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By Murl R., Schrock

THE use of diamond-drill holes at the Moctezuma Copper Co., Pilares mine, Pilares de Nacozari, Sonora, Mexico, for the ventilation of raises has resulted from experiments that were made in an effort to secure continuous and efficient ventilation in hot and gaseous raises. The method is particularly applicable in long raises and in raises run from development levels where the source of fresh air is distant. In order that the CO, concentration shall not become excessive, 30 cu. ft. per min. per man must be supplied to a heading for breathing purposes:' The quantity of air that will pass through a diamond-drill hole by

Jan 1, 1929

By William R. Barton

It is generally conceded as axiomatic that the aggregate producer and the average urban resident have mutually incompatible goals. The producer wants to be near his mass market and the average resident wants him as far away as possible. The traditional economic decision to mine local material to avoid transportation costs is increasingly challenged by zoning and planning bodies and by citizens' groups. This conflict generated much written and spoken word and offered opportunity for great flights of rhetoric. The pressure on the producer also has forced re¬appraisal of the need for the classic confrontation with indications that much of the conflict is more apparent than substantive. In New England several factors combine to increase complexity of the siting problem. It is a region of high land values and population density. Due to the proximity of metropolitan areas, even rural communities have experienced a prolif¬eration of environmental groups with concomitant enthusiastic regional planning and critical watch over land use. In New England where there is little public understanding of the necessity for, and the realities of, mineral production a sophisticated approach to aggregate production site location is required.

Jan 1, 1975

By C. E. Meissner

THE need for a practical method of deciding upon new alloy steel analyses to widen its markets was the problem facing the Chrome Steel Works at the beginning of 1927. In addition to determining the physical properties of the alloys under consideration, it was necessary to test their value in service. Small melts of steel in a crucible, from which test specimens could be made, were insufficient to determine more than physical properties, therefore actual castings and ingots had to be made before the products could be placed in service. Furthermore, to save further experimental work, it was desirable that these heats should be made on a "pilot" basis; that is, actually duplicating regular open-hearth furnace practice. Therefore a one-ton (2000-lb.) acid open-hearth furnace was built. FURNACE DESIGN There are no special features in the design of this furnace. Standard acid open-hearth practice is followed, with the reproportioning of the dimensions necessitated by its small size; notably, a considerably increased length in proportion to width in order to get proper flame propagation. After the first heat was made it was found that the checkers were insufficient. The checker chambers were increased in size and the checker openings were made much smaller. The furnace was started late in March, 1927, and made 65 heats. Experimental work was then stopped until the data accumulated could be worked up and summarized. The run was resumed in November and the furnace is still operating, having made at this writing 60 additional heats. For the last series of runs, there was some rebuilding of both the furnace itself and the checkers. Slag pockets were added. Fig. 1 shows details of construction; Fig. 2, the, general appearance. All designing was done by the operating and engineering staff. Construction materials were those available around the plant. The reversing valve is a standard butterfly type taken from an old furnace.

Jan 1, 1928

By C. E. Meissner

The need for a practical method of deciding upon new alloy steel analyses to widen its markets was the problem facing the Chrome Steel Works at the beginning of 1927. In addition to determining the physical properties of the alloys under consideration, it was necessary to test their value in service. Small melts of steel in a crucible, from which test specimens could be made, were insufficient to determine more than physical properties, therefore actual castings and ingots had to be made before the products could be placed in service. Furthermore, to save further experimental work, it was desirable that these heats should be made on a "pilot" basis; that is, actually duplicating regular open-hearth furnace practice. Therefore a one-ton (2000-lb.) acid open-hearth furnace was built. Furnace Design There are no special features in the design of this furnace. Standard acid open-hearth practice is followed, with the reproportioning of the dimensions necessitated by its small size; notably, a considerably increased length in proportion to width in order to get proper flame propagation. After the first heat was made it was found that the check-ers were insufficient. The checker chambers were increased in size and the checker openings were made much smaller. The furnace was started late in March, 1927, and made 65 heats. Experimental work was then stopped unti1, the data accumulated could be worked up and summarized. The run was resumed in November and the furnace is still operating, having made at this writing 60 additional heats. For the last series of runs, there was some rebuilding of both the furnace itself and the checkers. Slag pockets were added. Fig. 1 shows details of construction; Fig. 2, the general appearance. All designing was done by the operating and engineering staff. Construction materials were those available around the plant. The reversing valve is a standard butterfly type taken from an old furnace. The slag pockets and checkers are cleaned by digging down to arched openings which have been bricked up.

