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By R. D. Satterley

The Sherwood mine is an iron-ore mine owned and operated by the Inland Steel CO. in the Iron River district of the Menominee Range in Michigan. The property consists of an 80-acre tract in the village of Mineral Hills, about a mile north of the city of Iron River. The Ore mined is a reddish brown, nonbessemer hematite and limonite. The ore bodies lie in a large syncline having an east-west strike and are associated with a rock structure that is folded in a complex manner. The two 40-acre tracts run north and south and the shaft is at about the center of the north forty. The shaft is 1286 ft. deep, with the principal operating level at a depth of 1000 ft. from surface. A new level at a depth of 1200 ft. is in process of development. The ore body at the 1000 ft. level is east and west at approximately the line between the north and south forties and at this elevation is from 50 to 250 ft. wide and extends above the level at varying heights of from 50 to 250 feet. The 1000-ft. haulage drift extends south from the shaft, through the east-west ore body to the slates on the south side of the body. The ore body was intercepted in the drifting at about goo ft. from the shaft and was about 200 ft. wide at that point. A crosscut was extended to the east and west property lines along the south boundary of the ore body. This east-west ore body is laid out in areas go ft. wide for stoping with pillars 40 ft. wide. With this plan a stope is approximately 80 by 200 ft. and has a pillar on both sides that is approximately 40 by 200 ft. The ore has a height of from 50 to 250 ft. above the elevation of the level, therefore the height of the stopes varies in accordance with the height of the ore body at each location. From the crosscut, two main raises, each 6 by 7 ft., are put up into each 80 ft. stope area (Fig. 1). One of the main raises serves as a manway and the other as an ore raise, through which the ore is transferred from the stope to the tramming equipment. The original layout as for scraping of the ore from under the stopes into the ore raises. The ore raise is put halfway (east and west) along the 80-ft. stope area and usually is in iron formation. The manway and ore raise are on approximately 15-ft. centers. These main raises are Put up to a height of from 22 to 30 ft. above the level floor, then from the toll of the ore raise a scraper drift (8 X 8 ft.) is driven across the width of the ore body, which is in the neighborhood of 200 ft. From the scraper drift, mill raises are put up at 12½-ft. centers, staggered on opposite sides of the drift, and these are extended 23 ft. above the elevation of the scraper-drift floor. The sublevel connecting the top of these mill raises is called a mill "sub." Other sublevels are developed parallel to the scraper and mill subs at 25-ft. vertical intervals above the mill sub. The mill sub and the other subs above the mill sub are known as sublevels or dog drifts, and are 4 ft. wide and 6 ft. high. These dog drifts

Jan 1, 1946

By M. Bevis, E. O. Fearon, P. C. Rowlands

Fe-Ni and Fe-Ni-C martensite specimens have been deformed in compression at room temperature and the habit planes of operative deformation twins determined by two-surface optical trace analysis. The full orientations of the martensite crystals were determined from divergent X-ray beam diffraction patterns. The experimental results are in excelled agreement with predicted twinning modes. In particular, the habit planes of some deformation twins in bcc martensites are consistent with a "Type XI compound'' twinning mode with Kl, Kz, 71, 77~ elements given by {5, 8, 11) {ioi}, (33) (ill) Tetragonal derivatives of this mode are operative in bct martensites. UnUSUAL deformation twinning modes have been reported by Richman' to occur in Fe-Ni-C martensites with bcc and bct structures. The twin habit planes were determined by single-surface trace analysis (pole locus method) and the remaining twinning elements were determined from the geometry of twin-twin intersections. The indices assigned to the observed habit planes are (3 101, "(089)" and "{0, 1,13)" or "{1,2,7} ", and only in the first case do the twinning elements correspond to a predicted twinning mode. This is mode 1.4 in the paper by Bevis et al. The results presented in Ref. 2 indicate that previously unpublished bcc and bct modes should be operative in preference to the (130) mode and anomalous modes reported by Richman. In view of these results and the uncertainties involved in determining habit planes from single-surface trace analysis and twinning elements from twin-twin intersections5 a two-surface trace analysis of deformation twins in Fe-Ni and Fe-Ni-C martensites has been carried out. EXPERIMENTAL PROCEDURES Two alloys with compositions Fe-23 pct Ni-0.6 pct C and Fe-30.4 pct Ni were prepared from 4N pure materials by induction melting under a vacuum of 10"5 mm Hg. The alloys were homogenized at 1350'~ for 5 days. Both of these alloys are austenitic at room temperature with M, temperature - — 50"C. The aus-tenite grain size of the Fe-Ni-C alloy was approximately 300 to 400 p. The Fe-Ni alloy was remelted in a vertical tube furnace and the melt lowered slowly from the hot zone of the furnace to produce s ingle -crystal austenite specimens. Specimens approximately 10 by 5 by 5 mm were cut from the ingots and quenched to various temperatures below the Ms temperature. The specimens were elec-tropolished in a 10 pct perchloric acetic electrolyte before being deformed in compression at room temperature. Martensite plates which exhibited profuse deformation twinning were selected for analysis and the specimen polished on a second surface such that the two surfaces which contained the martensite plate enclosed an obtuse angle of approximately 145 deg. The specimens were then electropolished to reveal the traces of the deformation twins on both surfaces. The full orientations of the martensite crystals were determined using a divergent X-ray beam technique (Kossel line technique) employing an AEI SEM2 electron probe microanalyzer. Details of this technique which include a detailed description of the Kossel camera attachment to the microanalyzer used in the present experiments have been discussed elsewhere.3 Only additional details relevant to this investigation are discussed here. The specimens were mounted with one surface normal to the incident electron beam as illustrated schematically in Fig. 1. The martensite plates to be analyzed were located using the normal scanning equipment of the microanalyzer. The position of the electron beam and hence the position of the source of divergent X-rays generated within the crystal could be located to within 1 p. Exposure times of approximately 8 to 10 min were required for back-reflection Kossel patterns and it was found that useful diffraction patterns could be obtained consistently from heavily deformed martensite plates. A reference line (carbon contamination mark) produced on the two surfaces of the specimen by scanning the electron beam in a direction having a known relationship with the reference !ine in the X-ray camera enabled the full orientation of the crystals to be determined. Martensite plates with widths as small as 8 p could be oriented using this procedure. The Kossel line patterns were interpreted using the charts developed by Rowlands and ~evis~ as generally

