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By K. R. Janowski, H. Conrad

As part of a program to establish the effect of crystal imperfections on laser output, a detailed study was made of the dislocation structure of ruby crystals obtained from varioius sources. Using KHSO4 as an etchant, a detailed mapping of the dislocation structure on the (0001) and the (1150) planes was carried out. A less extensive study was made of the rhombohedral planes. The average dislocation density on the basal plane was 1.5 to 3 x 106 cm&apos;2 and on the (1120) planes -5 x 10&apos; cm-2. Howe7!er, considerable variation existed between areas on a given plane. The subboundaries in the basal plane tended to lie along [1100] type directions while those on the (1120) plane tended to lie: a) parallel to the basal plane, b) along truces of the (1101) plane, and c) along normals to the traces of the (0001) planes. The orientations suggest that mang of the dislocations lie on the (0001), (1701), and (1150) planes A Laue back-rellection nnalysis Of c-axis wander in one crystal revealed a gradual increase in misorientation from me end of the crystal to the other, with a maximum variation of 1°35&apos;. Slight etching of a mechanically polished crystal revealed mumerous polishing scratches, suggesting the existence of a plastic flow layer. It is recognized that chromium concentration, lineage, and other crystalline imperfections play a role in the efficiency of ruby lasers.&apos; Consequently, it is desirable to establish in detail the imperfection structure of ruby crystals to be used as lasers. Some exploratory work along this line had already been done by Barnes&apos; and Dils and Martin.&apos;" How- referred to as the morphological system, as described by Kronberg.4 The Miller indices given in the present paper will be based on this morphological system. In this system, the three common growth planes are the (0001), the {1120}, and the {1101}.5 Since these planes are of high stability, it is expected that they are most suitable for study by the etch-pit technique. In the present investigation, a detailed mapping of the dislocation structure was carried out for the (0001) and the (1130) planes. A less extensive study was made of the etch-pit structure on rhombohedra1 planes. An analysis of dislocations in sapphire (the ruby host) was first made by kronberg,4 who postulated that slip on the (0001) plane occurred by motion of dislocations with Burgers vector 1/3 <1150> and that these dislocations might dissociate into quarter partials. An etch-pit technique using boiling phosphoric acid was developed by Scheuplein and Gibbs6, 7 for identifying dislocations intersecting the (0001) planes and the planes near (2021) in sapphire. Al-ford and Stephens8 were able to reveal dislocations on the (0001) and {1120} planes in sapphire and ruby by etching in fused potassium bisulfate (KHS04) at 675°C. The termini of basal dislocations [(000l), <1120>] were revealed by Palmour and Kriegel9 on {1120} and {1010} by thermally etching sapphire. Bond and Harvey10 were able to decorate dislocations in sapphire by adding 1 pet by weight of ZrO2 to the Al2O3 powder used in the growth of the crystals by the Verneuil technique and then heating the grown crystals for 4 hr at 1500°C. They found, using optical transmission microscopy, that many dislocations had straight segments which were parallel to one of the six [1120] directions in the basal plane. Single dislocation lines curved smoothly from one alignment to the other, turning through angles of 60 to 120 deg. Also observed were isolated hexagonal loops, approximately 10 µ in diameter, the sides of which lay along the same directions as did the longer dislocations. MATERIAL AND ETCHING PROCEDURE The ruby specimens used in this study (all containing approximately 0.04 pet Cr) had been obtained from four different sources: Linde Co., Meller Co., Valpey Crystal Corp., and those grown at Aerospace Corp.* For the initial phase of this study, several specimens were cut from the slab shown schematically in Fig. 1 which had previously been cut from a disk boule obtained from the Linde

Jan 1, 1964

By J. Chipman, J. C. Fulton

Reduction of silicon from blast-furnace-type slags by carbon-saturated iron is a very slow reaction even under conditions of rapid stirring. Equilibrium under atmospheric pressure of carbon monoxide was approached from both sides, and the silica-silicon relation was established at temperatures of 1425º to 1700°C for slags containing up to 20 pct Al2O3. Silicon carbide is formed by reaction of graphite with high silica slags and the conditions for its formation were determined. Activities of SiO2 and of CaO in their binary solutions at 1600°C were determined. The effects on the activity of SiO2 were established for additions of 10 to 20 pct AI2O3 and of 10 pct MgO. STUDIES of chemical reactions between blast-*>-5 furnace slags and metal have been focussed on the transfer of sulphur from metal to slag. The important role which silica reduction plays in this process&apos; has led to a study of the transfer of silicon from slag to metal and the equilibrium represented by the equation SiO2 (in slag) + 2C (graphite) = Si + 2CO (g). [I.] The authors2 have published an account of the experimental method employed in attacking this problem and of three series of experimental results at 1600°C. This paper contains data on some additional compositions and extends the observations to cover the temperature range 1425° to 1700°C. The data are used to determine the activity of silica in the slags. As in the earlier experiments, the atmosphere was carbon monoxide which flowed through the furnace at 150 ml per min. In certain heats designated in the tables and figures, the gas was bubbled through the melt using a graphite tube. The latter increased in porosity during a run and in some cases bubbling ceased as indicated by cessation of surging. Some results were discarded for this reason. Experimental Results The new data are shown in Tables I through V and both old and new data are plotted in Figs. 1