By K. K. Clark, W. H. Somerton

The feasibility of reducing pressures needed to fracture formations by heating the borehole intensively has been investigated on a laboratory scale. Reductions in fracture pressures of heated Bandera, Berea and Boise sandstone cores were observed at various confining pressures, simulating reservoir stress conditions to a depth of 5,000 ft. Both permeable and impermeable borehole boundary conditions were tested. Three heating conditions were used in the tests. One group of samples was heated in a furnace to a temperature of 600C, cooled to room temperature., and then tested for fracture pressures. Two other groups of samples were heated with a borehole heater to maximum borehole wall temperatures of 625 and 792C, then cooled and tested. In all cases, heated samples showed substantial reductions in fracture pressures compared to unheated samples. The higher temperature borehole heated samples showed the greatest reduction, amounting to 70 per cent in the extreme case. In actual application to wells, it was concluded that fracture pressures may be reduced by 20 to 50 per cent by intensive borehole heating before fracturing. INTRODUCTION Hydraulic fracturing of oil and gas wells is an accepted and effective technique for increasing well productivity in many areas, but is notoriously unsuccessful in other areas. The technique becomes less interesting for production stimulation purposes when excessive fracture pressures are necessary. Formations which do fracture but require high pressures, probably form horizontal fractures. There is good evidence to indicate that some formations do not fracture at all, but yield plastically under high pressure. Thus, any technique which will make formations easier to fracture or will reduce fracture pressures is potentially interesting to petroleum producers. Horizontal fractures generally occur in planes of weakness, such as bedding planes, and at pressures equaling or exceeding the effective overburden stress. In the unusual case where planes of weakness do not exist, a slightly higher pressure must be applied to overcome the tensile strength of the rock. In general, high fluid-loss fracturing fluids are conducive to the formation of horizontal fractures.1 Vertical fractures, when they do occur, are initiated at pressures generally lower than overburden pressure and are favored by low fluid-loss fracturing fluids or impermeable boreholes. The pressure necessary to create a vertical fracture in an impermeable, elastic medium can be predicted from the so-called "thick-walled cylinder" equation: Pf = 2 or+ st.............(1) where pf = fracture pressure or = radial stress st = tensile strength of medium. This equation indicates that to fracture it is necessary to overcome a tangential compressive stress, induced by and equal to two times the radial confining stress, plus the tensile strength of the medium. If the above equation held strictly true for rocks, very little could be done to reduce fracture pressure. The tensile strength of the rock might be reduced somewhat, but any change in st would be small compared to the magnitude of the tangential stress at reasonable subsurface depths. Laboratory tests have shown that actual fracture pressures exceed values predicted from Eq. 1 by factors of 1.4 to 28.6.1,2 Eq. 1 thus does not properly describe the

Jan 1, 1966

By Robert L. Kidd

ONE time or another placer mining attracts the attention of a large number of people, because of the possible low initial investment, low operating cost, and quick returns. Much has been said about sampling and operating placer deposits but many are still inclined to jump at a property without giving much, if any, consideration to proper sampling. This statement does not apply to experienced mining engineers but I am hoping that readers will tear out this article and pass it along where it will do the most good and thus help to preserve the reputation of legitimate mining. It would appear that some people have an ability to sink test pits in high-grade gravel that is positively uncanny. One pit showing high-grade values proves to some property owners and promoters that the entire property will do the same. This is

Jan 1, 1937

By A. E. Powers

YEARS of experience and research have shown that molybdenum, tungsten, and vanadium are among the most useful and effective elements in augmenting the high-temperature strength of heat-treatable, ferritic steels. Grün, for example, in conducting an early survey of the strengthening effects of various elements in annealed, low-carbon steel at 750° and 930°F (400° and 500°C), found molybdenum and vanadium to be the most effective.' Holtmann comprehensively studied the strengthening effects of molybdenum and vanadium in quenched and tempered steels at 930°F as a function of microstructure and carbon content.' He found creep strength to be related to the saturation of the austenite at the austenitizing temperature, such that strength is increased by alloying until the y loop is reached—the terminus of complete austenitization. Beyond this degree of alloying, the inability to heat treat will result in lower strength, at least for low and moderate creep temperatures. Tungsten, owing to higher cost than molybdenum, has been little used in ferritic high-temperature steels. As a result, very few investigations have been made of its high-temperature strengthening ability, and reliable data on its effectiveness are not available. Grün rated tungsten far less effective than molybdenum at 930°F (500°C) on a weight-percentage basis.' However, Tammann has found that tungsten is as effective as molybdenum in raising the recrystallization temperature of iron. For this reason Smith, in a review of the subject, has advocated a re-examination of the influence of tungsten on the high-temperature strength of steel.' In any study of the effects of alloying additions on high-temperature properties, recognition must be made of the many variables involved, some of which can be controlled but have been ignored in innumerable investigations of the past, and some of which are difficult or even impossible to control. One cannot, for example, assign a definite strengthening index to any one alloying element, for this will be dependent upon the microstructure (heat treatment and mechanical treatment), testing con- ditions (temperature, time, and stress), and the complete composition of the steel (accompanying alloying elements, impurities, etc.). One ever-appear ing variable, hardness or tensile strength at room temperature, may be eliminated by heat treating all of the alloys to a single hardness level. An attempt may be made to control the variable of primary quenched structure by oil or water quenching the test specimen. The control of such variables may not be possible in cases of widely varying compositions, but at least such deviations should be recognized. Material and Procedure The alloys, made as 30-lb ingots in an induction furnace, all had the same base composition of about 0.18 pct C, 0.85 pct Mn, and 0.48 pct Si. They were grouped as follows: Mo Group—Containing up to 5.2 pct Mo. W Group-Containing up to 6.0 pct W. V Group—Containing up to 3.3 pct V. Mo-V Group—Containing up to 2.7 pct Mo and up to 1.4 pct V. Mo-W Group—Containing up to 2.5 pct Mo and up to 2.5 pct W. The steels, listed in Table I, were forged to ¾ in. sq bars and the material for the high-temperature, rupture-test specimens was further swaged to %-in. round bars.

