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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 K. R. Evans

The yield and fracture characteristics of a Ti-8 pct A1 alloy have been examined at room temperature as a function of exposure temperatures to 1700°F. Em-brittlenzent of the alloy is observed to occur following exposures below 1280°F, the critical temperature for formation of ordered Ti3Al particles. Presence of Ti3Al particles results in an increase in the yield and flout stresses which, for a given exposure time, are a maximum at the critical temperature. Embrittlement is attributed to the occurrence of plastically induced fracture controlled by the increase inflow stress. This interpretation is compatible with the decreasing ductility and increasing crack popagation rates observed with increasing exposure time below the critical temperature. Slip line observations and measurements of the gain size dependence of the yield stress indicate that the slip character of the alloy is not influenced by exposure treatments below the critical temperature. ThE potential usefulness of Ti-A1 alloys in structural applications has prompted considerable interest in studies of their microstructure and mechanical characteristics. The relatively recent surge of interest in the system and at the same time its complexities are indicated by the fact that despite numerous recent efforts14 there has not even been general agreement upon a phase diagram for the system. However, most of the Ti-A1 alloys of commercial interest range in aluminum content from 5.0 to 8.5 wt pct where most investigators achowledge the existence of the ordered Ti41 phase. Goldak and parrs confirmed the existence of Ti41 from X-ray diffraction analysis of the stoichiometric composition while Blackburn6 has determined its presence in commercial Ti-8 pct A1 alloys by electron diffraction analysis and dark field transmission microscopy. The embrittlement of Ti-8 pct A1 alloys was initially reported by Crossley and carew7 who showed that aging in the temperature range of 750" to 1110°F resulted in extreme brittleness. This behavior was later confirmed and correlated by Soltiss to the presence of long-range order in the commercial Ti-8 pct Al-1 pct Mo-1 pct V alloy (Ti-8-1-1) which exhibited anomalous creep and yield strength behavior at 950°F. This investigation was designed to study the mechanical characteristics of a binary Ti-8 pct A1 alloy in order to further examine the nature of the embrittlement behavior reported for the alloy system. EXPERIMENTAL PROCEDURE A Ti-8 pct A1 binary alloy was obtained from the Wah Chang Co. in the form of 0.050 in. sheet. The chemical composition of the alloy is given in Table I. Heat-treatment of the alloy was conducted in evacuated vycor capsules for the tensile sDecimens (1.25 bv 0.25 in. gage section) and in evacuated stainless steel jackets for the crack growth-rate panels. All specimens were annealed at 1700°F for 24 hr and water quenched before an exposure at elevated temperatures. Tensile testing was conducted at room temperature in an Instron testing machine equipped with a pushbutton control for instantaneous control of the cross-head motion. In this manner the change in stress, Aa, accompanying an increase in strain-rate by a factor of 10 was measured immediately after yielding. The crack growth-rate panels used were 3 by 18 in. with a precrack 0.33 in. long cut into the center of the panel. The panels were cycled at a maximum gross stress of 40,000 psi at a maximum to minimum tension stress ratio of 20. Cycling of the specimens was continued until the crack propagated the width of the panels. The crack length was intermittently measured during slow crack growth. EXPERIMENTAL RESULTS Correlation of Embrittlement to the Critical Temperature for Ordering. Blackburn6 has found that the critical temperature, T,, for the formation of Ti41 in a Ti-8 pct A1 alloy with an interstitial content on the order of that given in Table I is approximately 1245°F. The relationship between the ductility of the binary alloy used in this investigation and T, was estimated from tensile tests on specimens which were exposed for 24 hr at various temperatures followed

Jan 1, 1970

By C. A. Herbert

WITHIN the past few years, considerable development has been made in the coal-mining industry in Illinois and Indiana and it is the purpose of the authors to record its most important phases. Perhaps the two most striking features are the entry into the producing fields of certain large consumers of coal and the magnitude of some of the new operations. Mines are now being equipped by the. Chicago & Northwestern R. R., the Chicago,. Burlington & Quincy R. R., the Standard Oil Co., of Indiana, and the Union Electric Light & Power Co. of St. Louis. Besides, the U. S. Steel Corpn. has increased its coal-mining activities by commencing operations on its large holdings northeast of Benton, Franklin Co., Ill. Until recent years, the supply of coal has commonly been adequate and prices have generally been favorable, especially to large consumers, as the magnitude of their purchases enabled them to obtain satisfactory quotations from producers. In fact, most consumers of coal have believed that they could buy their fuel more satisfactorily in the open market than they could produce it themselves. The entry of the United States into the war was accompanied by various disarrangements of industrial conditions, and a demand for coal in excess of the supply seemed likely to be experienced for an indefinite period. Under these circumstances it seemed that the greatest assurance of a supply of fuel was its production by the operation of mines. In the case of the Chicago & Northwestern R. R. and the U. S. Steel Corpn. the production of coal has been carried on for many years with satisfactory results, but the C. B. & Q. R. R., the Union Electric Light & Power Co. and the Standard Oil Co. of Indiana had not been producers of coal, at least in this district. Apparently the most impressive feature of the newest mines is their capacity, for in some cases it is planned that production shall be in the neighborhood of 1000 tons per hr. There are mine, now in this field having nearly that output, but these were not planned for such

