydrogen. Work at Battelle Memorial Institute for the Navy Bureau of Aeronautics," done concurrently with the work described in this paper, provided information on the experiment converse to that described by Jaffee and Campbell in 1949. This work showed that the amount of line markings in unalloyed titanium could be decreased by vacuum annealing at elevated temperature. It was suggested that this was the result of removal of hydrogen; previously, the data of Jaffee and Campbell9 howed that the markings appeared when hydrogen was added. Recent work at Battelle for the Ordnance Corps, Department of the Army, which is described here, now establishes the identity of the markings as a titanium hydride phase. A sample of high-purity titanium rod 1/4 in. diam, 0.009±0.001 wt pct H (0.43 atomic pct), Fig. 1, was resistance heated in a 0.05 micron vacuum for 41/2 hr at 1500°C. During the early states of heating, the pressure increased to 150 microns, indicating that gases were being evolved. After cooling in vacuum, the iodide titanium rod was analyzed for hydrogen. Vacuum-fusion analysis showed 0.0006 ±0.0003 wt pct H (0.03 atomic pct) to be present.* Metallographic examination of the vacuum-annealed rod showed only a transformed ß structure. There was no evidence of fine line markings, Fig. 4. A section of this vacuum-annealed rod was cold rolled to 0.040 in. thick sheet and annealed for 1 hr

Jan 1, 1953

By J. A. Roberson, A. M. Adair

A plate of an 18 pct Ni (250 grade) maraging steel was homogenized, cold rolled, and re crystallized to produce a fine grained, homogeneous structure. The strength properties were then measured as a function of cold rollinq prior to aging. Material in the solution heat treated condition exhibited substantial strengthening from cold rolling. The strength increase and the rate of strengthening were proportional to the amount of cold work. This benefit, of up to 50 ksi, was retained during aging for 3 hr at 900°F. The structures produced by cold rolling and aging have been systematically investigated by transmission electron microscopy. The dislocation debris produced by cold work influences the morphology and distribution of precipitate particles as well as the rate of precipitation. An essential feature of the high nickel maraging steels is a martensitic transformation from y to a with the resultant production of a supersaturated bcc solid solution. This solution decomposes to form precipitates during a subsequent aging treatment.' The martensitic transformation in low carbon Fe-Ni alloys and the mechanical properties of these alloys have been studied extensively.2"4 The carbon content of maraging steel is kept very low to avoid the embrittling Fe3C phase and to minimize the acicular, or twinned martensite common to high carbon steels. Maraging steels are usually strengthened only by heat treatment, but cold working is also quite effective in this respect as was shown by Decker, Eash, and Goldman,' Kula and Hickey,5 and the present work. There is, however, some disagreement between the present results and those of Kula and Hickey. Because of their unusual combination of strength and toughness, the maraging steels are suitable for applications where reliability at high stress levels is required. The strength level is achieved through a combination of solid solution hardening,' substructure hardening,3 and precipitation hardening,6 of which the last is the most important. The toughness of these steels has been reported to be a result of the absence of brittle grain boundary precipitates,7 and the fact that the matrix is a massive martensite rather than a twinned martensite.' One of the problems in the manufacture and use of maraging steel has been its tendency to develop laminations in sheets parallel to the rolling plane.9,10 The present investigation was intended to determine whether or not an improvement in the microstructure of this steel by thermomechani-cal processing would improve mechanical properties. I) EXPERIMENTAL PROCEDURES A plate ($ in. thick) of an 18 pct Ni-250 grade maraging steel* was obtained and examined metallo- "Analysis, wt pct: 18.1 Ni, 8.6 Co, 5.3 Mo, 0.4 Ti, 0.07 Mn, 0.05 Cr, 0.05 Si, 0.002B. 0.1 Cu. 0.05 Al. 0.002 Zr, 0.016 C, balance Fe. graphically. It was found to contain laminations as described by previous investigators. This plate was then annealed at 2100°F for 24 hr to promote homogeneity. Pilot studies on small coupons had shown that this was the minimum temperature required to accomplish homogenization in 24 hr. Since this treatment caused excessive grain growth, the steel was subsequently cold rolled and recrystallized. Cold rolling was accomplished at about 5 pct per pass to a total reduction of 66 pct. The standard solution annealing treatment (1 hr at 1500°F) was adequate to cause recrystallization of the cold worked steel. The recrystallized steel was then cold rolled and ground by various amounts to produce specimens with uniform thickness and selected amounts of cold rolling. All tensile specimens with less than 90 pct cold reduction were ground to a thickness of 0.050 in. Specimens with 90 pct cold reduction were 0.025 in. thick. All

