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By Charles H. Eldridge, Robert J. Anderson

When a metal or alloy is poured into a mold, internal stresses are set up by the cont,raction in volume on passing from the liquid state at the temperature of pouring to the solid state at the ordinary temperature. It is assumed by the writers that all castings are internally stressed. That any casting is internally stressed is shown by its warping or distortion on maturing at the ordinary temperature, by warping or cracking on heating, or by cracking or distortion on machining. In the investigation made by the writers, the internal stresses in cast non-ferrous alloys were studied with a view to examining the effects of heat treatment (i.e., simple anneal) on the release of such stresses. Data have been obtained that show quantitatively that internal stresses can be removed by appropriate heat treatment. The investigation was limited to a study of the effect of heat treatment on stress release in cast rings of the 87:7:5:1 copper-tin-zinc-lead alloy. The writers are indebted to E. G. Fahlman for the method used for inducing stress in cast rings; to M. E. Boyd for microscopic examination and for other aid; and to G. B. Dalrymple for the chemical analyses. General Information While casting stresses have been discussed considerably, they have usually been treated in a general way, and so far as the writers know no measurements of the magnitude of internal stresses in castings have been reported. It has been suggested by many investigators, particularly in the case of steel castings, that castings are internally stressed, and that these stresses can be removed by annealing. In a casting, the internal

Jan 1, 1923

By Robert J. Anderson, Charles H. Eldridge

When a metal or alloy is poured into a mold, internal stresses are set up by the cont,raction in volume on passing from the liquid state at the temperature of pouring to the solid state at the ordinary temperature. It is assumed by the writers that all castings are internally stressed. That any casting is internally stressed is shown by its warping or distortion on maturing at the ordinary temperature, by warping or cracking on heating, or by cracking or distortion on machining. In the investigation made by the writers, the internal stresses in cast non-ferrous alloys were studied with a view to examining the effects of heat treatment (i.e., simple anneal) on the release of such stresses. Data have been obtained that show quantitatively that internal stresses can be removed by appropriate heat treatment. The investigation was limited to a study of the effect of heat treatment on stress release in cast rings of the 87:7:5:1 copper-tin-zinc-lead alloy. The writers are indebted to E. G. Fahlman for the method used for inducing stress in cast rings; to M. E. Boyd for microscopic examination and for other aid; and to G. B. Dalrymple for the chemical analyses. General Information While casting stresses have been discussed considerably, they have usually been treated in a general way, and so far as the writers know no measurements of the magnitude of internal stresses in castings have been reported. It has been suggested by many investigators, particularly in the case of steel castings, that castings are internally stressed, and that these stresses can be removed by annealing. In a casting, the internal

Jan 1, 1923

By J. E. Neal, C. T. Naber, O&apos

Silicon thin films were deposited by electron beam evaporation in an ultrahigh vacuum onto (0001) and (1102) sapphire substrates. Attempts were made to correlate the structural properties of the deposited silicon films with the following: 1) substrate orientation, 2) substrate surface condition, 3) substrate temperature, and 4) deposition rate. The substrate temperatures were varied between 500° and 1000°C and the deposition rates were varied between approximately 50 and 700A per min. The following sapphire surface treatments were investigated: 1) annealing in the ultrahigh vacuum at 1200° C; 2) heating in a hydrogen atmosphere at about 1350°C; and 3) etching with silicon by evaporating silicon onto sapphire substrates heated to 1200° or 1300°C. Twinned silicon films were occasionally formed at substrate temperatures between the 800° and 1000°C on unprocessed mechanically polished substrates; however, the reproducibility and crystallinity of these films were generally poor. Single-crystal and twinned silicon films were formed at substrate temperatures between 700° and 1000°C on substrates which were silicon-etched at 1300°C prior to deposition. Fiber texture films were formed at substrate temperatures between 500° and 700°C. The (111) lattice plane of the silicon single-crystal films was parallel to the (0001) sapphire plane and the (100) silicon plane was parallel to the (1102) sapphire plane. TECHNIQUES are well-established for growing single-crystal silicon thin films on sapphire by chemical methods.'-* Vacuum epitaxial techniques have a number of advantages over chemical epitaxial techniques in the fabrication of thin film microcircuits.5 One of the most significant is that vacuum techniques are compatible with well-established procedures for forming vacuum-deposited thin-film resistors, capacitors, and interconnections. Salatna, Tucker, and young6 recently reported the formation of relatively poor crystalline-quality silicon films on sapphire by electron beam evaporation in a vacuum of about 7 x 10-7 Torr. The purpose of the investigation described here is to determine the relationships between the structural properties of silicon thin films deposited on sapphire substrates by electron beam evaporation in an ultra-high vacuum and the following: 1) substrate orientation; 2) substrate surface condition; 3) substrate temperature; and 4) deposition rate. The substrate temperatures were varied between 500° and 1000° C and the deposition rates were varied between approximately 50 and 700A per min. Two orientations of sapphire were used: (0001) and (1102). The thick-nesses of the films were between 8000 and 16,000A. The morphology and crystallinity of the deposited films and substrate surfaces were investigated by optical microscopy, electron microscopy, and electron diffraction. APPARATUS AND PROCEDURE A cross section of the ultrahigh-vacuum thin-film evaporator is shown in Fig. 1. The bell jar is stainless steel and copper gaskets are utilized at all seals. High vacuum is attained by a 6001iter per sec ion pump and a titanium sublimation pump. Sorption pumps are used to rough the system to a pressure of about 10 µ prior to starting the ion pump. Vacuums in the 10- 10 Torr range were attained after baking the system at about 250°C for 8 hr. Vacuums in the 10-8 to 10-9 Torr range were maintained during silicon evaporation. The silicon charges, which were cut from p-type single-crystal boules which had resistivities between 2000 and 4500 ohm-cm, were placed in the water-cooled copper crucible of a 180 deg bent-beam 5-kw electron beam evaporator. During evaporation, a molten zone was formed in the top center of the silicon charge, and the portion of the silicon charge in contact with the water-cooled crucible remained solid. The substrate heater, which has been described