Jan 9, 1919

By Hollis L. Caswell

The present status of integrated circuits utilizing. superconductive switching. elements is reviewed with special attention given to fabrication techniques, methods for interconnecting completed circuits, and refrigeration requirements. Cryoelectronics has been largely an "inte- grated-circuit" technology since its conception because the switching speed of superconductive devices is attractive only when these devices are fabricated with thin-film techniques. It is true that cryotron circuits can be constructed from wires of appropriate materials (as indeed was done by Dudley Buck 1 in his early investigations) but these circuits will switch in times characteristic of milliseconds whereas similar circuits fabricated by thin-film methods have potential switching times of nanoseconds. Furthermore, cryo-electronic devices such as the cryotron lend themselves readily to fabrication by thin-film techniques since these components may be made from polycrys-talline thin films and are relatively insensitive to the presence of impurities (as measured by semiconductor standards). Therefore, during the past decade considerable effort has been devoted to developing techniques for batch fabricating circuit arrays containing superconductive switching elements. Technology had developed to the point several years ago that fabrication of cryoelectronic arrays containing up to one hundred devices was rather straightforward. However, larger arrays containing between lo4 and 106 components which are required for commercial development of cryoelectronics still pose very severe yield problems. Thus in a sense cryoelectronics found itself in 1962 at the point semiconductor technology finds itself today; namely, individual devices and small groups of integrated devices could be fabricated with acceptable yield and the outlook for building larger integrated-circuit arrays was bright. Unfortunately, problems associated largely with yield have made fabrication of these larger arrays difficult. Unlike semiconductor technology, cryoelectronics had to solve the problems of large-scale integration before it could become economically attractive. This has proven to be a sizable burden to bear. Since several reviews exist on superconductivity,2 superconductive devices,3 and cryoelectronic technology, no attempt will be made in this paper to summarize these areas. Instead a few specific topics will be dealt with in more detail. First, a brief description is given of selected superconducting switching and storage devices with special attention to several metallurgical techniques which improve the performance of these devices. Second, techniques used to fabricate cryoelectronic devices are described with emphasis on problems affecting yield. Third, techniques for interconnecting a number of cryoelectronic planes are described. And last, refrigeration of cryoelectronic components is discussed briefly since the low operating temperature of superconductive devices is an important consideration in this technology. SUPERCONDUCTING STORAGE AND SWITCHING DEVICES The basic superconductive switching device is the thin-film cryotron. The geometry of this device is attractively simple, since it involves only the intersection of two lines that are electrically insulated from each other. The switching element (gate) and control element (control) of a crossed-film cryotron are arranged as illustrated in Fig. 1. The material for the gate is selected to permit the gate to be switched from the superconducting to the normal (resistive) state by the application of a control current. Tin, which has a critical temperature (T,) of 3.7°K, is commonly used for the gate and the cryotron is operated at a temperature just below T, (for example, 3.5°K). The control material (normally lead, with T, = 7.2°K) is chosen so that the control is never driven normal during circuit operation. To improve cryotron operation, a ground plane, also of lead, is placed under all of the circuitry to act as a diamagnetic shield and improve the current-density uniformity across the width of various thin-film elements. Normally, line widths vary from 0.005 to ^ 0.020 in. and film thicknesses from 300 to 10,000A, although new fabrication techniques make narrower lines feasible. In fabricating cryotrons it is important that the edges of the gate elements be geometrically sharp to avoid undesirable switching characteristics associated with a thinner edge region, Fig. 2. One technique which has been used extensively to form patterns consists of placing a physical mask containing the film pattern between the evaporation source and the substrate and depositing through the mask. Film strips formed in this manner possess a penumbra at the film edges due to shadowing of the evapor-ant under the mask. Several techniques have been proposed for minimizing effects due to this penumbra. One of the more promising metallurgical techniques

Jan 1, 1967

By W. C. Browning

The influence of the hydroxyl factor is more damaging to formations penetrated and causes greater consumption of drilling mud additives than previously realized. This hydroxyl effect on clays is essentially independent of the cations present in the drilling fluid and thus differs from the base exchange reactions that have preoccupied mud chemistry with sodium and calcium bentonite concepts for nearly two decades. The new organic polyelectrolyte-con-ditioned muds haw made it possible to use materials other than sodium hydroxide to maintain the alkalinity of such muds. The properties of silicates, as indicated by their dissociation characteristics and buffering action, are such that they can control he pH and alkalinity of drilling muds at the desired level and, at the same time, minimize undesirable hydroxyl effects associated with sodium hydroxide. This use of silicate compounds is different and distinct from prior applications of silicates as deflocculants or shale preservers. Laboratory and field data presented in this report show that silicate compositions can be utilized to adjust the alkalinity of drilling muds and, at the same time, minimize hydmxyl-promoted clay cleavage. INTRODUCTION Studies for improving the efficiency of rotary drilling techniques must consider the chemistry of drilling fluids and of the formations being penetrated. The chemical aspects of drilling must be studied in conjunction with and in relation to the mechanical factors if, for example, penetration rates are to be optimized. Drilling fluid technology has been largely influenced by chemical reactions of the montmorillonite (bentonite) clay minerals. Most of the literature of mud chemistry concerns the properties of bentonite. Clays of the kaolin or illite type, which are nonswelling, are not generally regarded as sources of drilling mud problems. If these nonswelling shale clays are considered, they are commonly regarded as inert solids. Particularly noteworthy is the fact that the relation of surface and colloid chemistry to massive shale bodies has received only scant attention from drilling technologists. Clay studies reported in the drilling mud literature have dealt, for the most part, with the properties of clays in a finely divided state, and often in very dilute suspensions. Yet frequently during drilling, shale problems not related to the rheology of clay suspensions develop in massive non-bentonitic shale sections of zero or near zero permeability. This paper is concerned with surface chemical reactions that can influence the behavior of these non-bentonite clay masses in such a manner as to adversely affect drilling operations. Browning and Perricone1,2 have pointed out that some of the most troublesome shales to drill, such as the Atoka, contain no montmorillonites. They also pointed out that mud problems can frequently be mitigated by reduction of clay cleavage achieved by using drilling fluids with a minimum of available hydroxyl ions. If pronounced clay cleavage occurs during drilling, the borehole may soften, increasing the possibility of sloughing. In addition, the resulting increased incorporation of high-surface-area clay solids into the mud system can reduce penetration rates and necessitate greater chemical treatment. This increase of shale, of colloidal or near-colloidal dimensions, into the drilling mud is due to clay aggregate cleavage and not to base exchange or swelling reactions, such as occur with bentonites. Searle and Grimshaw3 point out the difference between cleavage or slaking reactions of nonswelling clays (such as illite and kaolinite) and the swelling of bentonite. They further state that the speed of slaking is increased in alkaline water. Eitel1 cites Salmang and Becker, who recognized that clay surface reactions impart plasticity and workability to clay masses. Their results clearly show that all liquids which contain hydroxyl groups in their molecules favor the workability of clays. The ancient technique of aging clays in the moist condition to increase their "workability" is evidence that these clay cleavage phenomena are of considerable importance. The same hydration cleavage that occurs during the aging of nonswelling clays for ceramic use also acts to break up cuttings and soften the borehole during drilling operations. The mechanism of cleavage of the crystalline aggregates of illitic, kaolinitic and other nonswelling clays; and the chemical means of controlling this cleavage are therefore of considerable significance to the drilling mud chemistry. Inasmuch as there is little reference in the drilling mud literature to the cleavage reactions of nonswelling clays, the structure of these clays and certain properties that relate to the mechanism of clay cleavage will be reviewed briefly. STRUCTURE OF NONSWELLING CLAYS Clays such as kaolinite, illite and montmorillonite are compmed of alternate layers of (1) silicon-oxygen tetra-