Jan 1, 1957

By M. S. Burton, J. E. Ritter

The sessile-drop technique was employed to study the effect of atmosphere and alloy additions of titanium and chromium on the surface tension and contact angle of nickel on sapphire substrates at 1500oC. Shear tests were conducted on the solidified sessile drops. The liquid-gas surface tension of pure nickel on sapphire was found to be 1747 ergs per sq cm, with a contact angle of 109.6 deg. The surface tension and contact angle were not affected significantly by the four atmospheres, vacuum, impure argon, pure argon, and pure hydrogen. Small additions of titanium and chromium decreased the contact angle and for the same alloy content, contact angles were lower in pure argon than in pure hydrogen. On the basis of shear-test results , the most probable cause for adhesion in the Ni-Al2O system is the formation of NiO:Al2O3 spinel at the interface; in the titanium alloy-Al2O3 system adhesion is thought to be due to a mechanical interlocking effect; in the chromium alloy-Al2O3 system adhesion is believed to be the result of solid-solution formation involving Cr2O3 and Al2O3 at the interface. In the last few years there has been considerable interest in reinforcing metals with ceramic whiskers. It has been postulated that, by reinforcing nickel with sapphire whiskers, high-temperature strengths su- perior to those of the best current nickel alloys can be achieved.' To fabricate a whisker-metal composite having optimum density and strength, understanding of wettability and adherence in the system is of prime importance. Although there has been some previous research in this a a need for further studies has been cited4-8 to understand better the effects of impurities, alloys additions, and atmosphere on the wetting and bonding of metals to metal oxides. The purpose of this research was to investigate the factors involved in the wetting and adherence of nickel/nickel alloys to sapphire (single-crystal a Al2O3). Sessile-drop experiments were conducted in vacuum, impure argon, pure argon, and pure hydrogen, to determine the effect of atmosphere on the contact angle of the metals on sapphire, and the liquid-gas surface tension of pure nickel. The solidified sessile drops were sheared from the sapphire substrate to measure the bond strength between the metal and ceramic. The alloy additions chosen were titanium and chromium since both enhance the wetting of nickel on sapphire.2, 3c Also, titanium is commonly used when brazing metal-to-ceramic seals to improve adherence. Since chromium forms an oxide that is completely soluble in sapphire, it was thought that the oxide could serve as a transition phase to promote bonding between the Ni-Cr alloy and sapphire. I) EXPERIMENTAL APPARATUS AND PROCEDURE The sessile-drop apparatus was a horizontal furnace, with an induction-heated, cylindrical molybdenum susceptor in a water-cooled Pyrex chamber. This apparatus was patterned after that of Williams and Nielsen.7 A flat Pyrex window at each end of the furnace allowed temperature measurements to be made at one