Jan 1, 1969

By T. E. Leontis

Data are presented in this paper to show that addition of zirconium to sand-cast Mg-Th alloys effects a marked decrease in the grain size of these alloys which is accompanied by a significant increase in the mechanical properties over the entire range of thorium content investigated. The beneficial effect of zirconium on the strength properties is maintained at elevated temperatures up to 700°F. In addition the alloys exhibit exceptionally high creep resistance at temperatures up to 600°F. Zirconium does not greatly improve the strength or creep characteristics of extruded Mg-'Th alloys. Cerium is not a desirable addition to Mg-Th-Zr alloys. IN a previous paper,' it was shown that addition of thorium to magnesium imparts exceptionally high resistance to creep at elevated temperatures both in the sand-cast and in the extruded states. The large grain size of the sand-cast alloys, however, renders them of questionable value as engineering materials. Sauerwald² showed that the grain size can be reduced effectively by the addition of zirconium. Results of a detailed study of the properties of Mg-Th-Zr alloys over a wide range of thorium content are presented in this paper. In accordance with the preliminary findings of Sauerwald,² it is shown here that Mg-Th-Zr alloys possess the unusual combination of fine grain size and a high level of strength and creep resistance at elevated temperatures. Some information on the effect of cerium on the properties of Mg-Th-Zr alloys is presented also. Effect of Th on Solubility of Zr in Molten Mg Sand-cast test bars and 3 in. diam extrusion billets of Mg-Th-Zr and Mg-Th-Ce-Zr alloys* were pre- pared according to the procedures described in the previous paper.' The zirconium was added in the form of sponge metal as supplied by the Bureau of Mines.' The maximum amount of zirconium that usually can be introduced into magnesium at the customary alloying temperature of 1350° F is about 0.7 pct. The analyses in Table I clearly show that thorium in- creases the solubility of zirconium in liquid magnesium. The tabulated zirconium contents are designated as "acid soluble;" this serves to distinguish these values from the small amount of zirconium that is not dissolved when the alloy is subjected to the action of a dilute acid solution. The "acid-insoluble" zirconium is believed to be ineffective as an alloying ingredient. The acid-insoluble zirconium contents ranged between 0.03 and 0.16 pct in the present alloys, the tendency was toward higher values in alloys with higher thorium content. In contrast to the effect of thorium, cerium decreases the solubility of zirconium in magnesium. This is shown by the three Mg-Th-Ce-Zr alloyst in- cluded in Table I and stands in agreement with the observations of Nelson and Stricter".' on Mg-misch-metal-Zr alloys. Effect of Zr on Grain Size of Mg-Th Alloys The potent effect of zirconium as a grain refiner has been demonstrated for many types of magnesium alloys.3-8 Of greatest importance to date are Mg-Zn-Zr alloys5,6,8 and Mg-mischmetal-Zr alloys.3,4,7 Zirconium has also been found to exert a marked grain-refining action on Mg-Th alloys. All the Mg-Th-Zr alloys cast in the form of ½ in. diam test bars have an exceedingly fine grain size (Table I). Reference to the previous paper' will show that binary Mg-Th alloys containing up to 6 pct Th are almost completely columnar even in the M in. diam reduced section of a test bar, and that 10 to 50 pct Th alloys have an equiaxed grain structure of 0.03 to 0.08 in. The data in Table I also show that addition of cerium to Mg-Th-Zr alloys has no effect on the grain size. In addition to increasing the strength properties, the fine grain size of Mg-Th-Zr alloys

Jan 1, 1953

By Lawrence A. Shepard, Richard H. Krock

A series of ductile, two phase W-Ni-Fe composites, sintered in the presence of a liquid phase, were tested in tension. Identical room temperature stress-strain curves were obtained for specimens containing from 80 to 92 wt pct W (58 to 75 vol pct W particles). The composites exhibited a maximum elongation of 29pct at room temperature, and 10.7 pct at 77 °K. The tungsten particles in the composite elongated by the same amount at these temperatures. Single-phase alloy specimens matching the composition of the composite matrix showed about one half the flow stress of the composites. The test results demonstrated that the mechanical properties of W-Ni-Fe composites are determined by the tungsten particles alone and are independent of matrix volume fraction or mean free path over the composition range studied. A general prediction of the plastic behavior of two-phase alloys or composites from a knowledge of the properties, size, shape, and dispersion of the individual components is not, at present, possible. The excellent theoretical and experimental exposition of the dispersion-strengthening problem applies only to cases where the stronger phase is present as fine particles and in volume fractions of a few percent. Materials in which the stronger phase constitutes a significant fraction of the volume, although of broad engineering interest, have received comparatively little analytical attention. This important group of materials is the concern of the present study. The degree of deformation of each phase during the plastic working of a ductile two-phase alloy has been established by Boas and Honeycombe,1 Clare-brough,2 and Clarebrough and berger.3 These workers empirically correlated the recrystallization temperature of the individual phases with the degree of cold work sustained by each. It was shown that the deformations of both phases are equal and therefore the same as that sustained by the composite whole when the volume fraction of the stronger phase exceeds 35 pct. At yielding, however, deformation begins in the weaker phase. A direct correlation of composite strength, as a function of deformation, with the strength of the individual components cannot be made. In the presence of stronger phase particles, the flow stress of the weaker matrix phase is enhanced by an additional hydrostatic component. The development and potential magnitude of matrix strength promotion upon straining has not even been experimentally established. Quantitative work in this field has been confined to ceramic-metal composites. These materials fracture at strains considerably below 1 pct by cracking of the ceramic particles.4-6 An indication of the potential constrained matrix strength capacity is given by the observation by Nishimatsu and Gur-land4 that the maximum tensile strength of a WC-Co composite is almost double the bulk strength of the cobalt matrix. The relationship between composite strength and geometry is ordinarily expressed in terms of the mean-free matrix path (referred to subsequently as MFP) between the stronger particles.* This factor takes into account both the volume fraction of matrix phase and the average size of the particles. unke17 concluded that composite strength varied directly with the decreasing log MFP, following the Gen-samer8 relationship. However, Gurland and ~radzil~ reported a strength maximum at an MFP of the order of 1 p in WC-Co composites. A ductile two phase metallic composite was chosen for the present investigation of composite mechanical properties with the view of avoiding the difficulties and ambiguities noted above. The W-Ni-Fe heavy alloy of Green, Jones, and pitkin9 contains 80 to 94 wt pct W, and nickel and iron in a weight ratio