Jan 1, 1965

By Edward B. Durham

The refinery at the San Francisco Mint takes the bullion purchased by the receiving department, and carrying Illore than 200 parts of precious metals in 1,000, or, in mint parlance, over 200 fine, and separates and refines the various metals contained therein, using electrolytic processes exclusively. Bullion containing silver is treated in cells charged with a nitric electrolyte. These cells produce fine silver and leave a residue rich in gold. The residue from the silver-cells, together with crude gold-bullion, is treated in cells having a chloride electrolyte. These produce fine gold and leave a residue containing silver chloride. The latter is reduced to the metallic state with zinc and is then treated in the silver-cells. The various waste solutions and the wash-waters, after being freed from the bulk of their precious metals, still contain copper and other metals. These are removed by scrap-iron, and are then treated in the copper-cells, having a sulphate electrolyte. These cells produce pure copper, and collect a residue containing lead, some gold and silver, and all the metals of the platinum gronp that were in the bullion. This residue is relatively small, and is melted into bars and stored until suEcient accumulates to warrant treating it for platinum, etc. The refinery occupies three large and three small rooms. The large ones are, a melting-room, 30 by 34 ft.; a cell-room, 39 by 46 ft.; and a wash-room, 30 by 33 ft. The small rooms are used as foreman's osce, laboratory, and , re- spectively. The methods here described are those in use in December, 1909, when notes for the present paper were taken.

Jan 1, 1912

By P. B. Bucky

IN previous publications1 it was shown that a scalar model of any weighty structure, where the stresses produced are mainly due to gravita-tional forces, will behave similarly to its prototype if the model and proto-type material are the same arid the model is placed in a centrifugal field of force (substituted for the gravitational field) which has been increased in the same proportion as the linear model scale is decreased. This principle is now being applied to a study of the behavior of mine structures in an attempt to determine the laws that control them. The importance of using models as a means of solving mining problems has been discussed previously, yet it may be well to reiterate here that this method of research and type of solution have been productive of far-reaching results in other engineering fields, notably those of hydraulics and aerodynamics. Model research is a comparatively cheap and time-saving method of checking practical experience. Up to the present time the best guide to what will happen underground has been the judgment of the man who has observed most and been longest acquainted with the local conditions. His experimental work underground is limited by the elements of cost, time and danger, while failure is seldom permissible. In a model, as in a test tube, all elements that affect the solution of a problem are under definite control. Each element may be varied independently at will and its effects noted. Experimental work on a small-scale model may there¬fore be performed at a comparatively small expenditure of time and money, instead of on a full scale in the mine, where the cost, time, and danger elements are necessarily great. The experiments described in this paper are to be considered as part of a progress report and a continuation of Technical Paper 425. One object of the research program is to make each experiment point the way for further research on that particular phase of the general problem. The results and conclusions presented here are therefore not to be considered as final, but are to be interpreted as

Jan 1, 1933

By M. M. Wong, D. E. Couch, G. M. Martinez

The Bureau of Mines electrowon hafnium metal with an average oxygen content of' 150 ppm at 700°C from an electrolyte containing 27 wt pct LiCl, 62 wt pct RbCl, and 11 wt pct HfC14. The average anode and cathode current efficiencies were 90 pct at anode and initial cathode current densities of 86 amp per sq ft. Haf-nium metal with an average oxygen content of 440 ppm was electrowon at 800oC from an electrolyte containing 90 wt pct KC1 and 10 wt pct HfCl4. The average anode and cathode current efficiencies were similar to those obtained in the LiCL-RbCl-HfCl, electrolyte. The chlorine gas given off at the graphite anode was vented through either a silica or a graphite tube to prevent cell corrosion. THE current method for the commercial production of high-purity hafnium is the thermal decomposition of Hfl4.1 The iodide method is not adaptable to continuous process techniques. Nettle, Hiegel, and Baker2 studied the electrorefining of hafnium from hafnium sponge containing 800 ppm oxygen. They failed to obtain hafnium with 600 ppm oxygen in their initial deposits and obtained AEC specification for oxygen only after 75 pct of the soluble hafnium had been removed from the electrolyte. Calculations using their data indicated this was approximately 4 lb of hafnium. The electrolyte was then used to produce approximately 3 lb of hafnium with a low oxygen content. However, no data are shown concerning the amount of anode material initially used or what percent of it was dissolved, therefore, results are not suitable for evaluation of a continuous operation. In general, it was not possible to consistently obtain low oxygen content metal with the electrolytes described by Nettle, Hiegel, and Baker. Wong, Hiegel, and Martinez3 investigated the electrorefining process for hafnium and showed that even by strict control of electrolyte composition only relatively low oxygen reduction could be obtained. The oxygen contained in the hafnium anode material tended to transfer to the cathode deposit and only a limited purification was possible. Both the "iodide" and the "electrorefining" processes depend upon hafnium sponge as a starting material. The sponge is normally produced by magnesium reduction of HfC14 ' and does not meet AEC specifications for hafnium metal. Since only 30 pct of the anode feed could be utilized3 in the electrorefining cells, the Bureau of Mines developed an electrowinning process. HfC14 was used as the feed material for the electro-winning process described in this report. Many of the electrolytes used in the electrorefining studies3 ap- peared to be suitable carrier-electrolytes for HfC14. However, in the initial studies on electrowinning, it was desirable to use electrolytes that had low solidus temperatures and could be operated over a wide temperature range to investigate parameters of the process. Therefore, electrolytes containing LiC1, NaC1, KC1, RbC1, CsC1, and HfC14, in various combinations were explored. EQUIPMENT Chlorinator. Hafnium carbide was chlorinated to produce HfC14 in the batch-type chlorination shown in Fig. 1. Chlorination temperatures were measured with a thermocouple placed in the center of the HfC charge. A flow meter was used to monitor the helium and chlorine. The exhaust side of the silica chlorina-tor tube was equipped with a flask for collecting organic material released during the initial heating of the HfC. The temperature of an internal heater, which extended from the HfC14 condensing flask to the hot end of the chlorinator, was adjusted to prevent the HfC14 from condensing before entering the collection flask. Helium and excess chlorine were exhausted through the lid of the collection flask to an aqueous NaOH solution. Sublimer. Initial studies were conducted using a sublimer, Fig. 2, made by placing a 13/8-in. OD nickel thimble 11 in. long, inside a 11/2-in. ID nickel bell 12 in. long, and locking it in place. This unit was loaded with HfC14 and partially immersed in the molten electrolyte for sublimation directly into the electrolyte. In another sublimer shown in Fig. 3, the HfC14 was contained in a "resin reaction flask". Quartz wool, previously heated to 600aC, secured between two nickel wire screens, was placed just above the HfC14 powder. The lid contained a vacuum outlet, a gage, an argon inlet, and an air-cooled pipe for condensing the HfC14. This sublimer was evacuated and heated. The sublimation temperature was not critical and the sublimer operated satisfactorily at all temperatures between 250" and 350°C. Electrolytic Cell. The electrolyte chamber, Fig. 4, was made of mild steel 8-in. schedule 20 pipe, 30 in. long. The exterior was metallized with a Ni-Cr alloy. The electrolytes were contained in a 16 gage nickel or iron liner with a nickel heat shield on top. The cell was heated by a resistance furnace. A 21/2-in. ID by 25 in. long air lock was connected to one port of a two-port cell cover assembly through a slide valve. The cover assembly of the air lock was electrically insulated from the cell and was equipped with a rubber sleeve that provided for the passage of the cathode lead. This allowed the cathode deposits to be removed and a new nickel cathode to be introduced without allowing air to enter the cell. A tube-rod assembly was bolted to the other port on the cell cover assembly and was sealed by a packing seal. The tube-rod assembly consists of a graphite