Jan 1, 1955

By J. H. Reynolds

The possibility of using a semiconductor, metal, semiconductor structure as a light detector is discussed. A brief theoretical argument is presented which predicts that this structure should have pho-toresponse characteristics similur to those of a metal-semiconductor diode, but with the magnitude of the current increased by a transistor gain factor. Measurements made at low frequencies on a particulur S-M-S device agree u~th this prediction aizd shoul that significant gain factors are attainable. The current gain is substantially higher than that predicted for the metal base transistor. It is proposed that the operation of these units is more closely approxinzated by the metal-gate model which predicts higher current gains. A qualitative discussion of its high-frequency performance indicates that detection of light modulated at microwave frequencies by a S-M-S deuice is possible. ADVANCES in the technology of chemical vapor deposition&apos;-4 have made it possible to form a S-M-S structure consisting of a thin metallic layer sandwiched between layers of single-crystal semiconductor. It has been proposed5-7 that, if the metal layer is made sufficiently thin, this structure could be used in an amplifier having several advantages over presently available transistors including higher-frequency performance. Calculations8 predict that the maximum gain-bandwidth product is nearly an order of magnitude higher than that of a bipolar transistor having the same geometry. Maximum oscillating frequencies in the microwave range are predicted. Electrical measurements made on S-M-S devices have shown that, although the theoretical frequency characteristics are not presently attainable, current and power gain is obtained. Another possible application of the S-M-S structure is its use in a light detector. This structure is a three-layer relative of the metal-semiconductor diode. Light detection using M-S diodes is the subject of considerable present interest. They are known to have a short response time and the ability to detect light having a photon energy below the semiconductor bandgap. In this paper, the mechanism by which S-M-S photocurrent is generated is discussed. An expression is derived which shows that the response characteristics are similar to those of a M-S diode, but that the magnitude of the current is increased due to the transistorlike behavior of the device. The results of low-frequency response measurements on a particular S-M-S structure are presented. These agree with the theoretical predictions and show that significant gain factors are attainable. A qualitative discussion of the high-frequency characteristics is presented which predicts that detection of light at microwave frequencies by a S-M-S device is possible. THEORY In this discussion, the photocurrent, iT, measured across terminals attached to the semiconductor emitter and collector layers is related to the current pho-toemitted from the metal base layer. The latter will be referred to as the primary photocurrent. The behavior of the primary photocurrent is determined by the characteristics of the two metal-semiconductor diodes. The diode photoemission is illustrated in Fig. 1, which shows the emitter and collector potential barriers. The metal thickness has been greatly exaggerated. As will be discussed later, the actual barrier shape can be much more complex, but this diagram illustrates the essential features of the photoemission process. An absorbed photon having an energy greater than the barrier height can emit an electron into the emitter or collector. Since this process is discussed

Jan 1, 1968

By John W. Chandler

INTRODUCTION Where and how to spend the exploration dollar to net the greatest economic return, has always been a major problem for mining companies. Minerals are being consumed in ever increasing quantitie: throughout the world and the mining industry is being called on to meet this demand. In view of the present world-wide nationalistic trend, it is becoming more difficult to determine where exploration money can be spent most profitably. An exploration project should be evaluated continually--in the beginning, to define economic targets and later to determine, as early as possible, whether a project has a chance of developing into a venture that will make a profit acceptable to top management. At this point, a decision can be made whether or not to proceed with plans to bring in a new mine. Exploration, evaluation, and development, therefore, must be fully integrated in transforming an exploration project into a profitable mining venture. DEMAND FOR LEAD AND ZINC Lead and zinc are metals basic to our economy and the demand for them will continue to rise as the population in the world increases. If the population continues to grow at the present rate, it will about double by the year 2000. The accelerating demand for these metals is shown by statistics. Nearly half as much lead (48%) and more than half as much zinc (55%) have been mined between the two world wars as in all pre- ceding history. (1)

Jan 1, 1970

By Donald L. Ball

The equilibrium of boron nitride, gaseous nitrogen, and solute boron was investigated in Fe, Fe-C, and Fe-Si rnelts by the Sieverts technique. The free ellergy of dissociation of boron nitride, involving -the implied standard states, was evaluated in the range 1400" to 1600°C as: ?Fo = 38,000 - 23.5T, cal ber mole The solute-solute interaction coefficients eBB, ecB , and eg were determined as 0.058, 0.22, and 0.078, respectively. BORON is an important ingredient in some types of steel and cast iron. It was employed at a particularly high level, 0.01 wt pet, in the as-cast malleable iron "CentraSteel".1 In this material, the boron concentration was critical in order to assure graphite, and avoid primary cementite, in the as-cast structure. One possible function of the boron in this alloy was the nucleation of graphite through the agency of boron nitride, which, conceivably, precipitated from the melt during cooling. A research program was instituted to assess the validity of this proposed mechanism. The first undertaking, reported here, was a determination of the thermodynamic possibility of boron nitride precipitation. Data available in the literature on the activity coefficient of solute boron in iron-base melts are inadequate for the present purpose. They will be reviewed briefly, however, so that comparison can be made, when

Jan 1, 1968

THERE is probably no phase of basic open-hearth steelmaking that is more of an art and less of a science than mold and pouring practice. It varies widely from plant to plant; consequently it is obviously impossible to set down in detail in this chapter the many ramifications of practice that are now used. The discussion that follows therefore attempts only to give an outline of the more common methods, with examples sufficient to indicate some of the variations in procedure, and to emphasize a few of the many underlying problems that face the open-hearth operator at the pouring platform. MOLD DESIGN Ingot molds are hollow castings of considerable wall thickness. Despite this, non-uniformity in wall thickness should be strictly avoided to prevent the setting up of unnecessary stresses. Molds are cast in various forms, sizes, and shapes depending on the type of steel, and on whether the ingot is to be withdrawn from the mold or the mold from the ingot. The selection of ingot-mold material, which is usually cupola or blast-furnace iron containing approximately 3.5 per cent carbon, 1 per cent silicon, 0.9 per cent manganese, 0.2 per cent phosphorus, and 0.07 per cent sulphur, is governed by numerous considerations that affect mold life. (1) Factors in addition to wall thickness that should be given primary consideration at the foundry are casting skin, shrinkage stresses, and casting temperature. Factors in Mold Design. In use, mold thickness is of greatest importance, for it must be heavy enough to freeze the molten metal rising in the mold and to reduce the temperature of the steel below that of the melting point of cast iron to prevent the inner surface of the mold from melting, which would result in the welding of the ingot to the mold. The skin of an ingot, which solidities immediately upon contact with the relatively cold walls of the mold, must soon become thick enough to contract on the fluid mass and pull away slightly from the sides of the mold. The separation of the ingot from the mold depends entirely on the heat taken up by the mold from the steel.