Jan 1, 1963

By J. D. Sullivan, Chester R. Austin, J. L. Nunes

IT is characteristic of most shales and surface clays that a bloated or vesicular structure is produced by burning to a sufficiently high temperature, usually about 150° to 200°F. above the normal maturing temperature used when making common brick of the same material. Early investigators recognized this phenomenon and attributed it to the release or expansion of gases in the vitrified piece. A number of theories have been advanced on the source of the gas and the manner in which it is evolved, but none has been proved conclusively. T. E. Jackson's1 early theory that bloating was due to oxygen evolved by ferric oxide in forming ferrous oxide or ferrous silicate was never substantiated by experimental work. Orton2 thought this explanation unlikely, as all clays should bloat at the same temperature, if this were true. He found that clays bloated in the range from 1100° to 1706°C. Also, many clays containing iron did not bloat at all, but fused. Orton also stated that oxidation of different types of carbon was a source of gas during burning of clay and should be expelled before vitrification was reached. Seger3 said that he had dissolved ferric oxide in glass with no apparent evolution of gas. Matson4 came to the conclusion that ferruginous days burned to vitrification in an oxidizing atmosphere did not have the iron in a ferrous form. According to Wilson,5 the gases responsible for bloating might be entrapped air, water, sulphur dioxide, carbon monoxide, oxygen, and hydrocarbons. He spoke of two causes of bloating: (I) incomplete oxidation below the vitrification temperature, and (2) decomposition of compounds and evolution of gases above the vitrification temperature. Decomposition of calcium sulphate was given as the principal example of the latter. Bleininger and Montgomery6 stated that if any time after vitrification began the evolution of gaseous matter was rapid enough so that pressure was produced within the day, bloating was certain to take place, and a vesicular structure would be formed. They recognized that initial structure, rate of heating, kiln atmosphere, and many other variables affected bloating. Work done by the Emergency Fleet Corporation7 indicated that clays suitable for bloating should be high in compounds of metallic oxides, carbon, sulphur, sodium, potassium, or other equivalent compounds, which either act as a flux or give off gas. Orton2 studied the chemical status of the carbon, iron, and sulphur in clays during various stages of burning, using a Huron shale that bloated under oxidizing or reducing conditions. He found that only a small amount of sulphur remained in the clay and that it could be detected only by dissolving the clay in hydrofluoric acid; i.e., the sulphur was in combination with the silica. Orton said that this sulphur was probably in the form of iron sulphide dissolved in the glass phase. The iron sulphide was not broken down by heating, but was dissolved in the basic slag or glass,

Jan 1, 1942

By J. W. Pamplin

SOME 40 miles east of Tampa is the center of the Florida pebble phosphate deposits. These are of Pliocene age and consist of several members of the Bone Valley formation.' Physically the phosphate-bearing formation, locally called the “matrix," varies in thickness from nothing to about 25 ft., usually averages 10 ft., and consists mainly of clay, quartz and phosphate pebble, the latter particles varying in size from colloidal to one inch diameter. The pebbles range in color from dark gray, through white, to cream, and in hardness from that which must be broken with a hammer to that which can be scratched with the fingernail. They assay from 66 to 80 per cent bone phosphate of lime (called "B.P.L." in the trade) or tricalcium phosphate (Ca3P2O8) and are secondary apatites. They form from 15 to 50 per cent by weight of the matrix. Over the matrix lies a varying thickness (5 to 40 ft.) of overburden composed of quartz and clay. Briefly, mining consists of two operations2 in open pits about 200 ft. wide by ½ mile or more in length. In the first step overburden is removed by large electric draglines at the rate of about 700 cu. yd. per hour. This is followed closely by hydraulic mining of the matrix at the rate of about 300 cu. yd. per hour. This matrix yardage is moved from an improvised sump in the floor of the pit to an ore-dressing plant, some- times a mile distant, by one, two or three "pit-cars" (pumping units) working on a 12-in. pipe line called the "rock line" to distinguish it from the high-pressure water line feeding the hydraulic giants, which is known as the “hydraulic line." This 700 yd. of overburden and 300 yd. of matrix-1000 yd. in all- amounts to about 1500 tons of material moved each hour. The over- burden is dumped into the mined-out pit alongside the pit being mined. The matrix moves to the ore-dressing plant at the rate of about 7 ½ long tons in 4000 gal. of water per minute-a ratio of approximately one to two by weight, or one to six by volume. This constitutes the "mill heads," but in the pebble district it is called the "matrix."

Jan 1, 1938

By Morris Cohen

STAFF: Editor, Gerhard Derge Acting Editor, Paul G. Shewmon Carnegie Institute of Technology Schenley Park Pittsburgh 13, Pa. Editorial Assistant, M. A. Redmerski Production Editor, Otto T. Johnson THE METALLURGICAL SOCIETY: Karl L. Fetters, President; J.S. Smart, Jr., Past-President; R. C. Cole, Vice President; R. L. Hennebach, Treasurer; R. W. Shearman, Secretary. PUBLICATIONS COMMITTEE OF THE METALLURGICAL SOCIETY: W. O. Phil-brook, Chairman; H. L. Bishop, R. C. Cole, J. N. Hobstetter, W. A. Krivsky, E. R. Morgan, F. C. Langenberg, R. E. Lund, Paul Queneau, L. L. Seigle, R. L. Smith, and T. V. Wainwright. Review and selection of manuscripts is under the supervision of the following Publicotions Committes: Extractive Metallurgy Division: R. E. Lund, Chairman; M. E. Wadsworth, Vice Chairman; D. H. Baker, Jr., G. M. Bell, R. E. Grace, W. A. Krivsky, R. A. Lewis, Charles O'Neill, A. W. Schlechten, R. E. Shinkoskey, and W. W. Stephens. Iran and Steel Division: T. V. Wainright, Choirman; H. L. Bishop, R. E. Boni, T. E. Dancy, D. W. Murphy, Johannes M. Uys, J. B. Wogstoff, and R. G. Ward. Institute of Metals Division: J. N. Hobstetter, Chairman; P. A. Beck, J. J. Becker, G. F. Bolling, J. 0. Brittain, J. B. Clark, R. S. Davis, A. E. Dwight, R. W. Fountain, L. R. Girifolco, J. C. Gurland, H. S. Kalish, E. L. Harmon, J. J. Heger, A. A. Hendrickson, L. Himmel, J. P. Hirth, H. H. Johnson, Jr., R. R. Nosh, T. N. Rhodin, R. H. Richman, H. C. Rogers, W. Rostoker, C. Sherby, C. W. Spencer, L. Slifkin, J. K. Stanley, R. Steinitz, H. C. Stumpf, and R. Swalin. Published bimonthly by the Americon Institute of Mining, Metallurgicol, and Petroleum Engineers, Inc., 345 East 47th Street, New York 17, N. Y. Telephone PLaza 2-6800. Subscription $30 peryear for non-AIME members in United States and North, South, and Centrol America; $32, foreign; $30per$7.50 for AlME members. Single copies, $6.00, single copies foreign, $6.50. The AlME is not responsible for any stotement made or opinion expressed in its publications ... Copyright 1962 by the American Institute of Mining, Metallurgical, and Petroleum Engineers, Inc. . . . Registered cable address, AlME New York . . . Indexed in Engineering Index, Industrial Arts Index, and Chemical Abstracts . . . Second class postage paid at New York, N. Y., and Ann Arbor, Mich. Volume 224, Number 4.