Jan 1, 1970

By William Phillips

This paper was prepared at the request of Capt. A. F. Lucas, Chairman of the Institute's Committee on Petroleum and Gas, as a preliminary discussion of the fuel oils which are used-in the Southwest, giving their composition, uses, prices, etc. Should it afterward be found desirable, these observations may be extended over a larger area. In this connection, J suggest that samples of all kinds of fuel oils, including various tars, from all over the country, be collected and sent to some testing laboratory for examination. We would be glad to extend the facilities of the laboratory of the Bureau of Economic Geology and Technology, at the University of Texas, for this purpose, and to do the work free of all expenses. After such examinations are made and the results classified and studied, it would be possible to suggest specifications under which fuel oils should be purchased, and to arrive at some conclusions respecting the uniformity of such specifications, with due allowance for climatic and other more or less local conditions. While such specifications do not now depart from a common standard anything like as much as specifications for illuminating oils do, yet there are discrepancies which should not exist. Some of these discrepancies are of lessening importance, due to the pronounced tendency to distill every crude oil that can be distilled with profit, and to send residues into the market for fuel oil. Residues from the refineries are of a more uniform chemical composition and physical nature than many of the crude oils that go direct into the fuel market. The lower flash-point oils are disappearing from the fuel trade and in place of them we are getting oils that do not flash below 220° F., nor burn under 320° F. In the Southwest, it has not been long since we had fuel oils that flashed at as low as 110° F., and the average flash point over a considerable period was 151° F. Now we are using oil that flashes at about 30° lower than the former burning point, which was 256° F. In other words, the flash point has risen, on the average, 100°; and the burning point about 75°.

Jan 6, 1914

BASIC open-hearth slags have no obviously unique features when compared with slags from other metallurgical operations. Open-hearth slags form and exist at temperatures ranging from 2500 to 3100 F (1370 to 1700 C). In the course of an open-hearth heat, the liquid slag changes slowly and continuously in chemical analysis-the changes reflecting changes in temperature, changes in metal analysis, solution of charged materials, solution of feed additions, and solution of refractories. Throughout the process, the slag performs specific primary functions that are responsible for a large measure of the technical and economic importance of basic open-hearth steelmaking. It is intended in this chapter to present the chemical reactions by which the primary functions of basic open-hearth slags are carried out, to describe the physical and chemical principles governing the extent to which these reactions proceed and can be regulated, and finally to indicate the limitations imposed on other features of the process by the fulfillment of the primary functions. Basic open-hearth slags perform three primary functions. First and foremost, the slag provides a means of removing phosphorus from the liquid metal. The ability to remove phosphorus is largely responsible for the widespread use of the basic open-hearth process. Second, the slag serves as a controlling factor in transferring the oxygen required for the other oxidizing reactions. Third, the slag provides a means of eliminating some sulfur from the open-hearth bath. While the extent of sulfur removal is not comparable to that of phosphorus, the ability to remove some sulfur contributes much to the value of the process. The effectiveness with which the slag fulfills its functions varies with temperature and composition. Since operating temperatures are fixed largely by the specified nature of the product and the equipment, .the speed and extent to which the functions of the slag are carried out depend largely on its composition. The necessity for regulating slag composition to accomplish the re-