Jan 1, 1969

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

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

Jan 1, 1948

By Hubert D. Keiser

An observation that is finding fresh confirmation in the leadership of the Allied Nations is "When the need arises, a man will rise to fill it." Much the same can be said of industrial advances in the current emergency. Events and circumstances have created needs for new materials and stimulated the inventive genius of man to fill them. The timely development of fibrous-glass manufacture, in relation to present war needs, is a case in point. Historical Summary It is probable that the first glass fibers produced by the hand of man were made about 4000 years ago. Successful production of glass fibers for manifold industrial and military uses, on any scale that could begin to meet present needs, has a history spanning a scant four years. The American processes that brought this young industry into being were developed between the years 1931 and 1937; research and manufacturing efforts were consolidated to create a commercially satisfactory operation as late as November 1939. In the 40 months that have elapsed since then, the industry has achieved commercial importance, supplied immense quantities of materials for pre-war industrial expansion to meet "defense" and Lend-Lease needs, and now has reached such importance that it has come largely under the control of the War Production Board for equitable distribution of its major products among 13 claimant agencies. It is worthy of note, however, that the Bureau of Ships, Navy Department, began exploratory use of fibrous-glass products in 1932, within one year after the date of the first experimental trials of the new process. This early recognition of the military importance of fibrous glass gave impetus to further research work. The foregoing rapid scanning of historical development must be filled in, lest the. impression be created that glass fibers were suddenly created out of thin air, with no antecedent trials, failures, and partial successes. The actual history of glassmaking is lost in antiquity. Probably some nomadic tribesman found molten glass beneath the bed of a hot campfire built upon sand of the proper composition. If his curiosity led him to poke the glowing, viscous pool with a green-wood stick or some handy weapon, it is probable that he drew a glass fiber on the end of his implement. Molten glass has much the same consistency as honey or molasses; it strings out tenuously. Thus, it is quite possible that the first form of glass shaped by man's action was a glass fiber. Artisans in glass, throughout the long Direc- of glassmaking as an art> drew fibers from fine rods for many decorative applications, including "spun glass" decoration in beautiful stemware and the charming millefiori (glass flowers encased in solid glass), and even to create the blood vessels in glass eyes to match a bleary mate.

Jan 1, 1948

By Jack White, Ralph B. Adair

THE Mufulira mine in Northern Rhodesia is 13° south of the Equator and at an altitude of 4100 ft above sea level. The concentrator was planned in 1930 to treat about 10,000 tons of ore per day, but before construction work had gone very far the size of the plant was reduced to 1500 tons daily capacity. Then just as the construction was finished in December 1931, the price of copper did not warrant the commencement of operations and it was not until October 1933 that the mine re-opened. Extensions have been made to the plant in 1034, 1935, 1936 1937, 1940 and finally in 1945. The capacity of the plant when the 1945 extension is completed will be 9500 short tons per day. The various extensions have been carried out with the minimum confusion and with great skill on the part of the designers and the erection engineers. In dealing with the plant in this paper it is proposed to give an outline of the present practice and only to refer to past practice when it is considered of particular interest. ORE TREATED The ore consists of fine grained highly siliceous, felspathic and argillaceous quartzites, the latter containing graphitic carbon and locally termed graywacke. These rocks are considerably altered near the surface and on the fringes of the ore body. Despite extensive efforts at mixing, both underground and in the crushing plant, many sudden changes of ore take place in the mill feed. The minerals are chalcocite, bornite and chalcopyrite with relatively unimportant amounts of malachite, native copper, azurite, covellite and chrysocolla. Gold is present at about 1 oz to a thousand tons, and silver at about 2 dwt. per ton. With increasing depth the nonsulphide content or, as commonly termed locally, the oxide content of the ore is decreasing and is now less than 0.25 pct against a high of 0.77 pct. CRUSHING OPERATIONS The first stage of crushing is done underground where the oversize from 12-in. grizzleys is fed to a 56 X 72-in. jaw crusher. At the Selkirk shaft two 10-ton skips deliver the run of mine ore at a rate of 450 tons per hour to the 400-ton capacity shaft bins. One 6 ft wide Ross chain feeder and one locally made air-operated finger feeder feed the two 30-in. McCully crushers. Ahead of each crusher is a 6 X 9-ft spring grizzley spaced at 5 in. on a 35° slope. Neither the Ross chain feeder nor the finger feeder is entirely satisfactory and for the wet, muddy and sticky ore that has to be handled it is planned to replace these with two 6-ft wide apron pan feeders. The McCully crushers deliver a minus 5-in. product, at a rate of over 500 tons per

Jan 1, 1947

By G. A. Parks, C. Kovacs

The xanthate flotation of chrysocolla can be improved by preheating the ore to 500° to 600°C. Addition of water vapor reduces the preheating temperature required. The temperature range of flotation improvement correlates well with the range over which dehydroxylation occurs. This observation is consistent with changes in the structure and properties of the mineral as detected by measurement of solubility, D.T.A., T.G.A., X-ray diffraction, and infrared absorption and with recently proposed explanations of chrysocolla flotation. A brief experiment by one of the authors several years ago showed that exposure of chrysocolla to high temperature water vapor resulted in increased recovery by subsequent xanthate flotation. Recently renewed interest in flotation and thermal properties of chrysocolla1-5 prompted further study of this effect and more detailed reporting.6 This paper explores the response of pure chrysocolla recovery in Hallimond tube flotation tests to the time and temperature of exposure and to changes in the water vapor pressure. Brief attention is paid to the effects of such treatment on accessory minerals. A single experiment with flotation of chrysocolla from an ore is reported. MATERIALS Chysocolla from Inspiration, Gila County, Ariz. was crushed in a porcelain mortar, cleaned of visible and magnetic impurities, and screened to obtain -48 +65 mesh (Tyler) material for flotation test work. The clean material contained 28.5% Cu and was identified as chrysocolla by comparison of X-ray diffraction, D.T.A., and infrared absorption data with that reported for similar materials by others.4,5,7 Samples of malachite and chalcopyrite, prepared in a similar way, were also identified by X-ray diffraction. The sources of these materials are unknown. A sample of chysocolla-bearing oxidized copper ore from Pima, Ariz. was crushed and dry ground for thermal activation and subsequent flotation. PROCEDURES Activation: The apparatus used for exposure of materials to high temperature water vapor is illustrated in Fig. 1. The furnace temperature was controlled to ±5°C Water vapor pressure was controlled by pre-saturating a nitrogen carrier stream with water in an adjustable thermostatted saturator (multiple bubbler). A water manometer and capillary orifice flow meter were used to measure nitrogen flow rates for activation and for subsequent Hallimond tube tests. Gas flow rate and temperature were adjusted prior to introduction of samples. Exposure time was measured from the moment of introduction of the sample to the furnace, hence includes an initial drop below the control temperature. After heating, the samples were removed from the furnace and stored at room temperature in air. Representative samples of activated materials were examined by X-ray diffraction and infrared absorption.