Jan 1, 1968

By R. E. Cech, J. H. Hollomon

KURDJUMOV and Maksimova reported experiments with manganese steels and high carbon steels' and with an Fe-Ni-Mn alloy' in which mar-tensite was formed isothermally over a range of temperatures. They found in some cases that mar-tensite formation could be suppressed by rapid quenching to liquid nitrogen temperature. From their microstructural observations of martensite formed isothermally, they concluded that the rate controlling step is nucleation rather than growth. Kulin and Cohen,3 in an attempt to reproduce these experiments, found that with a steel having the same composition as that reported by Kurd-jumov and Maksimova, the transformation to martensite was essentially complete above the temperature range of Kurdjumov and Maksimova's isotherms. The possible reasons for this disagreement were not considered. Recent papers by Das Gupta and Lement4 and Kulin and Speich5 report the formation of isothermal martensite in a high chromium steel and in an Fe-Cr-Ni alloy, but neither paper can be considered a verification of the original Kurdjumov and Maksimova results. Further, in neither case were the authors able to suppress the formation of martensite entirely. Because of the important bearing the Kurdjumov and Maksimova results have to an understanding of the mechanism of martensite reactions it was felt that an experimental investigation directly concerned with checking the validity of their results was in order. This paper describes the results obtained on the isothermal transformation over the temperature range from —79" to —196°C of an alloy of iron, nickel, and manganese. Experimental Apparatus A 15 lb heat of an alloy containing 73.3 pct Fe, 23.0 pct Ni, and 3.7 pct Mn was melted by induction and cast under argon. The ingot was forged to 1-in. bar and a portion rolled to 1/16x1 1/2-in. strip. This strip was pack-homogenized 300 hr at 1100" in a helium-filled sealed iron tube. The composition after homogenization was 73.2 pct Fe, 22.94 pct Ni, 3.73 pct Mn, 0.05 pct C, and 0.015 pct N. The strips were cut to 1/2-in. width for dilatometer and metal-lographic specimens. Only the center portion of the 11/2-in. strip was used in the present investigation. The dilatometer employed was similar in design to one described by Flinn, Cook, and Fellows." A concentric fused auartz rod and tube assembly with hooks for holding the specimen was mounted so as to transmit the specimen dilation to a 1/10,000 in., 1/10 in. travel dial gage. The dilatometer proper was mounted by means of extension arms to a counterweighted sliding member on a vertical standard. This method of mounting permitted rapid transfer of the dilatometer from the austenitizing furnace to the quenching bath and low temperature chamber. A small electrical vibrator on the dilatometer kept frictional effects of the quartz members at a minimum. The austenitizing unit was a vertical, molybdenum-wound, hydrogen atmosphere furnace maintained at a constant temperature ±3°C by means of constant power input. A 12-in. stainless steel jacketed copper liner having 1/2-in wall thickness acted to equalize the temperature in the hot zone of the furnace. This liner, closed at the bottom end and open at the top to permit entrance of the dilatometer and specimen, was kept filled with dry nitrogen gas. A chromel-alumel thermocouple was placed inside the tube to determine the temperature. The 4-in. dilatometer specimens in the chamber varied less than 1/2° across the specimen length except for a 1 1/20 drop at the end nearest the open end of the furnace. The low temperature isothermal holding bath was a double Dewar arrangement similar to one described by Turnbull7. The outer bath was filled with a refrigerant at a temperature lower than the desired holding temperature. The inner bath was filled with Freon "11" or "12" or a mixture of both, depending upon the holding temperature. This inner bath which tended to be cooled by the outer bath was kept at a constant temperature by introducing a small amount of heat with a manually controlled electric heater. Stirring was accomplished by bubbling dry air through the bath. A Leeds and North-rup type K potentiometer was used to measure the inner bath temperature as indicated by a five element copper-constantan thermopile. The bath temperature was maintained within ±0.2°C of the desired temperature by occasionally adjusting the heater current so as to keep the Leeds and Northrup galvanometer at zero deflection with a constant setting of the potentiometer. Isothermal tests were usually continued for 300 to 400 min and another reading made at approximately 1000 min if the bath, unattended overnight, had not deviated in temperature more than 5°C. Transformation curves are drawn dashed (Fig. 1) through the time region where temperature was not controlled precisely. Experimental Procedure Dilatometer specimens of 1/2x1/16-in. strip were cut to 41/2-in. length and holes were drilled for the quartz hooks with proper spacing to give a 4-in. measured length. A thermocouple consisting of 0.012-in. diameter chrome1 and alumel wires was spot welded to the specimen and threaded between the dilatometer rods to binding posts near the dial

Jan 1, 1954

By Raymond Brandon, Marshall Pease

A pattern for providing a large utility, The Detroit Edison Co., with an adequate fuel supply is outlined. From the standpoint of both fuel procurement and utilization, consideration has been given competition relationships, fuel properties as well as operating conditions that must be met to provide dependable service and reliable plant operation.

Jan 2, 1951

By G. S. Anderson

PRIOR to 1912 the only rail outlets for a large part of the coal regions of Carbon and Emery Counties. Utah, were over single-track lines of the Southern Utah R.R. and Castle Valley Ry. Companies, forming a route extending from Mohrland and Hiawatha, Utah, to a point of connection at Price, Utah, with what is now The Denver and Rio Grande Western R.R. Co. These two tap lines were poorly and inadequately constructed, with impracticable grades, lacking in equipment, and incapable of being put into condition to serve this region and handle the output of the mines, to say nothing of making possible increased production and new development. Any plans for the new development and increased output necessarily involved new and adequate railroad service. It was essential, not only that additional facilities be provided but that these facilities be in the form of an efficient and up-to-date railroad providing an independent through operation from the coal mines to a point of direct connection with the lines over which the product would be distributed to consumers. For this purpose, the Utah Railway Co. was incorporated in

Jan 1, 1948

By K. A. Johnson, F. H. Ellinger, C. C. Land

The Pu-In phase diagram has been determined by thevmal, filtvation, micrographic, and X-ray diffraction methods. This alloy system is characterized by 1) limited solubility of indium (-2 at. pet) in 6 and e plutonium, 2) one eutectic, and 3) fizte intermediate phases. Of the latter, PuIn and PuIn3 melt con-gruently, and Pu31n, Pu31n5, and the high-temperatuve 7 solid solution are formed peritectically. The n phase decomposes eutectoidally into Pu3In and PuIn at 865°C. Although the 6 phase can readily be retained to room temperature, it is not thermo-dynamically stable below 300°C. The Pu-In system has been investigated at Los Alamos as part of a study of the alloys of plutonium with the group 111-B elements. Constitutional studies of the Pu-A1 and Pu-Ga systems have been reported earlier.'9' To date, published information on the constitution of Pu-In alloys covers only the identification of the intermediate phases Pu3In and PuIn3.3,4 EXPERIMENTAL The alloys were prepared from five different lots of plutonium and three different lots of indium. The principal impurities in the plutonium differed from lot to lot within the following limits (in ppm): iron, 30 to 80; nickel, 30 to 60; aluminum, 15 to 30; silicon, 20 to 145: thorium, 50 to 300; carbon, 75 to 335; and oxygen, 20 to 135. Spectroanalysis of the indium showed that two of the lots were 99.9 wt pct pure, and the third was 99.98 wt pct pure. The alloys, in the form of about 5-g buttons, were prepared by arc melting in conventional equipment. The components for each alloy were weighed to the nearest 0.1 mg. In most instances the weight losses during melting were negligible; therefore confirmatory chemical analyses were usually not made. A few results are given in Table I. Density measurements of the arc-cast buttons were made by the liquid displacement of bromoben-zene. Heat treating was done in muffle furnaces controlled to +2°C. The alloy buttons were wrapped in tantalum foil before being enclosed in evacuated clear silica capsules. As a precaution against breakage, each capsule was in turn placed in a silica tube closed with a refractory wool plug prior to being charged into the furnace. For a moderately fast quench of the alloys, the outer protection tubes containing the capsules were plunged into cold water. Differential thermal analyses (DTA) in racuo were performed at heating and cooling rates of about 1.5°C per min. Tantalum crucibles having inverted thermocouple wells were used for most of the analyses. It was found micrographically, however, that the liquid alloys containing more than about 35 at. pct In reacted with the tantalum, resulting in contamination of the surface of the alloy. A search for a more satisfactory crucible material was not successful: tungsten and yttria behaved comparably to tantalum; thoria, vitrified magnesia, and vitrified alumina severely contaminated the alloys. So as to minimize the contamination by the crucible mater-