Jan 1, 1964

By Michael Perch, Charles C. Russell

Nearly 75 per cent of the total coke production in the United States in 1940 was consumed in blast furnaces. In 1939 the percentage was 69.9, and in 1938 it was 61.3. To produce a net ton of pig iron 1757 lb. of coke was required in 1940 and 1760 lb. in 1939. These figures indicate how dependent the production of pig iron is upon the production of coke. The rate of iron production is in some measure influenced by the physical and chemical characteristics of the coke. Blast-furnace operators believe that the physical properties of coke profoundly influence the operation of the furnace. Although it is true that coke of a wide variety of physical properties is used in American blast furnaces, the periodic fluctuation of these properties is one of the causes of irregularity of blast-furnace operation. The purpose of this paper is to present a broad picture of the physical properties of by-product coke that are generally considered to be important by coke-oven and blast-furnace operators. Effects of the kind of coal used, the preparation of the coal, the rate of coking, and other factors will be considered. Of all these variables, the kind of coal used is by far the most important. In this discussion use will be made of the method of coal classification by rank1 standardized by the American Society for Testing Materials. Although this method of classification has been an American standard for several years, it has not been given the attention it deserves, especially in connection with the selection of coal for production of coke. Despite the fact that many producers of blast-furnace coke are limited to some particular source or sources of coal, a good understanding of the principles involved in setting up the standard method of classification may lead to a better appreciation of the causes of changes in coke characteristics. Furthermore, where it is necessary to select new coals as a substitute for the original supply, this method of classification can be of great assistance. The data presented have been gathered from many sources, but no attempt has been made to cover the vast literature. Instead, some of the more important data have been selected to show how coals and their carbonization affect some of the physical properties of coke. It must be emphasized that generalizations concerning the behavior of coking coals are presented and that individual coals may vary widely from the specific behavior depicted. Nevertheless, it is believed that the data illustrate certain typical behaviors. Neither was it possible to cover completely all the factors involved; for example, the macroscopic constituents of coal are believed to have important effects on the physical properties of coke, but since there is little information available for such study, and since vitrain, clarain, and durain are neither

Jan 1, 1942

By T. E. Davidson, A. P. Lee

It is known from the early work of Bridgman that the two lowest-pressure transitions (I-II and II-III) are accompanied by substantial and abrupt changes in resistivity and Volume. However, unlike the temperature -induced allotropic transformations observed in such elements as lithium, cobalt, tin, and so forth, there is little actually known about many of the characteristics of the pressure-in&ced transitions. This current work involves an examination of the structural and transformation characteristics of the bismuth I-II and II-III transitions under hydrostatic pressures. The relationship of initial structure to the transformation pressure, rate, resistivity change, and resultant structure is discussed. It is shown that the transition pressure and transformation rate are independent of the presence of grain boundaries and associated anisotropy-induced deformation. An observed hysteresis in both the I-II and II-III transitions is shown. BISMUTH is one of the most interesting of the elements exhibiting pressure-induced polymorphs since it undergoes several allotropic transformations at pressures below 90,000 atm. It is known from the early work of bridman1,2 that the two lowest-pressure transitions (1-11 and 11-111) are accompanied by substantial and abrupt changes in resistance and volume. However, unlike the temperature-induced allotropic transformations observed in such elements as lithium, cobalt, tin, and so forth, there is little actually known about many of the characteristics of these pressure-induced transitions. It is the purpose of this work to examine some of the structural and transformation characteristics of the bismuth 1-11 and 11-111 transitions under hydrostatic pressures. Another interesting characteristic of bismuth is that, in its polycrystalline form, hydrostatic pressures of sufficient magnitude will induce severe progressive plastic deformation in the region of the grain boundaries.3 This deformation, which has also been observed in several other metals, is attributed to the high degree of anisotropy in the linear compressibility of bismuth, resulting in shear stresses in the grain boundaries when it is exposed to hydrostatic pressure. Most thermally induced allotropic transformations in metals, whether of the diffusionless ather-ma1 (martensitic) or isothermal nucleation and growth types, are dependent upon structure and prior history,4 viz., grain boundaries, deformation, and so forth. One logically wonders then whether the transformation characteristics of the pressure-induced polymorphs in bismuth might also depend upon initial structure, particularly with respect to the presence of grain boundaries and associated plastic deformation. In this investigation, the role of grain boundaries and plastic deformation on the characteristics of the bismuth I-II and 11-111 transitions will be established. The rather unique residual microstruc-tural changes associated with these transitions will be presented and discussed. The occurrence of a measurable hysteresis in both the I-II and 11-111 transitions will be demonstrated. The type of transformation mechanism based on the observed transformation rate will be discussed. EXPERIMENTAL PROCEDURE A) Apparatus. The hydrostatic pressure system utilized in this investigation is similar to that previously reported by Bridgman' and Birch and Robertson,5 and has been previously described.3 For the purpose of this work, the pressure medium utilized was a 1:l mixture of pentane and isopentane. Pressure measurement was by means of a manganin coil in conjunction with a Foxboro Recorder. The manganin coil was mounted in the bottom closure and inserted inside the pressure cavity. Based on calibration against a controlled clearance piston gage at approximately 10,000 atm, the estimated error in the pressure measurement was +2 pct. Assuming the nonlinearity in the pressure coefficient of resistivity between 10,000 and 28,000 atm to be not greater than 1 pct, then the estimated error in the range of the I-II and 11-111 transitions was +3 pct. B) specimen Material and Preparation. The bismuth utilized throughout this investigation was of

Jan 1, 1964

By Adrien Berton

For nearly a century, the Corocoro deposit has been renowned among geologists from the fact that it shares with the Lake Superior deposits of the United States the distinction of being the only important copper districts where native copper is worked commercially. The Corocoro district, named after the city of Corocoro, the capital of the Province of Pacajes (Department of La Pas, Bolivia) is situated towards the western edge of the Bolivian high plateau, at an altitude of a little over 13,000 ft. The most interesting part of the district is that immediately surrounding the city of Corocoro, at a distance, by railway, of 65 miles from La Paz and 210 miles to the Chilean port of Arica on the Pacific Coast. The climate of the region is that of the semi-desert, dry and relatively cold, with big differences of temperature between day and night. The rainy season begins in December and ends in March. The cupriferous lodes of Corocoro have not attracted as much interest as the silver mines of Potosi but they have been known for a long time. In the precolonial period the Indians were using the charquis, or natural native copper sheets, to make small brass objects, and the beautiful pale green brochantite of the outcrops as a source of green pigments. During the colonial period, the oxidized ores of the outcrops were worked for copper, which was used by the mint at Potosi for making currency. The native copper ores were neglected because the Spanish miners were at a loss how to concentrate and smelt them. About 1830, work by rudimentary processes, which were improved little by little, was begun on the native copper ores. The total production of native copper in the following hundred years amounted to 100,000 tons. This native copper, which was remarkably pure, was much appreciated on the European markets before the introduction of electrolytic copper. It was useful for making special ornamental brass or bronze and the granular concentrates of native copper, called barrilla, containing 80 to 85 per cent copper were suitable for making copper sulfate by direct oxidation and subsequent leaching with sulfuric acid.