Jan 1, 1967

By R. Maddin, A. Lawley

Single crystal Mo-Re alloys (99.99+ purity), grown by electron bombardment floating zone heating, were deformed in tension at temperatures from -196° to +200°C. At -196oC, Mo-6 pet Re and Mo-20 pct Re alloys deform by slip on systems of the type (101) [1ll]. Above-78oC, the operative slip system is of the type (112) [111], (101) [1ll], or (211) [1ll], depending upon orientation. Mo-39 pct Re alloys deform by twinning and slip at all test temperatures with a (112) [111] type twinning system. Critical resolved shear stress for slip and twinning is calculated as a function of composition and temperature, and the mechanical properties compared with those of less Pure polycrystalline alloys. HEXAGONAL close-packed rhenium, c/a = 1.615, is highly soluble in molybdenum. The bee molybdenum-rich solid solution extends to approximately 42 at. pet Re* at 2500°C,' Fig. 1. Recent work on polycrystalline alloys has shown conclusively that the addition of rhenium to molybdenum improves both the fabricability and the low-temperature ductility.2"5 Similar behavior is observed for W-Re and Cr-Re alloys.5 Initially, Geach and Hughes2 found that the Mo-30 pct Re alloy could be cold rolled to approximately 90 pct reduction before the onset of cracking. The remarkable effect of the rhenium was later confirmed by Jaffee et al. 3-5 with an optimum condition of maximum strength combined with maximum ductility at the Mo-35 pet composition. At higher rhenium contents, the ductility was severely impaired due apparently to the presence of brittle tetragonal s phase, Fig. 1. In the Mo-35 pet Re alloy, at least 110 ppm (parts per million) oxygen could be tolerated before cold workability was reduced. By comparison, the corresponding tolerance figure in unalloyed molybdenum is ~ 10 ppm. The effect of rhenium was considered in terms of: 1) a reduction in oxygen solubility in the lattice; 2) a redistribution of intergranular oxide from grain boundary films to isolated spheres; and 3) the promotion of twinning as a further mode of plastic deformation. Paradoxically, twinning is often associated with brittle fracture in iron and the silicon irons.6'7 Work to date has been restricted to are-melted and powder metallurgy polycrystalline material of ~ 99.9 pet purity. The present investigation was initiated for the purpose of determining the mode and crystallography of plastic deformation of a series of bee Mo-Re alloy single crystals over a wide temperature range from -196o to +200°C. Using high-purity zone-melted material (~ 99.99 + pet purity), the specific role of rhenium in the molybdenum lattice may be understood more fully since effects due to grain boundaries and oxide redistribution will be absent. At the same time, the study provides a comparison of single and polycrystalline deformation behavior. EXPERIMENTAL TECHNIQUES A) Preparation and Purity of Single Crystals. Alloys in the form of either are-cast and swaged or sintered and ground bar, 0.120 in. diam, were obtained from the Chase Brass and Copper Co. Alloy single crystals ~ 8-in. long were grown from the

Jan 1, 1962

By D. R. Parrish, T. M. Geffen, R. A. Morse, W. W. Owens

Laboratory studies of several factors affecting measurements of relative permeability were made using the three-section plastic-covered core technique. Results show that the core assembly, properly constructed, will perform as a single unit, and that the testing technique will, under suitable conditions of pressure gradient, gas expansion, and migration of partial water saturation, permit measurement of flow characteristics not affected by technique. Wettability equilibrium is readily established in cores exhibiting strong wetting preference to water or oil when initially saturated with water. Laboratory tests must be conducted so that saturation changes represent those that occur in the reservoir. Immediate implications of saturation history are (1:) that the possibility exists of increasing the displacement efficiency of solution gas drive reservoirs over the natural process, and (2) residual gas saturations following water flooding in gas or gas condensate reservoirs will be 15 to 50 per cent pore space rather than 1 to 11 per cent as generally believed. INTRODUCTION Solutions of petroleum reservoir problems pertaining to production performance require the use of true relative permeability characteristics. 'This relationship of fluid conductivity and saturation has been obtained by reservoir engineers in four ways, namely: 1. From past gross reservoir performance and the extrapolation of this data based on experience,1,2 2. By using published fluid flow relationships obtained in laboratory studies on general type porous materials,3, 4, 5 3. By attempting a mathematical derivation of flow behavior," using some experimentally obtained characteristics of reservoir rocks, and 4. By laboratory flow tests using representative rock samples of a reservoir. The first three methods listed above have shortcomings which make their use limited or questionable. Production characteristics of only certain processes are obtained from field data and these are not available at the beginning of a reservoir's producing life, at which time they are desirable. It is fortuitous if general fluid flow characteristics obtained experimentally have accurate application to specific field problems. Also, it is felt that at this time there is not sufficient knowledge of the flow behavior of oil, water, and gas in porous materials to enable applicable analytical description of this to be made based on other measured rock characteristics. Measurement of relative permeability in the laboratory offers the only direct method subject to adequate checking for determination of flow characteristics applicable to field problems. Primarily, this paper deals with laboratory experiments to establish the effects of several factors on the measurement of relative permeability and the practical significance of this knowledge. Laboratory flow tests are essentially model studies; the gross reservoir being represented by a small rock sample. It is necessary to simulate in experiments all the factors of importance which influence flow in the reservoir, as well a.; to eliminate effects peculiar to the model studied to assure measurement of valid flow behavior. Factors which affect the measurement and behavior of fluid flow in porous materials have been studied by other investigators.3,7,8,9,10,11 This analysis is not intended to evaluate fully all factors involved, but to establish first the validity of measurements by the technique used and then study flow behavior under various conditions. The studies reported are as follows: 1. Effects of pressure gradient, 2. Effects of gas expansion in tests conducted at near atmospheric pressure, 3. Effects of migration of a partial water saturation in oil-gas-water systems, 4. Effects If wettability characteristics of the system. and 5. Effects of the fluid saturation history. The first three are important only in laboratory measure-

