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By S. J. McCarroll

Fig. 1—The belt shown at right carries filter cake to mixing station over calciner. Crude ore conveyors appear in right background. THE Three Kids mine, some six miles east of Henderson, Nev., is in a typical southwest desert area, with high dry summer heat and cool to cold winter seasons. The manganese deposit was located during World War I.' During this period 15,000 to 20,000 tons of ore assaying up to 41 pct manganese were shipped. Interest in the deposit was not revived until the middle thirties, when experiments on the ore were initiated. Test work indicated possible recovery of only 70 pct by flotation, but in 1941 additional work was done at the Boulder City pilot plant of the U. S. Bureau of Mines and also by M. A. Hanna Co. As a result, the Manganese Ore Co. was formed and a plant utilizing the SO2 process was constructed. Numerous operation difficulties ensued, and the plant. was closed when the manganese situation in the country eased. In 1949 Hewitt S. West initiated negotiations to acquire the plant. In 1951 Manganese, Inc., was formed and contract entered into with the General Services Administration to supply 27 million units of metallurgical grade manganese in the form of nodules to the national stockpile. A second contract was made to upgrade 285,000 tons of stockpile ore. Test work was undertaken by the Southwestern Engineering Co. and likewise by the Boulder City pilot plant at the U. S. Bureau of Mines. Results obtained indicated the commercial feasibility of the flotation process. Construction of the plant, which is shown in Figs. 1 and 2, was started in June 1951, and operations on a break-in basis began in September 1952. Apart from the usual starting difficulties two major disasters caused serious setbacks, one a kiln failure in February 1953, and the other a fire that destroyed the flotation building in June of the same year. The nodulizing section of the plant resumed operation in November, and the flotation section in January 1954. The ore minerals are chiefly wad,* with minor amounts of psilomelane, and occur in sedimentary beds of volcanic tuff. The ore is overlain with beds of gypsum which outcrop or may be covered with surface gravel. Intermediate beds of red and white tuff occur frequently with lenses of red and green jasper and stringers of gypsum and calcite. Small amounts of iron are present; lead content averages about 1.0 pct and minute amounts of copper and zinc are found. Barite, celestite, and bentonite are present. Since these are made up of minute asicular crystals, moisture content is very high, averaging about 18 pct. Ore reserves have been estimated at 3 million tons averaging 18 pct Mn2 and up to 5 million tons after grade is dropped to 10 pct Mn. A good part of the orebody was stripped of overburden by the previous operating company . Approximately 50 pct of the ore, representing more than 60 pct of the manganese, can be mined by open-cut methods. A system for underground min- . ing has not yet been decided on. Open-cut mining with benches of 20 ft has proved satisfactory. Although the ore is soft and appears dry and dusty it has a certain resilience, probably due to the porosity and moisture which makes drilling and fragmentation difficult. Wagon drills have been abandoned in favor of the Joy 225-A rotary drill which will put down a 43/4 -in. hole at the rate of 2 ft per min. Holes are spaced in a pattern with 8 to 9-ft centers. Forty percent powder has been used, but better breaking to 2-ft size is obtained with low velocity bag powder of 30 pct strength. Loading is done with one 21/2-yd shovel, and cleanup follows with one D-7 bulldozer. The ore is hauled with Euclid trucks about 1000 ft from the pit to a blending pile, where the daily mine production is spread in layers by bulldozing until approximately one month's mill feed is accumulated. A new pile is then started and mill feed is drawn from the first pile by one 13/4-yd shovel and Euclid trucks, with a haul of approximately 500 ft. Mining is performed by an independent contractor with engineering and supervision by the company staff. Early test work indicated that the manganese could be floated with soap, a wetting agent, and fuel oil to give a recovery of better than 75 pct with a grade of 43 pct Mn. The concentrate when nodulized with coke would upgrade to 46 pct Mn or over, and the lead volatilized to 0.6 pct residual.