Jan 1, 1944

By A. W. Hahn

THE importance, to geologists and metallurgists, of knowing the mineralogical forms in which the silver occurs in ores has increased during the past few years, and will continue to increase as long as working hypotheses are converted into proved theories and rule of thumb methods are supplanted by scientific data. The methods used for the determination of these silver minerals have been: (1) careful study of polished ore sections with the microscope, and more recently, (2) metallurgical separation of the ground ore by flotation, tabling and panning, supplemented by microscopic examination of these separated minerals, and chemical analyses. During the course of some experimental work carried on by the testing department of the International Smelting Co. and myself, it becomes important to determine, in a general way, the distribution of the silver values in the minerals, as they exist in Tintic Standard and North Lilly siliceous oxidized lead-silver ores. These mines are in the East Tintic Mining District, State of Utah. In this work the second method was used. Differential flotation tests were run, separating the sulfide from the oxidized minerals, and the products were assayed. Various concentrates were carefully panned and the separated products assayed; special samples of ore were selected and assayed. The microscope was used only for the purpose of determining whether the individual minerals were in the so-called "free" state.

Jan 1, 1929

By C. A. Stickels

A heat of electrolytic iron, to whzch alunzinutn and nitrogen had been added, was hot-rolled, cold-rolled 90 pct, and recrystallized at temperatures from 500" to 700°C. Primary recrystallization textures appear to arise from competitive growth of two types of nuclei: 1) those having orientations belonging to the "usual" primary recrystallization texture found in riming steel, and 2) those with the {111} (110) ovientation. Development of a (111}(1 10) component in the primary recrystallization texture occurs only over a certain interval of isothermal recrystallizatzon temperatures when the material is supersaturated with respect to the precipitation of AlN. Lowering the degree of supersaturation depresses the temperature interval in which a (111)(110) component occurs. An elongated, &apos;pancake-shaped" recrystallized pain structure and a marked delay in the start of recrystallization were found in all specimens which were supersaturated with respect to A1N precipitation after cold work, regardless of their recrystallization texture. ONE of the consequences of killing low-carbon steel with aluminum is a significant change in recrystallization behavior. About 15 years ago, Solter and eatttiel showed that this behavior was largely controlled by aluminum and nitrogen in the steel. If complete precipitation of A1N was prevented before cold rolling, an increased "recrystallization temperature" was observed in subsequent. annealing, and the recrystal-lized grains were not equiaxed. Leslie et a1.2 studied this phenomenon in some detail and clearly demonstrated the relationship between A1N precipitation, recrystallization kinetics, and the development of "pancake-shaped" grains. It has also been known for some time that aluminum-killed steels, processed to produce elongated "pancake" grains, develop a (11 I}( 110) primary recrystallization texture. This texture has not been found in iron or low-carbon rimming steel as a primary texture4j5 but has been observed following grain growth in electrolytic iron.5 The present work was undertaken to study in more detail the effect of A1N supersaturation on recrystallization textures in iron. LITERATURE REVIEW The deformation texture in heavily rolled iron has been studied in detail by Bennewitz.~ The texture consists primarily of a partial fiber texture about a (110) axis in the rolling direction, designated here as fiber texture A. It includes the range of orienta- tions (111)[110] - (001)[ 110] - (11l)[110]. A weak secondary texture also is present.6 This is a duplex partial fiber texture about two (110) fiber axes located 60 deg from the rolling direction and 30 deg from the sheet normal. The range of this texture, designated here as fiber texture B, about the [101} fiber axis is (112)[110] - near (545)[252] - (211:1[011] *The range given here follows Bennewit~.~ A few pole figures from re-crystallized material indicate a broader range than this.&apos; However, the components which are strongest in the recrystallization texture are in this range.&apos;________________________________________________________ Primary recrystallization textures in unkilled steels can be accounted for by growth of members of fiber texture B present in the deformed metal.5 However, while members of fiber texture B dominate the primary texture, other orientations survive primary recrystallization as well. In particular, some {111}(110) members of fiber texture A must also grow during primary recrystallization, because a well-defined {1ll)( 110) texture develops during subsequent grain growth at 700°C.5 The unusual recrystallization behavior of deformed supersaturated solid solutions has been attributed to: 1) retention of the solute in solution,&apos; 2) formation of coherent, preprecipitation solute clusters prior to and during re~r~stallization,~ and 3) formation of a precipitate prior to and concurrent with recrystallization.&apos;~-&apos;~ When aluminum is supersaturated with iron, the difference in grain boundary mobility between general high-angle boundaries and certain special coincidence site boundaries is apparently eliminated.&apos; In aluminum-killed steels, precipitation of A1N can take place at ordinary subcritical recrystallization temperatures. The rate of precipitation increases with increasing aluminum or nitrogen contents.2&apos;13 There is some doubt, however, as to whether true precipitates form during the time at temperature needed to complete recrystallization. Leslie ef a1.2 found that precipitation in one steel was complete after about 100 min at 700GC, or after about 1000 min at 650GC, as measured by chemical analysis for AlN. Aoki et a1.,13 using internal friction for dissolved nitrogen, showed that a large fraction of the dissolved nitrogen was removed from solution within a few minutes annealing time at temperatures from 400" to 800°C. However , the rate of formation of AlN, as detected bv chemical analvsis. was much slower than the apparent rate of nitrogen removal. Hasebe,&apos;~~ using carbon extraction replicas, has identified A1N precipitates by electron diffraction in a 0.2 C steel, solution-treated at 1300°C and annealed 2 hr at 700°C. Borchers and kim,I6 also using a replication technique, observed precipitates after annealing treatments as short as 2 min at 640°C. However, Leslie et a1.&apos; state that no A1N precipitate can be seen while recrystallization is being inhibited in aluminum-killed steel.