Jan 1, 1951

By J. L. Huitt, IM. L. Slusser, E. E. Glenn, M. L. Slusser

This paper is concerned with: (1) an analysis and interpretation of the filtration characteristics of drilling muds on filter paper, and (2) an interpretation of early stage "filtration" on consolidated synthetic porous media, herein referred to as the "surge loss" period. It is shown that water-loss data can be plotted so as to disclose the existence of a surge period, a non-uni-form cake thickness period, and a constant pressure filtration period. The same periods occur during exposure of consolidated porous media to muds. The data on consolidated porous media indicate that: (1) mud particles bridge pores inside porous media during the surge period, and (2) "filter cake" build-up inside porous media occurs during subsequent filtration. The data also show that the surge loss increases with sample permeability and with pressure differential. INTRODUCTION The phenomenon of mud filtrate invasion into porous sands of a drilling well has been recognized for many years.1,2,3,4 Part of API Code 29 is concerned with the filtration testing of a drilling mud to yield data on this subject. In a number of investigations "plastering properties" of muds were studied,5,6,7,8 "plastering" apparently meaning cake forming properties. In such studies the characteristics (thickness, permeability, and toughness) of the cake and the rate at which the liquid phase of the mud was separated from the solids (filtration rate through the cake, commonly referred to as water loss) were measured or at least considered. From the standpoint of its effect on well productivity, filtrate invasion into "clean" water-wet pay sands (i.e., those containing no clay minerals) may not necessarily be a serious problem. An increase in water saturation in the zone immediately around the wellbore would cause loss of oil permeability. However, in such a system, much of this water may be removed by oil production. If filtration occurs into "dirty" sands (i.e., those containing appreciable clay minerals), clay swelling or other mechanisms may cause losses in oil permeability which cannot be overcome by oil production.'.' The phenomenon of mud particle invasion into the pay sand of a drilling well was considered to be a possible serious factor in the early days of rotary drilling. Opinions varied as to whether drilling fluids "mudded up" and "sealed" oil sands. In 1932 Rubel,9 Gill,10 and Parsons," in separate papers, presented differences of opinion together with some laboratory and field data concerning loss of oil production caused by drilling fluid invasion into pay zones. These papers, and the discussions, furnish a large background on this subject and indicate the very clear thinking of many men at a time when petroleum production technology was in its infancy. Jones and Babson,6 and Silent,' experimentally found that the depth to which mud particles invaded water-saturated sand packs was small (no greater than 1/16 in). Arn-quist12 in 1937 observed loss of water permeability due to invasion of mud particles into water-saturated, consolidated, sandstone samples. Williams and Cannon5 studied the filtra-tion properties of drilling muds and mentioned "sludge" which penetrated a short distance into the water-saturated filtering medium (sand pack). Byck5 concluded that formation permeability had no effect on the plas-

Jan 1, 1958

By D. C. Ganow, N. J. Grant, I. R. Silver, A. R. Chaudhuri

The structural changes that result when a metal is "cold worked" lead to higher values of yield and tensile strength on subsequent deformation at room temperature. Further it has been shown that the beneficial effects of cold work are, to some extent, maintained even at elevated temperatures. The strength improvements from prior cold work decrease as the temperature is raised and are a function of grain size, amount of cold work, and strain rate (or time of test). If recrystallization occurs during test, the cold-worked material may even show weaker stress-rupture properties than the annealed material. A question that arises is whether there is a significant effect on the subsequent creep behavior if the prior deformation is performed under creep conditions at high temperatures not too different than the test temperature, but at significantly different strain rates and for various amounts of strain. An answer to this question was sought in the work described in the present paper. PROCEDURE Two materials were used in this work, high-purity aluminum and a solid-solution A1 2 pct Mg alloy. The compositions are shown in Table I. The grain size in the two materials was 0.2 to 0.3 mm. The deformation for both the prestraining and for the creep tests was at constant stress. The procedure adopted was to prestrain a specimen to the desired elongation, remove the load, and cool to room temperature rapidly by spraying. the test bar with cold water. The specimen was then placed in a creep machine, held at the test temperature for less than 5 min, and loaded. Less than 15 min was required to bring the specimen to temperature. The creep stress was based on the diameter resulting from the prior strain, and values of elongation at fracture listed in the tables do not include the respective elongations resulting from the prestrain. The choice of the stresses for the prestrain and for the creep tests was made so that the tests could be completed in relatively short times to maximize any effect of prestrain. The specimens had a gage length of 1 in., the pure aluminum specimens having a diameter of 0.150 in. and the alloy specimens a diameter of 0.125 in. For pure aluminum, prestrain was at 400' and 600°F with testing at 500°F; for the alloy, prestraining was at 700°F and testing at 600°F. RESULTS The prestrain conditions for the aluminum and the alloy specimens are shown in Tables I1 and IV. In addition to showing the time required to achieve the desired prestrain, the expected rupture Lives are given which would have resulted had the creep prestrain been continued to fracture. The prestrained and the annealed specimens were then tested in creep to fracture; the temperature and stress for the pure aliminum specimens being 500o F and 1800 psi and for the alloy, 600' F and 3400 psi. The creep curves for the prestrained and the annealed specimens of pure aluminum showed all three stages of creep; hence in addition to the effect of prestraining on the rupture life, it was also possible to determine the effect of prestraining on the minimum creep rate, and on the time and elongation at the end of each of the three creep stages. The results are shown in Fig. 1 and Table 111.