Jan 1, 1955

By W. R. Kube, W. H. Oppelt

Five low-rank coals, including two lignites, a steam-dried lignite, and two subbituminous coals, were carbonized at 940°F, in a bench-scale carbon-ize~ with a nitrogen and hydrogen atmosphere, or both, using pressures from atmospheric to 1000 psig. Coking characteristics were observed, and yields, composition, and distribution of products from thermal degradation were determined. The influence of pressure on yields from coal carbonization, as well as on modification of coke structure, has long been of interest. One of the first attempts at pressure carbonization was that of Capps and Hulett1 who in 1917 investigated the effects of pressure ranging from 1 to 20 atm on yields at 600° C for bituminous and subbituminous coals. High-pressure carbonization has not been used commercially for coke production. More recently, carbonization at relatively high pressure in the presence of catalysts and hydrogen has been proposed for the production of metallurgical-grade coke from low-rank fuels.' Another approach involves carbonization at high pressure in a hydrogen atmosphere to produce heating gas and a solid residue that is gasified with oxygen and steam to produce the hydrogen required for hydro-genation 3-4 or, alternatively, is burned to raise steam. Recently it has been proposed to carbonize coal under pressure, burn the gas in a gas turbine to produce power, and use the solid residue as fuel in a conventional steam-power plant.5 Our primary interest in pressure carbonization stems from an attempt to determine the influence of pressure on the coking characteristics of lignite to design a high-pressure slagging gasifier using oxygen and steam. At atmospheric pressure, North Dakota lignite is generally considered to be non-coking. However, lignite and subbituminous coals have been reported to show increased agglomeration properties in steam, hydrogen, or carbon-dioxide atmospheres at elevated pressures, although no strong coke was formed at pressures up to 3000 psig.6 If lignite would agglomerate seriously under gasification conditions, proper flow of charge through the gasifier could not be maintained without installing a mechanical stirring device or destroying the caking properties of the coal before gasification. Data on the caking properties of lignite or other gasification fuel, as well as on yields of products under conditions simulating the carbonization process in the pregasification zone, were required in the preliminary design stage of the gasifier. No information was available concerning the influence of pressure and type of atmosphere on North Dakota lignite. To provide this basic data, the U. S. Bureau of Mines investigated the effect of pressure and a hydrogen or nitrogen-rich atmosphere on the yield of products and on coke formation in a bench-scale pressure carbon-izer. SCOPE OF INVESTIGATION Five coals were tested at 940°F in nitrogen and hydrogen atmospheres, at pressures from atmospheric to 1000 psig. Yields and composition of solid, liquid, and gaseous products were determined as a function of the process pressure, and attention was given to the physical appearance of the solid residue. Percentage of carbon converted, sulfur retained in char, and thermal balances were calculated. In this paper, primary consideration is given to carbonization in a nitrogen atmosphere and to data concerning the solid phase. The effect of pressure in hydrogen is presented in detail for only one lignite. More complete information on the response of the other coals in a hydrogen atmosphere will be published in a USBM Report of Investigations. DESCRIPTION OF EQUIPMENT Details of the bench-scale pressure carbonizer are given in Fig. 1. The body of the carbonizer was