Jan 1, 1967

By B. G. Klugh

(New York Meeting, February, 1913.) THE present paper represents a preliminary investigation of the ultimate structure of iron-bearing materials which have been subjected to heat treatment for the purpose of agglormer-ating the fines, or for chemical changes, or for both. Primarily, the object of this inquiry was to determine the relation of the microstructure of each of these materials-to its specific susceptibility to the action in the blast-furnace. It is hoped that this small contribution will serve as a nucleus to a line of research that will throw some light upon the debated questions involved in blast-roasting. Thus far the investigation has been conducted along lines of analogy of known principles, rather than through the applica-tion of any well-defined system of microscopic petrography. This is obviously due to the fact that we are dealing with materials which are the artificial products of a new art, rather than with the natural materials from which the essence of the pure science has been evolved. In the limited tune allowed for this research, it has been conducted upon those materials the chemical content of which allows of their use commercially in the blast-furnace.

Jan 5, 1913

By Clarence Zener

THE heat-treatment of steels will not pass from the stage of an art into that of a science until the mechanism of the phase transformations associated therewith is thoroughly understood. Such an under- standing will involve two distinct types of inquiry: (I) the direction of the trans- formation, (2) the speed of the transformation. The first type of inquiry, which involves only the equilibrium relations between the various phases in steel, is the subject of the present and of a succeeding article; the second type, which involves the kinetics of transformations, will be the subject of a later article. In order that a theory of equilibrium relations may give useful information, certain experimental data must be put into the theory. It is customary for these data to be of a thermal nature. In the system discussed here in steel such thermal data are very scarce. In medium-alloy steels another type of data might possibly be used. In these steels the concentration of solute is so low that it may be possible to use the approximation of dilute solutions, an approximation that leads to linear relationships. The unknown functions, or constants, may then be determined in some cases from the binary iron- alloy system and in other cases from the ternary iron-carbon-alloy system. From the functions, or constants, so determined may then be computed the equilibrium relations in more complex medium-alloy steels. In the present article this second approach is used, the necessary experimental data being obtained from the simple binary and ternary equilibrium diagrams. FUNDAMENTAL EQUATIONS In this article are derived the equilibrium relations between the gamma, alpha and cementite phases. In these three phases all important alloying elements except carbon form substitutional solutions, each alloying atom occupying a potential site for an iron atom. Carbon, on the other hand, forms an interstitial solution. In the face-centered cubic gamma phase, the interstitial positions are at the center of the unit cube and at the centers of the cube edges. There are therefore four interstitial positions per unit cell, or one per iron atom. In the body- centered cubic alpha phase, the interstitial positions are probably at the face centers, and also at the centers of the edges, which are crystallographically equivalent to the face centers. There are therefore six interstitial positions per unit cube, or three per iron atom. In cementite, (Fe, Alloy)3C, the carbon atoms may be regarded as filling all the lattice positions of a given type. The introduction of the fundamental equations of equilibrium has presented the author with a quandary. These equations are formally so similar to other equations that have been derived1,2 for related systems that they may appear to some readers

Jan 1, 1946

Dr. Rossiter W. Raymond, Past President, Honorary Member, and Secretary Emeritus, died suddenly of heart failure at his home, 123 Henry St., Brooklyn, N. Y., on the evening of Tuesday, Dec. 31, 1918. -He is survived by his&apos; wife and daughter. He was one of the founders of the Institute and its second President, his genius was the Institute&apos;s guiding spirit for 47 years. Dr. Raymond was one of the most versatile men of the day; an able mining engineer; brilliant professional writer; an editor of the first rank and an acknowledged authority on matters of literary style; a creator in the fields of fiction, children&apos;s stories, poetry and music; for more than 50 years the superintendent of Plymouth Church Sunday School. He will be revered by thousands as a loyal friend, generous giver, and admired colleague. His passing marks an epoch in mining and metallurgical literature. A memorial service will he held in the Auditorium of the-Institute Headquarters on Monday, Feb. 17, 1919, at 4.30 P.M.

Jan 2, 1919

By J. Chipman, T. Fuwa

The effects of various elements on the activity coefficient of carbon in liquid iron have been studied by two experimental methods: 1) equilibration with controlled mixtures of CO and CO2; 2) the solubility of graphite in the melt. Activity coefficient of C is increased by Al, Co, Cu, Ni, P, Si, S, and Srz. It is decreased by Cr, Cb, Mn, Mo, W, and V. THE thermodynamic properties of the iron-carbon binary system have now been fairly well established, although some uncertainty remains with respect to the exact location of some of the phase boundaries. The activity of carbon in ferrite and in austenite has been measured in the classic researches of R. P. smith&apos; while similar measurements by Richardson and ~ennis, and by Rist and chipman3 have established the values of the activity of carbon in liquid iron up to 1760°C. On the other hand, our knowledge of the effects of alloying elements on the activity of carbon in dilute solutions is restricted to Smith&apos;s experiments on systems Fe-C-Mn and Fe-C-Si in the austenitic range and to some more recent experiments of schwarzman4 in the a range. In addition there have been a number of determinations of the effects of various elements on the solubility of graphite in liquid iron, and from these the corresponding effect in saturated solution may be obtained. The purpose of the present study was to extend the investigation of the liquid system to include the effects of alloying elements upon the activity coefficient of carbon, principally in dilute solutions. Equilibrium measurements were made on the reaction C + co, = 2 CO (g) The prepared mixture of CO and CO,, diluted with argon, flowed over the surface of the liquid metal which, after several hours&apos; exposure to the gas, was quenched and anqlyzed. As in the earlier experiments, the principal experimental difficulty was in the deposition of carbon on the parts of the furnace at temperatures slightly below that of the metal bath. In order to minimize this difficulty, the ratio (Pco)2 /PCo2 was restricted to values not much higher than 100 atm, and correspondingly the carbon concentration in the metal seldom exceeded 0.30 pct. EXPERIMENTAL METHODS The method and apparatus were essentially the same as used by Rist and Chipman.3 The gaseous mixture consisting of highly purified CO, CO,, and argon, each controlled by a flowmeter, was led into the furnace and passed over the surface of the liquid-iron melt which was heated and stirred by high-frequency induction. One slight modification was made in that a molybdenum susceptor was placed outside the crucible for the sake of uniformity of temperature and to combat the tendency of carbon to precipitate on the crucible wall. Pure alumina crucibles approximately 25 mm ID were used. The charge consisting of about 30 g was made up of electrolytic iron, the alloying element to be added, and enough graphite to supply slightly more or less than the anticipated equilibrium carbon concentration. All metals used were of high purity. Metallic chromium, columbium, and vanadium were from special lots supplied by the Electro Metallurgical Co. Tin, copper, molybdenum, tungsten, cobalt, and nickel were of purest commercial grades. The electrolytic iron, after being cut to the proper size for charging, was prereduced by hydrogen at 850° to 1000°C to remove surface oxidation. The oxygen content of the reduced material was 0.002 pct. This treatment made it easy to control the carbon content of the initial melt. The charge was melted under the gas mixture to be used for the entire run. In some earlier melts the charge was melted under a stream of argon, but in this case some alumina was reduced from the crucible, and the aluminum thus absorbed in the melt was subsequently oxidized with the formation of a solid film of alumina on the surface of the melt. AS another safeguard against film formation, overheating of the bath was carefully avoided. All runs were made at a temperature of 1560°C. Under experimental conditions a charge of pure iron picked up 0.17 pct C in 3 hr and 0.23 pct C in 6 hr under an atmosphere for which the equilibrium concentration of carbon is 0.27. It is clear that the time required to reach equilibrium from an initially carbon-free melt would be very great. For this reason each experiment was started with a melt of known carbon concentration not far above or below the expected equilibrium value, and each melt was held at temperature for a period of at least 5 hr. Under such circumstances it was possible to chart the approach to equilibrium from both high-carbon and low-carbon materials. Temperature was controlled by frequent optical observation and adjustment and the metls were timed in such a way that the final 2 hr occurred during a time when electric power was steady; for example, 2 to 4 pm or after 11 pm. In melts containine volatile metals such as copper, tin, and mangane\e the time of holding was decreased somewhat in