Jan 1, 1960

By M. Ahlers, Horace Pops

Slip features and formation of stress-induced mm-tensite have been studied in single crystals of fl phase Cu-Zn alloys in the composition range between 43 and 48 at. pct Zn. A phenomenological analysis is presented in which the glide planes and directions that me associated with the inhomogeneous shear in the martensite phase are predicted from the experimentally determined habit planes and macroscopic shear directions. It is shozun that the resolved shew component of the applied stress on the secondary shew system is a very important factor governing the formation of stress-induced martensite. A comparison is made between martensitic phases that form spontaneously upon cooling and those phases that arise from deformation. THE ß phase in Cu-Zn alloys is unstable and may transform spontaneously by a martensitic transformation1-3 during cooling. Thermoelastic and burst-type spontaneous martensitic phases have been observed3 in the composition range between 38 and 42 at. pct Zn. Martensite forms in alloys containing more than 42 at. pct Zn only by cold work below a temperature denoted as Md.4 Thus, alloys near the equiatomic composition can undergo a stress-induced transformation upon cold work at room temperature. Although data have been obtained concerning the effect of temperature5 and composition4 on the deformation of polycrystalline Cu-Zn alloys, it is rather surprising that few detailed studies6-8 on 6 brass have been conducted for single crystals where the stresses acting on crystallographic planes can be determined. A thorough understanding of transformation mechanisms requires knowledge of the role of stress and plastic strain. With this objective in mind we have made an investigation of the stress-induced rnartensite transformation in single crystals of 0 brass in the composition range between 43 and 48 at. pct Zn. Particular emphasis is given to the morphology and crystallographic features and a comparison is made between stress-induced and spontaneous martensite. I) EXPERIMENTAL PROCEDURE Master alloys having nominal zinc contents of 43, 45, and 48 at. pct were prepared from spectroscop-ically pure metals (>99.999 pct) in sealed quartz tubes under a partial pressure of helium. These alloys were cut into several pieces and converted into small rods by melting in quartz tubes. Subsequently, single crystals were grown from the rods by the Bridgman technique and homogenized for 3 hr at 820°C. Lattice parameter measurements were made from stress- relieved filings. Compositions, as determined from the reported variation of lattice parameter with atomic percent zinc,' indicated a zinc loss of less than approximately 0.25 at. pct. The single crystals were electroturned using a phosphoric acid electrolyte into tensile specimens each with a center section of 1 mm diam and 1 cm gage length. Cylindrical specimens approximately 3 mm in diam and 10 mm long and spark-machined rectangular blocks approximately 3 by 3 by 10 mm were prepared for compression studies. These blocks were stress-relieved and electropolished to eliminate any stresses introduced by spark machining. Crystal orientations were determined to an accuracy of ±2 deg primarily from back-reflection Laue photographs, and sometimes by the rotating single-crystal method. Figs. 1 to 3 show the location of the tension and compression axes within unit stereographic triangles. All specimens were deformed plastically at room temperature using an Instron machine, the crosshead speed being 0.02 or 0.05 cm per min. A hard steel hemisphere was placed at the upper end of the compression samples to achieve uniaxial loading. Trace normals of slip markings and habit planes of strain-induced martensite were determined to an accuracy of ±4 deg from the rectangular blocks by a two-surface analysis and from cylindrical specimens by measuring the angle between the tensile (or compression) axis and the tangent to the surface trace for different rotations about the deformation axis. Slip

Jan 1, 1969

By G. W. P. Rengstorff, H. B. Goodwin

A method is described for purifying iron in batches of 150 Ib or more. Oxygen, carbon, nitrogen, and sulphur are removed from flakes of electrolytic iron by treatment in wet and then dry hydrogen. A special consumable-electrode arc furnace is used to remove hydrogen and to melt the flakes into ingots. A LTHOUGH iron and steel have been made and used for hundreds of years, there are still wide gaps in the knowledge of these metals. One reason for this is that the true properties of the base metal, iron, are not well known, simply because really pure iron has never been readily available for research purposes. Commercial iron, and even the "pure" irons which are readily available, are really complex alloys of iron with oxygen, nitrogen, hydrogen, carbon, sulphur, silicon, and other elements. With such materials, it has not been possible to prepare steels and other iron-base alloys in which all elements are present in exactly controlled amounts. It is an almost impossible task to determine the separate effects of each element or the intereffects of several elements unless alloys can be prepared containing only the desired elements. The ability to make controlled alloys of low impurity content is of far more than academic interest. It will assist in the study of such practical problems as the brittle failure of ship hulls or high pressure gas pipelines, determination of the factors which affect the deep-drawing properties of steels, the improvement of magnetic and electrical properties, the phenomenon of aging, and many others. In 1949, the American Iron and Steel Institute asked Battelle Memorial Institute to prepare high purity iron on a larger scale than had been attempted up to that time. The work was directed toward the preparation of iron as low as possible in carbon, oxygen, nitrogen, hydrogen, and sulphur because these are the contaminating elements of major interest in many research investigations on iron-base alloys. The initial goal was to prepare iron with not more than a few thousandths of one percent of each of these impurities. The purpose was to set up a "bank" of high purity iron on which researchers could draw when a supply was needed and which could be replenished from time to time as the stock became depleted. Of several possible purification methods, heating in hydrogen was selected as the most promising for the removal of both carbon and oxygen from solid iron. Metallic impurities are not removed by hydrogen. Electrolytic iron was chosen as a starting material because it is the most readily available form of iron low in metallic impurities. Decarburization and desulphurization were obtained by annealing in wet hydrogen at relatively low temperatures. The optimum temperature, as found by previous investigators,' is about 1400°F. This treatment also removes nitrogen. Deoxidation is favored by dry hydrogen and higher temperature. There is evidence from previous work that an anneal at 2200°F in dry hydrogen will deoxidize the iron, reduce the nitrogen content, and reduce the sulphur content. These facts suggested the use of a two-stage process: treatment in wet hydrogen followed by treatment in dry hydrogen. To carry out the purification process in a reasonable time, it was necessary that the iron be in thin pieces. Flakes of electrolytic iron were used. Because most research requires that the iron be in massive form, it was nccessary to melt the purified iron to consolidate it into large pieces. For this work, it was decided to prepare ingots weighing at least 15 lb by using a cold hearth arc-melting process in which the iron was melted in an inert atmosphere of argon or helium. This method of melting was selected because the water-cooled copper crucible is not wet by the iron and does not contaminate it as a refractory crucible would. Because the crucible is also the ingot mold, contamination from mold materials was also eliminated. Contamination from electrodes was avoided by making them from purified electrolytic iron. The melting procedure finally developed was an adaptation of the Kellogg electric ingot process.' Two batches of purified iron weighing about 100 Ib each in ingot form were prepared with analyses as shown in Table I. The remainder of this paper describes the procedures in more detail. In general, the process con-