Jan 1, 1962

By F. A. Schaufelberger

METAL can be recovered from a leach solution either indirectly by precipitation as a compound that is later reduced or directly by electrolysis, cementation, or chemical reduction, for example, with hydrogen. The widely used electrowinning processes suffer from the inefficiency of converting fuel to low voltage dc and from low current efficiency when complete recovery is attempted. With the batch hydrogen reduction methods outlined here, a unit of fuel converted to hydrogen in modern gas reform plants will make two to three times as much metal as can be obtained with electrolysis. With continuous hydrogen reduction, four to six times as much metal is obtained. Precipitation by reducing a metal salt solution with H, has been known for almost 100 years.' Commercial use of the process, however, awaited the research and development program initiated by Chemical Construction Corp. Within the last few years this program has brought about construction of commercial plants, listed on page 545. Ideal hydrogen reduction will precipitate a pure metal from a solution obtained by commercial leaching methods at a rapid rate without excessive temperatures and pressures. The metal precipitate will be of desired size and density, less than a percent left in the vessel being deposited on the wetted parts. The present discussion will outline Chemical Construction Corp.'s early development program and will discuss the chemistry and mechanics of reducing copper, nickel, cobalt, and cadmium from solution by H2. Work on selective reduction of nickel from cobalt has been described earlier.' The article is not concerned with precipitation of copper by disproportionation of cuprous solutions where yield is limited3 to 50 pet.* It does not discuss use of gaseous reducing agents other than hydrogen, such as SO2,4 which may lead to contamination with sulfur, or C0,5 which is considerably more expensive than hydrogen and produces a gaseous reaction product, CO2.† The article does not include use of other nongaseous reducing agents,7 which are far more expensive than CO. Also, notes on the history, engineering design, and performance of commercial plants have been given elsewhere.' It will be shown that a pure metal can be precipitated by reduction with hydrogen from solutions obtained by commercial leaching methods when the solution composition is controlled and proper acidity, proper metal ion concentration through controlled complex formation and hydrolysis, and adequate agitation for suspension of metal and for transfer of reducing gas are maintained. Reasonable temperatures and pressures are used which are less excessive than those used by previous workers, although they are maintained at above equilibrium values because of process kinetics.' It has also been found necessary to induce nuclea-tion and to control growth and agglomeration in order to get a powder product of suitable physical properties. Review of the Literature Muller, Schlecht, and Schubardt of I. G. Farben10 claimed a successive reduction of silver, copper, nickel, cobalt, and zinc from ammoniacal solution by applying progressively higher temperatures and hydrogen pressures. Metals produced by this technique were not pure, since no attempt was made to adjust the solution composition between the different reductions. It is doubtful that the zinc product contained any metallic zinc. The major contribution to the literature" was made by the Ipatiews,12 who prepared platinum, iridium, copper, nickel, cobalt, lead, tin, arsenic, antimony, and bismuth at often extreme conditions, up to the critical temperature of water (373ºC), at 1500 to 7500 psi, for as long as several days. Even cadmium and zinc were claimed to have been observed in trace quantities precipitated together with basic salts. In the early experiments solutions were not stirred. Extensive formation of oxides and basic salts occurred, probably long before sufficient hydrogen could be supplied in the unstirred liquid to decrease the metal concentration in solution by reduction. The importance of hydrogen pressure was

Jan 1, 1957

By W. C. Thurber, C. J. McHargue

THE deformation textures of metals have been extensively studied because of both the practical implications in metal fabrication and the fundamental insight into the behavior of the metal during deformation. Most studies have been concerned with pure metals or solid-solution alloys and relatively little attention has been directed toward more complex alloy systems. Barrett' and Brick2 have summarized work on multiphase systems in terms of the relative ease of deformation of the phases. Thus, if the phases are deformed about equally, each develops its usual texture for the particular fabrication process. This behavior has been reported for Ag-Cu and Cd-Zn alloys3 and 62-38 brass.4 A dispersion of hard particles in a softer matrix interferes with the normal reorientation of the matrix lattice during deformation, thus distorting the texture or in some cases preventing its development. Increasing the amount of carbon in steel has been observed to increase the randomness of the ferrite5 and almost random textures have been reported for A1-12 pct Si eutectic alloy wires.6 It has also been noted that if the second phase is platelike, deformation causes the platelets to line up in a common direction. The present work concerns a quantitative study of the deformation textures in aluminum-uranium alloys containing 5 or 13 wt pct U. Alloys in this composition range are characterized by a pure aluminum matrix in which the intermetallic compound UA14, is dispersed. From a fundamental standpoint, this system affords the opportunity for studying the effects of a dispersed phase on texture development in an essentially pure metal matrix. Further, these alloys are commonly used as fuels for nuclear research and testing reactors, and a knowledge of the preferred orientation may be useful in evaluating physical and mechanical properties as well as response to fabrication. The deformation textures in sheets of aluminum (99.996 pct), Al-5 pct U alloy, and Al-13 pct U alloy were determined for reductions in thickness of 90 pct by cold rolling. The texture in an alloy rod of the latter composition which had been extruded at 455°C was also determined. It can be readily ascertained from the aluminum-uranium phase diagram that the 5 pct U alloy is hypoeutectic and contains 7.3 wt pct (3.4 vol. pct) UAl4. The 13 pct U alloy has the eutectic composition and contains 18.9 wt pct (9.4 vol. pct) UAl,; however, because of nonequilibrium solidification conditions, the typical eutectic microstructure was not developed. The alloys were prepared by open-air melting in graphite crucibles. The slab castings were cropped and cold reduced 90 pct in thickness on a two-high mill. The slabs were reversed after each pass. Although the uranium-bearing alloys exhibited edge cracking, they could be cold rolled to the desired thickness. Small coupons 1.5 by 0.75 by 0.1 in. thick were sheared from the center of the rolled plate with a major axis parallel to the rolling direction. To obtain a specimen of sufficient thickness for machining of the X-ray diffraction specimen, three coupons were laminated into a single three-ply sandwich with a suitable adhesive. Adhesives which could be cured at room temperatures were selected in order to prevent recrystallization during bonding. Bondmaster M-648 (Rubber and Asbestos Corp.) was used for the pure aluminum laminate and Armstrong A-6 (Armstrong Products Co.) was used for the alloys. The alloy for extrusion was cast in a cylindrical graphite mold! cropped and machined into a 3-in. diam by 4-in. long billet. This billet was preheated to 455°C, upset to 3.125 in. diam, and extruded on a 700-ton hydraulic press at a speed of 6 fpm through a flat-face, 1-in. diam die. The rod was quenched with a water spray as it emerged from the die. Pole-distribution data were obtained using the diffractometric method described by Jetter and Borie.7 A spherical X-ray diffraction specimen of