Jan 1, 1960

By H. M. Davis, T. S. Jones

A methodfor accelerating the reduction of a metal oxide by carbon monoxide is to incorporate in the reduction system a nonreducible oxide which has a strong affinity for carbon dioxide and will remove it from the product gases. The effectiveness of this method was investigated by veduction experiments on specially prepared powder mixtures of SrO-ZnO ("active&apos;.) and SrCO3-ZnO (reference). Measurements at 750o and 677°C showed that the rate of reduction of ZnO by CO was as much as an order of magnitude greater in the presence of SrO. Under the experimental conditions at least as much CO2 as generated by veduction was absorbed when SrO was present. Enlzancernent of the veduction rate by SvO resulted from modification of&apos; gaseous djfitsion processes. THE specific question posed in this investigation was whether the reduction reaction CO + ZnO - Zn(g) + COa [1] could be accelerated by incorporating strontium oxide into the oxide charge undergoing reduction. Provided the strontium oxide was active enough, it could remove carbon dioxide from the product gases of ( 11 and form strontium carbonate, thus modifying Reaction [ 1] to CO + ZnO + SrO = Zn(g) + SrCO3 [ 2] This study of the comparative rates of Reactions [ I] and [2] can be considered as a particular case of a general scheme for increasing the rates of reduction of metal oxides by carbon monoxide. The limits of applicability of this scheme are indicated in the simple thermodynamic analysis presented in the appendix. This analysis primarily dictated the choice of the system, CO-ZnO-SrO, employed in this investigation. In addition, it was thought that a study of this scheme by use of the CO-ZnO-SrO reaction system would be of special interest in view of indications in the literature1-3 that the gaseous reduction of zinc oxide is prone to rate control by diffusion in the gas phase. The distinction should be pointed out between this and other efforts to increase the rates of reduction reactions through the use of addition agents. Additions of a wide range of chemicals have been tried with de- grees of apparent success varying from slightly negative to highly positive. The reasons for making additions have included the use of salts to produce effects akin to catalysis,4, 5 the alteration of the microstruc-tural details of the reduction process (including solid-state diffusion),6-8 and the conversion of an en-dothermic reaction into one that is exothermic or autogenous.9-12 The means used to attempt alteration of the enthalpy change has commonly been to incorporate lime into the charge, so that heat is generated by reaction with either of the reaction products, carbon dioxide or water. The use of lime in this way makes endeavors such as that of Edstrom and Wiberg similar to the present one. Furthermore, the basic principle behind the present work has previously been set forth.13 However, a detailed investigation of the type of reaction described in this work does not appear to have been conducted. EXPERIMENTAL PROCEDURE From both conceptual and physical standpoints, Reaction [2] is to be considered as the sum of two reactions:

Jan 1, 1968

By L. Seigle, R. Resnick

An inert-marker movement experiment indicates that the ratio of the intrinsic diffusion coefficients DC:DTa = 80:l in TaC at 2500°C. Measurements of the diffmion coefficient of carbon in nonstoichiometric TaC at temperatures from 1700° to 2700°C reveal that Dc increases with decreasing carbon content, but much less than expected from the probable change in vacancy concentration with carbon content. A diffusion process involving two simultaneously operating mechanisms is postulated, and shown to be theoretically feasible. The average value of the carbon diffusion coefficient is given by DC = 0.18 exp[(-85,000 ± 3000)/RT] sq cm per sec over the composition range 46 to 49.5 at. pct C. BECAUSE of their high melting points and hardness, the carbides of the IV, V, and VI group transition metals, along with those of uranium, have attracted considerable interest for applications at high temperatures. In these applications the reactivity of the materials is important, and, since rates of diffusion within the compounds influence reactivity, a knowledge of diffusion kinetics and mechanisms is desirable. While many investigations of the mechanical and electrical properties of these compounds have been made, only two fundamental investigations of diffusion in the carbides are known. Chubb, Getz, and Townleyl measured the diffusivity of carbon and uranium in UC, and Gel&apos;d and Liubimov2 measured the diffusivity of carbon and niobium in NbC. This paper describes an investigation of the diffusion of carbon in tantalum monocarbide and, in particular, the influence of carbon deficiency on this process. Tantalum carbide melts at approximately 3800°C, which makes it one of the highest melting materials known. The compound exists over a rather wide range of carbon Content.3-7 At the peritectic temperature, 3240°C, the phase extends from about 36 to 50 at. pct C. Although the compound can exist with a substantial carbon deficiency, the high carbon phase boundary remains near the stoichiometric composition over the entire temperature range; i.e., no carbon excess is observed. The structure of TaC is the NaCl type wherein carbon atoms normally occupy the octahedral sites in a somewhat expanded fcc lattice of tantalum. Decrease of the lattice parameter with decreasing carbon suggests that the removal of carbon introduces octahedral vacancies into the lattice. I) EXPERIMENTAL DETAILS AND RESULTS Inert-Marker Experiments. In a compound such as TaC the interstitial element would be expected to diffuse more rapidly than the metal. This was confirmed by an inert-marker experiment, following Srnigelskas and irkeendall.8 Ideally, the markers should be placed at the interface between a slab of low-carbon TaC and graphite, and their movement during subsequent inter-diffusion measured. Unfortunately, no solid could be found which is unreactive in contact with carbon at the high temperatures employed in these experiments. In order to circumvent this problem, a specimen was designed in which the markers consisted of several small canals running just below the surface of a tantalum slab. This specimen was prepared by machining grooves on the surface of the tantalum slab and then diffusion-bonding a thin plate of tantalum to the slab over the grooves. The surface of the plate was then ground down until the distance between the canals and surface was as small as possible (about 0.01 cm). Thus, the canals would lie entirely within the TaC phase after a short period of diffusion. The diffusion anneal consisted of immersing the metal sample in high-purity graphite powder and heating for approximately 10 hr at 2500°C under vacuum. At this temperature, the vapor pressure is sufficiently high and the transfer of carbon from graphite sufficiently rapid to allow the surface of the diffusion sample to attain the stoichiometric carbon concentration very quickly. Conclusions regarding the relative diffusion rates of carbon and tantalum in the compound layers (TaC and Ta2C) can be drawn from the location of the canals after the diffusion anneals. If the growth of the layers is governed mainly by the diffusion of carbon, as expected, the canals should remain close to the sample surface. If the diffusion of tantalum contributes appreciably to formation of the compound layers, the distance from the markers (canals) to the surface should increase. Fig. 1 shows, diagrammatically, the appearance of the specimens after diffusion, and Table I presents the depth below the surface at which the

Jan 1, 1967

By E. N. Silverman

Mixtures of&apos; MnS, ZFeO-SiO,, ZMnO.SiO, FeO, and MnO were prepared synthetically and used to study the planes ZFeO. Si0,-ZMnO.SiO,-MnS, FeO-MnO-MnS, and Fe0-MnO.Si0,-MnS in the quaternary system. Because the oxides and silicates studied flux MnS, inclusions in this system are partly liquid during rolling. Also, oxygen can modify subsurface inclusions through the reaction MnS + FeO—FeS + MnO during the soaking of ingots and the reheating of blooms and billets. NONMETALLIC inclusions in steel have long been a subject of study and conjecture because the properties of many steels are known to be affected by the the kind and amount of inclusions present. In free-machining steels, the type of inclusion present has a direct bearing on their machinability; in other steels, cleanliness is important and metallurgists direct their efforts at reducing the amount of non-metallic inclusions. Thus, the accumulation of knowledge about the formation and behavior of inclusions during all stages of steelmaking and processing should aid in improving the quality of steels and decreasing the production costs resulting from rejections. Previous studies of inclusions have contributed greatly to our understanding. The investigators have generally examined inclusions in finished steel. Boulger, Moorehead, and Garvey1 found that in Bessemer steels, low silicon contents results in globular sulfide inclusions with attendant high machin- ability ratings. Van vlack2 also found that low silicon contents favor inclusion morphology that contributes to high machinability ratings. Boquist and Stoll* found that silicon content affected sulfide inclusions and machinability of blooms, billets and bars more than did the finishing temperature in the range between 1775° and 2075°F. Carney and Rudolphy3 studied the distribution of sulfur, manganese, silicon, and carbon in a commercial ingot of resulfur-ized steel and found that silicon was concentrated near the bottom. This helped to explain why bottom cuts had poorer machinability. Klinger and Koch4 developed techniques for extracting inclusions from steel, and the technique devised by Gurry, Chris-takos, and stricker5 for extracting carbides can also be adapted to the study of inclusions. A different approach to the study of inclusions in steel consists of synthesizing the constituents that form inclusions and determining the phase eauilibria in the system that contains the constituents. The resulting phase diagrams can then be used to determine the effect of the various constituents upon each other. In this way, it becomes possible to under-