Jan 1, 1937

By Mark J. Klein

An internal-friction profile induced by nitrogen in Cr-35 at. pet Re was studied as a function of nitrogen concentration and heat treatment. From these studies, the solubility of nitrogen in this alloy was deduced. The equilibrium solubility of nitrogen in chromium does not appear to be substantially changed by the addition of 35 at. pet Re to chromium. However the precipitation kinetics of nitrogen from solution are significantly decreased. IT has been found that rhenium has a beneficial effect on the low-temperature ductility of Group VI-A metals. As part of a study to investigate the mechanism of this increased ductility, it was of interest to compare the concentration of interstitials retained in solution in one of these ductile bee alloys with the concentration retained in solution in its base metal. In this regard, a study of the solubility of nitrogen in ductile Cr-35 at. pet Re was selected for this phase of the investigation because the solubility of nitrogen in the base metal, chromium, has been determined over a fairly large temperature span.1-3 In addition, nitrogen is known to have an embrittling effect on chromium,4 but does not appear to have a comparable embrittling effect on Cr-35 at. pet Re.5 The experimental method selected for this study was the measurement of internal friction. This technique allows interstitial-concentration measurements to be made in bee metals that are sensitive to small changes in the concentration of a particular interstitial, even when other interstitials are present. When subjected to a cyclic stress, interstitial atoms in solid solution give rise to an internal-friction peak whose origin is the stress-induced ordering of interstitials into preferred lat- tice sites. These internal-friction peaks signify a dissipation of elastic energy which is greatest when the period of oscillation is comparable with the relaxation time for the attainment of an equilibrium interstitial distribution. For dilute interstitial binary alloys, the height of the internal-friction peak is proportional to the concentration of interstitial atoms in solid solutions. snoek8 has shown that the internal friction, Q-1, may be expressed by the equation where A is the relaxation strength, ? is the angular frequency, and 7 is the relaxation time. The relaxation strength, A, is proportional to the concentration of interstitials taking part in the relaxation if no interaction occurs between solute atoms. Thus values of peak height, which can be determined as a function of temperature, give a measure of the relative concentration of interstitial atoms in solution at various temperatures. If equilibrium has been attained, the peak heights will be proportional to the interstitial solubilities. The constant that relates A to the interstitial concentration can be determined experimentally from peak-height measurements as a function of concentration for specimens of constant crystallographic orientation. When this has been done, Q-1 measurements can be used to determine solubilities according to the equation C - KQ-1 where C is the soluble interstitial concentration and K is a constant that can be experimentally determined. It is possible, however, that lattice imperfections may bind interstitials to certain lattice sites. Internal friction does not measure these interstitial~, but does measure interstitials which are free to move in the normal lattice sites when subjected to an applied oscillating stress. On alloying the pure metal with a substitutional

Jan 1, 1965

By James W. Dally, William L. Fourney, Anders Ladegaard Peterson

For the past 3 years, the authors have been conducting research sponsored by the National Science Foundation (RANN) to improve the process of excavation by drilling and blasting. The approach followed has been experimental where the development of stress waves and fractures initiated at the bore hole have been investigated in order to obtain a complete understanding of the dynamic fracture process. The second step in the approach has been to introduce modifications in the drill and blast procedure which will permit closer control of the fracture process. The laboratory investigations involve high speed photography where the dynamic fracture process is recorded with a Cranz-Schardin 1, 2 multiple-spark camera. The camera is equipped with 16 spark gaps which are pulsed at 25 K volts to produce an intense but very short (0.5 sec) flash of light. The camera is capable of recording 16 photographs of a dynamic event at framing rates which can be varied from 30,000 to 1,500,000 frames per second. The exposure time is sufficiently short to stop motion associated with detonating explosive charges and to make visible the details of the fracture process at a bore hose. The bore hole in a massive intact rock formation is modelled with a two dimensional plate containing a circular hole to represent the bore hole. The model material employed is a transparent polyester known commercially as Homalite 100.* This polymeric material is extremely brittle as evidenced by its extremely low fracture toughness of [ ]. The fracture toughness is a measure of the ability of a material to resist the propagation of flaws or small cracks. In comparison, Schmidt3 has recently measured the fracture toughness of Salem limestone and determined [ ]. Thus, the Homalite 100 should closely model the brittle nature of rock where fractures occur at small flaws and propagate without any apparent plastic deformation. Homalite 100 is also birefringent, which indicates that it becomes optically anisotropic when subjected to either static or dynamic loads. Circularly polarized light is transmitted through the loaded Homalite 100 model in a polariscope4 and the birefringence produces an optical interference pattern which is called a fringe pattern. For dynamic photoelasticity, the multiple-spark camera is equipped with polaroid filters to produce the circularly polarized light required to generate the photoelastic fringe patterns. An example of a singlespark frame showing a fringe pattern from a typical experiment is presented in [Fig. 1]. The photograph was taken 0.000072 sec (72 sec) after the detonation of the explosive charge. The circular fringes are due to the outgoing dilatational or P type stress wave and travel with a velocity of 85,000 in. per sec (2260 m/sec) in the Homalite 100. The P wave is followed by a second lower velocity stress wave known as the shear or S type wave which propagates at a velocity of 49,000 in. per sec (1245 m/sec). In the local neighborhood of the bore hole, several radial cracks are visible. These cracks propagate at essentially a constant velocity of 15,000 in. per sec (380 m/sec) prior to arrest. The fringes about the crack tips and in the local region of the bore hole are primarily due to the residual gases contained in the bore hole after the explosive charge was detonated. Sixteen frames similar to this one are recorded during the experiment to give full field visualization of the dynamic event at 16 discrete times over its duration. The fringe order number N is related to the difference in the principal stresses of and 02 according to a stress optic law4: [ ] where f0 = material fringe value, and h = model thickness. The wholefield dynamic-fringe patterns provide a basis for simultaneously observing the interaction between propagating cracks and the stresses which drive these cracks. Fracture Control Experiments Improvements in the efficiency of the drill and blast procedures must involve close control of the fracture process following the detonation of an explosive charge in a bore hole. By control it is implied that the number of cracks initiated and the location of each crack on the wall of the bore hole can be specified. Control also, involves orienting each crack and maintaining the crack path and velocity until the specified crack length is achieved. If the entire fracture process can be controlled, then rounds can be designed to optimize volume removed. fragment size and minimize costs. One area of blasting where fracture control is vitally important is in underground excavation where the strength and stability of the rock walls must be maintained and smoothness and precision of the walls must be achieved. The smooth blasting method is one of the most commonly employed procedures for achieving some degree of fracture control. In smooth blasting, the central region of material is first removed, and then the final row of closely spaced undercharged or cushioned holes are fired to remove the final volume and produce a smooth wall. In some instances, unloaded or dummy holes between the loaded holes are recommended to guide the fracture plane. This investigation pertained to an evaluation of 3 features of the smooth blasting process. These included (a) the effect of stress reinforcement on fracture by simultaneously firing 2 charges; (b) the influence of a dummy hole on control of the fracture planes between 2 simultaneously fired charge holes; and (c) the influence of dummy hole spacing on fracture plane control.