Jan 1, 1961

By Mark J. Klein, A. H. Clauer

The Snoek peak induced by solute nitrogen in chromium was studied. A rapid quenching rate is required to maintain nitvogen in solution in sufficient concentrations to be detectable by internal-friction techniques. An expression for the tertninal solubility, c, of nitrogen in weight pevcent is log,,c = -4130/T + 1.05. The diffusion of nitrogen in zhromium determined by internal friction is SNOEK&apos;S theory1 of the stress-induced ordering of interstitial solute elements in the bcc lattice has been used to deduce information on the behavior of interstitials in a number of bcc metals. Under the proper conditions, internal-friction studies of "Snoek peaks" permit a determination of diffu-sivities, solubilities, and precipitation kinetics of interstitial elements. Internal-friction studies of this type have been reviewed by Nowick,&apos; Wert,= Powers and Doyle,4 and others. The behavior of interstitial atoms in Group V-A metals and iron has been rather extensively studied in this way, and results are reasonably self-consistent. However, the study of interstitial elements using internal friction in the Group VI-A metals—tungsten, molybdenum, and chromium—is much less clear. Thus far, internal-friction studies have revealed a peak in tungsten tentatively attributed to carbon5 and a peak in chromium associated with nitrogen. In chromium, Bungardt and Preisendanz6 observed a peak at 180°C for a frequency of -1 cps which they associated with nitrogen. More recently, de Morton7 determined the diffusivity of nitrogen in chromium from studies of a nitrogen peak at 160°c for a frequency of -1 cps and from elastic aftereffect measurements. The objective of this investigation was to study the nitrogen peak in greater detail, to obtain information on the behavior of nitrogen in chromium. This information is of interest because nitrogen is known to have an embrittling effect on chromium. However, efforts to deduce the role played by nitrogen in influencing the mechanical behavior of chromium are handicapped by lack of information on its behavior and distribution in chromium. EXPERIMENTAL PROCEDURES The chromium wires used in these studies were drawn from arc-melted iodide-chromium stock that had been extruded, swaged, and centerless-ground to 0.25 in.-diameter rods. Internal-friction measurements were made in an evacuated torsional pendulum similar to the one described by Kê,8 utilizing wires 0.048 to 0.030 in. in diameter. A series of analyses of the fabricated wires showed that the major impurities were approximately 100 ppm Fe and a combined interstitial content of 50 to 100 ppm C, O, N, and H. A typical interstitial analysis of the wires after processing is 15 ppm C, 2 ppm H, 50 ppm O, and < 5 ppm N. Interstitial concentrations were determined by vacuum-fusion, combustion-conductimetric, and micro-Kjeldahl analyses. In general, vacuum-fusion analysis yielded a lower nitrogen concentration than did micro-Kjeldahl analysis. For consistency, the nitrogen concentrations reported in this investigation were determined by the former method. Heat treatments were carried out by encapsulating the specimens in quartz tubes containing argon. The fabricated wires were annealed at 1150°C for 1 hr to produce uniform recrystallized grain size of about 0.015 mm. Nitrogen and carbon were added to the fabricated wires by encapsulating them in a quartz tube containing ammonia or methane, and annealing them at 1150°C for 48 hr. In some instances, carbon was also introduced by coating the wires with carbon black prior to annealing.

Jan 1, 1965

By M. J. Blackburn

The ordering reaction and its effect on dislocation awangements in alloys based on Mg3Cd has been studied by transmission electron microscopy. Ordering occurs by the growth of ordered domains until they impinge, thus producing an antiphase domain boundary (APB) network. In the disordered alloy, slip is confined to the basal plane while in the ordered condition slip occurs on the basal, prism, and pyramidal planes. Superdislocations have been identuied on these planes and their origin and expected spacings are discussed and compared with the observations. The APBs between the dislocation pairs have been resolved and have proved useful in the identification of superdislocations. Observed slip structures are related to the changes in mechanical properties produced by ordering. At high temperatures magnesium and cadmium form a continuous series of disordered hexagonal solid solutions. At lower temperatures three ordering reactions occur; two based on the compositions Mg3Cd and MgCd3 are transformations to DO19-type superlattices while alloys based on the composition MgCd transform to a B19 lattice. The ordering process in alloys near the composition Mg3Cd has been studied by several workers&apos;-4 and it has been shown to be a first-order type transformation. The variation of mechanical properties with the degree of order and temperature have been reported in some detail by Stoloff and Davies.5 It was shown that the disordered alloy was stronger and less ductile than the ordered alloy, the differences being attributed to changes in the number of active slip planes and to the presence of short-range order in the disordered alloy. In this investigation the ordering reaction has been qualitatively followed by transmission electron microscopy. The effect of this ordering transformation on the dislocation arrangements in deformed specimens has been studied in some detail. EXPERIMENTAL Alloys were melted in a stainless-steel crucible under a flux and chill cast; both melting and casting were performed in an argon atmosphere. The ingots were hot-rolled and tensile specimens cut from the resultant strips. Some large-grained specimens were grown by the strain-anneal method. Heat treatment was carried out in inert atmosphere furnaces and in oil baths at the lower temperatures. Thin foils for transmission electron microscopy were produced by electropolishing in a cooled solution (temp = -40°C) containing 50 cu cm of nitric acid and 180 cu cm of methanol at a potential of 10 v. The ordering reaction was followed by transmission electron microscopy, X-ray, and electrical resistivity measurements but this paper will be concerned chiefly with the electron-microscopy results. Much of the work described in this paper was performed on an alloy containing 26.10 at. pct Cd although two of the micrographs are taken from work on a 25.55 at. pct Cd alloy. RESULTS i) The Quenched Solid Solution. Specimens which were quenched from above the critical temperature showed diffuse superlattice reflections. sark-field analysis on these reflections failed to reveal the presence of ordered domains and thus the superlattice reflections were assumed to arise from regions of short-range order. Quenching from temperatures of 400" to 450°C results in the formation of dislocation loops and helical dislocations. Such defects have been observed and analyzed in detail in several hexagonal metals6-&apos; and as the crystallography and interactions of such loops appear similar in this alloy they will not be reported in detail. Fig. 1 shows a high density of loops