Jan 1, 1962

By R. A. Vandermeer, J. C. Ogle

The preferred crystallographic orientations (texture) developed in randomly oriented, poly crystalline niobium during rolling were studied by means of X-ray diflraction techniques. The evolution of texture at both the surface and center regions of the rolled strip was carefully examined as a function of increasing defamation throughout the range 43 to 99.5 pct reduction in thickness. Certain aspects of the center texture development in niobium are in agreement with the predictions of a theory by Dillamore and Roberts, but others cannot be explained by the theory in its present form. Above 87 pct reduction by rolling, a distinctly different texture appeared in the surface layers which was unlike the center texture. The present results are compared with previous results obtained from other bcc metals and alloys. RANDOMLY oriented, poly crystalline metal aggregates when plastically deformed to a sufficiently large extent develop preferred orientations or textures. In a recent review article, Dillamore and Roberts1 pointed out that the nature of the developed texture may be influenced by a large number of variables. These include both material variables such as crystal structure and composition and treatment variables such as stress system, amount of deformation, deformation temperature, strain rate, prior thermal-mechanical history, and so forth. From a practical point of view, the control of preferred orientation may often be important for the successful fabrication of metals into usable components. During the past few decades many experiments have been devoted to the study of deformation textures. This work, however, has been confined in large part to metals and alloys that have an fcc crystal lattice. By comparison, bcc metals and alloys have received much less attention, and consequently our understanding of preferred orientations in these materials is only shallow. This state of affairs worsens when it is realized that almost all of our present howledge about this class of materials derives from studies on irons and steels.&apos; The bcc refractory metals, which are relative newcomers to the industrial world, have, on the other hand, been given at best only passing glances in the area of texture development. Our understanding of the evolution of preferred orientations in bcc metals can only remain fairly limited until systematic studies of metals and alloys other than the irons and steels have been carried out and the influence of the many variables has been determined. To that end a program was initiated to investigate in detail texture development in niobium. The present paper reports some of the results of this study. Textures were determined at both the center and surface of strips rolled variously to as much as 99.5 pct reduction in thickness at subzero temperatures. Emphasis in this paper is on texture description and on texture evolution during rolling to progressively heavier deformation. EXPERIMENTAL PROCEDURE The niobium was purchased from the Wah Chang Corp. as a 3-in.-diam electron-beam-melted billet. Chemical analysis indicated the impurities to be less than 300 ppm Ta, 40 ppm C, 10 ppm H, 170 ppm 0, and 110 ppm N. All other impurities were below the limits of detection by spectrochemical analysis. This large-grained billet was fabricated into specimen stock so that a fine-grained randomly oriented grain structure resulted. This was accomplished in three deformation steps alternated with recrystalli-zation anneals of 1 hr at 1200°C in a vacuum of low 10"6 Torr range after each deformation step. The first step was to alternately compress the billet 10 to 20 pct in each of three orthogonal directions. The second step was to compress in only two directions 90 deg apart to produce a 2-in.-sq bar. The final step was to roll this bar 50 pct to give a 1-in. by 2-in. cross section. After the final anneal, metallo-graphic examination showed the material to have an average grain size equivalent to ASTM No. 5 at 100 times (i.e., 0.065 in. diam). Specimens cut from the center and edges of this bar gave no indication of detectable preferred orientation when examined by X-ray diffraction. Samples 1.5 in. long, either 0.625 or 0.750 in. wide, and approximately 0.400 in. thick were machined from this fabricated ingot. The surfaces corresponding to the rolling planes were ground so as to be parallel. The samples were chemically polished in a solution of 60 pct nitric acid and 40 pct hydrofluoric acid (48 pct solution) prior to rolling to remove any cold work introduced in the machining operations. Rolling was accomplished with a 2-high hand-operated laboratory rolling mill that had 2.72-in.-diam rolls. Prior to operation, the rolls were polished with 600 grit paper, cleaned with acetone, and then soaked in a container of liquid nitrogen for several hours. The samples were also soaked in liquid nitrogen prior to rolling and were recooled between each pass. While some slight heating of the samples occurred during rolling, this procedure maintained the sample temperature well below 0°C at all times. The samples were rolled unidirectionally, and the rolling plane surfaces were not inverted during any phase of the operation. The draft per pass averaged between 0.010 to 0.012 in. After 96 or 97 pct reduction the draft was reduced to 0.001 to 0.002 in. per pass. Samples were rolled to various reductions in thickness between 43 and 99.5 pct.

Jan 1, 1969

By C. E. Richards, C. N. Reid

The influence of preferred orientation on the incidence of defbrtnation tuinning has been studied. High-purity iron with almost vandonz grain orientation was cotnpared uitll iron of the sa)ne grain size and composilion lza,ing a strong (110) fiber texture. As expected from published work on single crgslfls, /he ))lean stress for the onset of luitzning-, and the l,olu)nt. fraclion of twinned nzaterial obserlled in lension differed fron the 1-a1ue.s it2 co?nPression for tnolerial with a slrong texlure. The llinning stress of "rctndorrl " )zalerial did not 17ary with the sense of the aPPlied unin.via1 stress, but sirprisinglg the incidence of 1c)i)zning- was about three 1i))zes greater ill conzp?&apos;ession Illon in lension. These results (Ire attributed entirely to ovienbation and may be nderslood in ler?ns of the shear slresses acting on the allowed twinning syster)is. J. HE twins most commonly formed in bcc metals may be described as regions of the crystal in which a particular set of (112) planes is homogeneously sheared by 0.707 in the appropriate ( 111) direction. A similar twin-related crystal could be produced by a shear of 1.414 in the reverse (111) direction but twinning by this large displacement has never been reported. Thus, twinning is unidirectional and a shear stress which produces twinning does not do so when its sense is reversed. The sense of a shear Stress is reversed when the loading is changed from tension to compression, or vice versa. Consequently, for a given orientation of a crystal relative to a uniaxial stress, only a fraction of the twelve (112) twinning systems are geometrically capable of operating in tension, and the remaining systems may operate only in compression. Therefore, when twinning is involved, there are expected to be differences in behavior between crystals tested in uniaxial tension and those tested in compression. This has been verified experimentally by Reid et 01.&apos; and Sherwood el al.,&apos; although a critical stress criterion was not encountered. Furthermore, twinning stresses in colmbium," tungten, tantalum,&apos; irn,&apos; i-Fe,\ nd molybdenum7 single crystals have been shown to depend critically on orientation, although again twinning did not occur at a critical value of the macroscopic shear stress. However, when twinning occurs, it generally does so on the most highly stressed systems, 1--4&apos;6&apos;8&apos;9 implying that the stress level does have some relevance to twin formation. In view of the large orientation dependence of twinning in bee single crystals, it might be expected that such an effect would be present in poly crystalline material which possesses a recrystallisation texture. Indeed, riestner" showed that the twinning stress in tension is very orientation-sensitive it1 <&apos;grain-oriented, silicon-iron;" this material possessed a very strong t c m^ii a nnr x_____k . i-_ii__ ri_______j. _x r»i_._:__i preferred orientation obtained by secondary recrystallisation. Reid et a/.&apos; observed a marked difference in the tensile and compressive yield stresses of polycrys-talline columbium which was rationalised in terms of the effect of a preferred orientation on twinning. No other such illformation is known to the authors. Several investigations of twinning in polycrystalline bcc metals have been reported in which the possible existence of a preferred orientation was not even mentioned. It is the purpose of this paper to show that there is a strong effect of texture on twinning in polycrystalline iron, and to poilt out the difficulty in eliminating preferred orientation in recrystallised metals. 1. EXPERIMENTAL METHOD Material and Specimen Preparation. Low-carbon, high-purity iron was obtained from the National Physical Laboratory in the form of $-in. diam rod which had been cold-swaged from a diam of 1 in. The composition of the material is given in Table I. The as-received bar was cold-swaged directly to 0.185 in. diam from which cylindrical tensile and compression specimens were machined. Specimen geometry is illustrated in Fig. 1. The gage length was 0.30 in. long and 0.10 in. diam; it should be noted that, apart from the extra heads which are necessary for tensile loading, the geometry and dimensions of the two types of specimen are identical. The specimens were heat treated either by sequence A or B outlined in Table 11. The essential difference between these two treatments is that in one case the material was repeatedly cycled through the y- to a-phase change in order to produce grains of almost random orientation ("random" iron)