Jan 1, 1965

By Jack L. Taylor

Single crystals of tungsten grown from powder -metallurgy swaged rod by high-temperative annealing were deformed in tetzsion at temperatures from 2500 to 5000 OF. Orientation of specittzen tensile axis, strained ratrix, and defortnation bands was determined optically by reflections from {110} etch pits. Slip traces meve analyzed and slip directiorz determined. Results indicate that {110}(111) , {112){111}. and {123}{111} type slip occur in tungsten over the terpe,atuve rutge investigated. Slip is orientation-dependent occuvring on tCuzt cowzbitzation oj slip plane and direction which has the highest critical resolved shear stvess. Overshooting appears to be a general occurrence between 2500' and 5000°F. Dejornatiotz bands show rotation in a directiorz opposite to the yotation of tire tensile axis. DEFORMATION in bcc metals has been reviewed by Maddin and Chen,' Keh and eissman,' and most recently by Nabarro, Basinski, and olt. From these reviews it is apparent that some disagreement exists concerning the crystallography of plastic deformation in bcc metals. There is evidence of noncrystallo-graphic slip, the crystal slipping on or near a noncrys-tallographic plane in the (111) zone for which the resolved shear stress is greatest. There is also evidence to indicate that crystallographic slip occurs in the planes of the same zone. On the other hand, there is agreement that the close-packed direction ( 111) is the slip direction in bcc metals. Further, in pure bcc metals4 and alloys the operative crystallographic slip systems are strongly dependent on temperature. Both crystallographic and noncrystallographic slip have been reportede8 in tungsten, a group VIA bcc metal. The number of slip families which operate in the case of crystallographic slip, however, has been in question. The earliest study by oucher' of single-crystal tungsten tested at temperatures of 1800 to 5000°F led to the conclusion that only {112)( 111) type slip operated. Raymond and eumann recently reached the same conclusion studying deformation by rolling of plasma-flame single crystals at 1800°F; however, only [loo] and [110] orientations were studied. Work at high temperatures by Leber and Pughs on tungsten and Chen and Maddin9 on molybdenum suggests that conjugate slip on nonparallel (110) planes may account for the slip on (112) and (123) planes found by others working with bcc metals. The present paper presents evidence of three families or modes of slip in tungsten in the temperature EXPERIMENTAL PROCEDURE Material. Standard microtensile specimens, 0.160 in. diam by 0.640 in. gage length, were ground from commercial powder-metallurgy swaged tungsten rod, type MK. Specimens were heated for 10 min at 5150°F or 6 hr at 5400°F in vacuum of 1 x 105 Torr or less to produce single crystals. At the lower temperature an average of two out of seven specimens heated in a group developed single crystals throughout the specimen length as shown in Fig. 1, the rest remaining wholly poly crystalline. At the higher temperature, three or four out of seven specimens became single crystal. Other than to note that Laue X-ray photographs generally showed sharp, well-defined spots, no attempt was made to assess crystal perfection or measure dislocation density. Two tungsten single crystals grown from the melt by the plasma-arc process were also used for this study.

Jan 1, 1967

By B. R. Coil

One executive writing on capital equipment re- placement stated: "We keep extremely close watch over repair costs and when we reach the point of uneconomical operation, we replace the machine." This, of course, is most elementary, but a very basic premise. However, it seems advisable to set forth several equally obvious principles, so that all discussion will start from the same point. The subject is intended to deal with the point of uneconomical operation due to high operating costs in general and high repair costs in particular. This is not concerned with a similar problem of high costs because of an insufficient number of machines, be they trucks, shovels, or dozers.