Jan 1, 1956

By R. F. Decker, J. R. Mihalisin

Phase transformations in a series of relatively pure nickel-titanium-aluminum binary and ternary alloys were studied. The purpose was to clarify age-hardening mechanisms, especially in predominantly titanium-hardened compositions. The micro-structures, as determined on a tirne-temperature-transformation basis by light and electron microscopy and X-ray diffraction, were correlated with room temperature and hot hardnesses. The stability of the ? phase under nonequilibrium conditions was shown to be important in this system. MANY of the alloys designed for service in the temperature range of 1200oF (649°C) to 1800°F (980°C) are austenitic iron-.nickel base alloys hardened by titanium and aluminum. The control of high-temperature properties of these alloys has been advanced by the knowledge of microstructures provided by the equilibrium diagrams of Taylor and Floyd1 and Taylor.2, 3 However, there have been notable instances in the predominantly titanium-hardened alloys where the correlation of properties with microstructure and the equilibrium diagrams has been unsuccessful. Mihalisin and carrol14 have noted in electron microscope studies of commercial alloys of this type that a fee phase precipitates and transforms to 71 (hep Ni3Ti) in stress-rupture tests. Bieber5 found that this fee phase remained in nickel-chromium-titanium alloys even when aluminum was excluded. Age hardening in iron-nickel base alloys has been attributed to 71 by Craver, Aggen, and Dyrkacz,6 ? (fee phase based on Ni3Al) by Beattie and Hagel7 and both phases by Lena,8 Yet, on the basis of the nickel-titanium-aluminum equilibrium diagram of Taylor and Floyd shown in Fig. 1, it has been difficult to understand the presence of ? in low aluminum alloys. Furthermore, the characteristic coarse dispersion of 71 has not been commensurate with the marked aging response. Also there have been anomalous grain boundary transformations in the iron-nickel base alloys which have markedly reduced notched stress rupture properties.9 These have been related to boron content and residual cold work. Recently, Golubtsova and Mashkovich,10 Buckle and Manene11 and Bogaryatskii and Tyapkin12 have noted a transition fee Ni3Ti in titanium-hardened nickel-base alloys. It was apparent from the latter studies that a transition phase was important in the nickel-titanium-aluminum system. The purpose of this study was to identify this phase, measure its transformation characteristics and kinetics, and attempt to correlate these findings with age hardening. It was believed that time-temperature-transformation studies, because of their adaptability to transition phases, would be of most value for this investigation. Accordingly, this method of study was undertaken on relatively pure nickel-titanium-aluminum alloys with controlled variations in boron content and residual cold work. EXPERIMENTAL PROCEDURES Compositions were selected from the nickel-titanium-aluminum isothermal section of Taylor and Floyd.' These compositions are plotted in Fig. 1 and were selected to obtain equal aging potential of approximately 15 pet by volume of the equilibrium precipitate. The materials were vacuum melted from

Jan 1, 1961

By Paul A. Beck

While the recrystallization properties of most of the practically important metals are known in considerable detail, those of lead are still relatively little known in spite of some valuable contributions in recent years. This condition probably is due to the difficulties involved in applying the usual methods of investigation. The microscopic method of discriminating between the elongated deformed grains of the familiar lenticular shape and the equiaxed recrys-tallized grains is not successful with fully refined lead: lenticular grains cannot be seen under the microscope with any amount of deformation whatever. The probable cause of this phenomenon was recognized long ago. With large deformation the time of recrystallization at room temperaturc is shorter than the period of preparation of the specimen for microscopic examination; with small deformation the deformed grains arc not noticeably elongated in the first place. However, with small cleformation the migration of the grain boundaries can be observed. The only systematic investigations yet published on the recrystallization of lead, those by R. S. Russell, 1-3 make use of the conditions just described. Russell considered recrystallization completed when no further migration of boundaries could be observed. A serious objection to this method arises from the fact that there is no evidence to support the view that grain-boundary migration results only from recrystallization. If such migration does result from other sources after recrystallization is completed—after the deformed grains have been replaced by new unde-formed grains—the "time of recrystallization" as determined by Russell may be too long. Furthermore, the method is subject to the above.-mentioned limitation because of the unavoidable period of preparation of the sample. According to Russell, even with the fastest technique 5 min. clapsed before the deformed samples were ready for examination, and consequently no reearystallization time shorter than 5 min. could be measured. The lack of success in detecting the recrystallization of lead by the usual hardness or tensile tests is again caused by the time necessary for

Jan 1, 1940

By M. J. Sinnott, J. C. Tobin

A technique for collecting and analyzing the small quantities of gases desorbed on heating metal powders has been developed. The gases collected from copper powders of various types of manufacture have been analyzed and found to contain H2O, CO2, and SO2. The quantities of gas evolved vary from 0.05 to 1.0 milliliter, under standard con-ditons of temperature and pressure, per gram of powder. The source of the gas has been traced to surface films on the powder. IN the powder-metal industry it is recognized that the presence of gases in the powders affects the processing and resulting properties of the sintered product. Several sources for these gases are possible; namely, chemical reactions with the sintering atmospheres, inclusion. of gases during compacting, dissolved gases in the metal powder, and sorbed gases on the metal-,powder surfaces.1, 2, 3 Since the quantities of the gases are relatively small, no direct attempt has been made to collect and analyze them nor to isolate their source. The work that has been done has been concerned with the dimensional stability of the product., its porosity, density, or mechanical and physical properties. Conclusions are drawn on the basis of studies of this type as to the probable effect of gases on these properties. 4, 5, 0 In the present investigation an attempt has been made to obtain quantitative data on the analysis and quantities of gases associated with copper powders manufactured in different ways. Apparatus was constructed to collect gases evolved from loosely packed powder samples on heating to various temperatures under a reduced pressure of helium. The powders were outgassed thoroughly in a vacuum prior to starting the collection operation in order to eliminate mechanically entrapped gases and to prevent the possible reaction of the powder with these gases. Experimental Procedures Materials—The major portion of the investigation was carried out on a single heat of water-atomized, oxygen-free, high-conductivity copper. A relatively coarse fraction of this powder, 20 to 100 mesh, was tested as a mixed powder and as various graded size fractions. The microstructure of this powder showed equiaxed grains and was clean and free from oxides and oxide eutectic. The particle size, as determined