Jan 1, 1979

By Waynw P. Myers

Electricity, as normally thought of by a layman's definition, is a man- made force that has no color, no odor, is not visible, cannot be heard, yet man can control it and make it perform his work for him. A person dare not cheat on it, for it then will master him. It has laws, rules, and regulations that must be strictly adhered to if it is to perform in the safest, most reliable, and economical manner. In the early days of coal mining the power source was man himself. Later man employed steam and compressed air to reduce his muscle power requirements needed in the extraction of coal. By the turn of the 20th century, electricity had become the most common source of power for the coal industry. Steam and compressed air are still found in use in isolated cases today for specific applications; however, they are rapidly being replaced by electricity. Electrical laws and regulations are quite extensive and somewhat complex for the proper generation and use of electricity. Man-made laws, rules, and regulations for the care and handling of electricity, designed for man's safety, are rigidly enforced by Mine Safety and Health Ad- ministration (MSHA), as well as the various state departments of mines. BACKGROUND Until recent years most underground coal mines used direct current exclusively for all movable and for some stationary equipment. Exceptions were for permanent pumps, tipple, fan, and outside buildings where alternating current was normally employed. However, over 60 years ago there were mines with dual systems. Alternating current was used for cutting machines, drills, and pumps ; direct current for locomotives. There were no other machines as these were hand-loading mines. Later, during early mechanization in the 1920s, alternating current was also applied to loading machines at some mines. Use of dual electric systems then declined until fairly recently when

Jan 1, 1981

By S. W. Kennedy, M. Cohen, L. D. Calvert

A high-temperature X-ray diffraction method has been used to study the composition and the kinetics of formation of oxide scales at 800 °C on iron and pure iron-nickel alloys containing 25.6, 75, and 84 pct Ni. This technique has been combined with chemical analysis, and metallographic examination, and determination of weight-gain time curves. The 75 and 84 pct Ni alloys form a three-layer scale consisting of porous NiO next to the metal, a spinel layer, and a -Fe203 on the outside. The 25 pct Ni alloy produces a scale of spinel and Fe,O,, in which, however, traces of FeO and NiO are observed near to the metal. On all the nickel alloys the Fe20, is orieqted with the (107) plane parallel to the surfaces. The exposed surface of the Fe20, is in the form of needles whose axis is perpendicular to the specimen surface. Although the alloys show nonuniform oxidation rates as determined from weight-gain measurements, the X-ray data indicate that the total scale and individml scale layers grow uniformly and parabolically. Deep subscale formation is observed in the 25 pct Ni alloy only. This appeared to be associated with cracking of the scale and penetration of the oxygen along the grain boundaries. ALTHOUGH film formation has been investigated by high-temperature electron-diffraction techniques, for example by Gulbransenl and Hickman,' there have apparently been no investigations of scaling by methods permitting study of the scale during formation at the reaction temperature. Methods involving quenching of scales are based on the working assumption that no significant changes take place on cooling. In general, however, compounds may disproportionate, and the different thermal expansions of constituents may cause disruption of the scale on cooling. In particular, some scales cannot be quenched without spalling. These difficulties are avoided by X-ray diffraction examination at the temperature of formation. High-temperature diffraction procedures also provide a method of determining the kinetics of formation of each phase separately in a multicom-ponent scale. The scaling of binary iron-nickel alloys Containing up to 30 pct Ni at a temperature of 800°C has been studied by Benard and Moreau using cooled specimen~,~ whilst Foley, Druck, and Fryxel14 made an electron-diffraction study of cooled oxide films formed on an iron 42 pci Ni alloy. In the present work the scaling of three iron-nickel alloys containing 25.6, 75, and 84 pct by weight of nickel, at 800°C in oxygen at atmospheric pressure was investigated by X-ray diffraction techniques using a Norelco Geiger-counter diffrac-tometer. The oxidation of iron was investigated for comparison. In addition normal metallographic examinations and chemical and X-ray analyses were made, and weight-gain time curves obtained. EXPERIMENTAL Materials Used—The analyses of the alloys were as follows: Percent by Weight Nickel 25.6 pct 75.4 pct 83.2 pct Aluminum 0.004 None found 0.002 Cobalt — 0.05 0.05 Chromium — — — Silicon — 0.01 None found Carbon All less than 0.05 pct The chemical analysis of the iron was: Al— 0.002 i 0.001 pct, C—0.007 pct, Cr less than 0.001 pct, Cu—0.002 ± 0.001 pct, Mn—0.003 ± 0.001 pct, N—0.0002 pct, 0-0.013 pct, P—less than 0.0002 pct, Pb—not detected, S—0.0017 pct, Si—0.0095 pct, Sn—none detected. This iron was the same as that used to make the alloys and was obtained from the same source-Vacuum Metals Corp. One specimen of pure iron from The British Iron and Steel Research Association was used for preliminary investigation. Specimens were machined to shape from cast rod, polished (the final polish being with 600 Sic paper), degreased in trichlorethylene, washed, and etched. The etching and neutralizing procedure adopted by