Jan 1, 1968

By R. J. Cowan

In the art of cementation a controversy has been going on for years as to whether solid or gaseous carbon is the active agent in carburizing steel. More recently opinion has crystallized into a compromise to the effect that carburizing is a gaseous cementation of steel in the presence of solid carbon. The importance of this theory has a particular bearing upon the work to be described. In the art of carburizing, use is commonly made of such words as &apos;&apos; Catalytic" "atomic," "nascent" and "activated," in an attempt to describe the process. These words are used indiscriminately. For instance, the gas is said to be "catalytic" and carbon to be "nascent," or vice versa. In using these expressions, investigators are confirming the acceptance in their minds of the active conditions of both gas and carbon in the cementation process. Reactions at Metal Surfaces For many years attention has been directed to the consideration of reactions that occur on metal surfaces. When heated under the influence of certain metals a complex gas will be broken down and one of the constituents thereof usually found to be very active chemically in combining with the metal. The process of nitriding is an example of this type of reaction. In this case ammonia (NH3) breaks down at the metal surface and the nitrogen which is thus liberated reacts with the surface to form a nitride whereas molecular nitrogen not so liberated is inert and will not react chemically with the metal. Whatever term or terms may be used to describe such reactions, the fact remains that they are highly important and valuable in their commercial applications. Since the metal surface plays so important a part in these reactions, it is evident that any accumulation of reaction products thereon interferes seriously with the progress of the reactions. For this reason it is desirable to remove all such products as rapidly as they are formed in order to bring active gas into contact with the surface and thus facilitate the desired reactions. This is accomplished best by conducting the operation in a continuous manner, thereby causing the reacting gases to pass over the

Jan 1, 1931

By J. J. Crowe, G. K. Burgess

PAGE, Theories of Allotropy of Iron........... 667 Previous Determinations of A2 and A3 in Iron........ 668 Critical Ranges as Determined by Expansion........ 669 Critical Ranges by Thermoelectric Observations........ Critical Ranges in Terms of Crystalline Structure....... 670 Critical Ranges as Detected by Mechanical Methods...... 670 Critical Ranges as Given by Electrical Resistance....... 670 Critical Ranges as Determined Magnetically........ 671 Critical Ranges as Determined Calorimetrically....... 673 Critical Ranges from Heating and Cooling Curves....... 673 Summary of Previous Determinations of A3 and A2 in Pure Iron.... 675 The Present Investigation............ 678 The Experimental Method........... 678 Arrangement of the Apparatus.......... 679 The Measurements of Temperaturc.......... 679 The Measurement of Time........... 679 The Thermocouples............ 681 Temperature Scale and Calibration of Thermocouples...... 681 Mounting of Thermocouples........... 682 The Differential Measurements.......... 682 The Furnace and Accessories........... 682 The Vacuum Pump............ 683 The Energy Supply............683 The Rheostat.............. 683 The Materials Investigated............ 685 Sources of Iron............. 685 Preparation and Description of Samples......... 685 Chemical Analyses............. 687 The Experimental Data............. 688 The Preliminary Observations.......... 688 The Definite Series............691 Description of the Plotted Curves........... 694 The Effect of Plotting at Different Intervals........ 694 The Curves of the Definite Series.......... 695 . The Curves of the Preliminary Series......... The Existence of Other Critical Ranges......... 696 The Microscopical Examination........... 697 Summary and Conclusions............ 698 Bibliography............... 701 Plates................ 704