Oct 1, 1955

By A. B. Kinzel

INCLUSIONS of nonmetallic matter have long been recognized as objectionable in steel. A complete theory of the effect of inclusions, which is consistent with that held today, was outlined in Howe's early work on the metallurgy of iron and steel. This, by analogy with wrought iron, assumes that the observed difference in properties between specimens with and across the direction of rolling is due to slag enclosures. The fact that wrought iron shatters in ballistic test is taken as an indica-tion of the effect of inclusions on the resistance to shock. Since that time, the development of the metallurgical microscope has shown the discrete nonmetallic inclusions in the iron matrix. Better resolution has, brought out the great number of inclusions, and the availability of good microscopes in steel mills and laboratories has constantly drawn more attention to these inclusions. This recent emphasis on inclusions has resulted in studies as to their nature, mode of formation and methods of elimination, such as the work of Wohrman, Dickenson, Herty, and Schenk. In addition to the published works; many investigations which have not been published' have been carried out by private laboratories and have resulted in the application of specifications intended to limit inclusions in many qualities of steel. Nevertheless, a clear understanding of the quantitative effect of inclusions on the physical properties of steel has been lacking. It is the specific purpose of this study to determine the quantitative effect of inclusions on the physical properties of steel that determine its behavior under dynamic stress. The study has shown that a quantitative relation between counted inclusions and impact strength does exist in normalized steels, but 'that factors other than readily visible inclusions are equally important. Furthermore, the present commercial rating of inclusions is unreliable with respect to counted inclusions or dynamic tests, and this study indicates that the tensile-impact test is more reliable as a measure of quality.

Jan 1, 1931

By R. A. Graham, R. W. Rohde

The effect of hydrostatic pressure upon the austenite start temperature of a commercial Fe-28.4 at. pct Ni-0.5 at. pct C alloy has been determined. For pressures to 20 kbar, the austenite start temperature decreased from its atmospheric pressure value of 380°C at the rate of about 4°C per kbar. These data are analyzed by two different thermodynamic approaches; first, considering the transformation as an isothermal process, and second, considering the transformation as an isentropic process. It was found that both these approaches fit the experimental data equally well. The effect of hydrostatic pressure upon the austenite start temperature is best described by considering the mechanical work done during the transformation as that work obtained by multiplying the applied pressure with the gross volume change of the transformation. It is widely recognized1 that strain has an important effect on the initiation of martensitic transformations.* For example, the martensite start tempera- *In this paper, use of the term martensitic transformation implies the reversal of martensite to austenite as wen as the formation of martensite from austenite. ture, M,, may be increased by plastic deformation. Similarly, plastic deformation is observed to lower the austenite start temperature, A,. The effect of uniaxial stress on the M, of iron-nickel alloys has been studied by Kulin, Cohen, and Averbach.2 They found that the martensite start temperature was significantly changed by stresses well within the elastic region. Moreover, the effect of tensile and compres-sive stresses differed. These effects were explained in terms of the interaction of the applied stress with both the dilational and shear components of the transformation strain. The magnitudes of the influence of uniaxial tension, compression and hydrostatic pressure on Ms were measured in 30 pct Ni 70 pct Fe by Pate1 and Cohen.3 Their thermodynamic calculations and similar calculations by Fisher and Turnbull4 predicted the experimental results when the transformation was assumed to occur isothermally at some fixed driving force. This driving force was assumed to be supplied by a combination of the chemical free energy difference between the austenitic and martensitic phases and the work performed during transformation by the applied stress. More recently, Russell and winchel15 reported the effect of rapidly applied shear stress on the reversal of martensite to austenite in iron-nickel-carbon alloys. They performed a thermodynamic analysis of this transformation based upon the assumption that the re- versal occurred adiabatically. They concluded that the applied shear stress did not significantly interact with the transformation strain and thus did not assist in inducing the reversal. Rather they concluded that the reversal was effected by localized strain heating which resulted from the gross local shear deformation of the experiment. In either the adiabatic or isothermal analysis it is necessary to compute the work performed by the interaction of the applied stress and the transformation strains. In the case of hydrostatic pressure this interaction has been treated by two different methods. In either case the applied pressure is assumed to remain constant during the transformation. In one treatment the applied pressure is assumed to interact directly with the dilatational strain associated with the formation of an individual martensite plate.3'4 This local strain has been measured at atmospheric pressure in iron-nickel alloys by Machlin and Cohen.6 In the above treatment this local strain is assumed invariant with temperature and pressure changes. In the other treatment the applied pressure is assumed to interact with the gross volume change of the transformation.7,8 The usefulness of this latter treatment has been demonstrated by Kaufman, Leyenaar, and Harvey7 who calculated the effects of pressure upon the martensite and austenite start temperatures of Fe-10 at. pct Ni and Fe-25 at. pct Ni alloys. Excellent agreement was obtained between their calculations and their experimental data on an Fe-9.5 at. pct Ni alloy. However, this treatment suffers from the fact that the data required to calculate the volume change of the transformation (i.e., the initial specific volumes, the thermal expansion and compressibility data for both the austenitic and martensitic phases) is, in general, not available for any material except pure iron. Thus the calculations of Kaufman et al.7 were necessarily performed by assuming that the volume change of the martensitic transformation in the iron-nickel alloys was that same volume change occurring during the a-? transformation in pure iron. While this approximation may suffice for very dilute alloys it is likely to be inaccurate in high nickel alloys. We have performed measurements of the effect of hydrostatic pressure to 20 kbar on the A, temperature of an Fe-28.4 at. pct Ni-0.5 at. pct C alloy. The composition is similar to the alloy used by Pate1 and Cohen3 to determine the effect of pressure upon the M, temperature. The present measurements permit calculation of the interaction between the applied pressure and the transformation strain. Additionally, measurements have been made which allow precise determination of the gross volume change of the transformation. The data allow direct comparison between the alternate hypotheses of the interaction between the applied pressure and a dilatational transformation strain characterized by either the formation