Jan 1, 1970

By H. B. Pulsifer

THIS paper outlines a method for rapid production of flat, granular surfaces on many of the medium hard- alloys. Grinding wheels and fabrics on wheels are not used; the purpose is accomplished by rubbing on emery papers and on fine abrasives heaped on hand boards. The advantages of chromic acid in the solutions for etching nickel and copper alloys are explained. Steels may require an additional step in the mechanical smoothing owing to the lack of an ideal etchant. Chloric acid is a. rapid etchant for steel. Thirty-six photomicrographs illustrate the steps in smoothing and the clearness with which structures can be rapidly exposed. New structures in low-carbon steels are announced. After a little practice, many materials can be finished within two or three minutes. CUPRONICKELS In the previous papers on the metallography of copper and magnesium which the author has presented to this Institute there has been an attempt to demonstrate the advantages of minimizing the mechanical work of surfacing and depending more on chemical means; the present paper continues the research into a still wider field. It is obvious that the success of a chemical method for smoothing and etching depends almost entirely on the use of a suitable reagent. The etchant must remove the surface granulations and outline the structures. A very special type of reaction between reagent and metal must be displayed. Promiscuous corrosion or the intensification of the attack about local couples -formed by impurities or components must not be in evidence. There are several reagents that etch highly polished nickel and copper alloys very well. Acidified ferric chloride solutions, concentrated and dilute nitric acid solutions, and nitric and acetic acids in acetone have -been used to good advantage for the purpose. However, if they are expected to etch rather deeply-that is, smooth as well as characterize-the results may be far from satisfactory.

Jan 1, 1928

C. H. GIBBS,* Salt Lake City, Utah (written discussion?).-The development of the coke industry in Utah had a somewhat checkered career for the first 50 years of its existence. About 1851 the iron-ore deposits of the Iron Springs and Iron Mountain mining districts, Iron County, Utah, were discovered. Obtaining of iron for any purpose at the time involved its shipment from the Mississippi River to Utah, by ox-team. The result was that the pioneers endeavored to establish an iron-ore smelting plant at what was known as Iron City, near Iron Mountain, in Iron County. Bishop Taylor, who headed this enterprise, endeavored to obtain coal suitable for coking from the Cretaceous coal measures immediately east of Cedar City. A few small bee-hive coke ovens were built and lump coal was coked in these ovens; or, more properly speaking, the volatile matter of the lump coal was driven off, leaving a carbonized material that retained the lump shape, although it was extremely, friable. This coke was promptly found to be unsuitable for iron smelting, not only clue to its extreme weakness under smelting load, but also because of the fact that it carried about 5 per cent. sulfur. Bishop Taylor made a very thorough and exhaustive effort, covering a considerable period of years, to locate suitable coking coal in Southwestern Utah. His efforts failed, and in 1902 the iron-ore property acquired by him and associates was sold to the Colorado Fuel and Iron Co. It is of interest to note in this connection that the iron-ore deposits of Iron County at present are the largest known undeveloped iron-ore fields on the North American continent. It is safe to say that there is at least one hundred million tons of 55-per cent. ore actually in sight. In 1878, coal-mining operations were begun at Connellsville, in Huntington Canyon, Emery County. Ten 8-ft. bee-hive ovens-were constructed, and lump coal was carbonized for use in the Salt Lake valley smelters. The product was somewhat better than that obtained in Iron County, but was still of very inferior quality. It had to be hauled a distance of 140 miles by wagon to the two small smelters then operating in Salt Lake valley. These ovens operated only a very short time. The next move toward obtaining a coking coal was in 1889 when the Pleasant Valley Coal Co. began operations in the Castle Gate district. In 1890, eighty 8-ft. bee-hive ovens were constructed, and the product. was used until 1900 in all the smelters in Utah. During this period, 124 additional. ovens were constructed, making a total of 204 bee-hive ovens operating at Castle Gate.

Jan 11, 1918