Jan 1, 1914

By Harold V. Pearce

THE plant of the Boston and Colorado Smelting Company, at Argo, Colo., has not received special notice in technical or scientific publications for some time past. Dr. Peters&apos; described the development of the reverberatory furnace used at Argo, but he dealt almost exclusively with the increase in size and the corresponding gain in capacity. Since the time of his description the changed conditions have led to a modification of the former smelting practice, and this modification forms the basis of the present paper. No claim is made for personal originality in connection with these improvements, in fact, I expressly disclaim any such intention, since nearly every successful step in the development of the process now used at Argo has been suggested by Dr. Richard Pearce at various times during the past 12 years, though perhaps in a different manner of exe¬cution. These changes would undoubtedly have been made before, were it not that many problems of greater metallurgical importance and momentarily of more urgent need thus had to be solved. At the time to which I refer, economic conditions demanded mainly an increase of capacity. Moreover, the determination and placing of exact losses in smelting were not so imperative as they now are. About three years ago, the time of my appointment to the - management of. the works of the Boston and Colorado Smelting Co., the great decline in the price of copper occurred; the Argo plant, in common with all others engaged in the copper-industry, lost heavily in the market-value of the stocks on hand of ore and furnace-products. Added to this was the unfortunate circumstance that the copper-supply of the Argo plant from general sources began to dwindle; in fact, it almost entirely ceased. About this time the smelting-plants in the west-

Mar 1, 1905

By Edward Steidle

NOTHING stands still. We go forward or backward. As a distinct group of educators, our immediate concern is with the preparation of young men and women for participation in the mineral industries on a professional basis. I am going to have the temerity to state emphatically that for a long time as a whole we have remained in a static position. We are now confronted with the choice of advancement or retrogression. We are still in a position to make a choice and our ultimate fate will depend upon the course of action we elect to follow. Wars always have been great catalysts and social and scientific advancements mature under military stimulus in a fraction of the time normally consumed. We are now in the midst of a conflict whose duration or final effects cannot be estimated. However, the record already demonstrates that this war as never before is a mineral war in which incentive, strategy, and outcome are directly related to the distribution and utilization of minerals. The inevitable result is a drain upon our resources now reflected in our record production and a growing public consciousness of the importance of minerals in the role of human affairs. It is evident that we must seek new mineral frontiers beyond the borders of the United States and take every other precaution to forestall a "have-not" position in the world.

Jan 1, 1943

By E. J. Roberts

Hydrometallurgical operations and many processes in the chemical industry require the separation of dissolved material from solids. One of the decisions which has to be made in designing a plant for such an operation is how this separation shall be accomplished. It is the purpose of this paper to set forth the general principles involved in such separations in order that the optimum flowsheet and equipment may be selected. Countercurrent Decantation with thickeners, or C.C.D. as it is generally known, is, of course, only one way of separating dissolved material from undis-solved. Other unit operations such as Countercurrent Filtration (C.C.F.) and Countercurrent Classification (C.C.C.) are widely used. The same principles may also be applied to centrifugally assisted decantation, as with nozzle or solid bowl type centrifugals, and to separations using screens, either gravity or centrifugal type. Another form of countercurrent washing applicable to granular materials is practiced in continuous ion exchange operations. All of these operations depend on two basic principles, or combinations of these principles. In C.C.D. type operations the solids are concentrated into the smallest possible volume of liquor in each step and the resulting pulp then mixed with a more dilute liquor and recon cen trated, the liquors separated from the more dilute stages being advanced to dilute the more concentrated liquor in a previous solids concentrate, as is shown in Fig. I. This will be termed the solids concentration-solution dilution principle, or, for short, solution dilution. countercurrent decantation- type circuits. The second principle used in washing solubles from solids is that of solids concentration-solution displacement. Here the solids are restrained from movement while wash solution is forced through the interstices in a counter-current fashion, as in Fig. 2. In the simplest application of this principle, once the "cake" has been formed, it is not disturbed until the operation is completed. In continuous filters, the cake is merely moved under the successive washes and the displaced filtrates removed as in Fig. 4. Often, however, after one or two displacement washes, a repulping step is introduced, the cake reformed and washed again. This procedure brings the solution dilution principle into play as well as the solution displacement principle. Merely thickening ahead of filtration, which is often operationally advantageous, does not hybridize he circuit. It still can be a pure solution displacement operation from the washing