Jan 1, 1970

Secretary's Note.—M. Beutter's introductory remarks, being simply a resume of Prof. Hone's paper, with a disclaimer of the intention to criticise it, have been omitted, and the remainder of this contribution lias been somewhat condensed, without omitting any essential statements. M. Beutter, St. Etienne, France* (communication to the Secretary†):—Professor Home observes in $ 47 of his paper that the fluid-compression of steel diminishes piping by forcing the liquid and viscous metal, while the pipe is forming, or after it has formed, into the space so produced, and proceeds to study the different systems of compression, alluding to the four principal ones: Whitworth, Illingmorth, Williams, and Harmet. Having now for many years had in use the last-named process, that of M. Henri Harmet, I shall attempt to explain the difference between it and the three other processes named. The Harmet process does not diminish or close the pipe after it has formed, but prevents it from forming; and I may add that it is the only one of the four which produces this result. It is true that the Whitworth process diminishes the interior hollow by localizing it around the central axis of the ingot, which is afterwards removed by trepanning the hollow forg-ings for which this process has been, I believe, always exclusively used. The Illingworth and Williams processes, which, in my opinion, are applicable only to small ingots, of necessity allow the pipe to form, or (to use the expression of Professor Howe) they allow the surface of the lake of liquid steel to sink into the interior of the ingot, since they involve waiting a certain time before compression, so that (in the case of the Williams process) the ingot-mold can be opened without danger, or (in the case of the Illingworth process) so as to avoid the risk that the steel may enter the grooves produced by the removal of the distance-bars. During the time required for

Jan 1, 1908

By Margaret V. Doyle, R. W. Powers

INTERNAL friction measurements, carried out as functions of temperature, have been used extensively to obtain data on the mobility of interstitial impurities in the Group V metals, vanadium, colum-bium, and tantalum. For a given frequency of vibration, the internal friction peaks at a temperature which depends principally on the particular metal and interstitial impurity. Relaxation times and corresponding activation energies are determined from a knowledge of such peak temperatures for various vibration frequencies. These experimentally determined relaxation times are related to the average jump frequencies of the interstitial impurity atoms using the model of Snoek1 and can be related to macroscopic diffusion coefficients as shown by Polder.2 In these Group V elements, the impurity whose diffusion rate has been the hardest to pin down by the internal friction technique has been carbon. The experimental difficulties associated with carbon in these metals were not evident in the earliest investigation carried out by Ke on tantalum.3 The carbon peak was thought to occur near 140°C at 1 cps on the basis of the enhancement of the height of this peak measured on specimens which had been coated with colloidal carbon and heated in vacuo at 1200°C. Presumably on the basis of similar evidence, Wert identified a peak also near 140°C at 1 cps in columbium with the diffusion of carbon in that element." Later, the difficulties of admitting carbon into these metals without simultaneously introducing oxygen were pointed out by Ang and Wert- nd by Marx, Baker, and Sivertsen.6 The latter suggested that the peaks attributed to carbon in tantalum and columbium might arise from the presence of traces of oxygen, adventitiously absorbed during the carburization. Such suspicion was confirmed, at least in tantalum, when it was found possible to prepare C-Ta specimens without contamination by oxygen. Specimens sci prepared, aitnough high in carbon content, gave no evidence of a damping peak near 140°C.7 Recently, the authors were able to Show that carbon diffuses in tantalum at a rate only slightly slower than that for nitrogen.8 At 1 cps, the carbon peak lies at 34g °C, only a few degrees above the nitrogen peak at 341 °C. However, the two could be distinguished by a different stability with respect to aging at 400 °C and by a different dependence of peak height on solute concentration. On the basis of this experience with tantalum, it appeared worthwhile to search also for a carbon peak in columbium in the vicinity of the nitrogen peak." In this report, the location of the carbon peak in columbium is made known and its position relative to the nitrogen peak is compared with the authors' previous findings in tantalum. In addition, a second source of internal friction seen only in specimens of very high purity is discussed. Columbium was obtained from the Fansteel Metallurgical Corp. as either 30 or 40 mil diam stock. The manufacturer gives the following analysis as typical of this material: 99.4 pct minimum Cb, 0.10 pct maximum C, 0.5 pct Ta, 0.04 pct Ti, 0.02 pct maximum Si, and 0.01 pct maximum Fe. Wires 13 in. long were degassed near 2200°C for 4 to 16 hr in an ultimate vacuum better than 2.10" mm Hg in an apparatus described previously.8 This treatment reduced the oxygen and nitrogen damping peaks observed in the as-received material to less than 0.0001 in Q-1 above background. The latter value varied somewhat from specimen to specimen in the range 0.0004 to 0.0008, depending chiefly on the wire diameter and the frequency of vibration. The high temperature evacuation left the specimen in a rather soft condition, the Vhn being reduced to 40 from a value near 121 in the as-received conGi~iur~. The gram size was about the diameter of the specimen. After outgassing, specimens were loaded with carbon by absorption of such hydracarbons as methane or ethylene. Carbon was also added to a few speclrnens by heating them at