Jan 1, 1969

By Ralph H. Kress

Haulage trucks designed expressly for mine service were introduced about 35 years ago. The first models to arrive on the scene hauled about 15 tons and easily outperformed the modified highway trucks that had been used for mine haulage up to that time. The specialized trucks obviously had a greater potential for mining operations, and once this fact was recognized, competitive pressures in both the mining and truck-manufacturing industries drove the new vehicles toward ever more efficient performance. Many changes in truck design resulted, but not all were accepted for a variety of reasons. Often poor design was the reason, but in many cases progressive and promising designs were abandoned only because the truck manufacturer lacked the courage, tenacity, or vision to capitalize on them. There were also many cases where a promising new design was abandoned because the prototypes were initiated in a mine where personnel lacked the patience to prove them. Truck Startup Takes A Team Ally new truck that is to succeed and contribute to lower mine haulage costs must be well conceived and launched with good teamwork on the part of all concerned. Naturally, it is essential to have a good overall truck design at the outset to make this teamwork worthwhile. The design must include the selection of the right components manufactured by suppliers that can each do justice to their particular part of the new truck. Last, but by no means least, the mine operator where the truck is to be launched must be enthusiastic in his efforts to see the new vehicle through to successful operation.

Jan 1, 1971

By F. L. Wolf, G. E. Alderson

The reclaiming of nietallics from slag and sweepings is of vital interest to every brass-foundry man, but the first cost and interest on the investment often make it prohibitive for the small foundry to enjoy the benefits of a complete concentrating unit. However, the cost of such an installation and the cost of operation are often exaggerated to such an extent that foundries dispose of their copper-bearing refuse at a great loss. This paper gives the actual costs and the returns obtained in the reclaiming plant used at the Ohio Brass Co. This plant is divided into two sections—the preliminary treatment and concentrating departments. In the former occur the screening, crushing, and hand-picking operations whereby rnost of the non-valuable and heavy metallic constituents are removed. The concentrating equipment consists of an elevator, bins, jigs, ball-mill, arid table. Preliminary Treatment The preliminary treatment varies, depending on whether the material is slag, foundry-floor sweepings, or brass casting-floor sweepings. Material of the first class is produced from the operation of six Steele-Harvey and six Schwartz furnaces; the second class represents sweepings from an average of thirty-five. molding floors per day: and the last is collected around the sprue cutters and knockout bench. Treatment of Slag The slag, analyzing 30 to 40 per cent. copper, is passed over a 3-mesh yard screen (Tyler standard mesh) stretched on an inclined rectangular frame supported by adjustable legs and equipped with a hopper which directs the undersize into a barrow. The fine product, running about 20 per cent. copper and representing 20 to 25 per cent. of the total slag, goes direct to a storage bin for concentrator frrd. The oversize is charged into a No. 2 Hill cinder crusher, which takes 175 to 200 11,. to a charge and crushes two charges per hour through &apos;a No. 4 riddle. It

Jan 1, 1921

By J. L. Lamon, W. Crafts

Study with the electron microscope of the carbides in vanadium-chromium-molybdenum steels after tempering up to 1000 hr at 600 teelsto 1400°F confirmed that alloy carbides are formed at the secondary hardening temperature by decomposition of the plate-like iron carbides. It was also demonstrated that vanadium carbide persists as much smaller particles than do chromium- or molybdenum-bearing carbides. Conditions conducive to the formation of fine vanadium carbides are indicated to be favorable for high temperature strength. IN order to determine the effects of long exposure at high temperatures on vanadium-bearing steels, a survey has been made of their hardness and carbide structure. The behavior of carbides in the tempering of martensite was studied by X ray diffraction and electron microscope examination of electrolytically extracted residues in greater detail than in an earlier investigation.&apos; A group of steels with about 0.25 pct carbon and containing chromium up to 5 pct, molybdenum up to 1 pct, and vanadium up to 1 pct was either quenched or annealed and then tempered for peri- ods of up to 1000 hr at 600° to 1400°F. Their hardness after tempering agreed with the Hollomon and Jaffe2 relation of equivalency of time and temperature, and with the degree of secondary hardening predicted from composition. Further, the appearance of the carbides indicated that the time and temperature equivalency relation was also applicable to the degree of carbide development. The mechanism of the tempering of martensite was demonstrated more clearly than in the earlier study. It was confirmed that carbides develop from martensite as poorly defined plates of the Fe2C type carbide followed by thickening of the plates with a transition to Fe,C. Finally the plate structure deteriorates into a lacy mass from which alloy carbides emerge as chunky particles that grow slowly with further increase of tempering temperature. Vanadium carbide derived from tempered martensite was found in characteristically small particles that tended to grow very slowly. The addition of chromium or molybdenum to a steel with predominant vanadium carbide tended to introduce carbides of a chromium- or molybdenum-bearing type having a somewhat larger particle size. The Cr,C, type carbide particles were larger than those of the V,C, type and somewhat smaller than the carbides in steel containing both Cr,C and M,C. Similar observations were made of the carbides in the pear lite of annealed steels. It appeared that high temperature properties would be benefited by the retention of the fine carbide particles that result from a balance of composition and heat-treatment designed to produce a maximum amount of vanadium carbide. Procedure: The steels used in the investigation were made with a base of Armco iron in a high-frequency furnace at the Union Carbide and Carbon Research Labs., Inc. The steels were deoxidized

Jan 1, 1951