Jan 1, 1959

By R. J. Teitel

The Pb-U system has been investigated by X-ray, thermal, and microscopic analyses. Two pyrophoric intermetallic compounds were found; UPb3 and UPb. The crystallographic structure of UPb3 is reported. A phase diagram was drawn for the system showing three eutectic and one syntectic reaction. AS a result of the interest in uranium metallurgy, the Pb-U system has been investigated. Practically no quantitative data appears in the literature concerning this system. Chips were machined from bars of high-purity uranium obtained from Ames Laboratory and the lead was 99.9 pct high-purity grade from Bel-mont Smelting & Refining Co. The chips were pickled in a dilute solution of nitric acid, washed in water and then alcohol, and dried prior to use for alloy preparation. The lead was simply sheared from the pig with no chemical treatment. Lead and uranium were placed in graphite crucibles and loaded into an induction furnace. In order to insure good alloy formation, the lead was loaded above the uranium to wash the uranium chips into the melt. After the charge was held at the desired preparation temperature for 15 to 30 min, it was allowed to cool in the crucible. The entire melting operation was accomplished in an atmosphere of helium (double charcoal refined grade with no further treatment). The procedure was to outgas the furnace by first evacuating it to 10-% m Hg, admitting helium to 8 psi gage and sealing the furnace off. An optical pyrometer focused on the charge through a hole in the crucible cover was employed for thermal measurements. Usually a 50° to 70°C correction had to be added to the instrument value. The most suitable alloy preparation temperature range was found between 1220° and 1280°C. The reason for this is explained by the phase diagram, Fig. 7. Above this temperature range two liquids separate preventing alloy formation upon cooling at a reasonable rate. On the other hand, below this range the uranium was coated with a layer of both compounds and the rate of reaction was diminished by diffusion processes. The fastest reaction rate occurred in this narrow range of temperature because only UPb coats the uranium and it evidently does not protect the uranium as well as UPb,. Apparatus and Procedures The development of apparatus and techniques was determined by the peculiar properties of the alloys. All the alloys were pyrophoric, and, consequently, the alloys had to be protected from oxidation by handling them in a dry box or inert atmosphere furnaces. Other operations performed outside of this equipment had to be done as quickly as possible. Thermal analyses were conducted in a furnace and arrangement similar to that described by the author.' Graphite crucibles having a thermocouple well in the bottom were used to contain the prepared alloy, while a helium atmosphere protected the melt from oxidation. The furnace was outgassed by pumping a 10-3 mm Hg vacuum until the furnace had been heated to 400°C. Then helium was admitted and the furnace sealed under 8 psi gage pressure. Chromel-alumel thermocouples were used for thermal analysis and the accuracy of thermal measurements was judged to be k 5°C. At first, an attempt was made to use platinum-platinum-rhodium thermocouples, but evidently lead vapors attack it at high temperatures forming low melting alloys. These lead vapors had direct access to the thermocouple because the graphite crucible was porous. A differential temperature control was used to regulate thermal analyses. The thermal gradient across an alundum tube placed between the windings and the crucible was maintained constant by a controller.

Jan 1, 1953

By Frank S. McKenna, Myron A. Pugacz, Clarence E. Jackson

The notched-bar impact test has proved worth while in certain applications as a test for control of the quality or the heat-treatment of steel. In view of the serious thought that even so simple a test as the tensile test has evoked during the past years, it is not surprising that the notched-bar impact test should have inspired much experimental work. The correlation of the great number of data has not been successful, although steady progress has been made toward a better understanding of this test. Progress has been made, especially, in the isolation of the effects of the many variables that control the results. The stress system imposed on a specimen with a notch is so complicated and unpredictable that no simple evaluation of the test has been possible. Because of this uncertainty, many engineers hesitate to use the notched-bar impact test, and this attitude may well continue until a better understanding of the effects of the variables in the test are available. Without doubt, any additional data that may be presented should be accompanied by as complete a description of the - details of test as possible, so that it may be added to the general fund of information that is to serve as the basis for the theory of the notched-bar impact test. The many factors that may control the test results have been pointed out repeatedly. The importance of the effect of the type and dimensions of the specimen has been recognized. Most of the work reported by American investigators has been limited to small-size (Charpy) specimens tested as a simple beam; two types of notches (Fig. I) have been commonly used. Considerable data have also been presented by many investigators using a small specimen tested as a cantilever beam; this commonly being called the Izod test. Here, however, a shear component of stress makes the stress system even more complex than that present in the simple beam type of test, hence a critical analysis is difficult. Much work has also been reported by investigators using the tension type of specimen. A second variable that must be controlled is the testing temperature. An indication that the tests were performed at room temperature is not sufficient, since in many steels a variation of a few degrees in the actual testing temperature may be sufficient to give an entirely different range of results. Other factors, such as the composition and condition of the material being tested and the speed of test, should also be considered. Many investigators have contributed data and theoretical discussions on the problems encountered in notched-bar test-

Jan 1, 1944

By Robert I. Jaffee, Bruce W. Gonser

Recent work by the authors1 established the constitutional diagram of the gold-germanium system. Of particular interest in the simple euctectiferous system was the eutectic alloy at 12 per cent Ge, which melted at the very low temperature of 356°C and had an unusually fine eutectic structure. It was of interest to investigate the effect of third components on the melting point, microstructure, and strength of this alloy, and in this work the effect of copper is discussed rather completely, with less complete data on the effects of silver, platinum, nickel, and zinc. No reference in the literature has been found on ternary systems with gold-germanium. Materials and Alloying -4s in the work on the binary system, gold-germanium, the germanium used in this investigation was the cyanide-reduced, 99.9 per cent product, produced by the Eagle-Picher Co. Fine gold sheet from Handy and Harmon, pure platinum sheet from the American Platinum Works, copper magnet wire from Anaconda, nickel wire assaying 99.6 per cent Ni. and electrolytic zinc were used for alloying. To make the alloys, charges of the pure metals were put in a graphite crucible and melted in an atmosphere of pure hydrogen. The temperature was brought up to about I050 C, and the metal allowed to freeze in the crucible. Except with alloys containing zinc, melting losses were consistently less than a few milligrams. When zinc was a component, weight losses up to half the weight of the zinc present were incurred, and the alloy compositions were calculated by assuming the entire weight loss to be that of zinc. In all other cases the composition of the alloys could safely be taken as that of the charge. The ingots, as solidified, were segregated, the heavier primary phases being at the bottom and sides of the ingots and the lighter primary phases at the top, but by sectioning the entire ingot vertically for microscopic examination, this factor could be evaluated. Ingots for microscopic examination weighed one gram, and those for thermal analysis weighed three grams. Thermal Analysis In this work, the eutectic line composition and thermal arrests were desired, and, after microscopic exploration had established the position of the eutectic line at a given percentage of the third component, a thermal analysis was run at that composition. The method of thermal analysis used in the work on the binary Au-Ge system,' a method particularly well suited for a small quantity of material, was used in this work. Clearly defined arrests were obtained, although with some alloys containing platinum