Jan 1, 1961

By M. L. Samuels

In a recent paper the authors1 discussed the reactions to tempering of hypereutectoid steels quenched from 1000&apos; C., as revealed by studies on changes in hardness, electrical resistivity, coercive force and metallo-graphic charaeteristics. The data, which related specifically to tempering temperatures of 550°, 650" and 710" C., for varying periods up to 125 hr., permitted the following important conclusions: 1. The rate of softening of the steels at the temperatures examined differed markedly although the steels were of "similar" general chemical composition. 2. The temperature of maximum softening spread over a range from about 650" C, to the eutectoid inversion. 3. The divergence in behavior was ascribed, at least in part, to a profound difference in tendency to graphitization on annealing at sub-eutectoid temperatures. 4. The mode of spheroidization of the cementite, as revealed by metallographic studies, also varied considerably in the different steels. The purpose of the present investigation is to pursue further the study of the difference in behavior of the steels on tempering over a wider range of temperature. To this end the following features have received particular attention: 1. Comparison of temperatures of maximum softening of the various steels when the quenched samples were tempered for 5 hi. and for 75 hr. followed by furnace cooling and by quenching in water. 2. Deterrnirlation of the effect of quenching temperature, prior to tempering, on the rate of softening. Quenching temperatures of 1000" and 850" C. were used in the investigations. 3. Examination of the microstructures obtained with the various heat-treatments employed.

Jan 1, 1938

By M. L. Samuels

In a recent paper the authors1 discussed the reactions to tempering of hypereutectoid steels quenched from 1000&apos; C., as revealed by studies on changes in hardness, electrical resistivity, coercive force and metallo-graphic charaeteristics. The data, which related specifically to tempering temperatures of 550°, 650" and 710" C., for varying periods up to 125 hr., permitted the following important conclusions: 1. The rate of softening of the steels at the temperatures examined differed markedly although the steels were of "similar" general chemical composition. 2. The temperature of maximum softening spread over a range from about 650" C, to the eutectoid inversion. 3. The divergence in behavior was ascribed, at least in part, to a profound difference in tendency to graphitization on annealing at sub-eutectoid temperatures. 4. The mode of spheroidization of the cementite, as revealed by metallographic studies, also varied considerably in the different steels. The purpose of the present investigation is to pursue further the study of the difference in behavior of the steels on tempering over a wider range of temperature. To this end the following features have received particular attention: 1. Comparison of temperatures of maximum softening of the various steels when the quenched samples were tempered for 5 hi. and for 75 hr. followed by furnace cooling and by quenching in water. 2. Deterrnirlation of the effect of quenching temperature, prior to tempering, on the rate of softening. Quenching temperatures of 1000" and 850" C. were used in the investigations. 3. Examination of the microstructures obtained with the various heat-treatments employed.

Jan 1, 1938

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

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

Jan 1, 1970

By D. J. McAdam, R. W. Mebs

THE technical cohesive strength of a metal means, not the interatomic forces, but the technically estimated resistance to fracture. An example of such resistance to fracture is the so-called "true" breaking stress of a tension-test specimen, the breaking load divided by the sectional area at fracture, According to the prevalent view,8 the technical cohesive strength of a metal, in any particular state as regards prior mechanical treatment and heat-treatment, is determined by a limiting value of the algebraically greatest principal stress.? In two previous papers by one of the authors,14,15 evidence was presented to indicate that the technical cohesive strength of a metal cannot be represented by a single stress value, but that it comprises an infinite number of values corresponding to the infinite number of possible combinations of the principal stresses. The technical cohesive strength thus comprises an infinite number of technical cohesion limits, each representing a specific stress combination at fracture. In the same papers, evidence was presented that plastic extension causes a continuous increase in the technical cohesive strength. Such a variation is not in accordance with the prevalent view that plastic extension first increases, then decreases, the technical cohesive strength.8 In this paper, as in the previous papers,14,15 the algebraically greatest principal stress will be designated S1, the least principal stress will be designated S3, and the intermediate principal stress will be designated S2. The technical cohesion limit under polarsymmetric tension (S1 = S2 = S3) will be termed the "disruptive stress," and the cohesion limit under unidirectional tension (S2 and S3 = o) will be termed the "severing stress." 14,15 The yield strength of a metal, according to the generally accepted theory, may be represented by a constant value of the shearing-strain energy, and hence by a constant value of the sum of the squares of the three principal stress differences. In the second of the previous papers,"&apos; however, evidence was presented that the limiting yield stress decreases with increase in the volume stress, the algebraic average of the three principal stresses.* This influence of the volume stress becomes prominent only in the field of triaxial tension. In the two previous papers, 14,15 diagrams were presented to show the variation of the technical cohesion limit, yield stress, and ultimate stress with the combination of principal stresses and with plastic extension. Those diagrams and the resultant conclusions are derived from experimental data obtained by other investigators. They are based chiefly on results of tension tests