Jan 1, 1961

By R. Ruka, E. A. Gulbransen

ONE of the interesting questions in the understanding of the reaction of iron with oxygen is the kinetics and the mechanism of the crystal structure changes occurring in the formation and breakdown of the oxide film. In a recent work' the electron diffraction method was applied to the study of the solid phase reactions occurring internally in the oxide film on iron above 570°C and under vacuum conditions. The higher oxides which form under oxidizing conditions are reduced internally by the diffusion of iron into the oxide leaving the bulk of the oxide FeO (wüstite). This paper will ask and consider three questions. These are: (1) What is the nature of the depressed transformation temperature for the Fe3O4 + Fe ? 4Fe0 reaction in very thin oxide films?; (2) What is the mechanism of the forward transformation of Fe3O4 to FeO at or near 570°C?; and (3) What is the nature and mechanism of the slow reverse transformation of FeO to Fe3O4 and Fe below 570°C?. The composition of the oxide as well as its electrical, magnetic, mechanical, and chemical properties depends upon the kinetics of these reactions. Literature Survey The composition of the oxide film on pure iron at temperatures up to 570°C has been studied extensively by the electron diffraction method. Nelson,' Jackson and Quarrell,3 and Gulbransen and Hick-man4 have shown that Fe3O4 and a-Fe2O8 are formed in the temperature range of 225" to 570°C, while y-Fe,O, is most likely formed below 225°C for long oxidation times. In very thin films FeO (wüstite) is found to form at temperatures as low as 400°C, and its existence depends essentially upon the extent of initial oxidation. FeO is not found in the formation of thick films below 570°C. On the other hand X-ray5 and micrographic8 studies show that Fe3O4 may exist as the main com-ponent in thick films up to 650°C, well above the equilibrium temperature for the reaction Fe3O4 + Fe ? 4FeO. The kinetics of the forward and reverse trans-formations of the reaction Fe3O4 + Fe ? 4FeO have not been studied in oxide films. However, analysis of the data of Gulbransen and Hickman4 on the heat-ing and cooling of iron oxide films of known thick-nesses in high vacua shows that the forward reac-tion proceeds rapidly even for a 4000 A film at temperatures of 575" to 600°C. The reverse reaction or decomposition of wüstite (FeO) occurs at tem-peratures below 450°C. Thus, there appears to be a slow process in the decomposition which is in marked contrast to the rapid forward reaction. The decomposition of a bulk sample of FeO free from the metal has been studied extensively by Chaudron, Benard and coworkers using the micro-graphic, X-ray diffraction, and thermomagnetic methods. Chaudron7,8 first observed that the de-composition of FeO (wüstite) below 570°C occurs by the reaction 4FeO ? Fe3O4 + Fe. Later work by Chaudron and Forestier5 showed that the rate of de-

Jan 1, 1951

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 H. C. George

IN 1901, the United States produced 69 million barrels of crude oil, which was 41.4 per cent of the world production. By 1931, these figures were 850 million barrels and 62.1 per cent respectively, showing over this period of thirty years an increase in the world production of 718 per cent. Over this same period, the increase in the demand for technically trained men in the industry was much greater than the increase in crude-oil production. This was because much of the increase in production was secured by the application of geology and physics to the finding of new oil fields and by .the development of engineering methods and equipment incident to drilling and producing from much greater .depths.

Jan 1, 1934

By Everett G. Trostel

AMERICAN industry in 1943 emerged from the construction phase into the production phase, and American military operations passed from preparation into full action in the many theaters of the global war. The increased pace of industrial and military operations has resulted in a tremendous demand for petroleum and its products.

Jan 1, 1944

By C. F. Austin, J. K. Pringle

In view of the many large differences in structure and physical properties between metals, plastics, and most rock materials, a continuing experimental petrodynamics program has been established at the U. S. Naval Ordnance Test Station. Part of this program has been to test the concepts of scabbing and corner fracturing in a typical coarse-grained rock. The results of this experimental program have shown very interesting differences between the formation of fractures in rock and the formation of fractures in other materials such as metals and plastics which are often used as models for rock behavior studies. The gross behavior of coarse-grained brittle solids subjected to impulsive loading has undergone abundant empirical investigation by mining organizations and by mineral research agencies. However, the fundamental response of coarse-grained brittle solids to impulsive loading has received little detailed study in recent years. The various organizations and agencies interested in rock-like brittle materials have broadly applied the stress wave interaction concepts of the impulsive failure of metallic samples to the behavior of granular brittle materials 1,2,3,4,5,6 Using the metallic sample approach for stress wave induced fracturing, an understanding of the fracture behavior of brittle materials which perhaps is not always appropriate has been established in the open literature. An excellent summary of the metallic sample approach is provided by ref. 7 which summarizes the general beliefs prevalent at present regarding the impulsive behavior of brittle solids: "In the blasting process, rock breakage occurs primarily from tensile slabbing, beginning at a free face and proceeding back toward the explosive charge. The tensile waves formed by the reflection of incident compressive waves from two or more free faces will collide in the rock, resulting in large localized tensile stresses. Fractures formed by the interaction of these tensile stresses begin at the apex of the free face planes and progress back into the rock as shown...."