Jan 1, 1958

By R. H. Frazier, C. H. Lorig, F. W. Boulger

THERE are both advantages and disadvantages in using semikilled steels in place of killed steels. One advantage of semikilled steels is they provide a higher ingot-to-product yield. This is especially important to shipbuilders in the time of national emergency. By using semikilled steel for hull plates, a greater number of ships can be built from the same ingot tonnage. Consequently, if semikilled steel plates are adequate for the service, the advantage of the higher product yield is of considerable importance. The performance of hull plates is closely associated with the ductile-to-brittle transition temperature, and this, in turn, is strongly dependent on the composition of the steel. A low transition temperature is desirable because it indicates that the steel is less likely to fail at low ambient temperatures. Killed steels are known to have lower transition temperatures than semikilled steels. It is believed that the better qualities of killed steels in this respect are due mainly to the low oxygen contents of the steel. The principal deoxidizers, aluminum, silicon, and manganese, lower the oxygen content. Fundamental studies1-0 have shown that the oxygen content remaining after the addition of one of these three elements is influenced by the residual amounts of the other two present. In the current study, therefore, various amounts of silicon and aluminum were added to steels containing different manganese contents for the purpose of studying the influence of silicon and aluminum on the notched-bar properties of hot-rolled steels. Eleven types of steels were studied. The temperature at which the plates were finish-rolled was care- fully controlled at 1850°F. All plates were rolled to ¾-in. thickness, and tested in the as-rolled condition. The total number of steels prepared in the laboratory for this study was 95. The charge consisting of ingot iron and ferrosilicon equivalent to 0.10 pct Si was melted in a 200-lb induction furnace under an atmosphere of argon to insure low, uniform nitrogen contents. After the charge was melted and the desired temperature was obtained, the melt was partly deoxidized with either silicomanganese or aluminum. Aluminum was used for this purpose in only those steels with very low silicon contents, where the finished steel was to contain some aluminum. This initial deoxidizing addition was made to obtain consistent recoveries of subsequent additions of ferro-manganese, ferrosilicon, and aluminum. Carbon, in the form of graphite, was added, either just prior to tap or to the final aluminum addition. The entire heat of 200 lb was poured directly into a 6x6-in. big-end-up mold. The ingots of semikilled steel were capped with a steel plate. The killed steels, on the other hand, were poured with a hot-top containing 14 pct of the total volume of the ingot. The ingots were processed by heating to 2250°F, followed by forging to slabs 1¾ in. thick and 6 in. wide. After reheating to 2250°F, the slabs were rolled to 0.9-in. gage, using reductions of approximately 1/6 in. per pass. The 0.9-in. thick plates were immediately recharged in a furnace held at 1850°F. After 30 min in the furnace at 1850°F, the plates were rolled to ¾ in. in one pass. Following the final pass, the plates were stacked on edge on a brick floor, with a brick separating one from another, where they were allowed to cool in air. Drillings for chemical analysis were taken from the plates at locations corresponding to the top and bottom of each ingot. Carbon, manganese, silicon, phosphorus, sulfur, and nitrogen contents of each sample were determined. The analysis of each steel was carefully controlled so that the contents of

Jan 1, 1957

By H. M. Howe

One of the little electric-resistance magnesia crucible-furnaces which I designed for the metallurgical laboratory of the School of Mines of Columbia University is shown, in vertical section, of full natural size, in Fig. 1. The furnace itself consists of two semi-cylinders of magnesia, A A, with a cover of magnesia, B B, and a stopper, C, also of magnesia, which is perforated to admit the leads x, y, of a thermo-electric couple. The furnace is heated by the spiral of platinum wire, D D (entering and departing by the openings I and K), through which a current of any desired strength is passed, bringing it to incandescence, and to ally temperature, not above the melting-point of platinum, which may be desired. The crucible E, actually used in this furnace, is also of magnesia. I have heard of some unsuccessful attempts with furuaces of this general class, and have attributed these failures to the use of siliceous materials for the furnace-walls. Of course such a result is what we should expect, since the silica mould readily be electrolysed by the current, and the resultant silicon would alloy with platinum ; but the little magnesia furnace here clescribed seems to be wholly free from this objection. It has been used in my laboratory for some months, and temperatures, up to 1100° C, hare been developed in it over considerable periods of time—indeed, it has been held near 1400" C. for about two hours—without indication, thus far, of any deterioration of the platinum. In the section shown, F F are a pair of steel disks, bound together, but separated by the inverted thermo-junction, N, of the thermo-electric pyrometer. This arrangement permits the observation and record of retardations in the heating and cooling of F F—in other words, the determinations of the cooling curves of steel—a purpose for which the furnace has been used, thus far, more frequently than in the determination of meltingpoints. For the latter operation, however, it is also suitable.

Jan 1, 1902