Jan 1, 1943

By H. Warlimont

In recent investigations of the microstructure and crystallographic features of martensite by electgon microscopy,&apos;, &apos;9 thin films (about 50 to l000A in thickness) have been used as specimens. It was found by W. Pitschl ,2 that the lattice relationships between the supefiaturated fcc matrix and the martensite formed in thin films were different from those observed after transformation in bulk material. Corresponding differences are to be expected in nucleation and in reaction kinetics in the formation of martensite in thin films, but have not yet been reported. In the course of the preparation of thin films of iron-base alloys for investigation by electron microscopy it was observed that films of alloys, whose martensite formation temperature (Ms) was below room temperature, formed martensite during the electrochemical thinning process conducted at or above room temperature. The chemical compositions and approximate bulkM,-temperatures of the three materials on which this observation was made are listed in Table I. The alloys were vacuum-melted from high-purity metals. It can be seen that the effect occurred in alloys that contain only sub-stitutional components as well as in alloys that contain carbon as an interstitial solute. The appearance of martensite formed in this manner is illustrated in Fig. 1, which represents a polished film of alloy No. 1 at XI0 magnification. The martensite plates can be recognized by their surface relief effects. In addition, heavily strained (owrinkledo) austenite can be seen near the edges of the film where it is thinnest and transformation stresses or strains can be accommodated by plastic deformation most easily. The nonuniform distribution of transformed regions is probably due to the varying thickness of the film. These observations were common to all materials listed in Table I. Because effects of supercooling or straining can

Jan 1, 1962

By B. W. Christ, P. M. Giles

Mossbauer effect measurewents have been made on I-mil-thick foils of commercial 1 wt pct C steel and Fe-2 wt pct C alloy. The experimental method required about 3 to 5 vol pct of a phase in the nzultiphase steel sample for detection. Room-temperature Md&apos;ssbauer patterns obtained on austenitized atid quenched samples exhibit fifteen, and possibly twenty-one, lines. A sharp parama&tetic singlet and a quadrupole doublet, poorly resolued from the singlet, are attributed to austenite. Remaining lines are due to tnartensite. Accurate evaluation of austenite line paranzeters is not feasible if significant amounts of other phases such as carbides or martensite occur simultaneously with austenite. This is demonstrated by comparison of hyperfine interactions determined for austenite in multiphase high-carbon samples with those reported for Fe-C austenite in a nearly 100 pct austenitic sanple. Lines from carbides are incompletely resolced from austenite lines, as demonstrated by comparison of austenite line positions with carbide line positions calculated frow published values of hyperfine interactions. One martensite line overlaps an austenite line in the pattern for commercial 1 wt pct C steel. Results of this study suggest that the usefulness of e M6ssbauer pectroscopy for quantitatizle analysis of austenite in bulk samples of quenched and tempered high-carbon steels is restricted by poor resolution. Use of Mossbauer spectroscopy for phase identification and for evaluation of atomic and electronic structures appears quite feasible, however, The Mijssbauer effect has been widely discssed,&apos;-and e Mossbauer effect measurements have been reported on materials of metallurgical interest.7"20 In particular, it has been proposed that e Mossbauer patterns of commercial steels and laboratory-made Fe-C alloys, in the quenched condition, are composed of lines originating in two phases, Fe-C austenite and Fe-C martenite.-&apos; Evidence accumulated in this study demonstrates that three absorption lines found in the central region of the e Mossbauer pattern obtained on quenched steels are attributable to retained austenite. This interpretation is supported by parallel decreases in the intensity of these three lines caused by subambient cooling of commercial 1 wt pct C steel samples after water quenching to room temperature. A second result of this study is to clarify effects of line resolution and sensitivity in the Mossbauer patterns of multiphase steels on the accu- rate determination of austenite line parameters. Experimental line widths (full width at half height) are generally 1.5 to 3 times larger than the natural line width of 0.19 mm per sec. At least two lines, and sometimes more, from a single phase such as cementite (Fe3C), other carbides, martensite, and austenite fall in the energy band, i0.85 mm per sec. hhis band width is employed simply for convenient reference. It represents approximately the energy interval between the + 112 to 112 transitions in ferrite and is expressed as the velocity needed to Doppler shift a 14.4kev 7 ray to the aforementioned ferrite energy levels. This energy band is referred to as "the central region of the Mossbauer atttern:: in this paper. Hence, due to the large number of lines from different phases in a multiphase steel falling in a relatively narrow energy band, absorption lines from the different phases may overlap. Analysis of available data, presented below, indicates that this occurs to a significant extent for phases which commonly occur in quenched and quenched and tempered high-carbon steels. One consequence of limited resolution in the Mdssbauer patterns from multiphase steels is difficulty in accurate determination of such line parameters as position, width, and intensity. In fact, it appears that quantitative analysis for retained austenite in quenched and tempered high-carbon steels is not practical with the present experimental method. Line resolution is influenced to some extent by sensitivity. We point out below that atom or volume fractions of less than about 3 to 5 pct are not detected by the present experimental method. Thus, the presence of a multiplicity of phases does not always lead to impaired resolution. Finally, we report in this paper room-temperature MGssbauer parameters determined for austenite in a freshly quenched, commercial 1 wt pct C steel and in a freshly quenched laboratory heat of an Fe-2 wt pct C alloy. These parameters are compared with others reported in the literature. Three types of hyperfine interactions are detectable in a Mossbauer effect measurement: isomer shift, quadrupole interaction, and magnetic dipole interaction. These interactions are evidenced by one, two, and six line patterns, respectively.&apos;-4 More than one type of interaction has been reported in certain metallurgical phases thus far studied by this method. Isomer shift is the experimentally measured displacement of line position from some arbitrarily defined reference position. In the case of a multiline pattern, isomer shift is given by the displacement of the centroid (center of gravity) of that pattern from the reference position. All isomer shifts measured at finite temperatures contain a second-order Doppler effect characteristic of that temperature. The isomer shift is related to the total s electron density at the nucleus, becoming more negative with increasing s electron density. The first-order quadrupole effect arises from the interaction between the nuclear quadrupole moment and any axially symmetric electric field gradient in