Jan 1, 1960

By Henry S. Rawdon

Any laboratory corrosion test, as judged from the practical point of view, is valuable only to the extent that it foretells what will, in all probability, occur in service. Such a test is most properly to be considered as a "pilot test," that is, a test which indicates the direction along which action may be expected to occur rather than as a truly quantitative test which would be expected to tell just how and to what extent the action would proceed. It is generally recognized that laboratory corrosion tests should be chosen with particular reference to the character of service expected for any particular type of metal, so far as it can be foreseen. Even when this requirement has been fulfilled, however, the question whether the actual service behavior of the material is in accordance and general agreement with predictions based upon such laboratory tests is always a pertinent one. The information available in the technical literature bearing on the correlation between the service behavior of metals from the corrosion standpoint, and the results obtained with the same metals in laboratory corrosion tests is very meager, indeed. The results which have been obtained in weather-exposure tests of sheet duralumin which form the body of this discussion are presented here for their bearing on this particular phase of the problem of the corrosion resistance of duralumin. Although the exposure tests described have not been completed, in the sense that all of the tests, in the series initially laid down, have been accomplished, still the general trend shown by the results is so clear that a number of definite conclusions at this stage (subject, of course, to possible modification in the light of later results) are believed to be warranted. The tests reported upon here form one phase of the much broader study of this subject which has been under way for several years in which the National Advisory Committee for Aeronautics, the Bureau of Aeronautics of the Navy Department, the Army Air Corps and the Bureau of Standards have cooperated. Without the very helpful cooperation of the manufacturers in furnishing the

Jan 1, 1929

By F. N. Rhines, A. H. Grobe

At elevated temperatures the oxide of silver is unstable in the air at atmospheric pressure, consequently no external oxide scale forms upon pure silver under conditions of high-temperature annealing. When small quantities of certain alloying elements are present in the silver, the formation of a thin external scale is possible1 and in addition there may form a subscale composed of the oxide of the solute element precipitated within the body of the silver. Norbury2 and Leroux and Raub3 have reported internal oxidation (subscaleiormation) in alloys of silver with 2, 7.5, and 30 per cent of copper. The presence of the subscale is believed to be responsible, at least in part, for the objectionable "fire mark" in Sterling silver.4 Several other alloys of silver, after oxidizing heat-treatments, are known to exhibit undesirable polishing characteristics that may be the result of internal oxidation. Except for the absence of an external scale of silver oxide, it is to be anticipated that silver alloys will prove to be very similar in their oxidation behavior to the alloys of copper, the oxidation characteristics of which have been studied in some detai1.5,6 The present research confirms this anticipation. The oxidation of a series of 20 dilute alloys of silver has been studied metallographically; some types of subscale not encountered among the copper alloys have been found. Instances of internal oxidation in alloys of most of the metals of the 1-6 and VIII groups of the periodic system are on record, but evidence of this type of oxidation in alloys of the metals of the intermediate groups is lacking. A number of the metals of the intermediate group, among them cadmium, lead, tin, and zinc, appear to provide the conditions essential to subscale formation; i.e., they form oxides with a relatively low negative free energy of formation, they dissolve other metals that form more stable oxides, and, presumably, oxygen will diffuse through them. In a study of 40 alloys of these white metals only a few cases of internal oxidation have been found. The probable reasons for this difference in behavior will be discussed presently. Experimental Procedure Silver.—The silver alloys employed in the oxidation studies were prepared in heats of 30 grams each from high-purity silver (99.993 per cent Ag)* and the purest avail-

Jan 1, 1942

By E. S. Davenport, R. J. Hafsten, R. A. Grange

The influence of austenite grain size upon the hardenability of steel is now fairly well understood; for a given austenite, increasing the grain size increases the depth of hardening, and, since the latter is related to the fundamental transformation behavior of austenite, it follows that austenite decomposition must, in general, be retarded by increasing grain size. To date, very little quantitative information bearing directly on the effect of austenite grain size upon the progress and result of isothermal transformation has been published. Davenport and Bain,' and Zimmerman, Aborn and Bain2 have shown that the transformation is definitely relarded in hypereutectoid plain carbon steels at temperatures near and above the upper region of rapid transformation (Ar') when the austenite grains are increased in size. More recently, Griffiths. Pfeil and Allen3 investigated the effect of a series of different heating (austenitizing) temperatures upon the progress of isothermal transformation at 1245°F. (675°C.) and 745ºF. (395°C.) in a medium carbon, low-alloy steel; while they reported no measurements of the austenite grain size, their results suggest that increase in grain size retards austenite transformation at 1245ºF. (675°C), but has no appreciable effect at 745°F (395°C.). The present paper deals with the results of a study of the isothermal transformation of austenite in a single commercial alloy steel in which a fine-grained or a coarsegrained state was induced by appropriate selection of austenitizing temperature. The data are presented in the form of isothermal transformation diagrams, which clearly show that coarsening the grains lessens the rate of transformation of austenite to proeutectoid ferrite and pearlite—that is, at temperature levels above about 1000°F. (540°C.)—but is without significant effect at lower temperatures where the transformation product is bainite or martensite. It is to he presumed that the foregoing statement would be valid, at least in principle, for many other grades of steel; it confirms again that the mechanism of transformation to a pearlitic structure is fundamentally different from that to an acicular structure. A number of related observations were made, which seem to throw some light on the transformation process as it takes place in alloy steels, particularly at temperature levels in the vicinity of 1000ºF. (540°C.), where the transition from lamellar to acicular structures occurs. These data are discussed briefly in the belief that they may lead to a clearer understanding of this aspect of the subject. Material and Experimental Procedure The material investigated was a cornmercial grade of S.A.E. 4140 steel, the composition of which was as follows: 0.37 per cent C, 0.77 per cent Mn, 0.98 per cent Cr and 0.21 per cent Mo. It was hot-rolled to a 1 1/8-in. round bar and normalized by air-cooling from 1800°F. (980°C.).

Jan 1, 1941