Jan 1, 1964

By G. L. F. Powell, L. M. Hogan

Samples of silver and Ag-O alloy, 0.12 wl pet, have been undercooled to a maximum of 250°C by melting in a slag of commercial soda-lime glass. Grain refinement occurred in undercooled silver samples whet the degree of undercooling exceeded a critical value which varied from 153° to 175°C, depending on the melting conditions. In under cooled Ag-0 samples, the grain structure was fine and equiaxed at degrees of undercooling larger than 50oC. The grain refinement in both silver and Ag-0 alloy samples was the result of recrystallization during or immediately following freezing. When the oxygen content was increased, the degree of undercooling at which recrys-tallization occurred decreased. INVESTIGATIONS of the grain structure of strongly undercooled melts of nickel and cobalt1-4 have revealed a change in room-temperature grain size as a function of undercooling. At large degrees of undercooling the grain size was fine, but samples undercooled by smaller amounts exhibited a much coarser grain structure. The reported undercooling at which this grain size transition occurred in nickel varied between 140° and 175'C undercooling. Walker1 attributed this grain size change to an enhanced nuclea-tion rate. Cavitation at a melt-solid interface was con sidered to generate a high-pressure wave. This, in turn, produced nucleation in the liquid as a result of an elevated freezing point increasing the undercooling in the melt. A similar grain size transition was observed in undercooled silver between 133° and 153°C undercooling.5 The grain size transition in undercooled silver was found to be dependent on oxygen con tent. More recently, the authors5 showed that a simila transition occurred in an undercooled Cu-O alloy, =0.08 pet, but not in oxygen-free copper undercooled more than 200°C. The transition was attributed to recrystallization of the solidified grains after freezing; the greater the degree of undercooling, the more complete the recrystallization. In view of this effect of oxygen content on grain size transition, additional undercooling experiments were carried out with silver. This paper reports the results of these studies on silver and Ag-O alloys undercooled more than 50°C EXPERIMENTAL 1) Materials. The silver was supplied by Matthey Garrett Pty., Ltd., as fine granulate, silver mini' mum 99.9 pet, and contained the trace impurities shown in Table I. The silver originated from two different refineries resulting in the difference in impurity content shown in Table I. However, the variation in composition had no discernible effect on the undercooling behavior. Samples of the silver were melted in an open-ended, vertical, cylindrical resistance furnace, wound with Kantha] wire, the temperature of which was adjusted by means of a variable transformer, 2) Melting and Undercooling. Bulk samples of silver were undercooled by preoxidation and removal of oxidized impurities in a glass slag, as previously described,5 but special precautions were necessary to allow comparison between oxygen-free silver and a Ag-O alloy. Samples free of oxygen were prepared by first melting weights of granulate varying from 200 to 400 g in clean vitreous silica crucibles in air for approximately 15 min. A sample was frozen and a complete glass cover formed over the top surface by adding granulated glass rod. The sample was then melted and frozen several times until no oxygen bubbles appeared in the glass slag. Subsequent thermal analysis indicated that the oxygen content was negligible. The glass acted as an efficient barrier to absorption of oxygen from the air. Ag-O alloys were prepared by melting silver in silica crucibles with only a partial cover of glass, so that the top surface of the melt was in contact with air. The silver was held molten in air for several hours at 1000° to 1050°C, in the expectation that the oxygen in solution would reach equilibrium with atmospheric oxygen. At T - 1000°C the equilibrium oxygen content was calculated as 0.12 wt pet O, using Sieverts law and Eq, [l], derived by Mizikar et a!.7 from the data of Sieverts and Hagenacker8 for 1 atm of oxygen:

Jan 1, 1970

By Claudio Barsotti, Leonard C. Kitchener

INTRODUCTION The major mining operations of the Ontario Division of Inco Metals Company are located in the Sudbury Basin. Around the rim of this geo- logical formation and within thirty miles of the City of Sudbury are located eleven under- ground mines of the Company, two of which are on standby. In addition, there is one mine in Northwestern Ontario at Shebandowan, 1127 Km (700 miles) northwest of Sudbury. In 1980 ore production totalled 10.6 x 10 mt (11.7 million tons) from underground and open pit operations. Underground mining methods include blasthole, vertical retreat, sublevel caving, cut-and-fill and undercut-and-fill. Inco reduction operations in the Sudbury area include two concentrators, one smelter, one iron ore recovery plant, a copper refinery a nickel refinery and a high purity plate rolling mill. As well, an electrolytic nickel refinery is located on Lake Erie at Port Colborne, near Niagara Falls. In the Sudbury District, Inco employees total 14,278 of whom 6,490 are directly employed in the mining and milling operations. I T H DRILLING Background - The use of long blastholes in underground operations at Inco mines in the Sudbury district started more than 30 years ago at the Levack Mine. At that time, diamond drills were used to mine ore by open stoping and for pillar recovery. By the 1950's the concept had been extended to all mines. The heavy percussion drill was also developed by this time and led to the acceptance of the 11.43 cm (4% in.) piston diameter drifter for bar and arm and fan drill applications. The use however, of the nominal 5.08 cm (2 in.) diameter blastholes drilled radially from the level or sublevel causes overdrilling at the periphery. As well, the smaller holes are subject to hole deviation, are difficult to drill beyond 38.1 m (125 ft.) and require the need for several drilling horizons or sublevels. The need to improve blasthole stoping techniques and our experience with large diameter blastholes in open pit mining led to the trial of an I T H drill in 1962. The test was unsuccessful however, since button bits were not available at the time and the tungsten carbide cross bit experienced short life, gauge loss, and caused hole deviation and poor penetration rates. The successful application of the raise borer beginning in 1964 and the ability to drill 20.32 cm (8 in.) to 27.94 cm (11 in.) diameter pilot holes led to the development of smaller, more portable, rotary drills for blasthole mining. As a result, a mining method based on the use of 20 cm (7 718 in.) diameter, 60.1 m (200 ft.) long rotary blastholes was developed and tested successfully at our C.C.North Mine. Between 1972, when the method was first tested, and 1973 over 4572 m (15000 ft.) of large holes were drilled with rotary drills in the recovery of 175,960 mt (194000 tons) of primary stoping and pillar remnants. A new mining method was developed that made use of this large hole drilling equipment. Instead of longhole carbide drilling by bar-and-arm or fan drill set-ups at sublevels, the large diameter holes were drilled full stope height from the top sill, which substantially reduced development, drilling time and improved fragmentation. These drills however, were difficult to move and set-up, and had a high maintenance cost and low productivity. The capital outlay and extensive technical development required to improve these drills and reduce the high drilling costs however, was economically unattractive. New development in drills and button bits INTRODUCTION The major mining operations of the Ontario Division of Inco Metals Company are located in the Sudbury Basin. Around the rim of this geo- logical formation and within thirty miles of the City of Sudbury are located eleven under- ground mines of the Company, two of which are on standby. In addition, there is one mine in Northwestern Ontario at Shebandowan, 1127 Km (700 miles) northwest of Sudbury. In 1980 ore production totalled 10.6 x 10 mt (11.7 million tons) from underground and open pit operations. Underground mining methods include blasthole, vertical retreat, sublevel caving, cut-and-fill and undercut-and-fill. Inco reduction operations in the Sudbury area include two concentrators, one smelter, one iron ore recovery plant, a copper refinery a nickel refinery and a high purity plate rolling mill. As well, an electrolytic nickel refinery is located on Lake Erie at Port Colborne, near Niagara Falls. In the Sudbury District, Inco employees total 14,278 of whom 6,490 are directly employed in the mining and milling operations. I T H DRILLING Background - The use of long blastholes in underground operations at Inco mines in the Sudbury district started more than 30 years ago at the Levack Mine. At that time, diamond drills were used to mine ore by open stoping and for pillar recovery. By the 1950's the concept had been extended to all mines. The heavy percussion drill was also developed by this time and led to the acceptance of the 11.43 cm (4% in.) piston diameter drifter for bar and arm and fan drill applications. The use however, of the nominal 5.08 cm (2 in.) diameter blastholes drilled radially from the level or sublevel causes overdrilling at the periphery. As well, the smaller holes are subject to hole deviation, are difficult to drill beyond 38.1 m (125 ft.) and require the need for several drilling horizons or sublevels. The need to improve blasthole stoping techniques and our experience with large diameter blastholes in open pit mining led to the trial of an I T H drill in 1962. The test was unsuccessful however, since button bits were not available at the time and the tungsten carbide cross bit experienced short life, gauge loss, and caused hole deviation and poor penetration rates. The successful application of the raise borer beginning in 1964 and the ability to drill 20.32 cm (8 in.) to 27.94 cm (11 in.) diameter pilot holes led to the development of smaller, more portable, rotary drills for blasthole mining. As a result, a mining method based on the use of 20 cm (7 718 in.) diameter, 60.1 m (200 ft.) long rotary blastholes was developed and tested successfully at our C.C.North Mine. Between 1972, when the method was first tested, and 1973 over 4572 m (15000 ft.) of large holes were drilled with rotary drills in the recovery of 175,960 mt (194000 tons) of primary stoping and pillar remnants. A new mining method was developed that made use of this large hole drilling equipment. Instead of longhole carbide drilling by bar-and-arm or fan drill set-ups at sublevels, the large diameter holes were drilled full stope height from the top sill, which substantially reduced development, drilling time and improved fragmentation. These drills however, were difficult to move and set-up, and had a high maintenance cost and low productivity. The capital outlay and extensive technical development required to improve these drills and reduce the high drilling costs however, was economically unattractive. New development in drills and button bits

Jan 1, 1981

By W. S. Owen, T. Bell

The variation ox the M, temperature with nitrogen concentration has been determined experinzentally. The free-energy difference between martensite and the parent y Phase at the M, temperature,is computed for Fe-C and Fe-N alloys using recent solubility data fm the Fe-N system and new data for the variatim of the activity of the interstitial solute in the y phase with concentration. The "geometric model" is considered also, and the importance of Zener ordering is noted. It is shown that, although at any chosen concentration the M, temperature and the activity are significantly different in the two alloy systems, the free-energy duference is the same within the limits of experimental error, thus Providing further evidence of the very close similarity between the martensite transformations in Fe-C and Fe-N alloys. TO estimate the free-energy change accompanying the martensite transformation in binary interstitial alloys it is necessary to know the variation of the Ms temperature with interstitial solute concentration and some thermodynamic data. To date, only the Fe-C system has been discussed because this was the only system for which relevant experimental measurements had been made, but even in this case the data available are insufficient to permit the free-energy change to be calculated without the assumption of a thermodynamic model for the y and a solid solutions. The last comprehensive discussion of the topic was by Kaufman, Radcliffe, and Cohen in 1.960.' Since that time Ellis, Davidson, and Bodsworth2 have redeter mined the activity of carbon in y iron as a function of carbon content and Atkinson3 has measured the activity of nitrogen in Fe-N austenite. The M, temperature of Fe-N is reported in the present paper. Thus, it is now possible to assess the free-energy change accompanying the martensite transformation in Fe-N alloys on the same basis as that used to assess the change in Fe-C alloys. This comparison is of interest because it has been shown recently4 that the structure of the martensite formed in Fe-N alloys changes with nitrogen content in a manner exactly parallel to the changes in structure in the Fe-C series and that the strain associated with the distortion of the iron lattice in either the y or the a phase by a nitrogen atom is the same as that produced by a carbon atom. The work to be described tests whether or not this close similarity between the two systems extends to the thermodynamics of the martensitic transformations. THE Ms TEMPERATURES FOR Fe-N MARTENSITE Bose and Hawkes5 found the M, temperature of a eutectoid Fe-N alloy (2.35 wt pct N) to be 35°C. This single determination appears to be the only reported measurement of the M, temperature of an Fe-N alloy. In the present work a series of more than fifteen different specimens containing between 0 and 2.7 wt pct N were prepared by controlled nitriding of pure iron wire 0.010 in. diam. The iron was supplied by B.I.S.R.A. and was designated ACN2. The nitriding procedure, all aspects of which are critical, has been described in detail elsewhere.' Each specimen was quenched from the austenitizing temperature into brine at 20°C without removing the specimen from the ammonia-hydrogen atmosphere in which it was ni-trided. Alloys containing up to 0.7 wt pct N were fully martensitic. More concentrated alloys contained progressively more retained austenite. At 2.34 wt pct N the specimen was completely austenitic after quenching to 20°C. The M, temperature of specimens with M, temperature above 20°C (those containing less than 2.35 wt pct N) was determined by a thermal arrest technique. For the more concentrated alloys (2.35 to 2.70 wt pct N) the variation of electrical resistance with temperature down to -196°C was used. Unlike Fe-C alloys, Fe-N alloys cannot be reaus-tenitized in vacuum or in an inert atmosphere because excessive loss of nitrogen occurs, even in short times, at high temperatures. Consequently, the alloys were reaustenitized in the ammonia-hydrogen mixture corresponding to nitrogen equilibrium at the temperatures used. The specimens were quenched in a blast of cold hydrogen gas and the temperature, measured by a thermocouple welded to the 2.5-in.-long specimen, was recorded on a high-speed recorder. In the first experiments 0.002-in.-diam chromel-alumel wire thermocouples were used but it was found that these gave erroneous readings due to contamination by nitrogen pickup. Pt-Pt 13 pct Rh wires of 0.002 in. diam were found to be satisfactory. Further details of the experiments are available in the original thesis.7 The cooling rate was controlled between 400° and 4000°C per sec by adjusting the hydrogen pressure. At least three cooling curves were obtained for each composition. Each specimen was analyzed for nitrogen after quenching. To aid in the accurate location of the martensite arrest points a series of cooling curves of 0.010-in.-diam platinum wire were recorded and these, which did not contain an arrest, were used as comparison standards. Four specimens, each in the form of a 3-in. loop, were prepared from each of the four alloys containing more than 2.35 wt pct N. One specimen was used for nitrogen analysis, one for X-ray diffraction, and two for resistance measurements below 20°C. The latter specimens were quenched from 20°C into constant-temperature baths at selected lower temperatures.