Jan 1, 1946

By R. M. Goldhoff, J. W. Spretnak

IN connection with establishing the mechanism by which boron enhances the hardenability of heat treatable steels, this research work has been undertaken. Spretnak and Speiser1,2 indicated the need for studying high temperature effects in boron steels and, chief among these, the possible nonhomogeneous distribution of boron in iron grains. The occurrence of this effect may be of significance in explaining the mechanism of enhancement of hardenability by boron as well as in explaining some of the observed characteristics of boron steel behavior. Primarily, then, this work was undertaken a) to establish whether or not boron is adsorbed in ? solid solutions, b) to determine the nature of any such adsorption effects, i.e., whether it is positive or negative as well as its temperature coefficient, and c) to establish the magnitude of any such effect. While past literature on this specific subject is rather sparse, a number of authors have postulated that the observed temporary loss of hardenability as well as the occurrence of the characteristic boron constituent in these steels is due to the nonuniform distribution of boron in austenite.3-0 Experimental Techniques Attempts to obtain direct evidence of grain boundary effects in boron steels have been unsuccessful, both by using the technique of Chalmers' in which small concentrations of solute made radioactive by irradiation are detected by autoradiographic means, and by means of studying cleavage fractures with the electron microscope.8 The present study attempts the design of experiments to measure the effect of boron on properties which might be subject to variation because of grain boundary effects. A study of the theoretical aspects of grain growth indicates that such studies may be pertinent. Another valid type of experiment involves the measurement in polycrystalline materials of properties whose values depend on the bulk of the grain, being negligibly influenced by grain boundaries. One such property is the lattice parameter measured by X-ray diffraction. The value of the lattice parameter reflects the average conditions which obtain over a relatively large number of atoms in a fixed crystalline lattice, and is therefore mainly dependent on bulk grain material. Measurement of lattice parameter variation over the austenite range of temperature in iron with and without boron should give evidence of concentration variations in the bulk grain material if adsorption effects are significant. Finally, use is made of the Grange and Garvey3 metallographic test for boron in which a characteristic boron constituent can be developed in austenite grain boundaries. The very occurrence of this phenomenon suggests grain boundary enrichment of boron and a study of the occurrence of this constituent and its varation with austenitizing temperature should help to detect grain boundary effects. Materials One lb ingots of pure iron, an Fe-B alloy, an Fe-C alloy, and an Fe-C-B alloy were prepared. The melting technique included an initial melting of electrolytic iron under vacuum followed by remelting under hydrogen and, finally, by melting under argon with the addition of alloying elements. Each ingot was surface machined between meltings. A quantitative spectrographic analysis was obtained on the pure ingot and, in addition, oxygen, nitrogen, and carbon were determined by the vacuum fusion method. The results of these analyses are given in Table I. Semiquantitative spectrographic analyses were performed on the alloy ingots, while quantitative determinations of the alloy additions were obtained. These analyses are shown in Table 11. The cast ingots were subjected to the following sequence of fabrication treatments: forging to 5/8 in. bars, swaging to 0.300 in. rounds, drawing to 0.025 and 0.010 in. wire. Portions of each ingot were retained in the swaged condition for experimental use. For certain portions of the work typical commercial steels were used. These materials were supplied by the Republic Steel Corp., Canton, Ohio, and correspond to AISI 8640 and AISI 86B40 compositions.

Jan 1, 1958

By Ralph H. Sweetser

C. H. Herty, Jr.,* Pittsburgh, Pa.—The procedure of the Open-hearth Committee in studying pig iron and its effects in the open hearth has been influenced by the discussions at a number of meetings of the Committee. Those discussions have brought out very clearly that, as far as the pig iron affects the quality, interference with quality should be studied for three definite types of pig iron: (1) low-silicon pig iron; (2) very high-silicon pig iron, and (3) pig iron in which the silicon varies over a wide range in short periods, such as 1 to 3 casts. If a blast furnace is running steadily at 1 per cent. silicon iron, practically no trouble, apparently, can be traced to the pig iron, but if the furnace is running from 0.60 up to 1.50 and then down to 0.60 per cent., trouble in the finished steel is liable to occur, which apparently cannot be corrected by open-hearth practice. The open-hearth discussions have shown fairly definitely that with low-silicon iron, one type of trouble arises on the finished product and with the high-silicon iron, another. Our experience with the high-silicon is that any non-metallic matter that may form, does so when the pig iron is put into the furnace; that is, silicates immediately form which are more liable to be high in silica than would happen with low-silicon iron. I think Mr. Sweetser will agree with mc that if a blast furnace is very irregular, different points in the furnace may be working irregularly, in other words, the temperature gradient all over the furnace is not uniform. We have made many determinations on that point which indicate that a varying blast furnace causes hot spots and cold spots around the periphery. If some of those spots are hot and others are cold, although the analysis of the iron may be constant, then unreduced oxide, going through on one side, and good iron, going through on the other, may mix and the average temperature of the iron is not high enough to get completely rid of the oxides. So our method of attack is directly on that point; it is our first objective. Our procedure is to make a number of consecutive tests at a blast furnace, making a detailed study of the blast-furnace operation. We follow the iron through the open-hearth heat, being sure that we know just how much of each kind of iron gets into the open hearth, and take out specified samples at certain times during the process, working on steel where no finishing additions are made in the furnace. We are following one type of iron through the billet mills and into the finished product and are obtaining complete data on open-hearth operation. Besides the ordinary chemical analyses on the samples of slag and metal, we are running samples of the plg iron, the steel in the open hearth and the finished steel, for oxygen and silicates. We are going to find out whether or not oxygen, as silicates or oxides, has some connection with this trouble. If not, we may be able to locate some other factor. I am inclined to think that the silicates are responsible, judging from the indications in some other work that the Bureau of Mines is doing in Pittsburgh. We have laid out a very definite method of testing to which we will adhere at the different plants where we are now working and at others in the future. Our aim is to

Jan 1, 1927