Jan 1, 1968

By C. G. Dunn

Many investigations have been made concerning the nature of plastic deformation and recrystallization of metals either in the form of polycrystalline materials or in the form of single crystals. However, similar work on strip materials composed of large grains as thick as the strip and having a high degree of preferred crystal oricntation has not been carried out extensively. Such materials are ideal not only for studying changes in single crystals under conditions where neighboring grains exert their influences but also for obtaining the complete change in orientation of a grain or group of grains during the course of plastic deformation and subsequent recrystallization. It is well known that first-order twins in the form of Neumann bands may be produced during the plastic deformation of ferrite crystals, particularly silicon ferrite crystals.l--3 Furthermore, it is believed that high-stress regions exist in the vicinity of these bands4. and it has been said that these banded regions would be active in recrystallization processes.5 Evidence to that effect in silicon ferrite had already been presented with photomicrographs showing recrystallization among Neumann bands during the early stages of recrystallizati~n.5,6 Similar evidence has been published for nonferrous metals5 Merging of fine parallel deformation bands during annealing to form twin bands of ordinary appearance in alpha brass has been offered not only as evidence of an active role being played by mechanical twins but also as evidence that such twins cxisted prior to recrystalli~ation.5 Proof that such new structures are actually first-order twins, or even twins at all, must be considered inadequate except for those cases where X-ray or optical data have been obtained to confirm the necessary oricntation relationship.7,8 A. B. Greninger has obtained excellent iliustrations of annealing twins in alpha iron.9 Since his material was previously deformed by sharp hammering, a process that generally induces Neumann bands, the existence of twins following annealing might suggest that they grew from mechanical twins. In connection with transformation twins in alpha iron, Greninger reported that grains containing three twin orientations occurred rarely and none occurred with three sets of well-developed bands.9 The work at the Pittsfield Works Laboratory on silicon ferrite in the form of large-grained strip steel deformed a small amount and annealed to partially recrystallize the grains indicates that the deformation and recrystallization processes that lead to annealing twins involvc more complex processes than the formation of Neumann bands and the growth of such bands. Some groups of annealing twins are obtained which contain not only three well-developed twin structures, but also contain more than one order of twins. It is also shown, in one case, that an interrelated group of twins did not arise within one single parent grain and, furthermore, orientations of grains in the group

Jan 1, 1944

By G. V. Smith, C. O. Tarr, R. F. Miller

Metallurgists have long recognized that the Fe3C type of carbide is not a stable phase in steel and that, given sufficient time, it will decompose with formation of graphite, at least at temperatures below about 205o°F.l This slow graphitization has never been of much concern to users of hypoeutectoid steels, for in such steels it occurs only during prolonged exposure at elevated temperature. In I935, Kinzel and Moore2 reported complete graphitization of the carbide of a 0. I5 per cent carbon steel that had been in service for about 26,000 hr. at a temperature of about IIOO° to 12oo°F. In discussing this paper, both Mathews and Sauveur reported graphitization of slightly hypereutectoid steels. In an extensive investigation of graphitization of high-purity iron-carbon alloys, Wells1 produced graphite at 7oo°C. (I29o°F.) in an alloy containing as little as 0.13 per cent carbon, and showed that graphite may precipitate either directly from austenite or from decomposition of Fe3C, also that the rate of graphitization tends to increase with increase of temperature, at least up to a certain point, and with the presence of graphite nuclei. Jenkins, Mellor and Jenkinson,3 studying the structural changes in carbon steels ranging from 0.17 to 1.14 per cent carbon during stress-rupture tests at elevated temperature, observed graphite in an aluminum-killed 0.40 per cent carbon steel after fracture in 4340 hr. under a stress of 17,900 lb. per sq. in. at 840°F. and in many of the steels after stress-rupture test at 12oo°F. Unstressed samples of 0.60, 0.86 and I.I4 per cent carbon steels graphitized almost completely in 360 hr. at I200°F., while samples of 0.90 and 1.10 per cent carbon steels were somewhat graphitized, and steels containing 0.24, 0.40 and 0.57 per cent carbon showed no sign of graphite. Their results indicate that graphite precipitates the more readily the higher the carbon content, though this tendency is influenced to some extent by deoxidation practice, and that graphitization is accelerated by plastic deformation. In a recent study of the graphitization of 0.4 and 0.6 per cent carbon steel strip, made by M. A. Hughes,4 it was found that at I2oo°F. graphite appeared after 72 hr. in aluminum-killed but not in silicon-killed steel, and that the rate of graphitization was increased by slow cooling from I6oo°F. and by cold-working prior to annealing. In creep tests of normalized 0.15 per cent carbon steel killed with 2.5 Ib. of aluminum per ton, at the U. S. Steel Corporation laboratory at Kearny, spheroidization and some graphitization were found after 3000 hr. at Iooo°F. under a stress as low as 1500 lb. per sq. inch. .

Jan 1, 1944