Jan 1, 1969

By D. T. Peters

To better understand the Co-Mo interaction in maraging steels, the precipitate reversion behavior of two alloys closely related to the 18 Ni-250 grade has been studied. The reversion experiment is carried out by aging at a low temperature, T1 (preaging temperature), for some period of time and then aging at a higher temperature, T2, called the reversion tetnperature. Aging and reversion tmzperatures used were in the range 679 to 1050°F. Reversion is marked by an increase in resistivity and a decrease in hardness. It teas found that some reversion took place at all cornbinations of T1 and T2 but that pronounced rei3ersion resulfed for an Fe-18 pct Ni-5 pct Mo-8 pcl Co alloy exposed to preagz&apos;ng&apos; temperatures at or below 850°F. A similar alloy not containing cobalt exhibited no such discontinuity in reversion behazior with TI. The results are interpreted in terms of two types of molybdenum-rich precipitates, one nucleated at dislocations and the other in the matrix. The increased supersaturation resulting from the presence oj- cobalt evidently provides sufficient energy to nucleate the matrix phase. Anotnalies in the temperature dependence of the resistivity decrease and peak hurdness in the Fe-18 pct Ni-5 pct Mo-8 pct Co alloy as a junction of aging temperature noted in an earlier study are explained on this basis. SEVERAL investigatorsl-9 have studied the details of the precipitation-hardening reaction of the maraging steels. The great majority of conclusions have been drawn from transmission electron microscopy. Although much has been learned by this method, alternate techniques which will give information earlier in the aging process are needed for a more complete understanding of the series of precipitation steps probably involved in the aging reaction. Peters and cupp&apos;&apos; have recently shown, by measurements of the kinetics of precipitation, that an 18 pct Ni-Co-Mo maraging steel shows a distinct change in precipitation behavior at aging temperatures below 850°F. The results were interpreted in terms of the existence of two types of precipitates at low aging temperatures. Only one precipitate type is present at higher aging temperatures. The low-temperature precipitate was found to be a particularly effective strengthener. No evidence for the additional low-temperature precipitate was found in the 18 pct Ni-5 pct Mo ternary alloy. The object of this investigation was to study further the characteristics of the two types of precipitates and the interaction between cobalt and molybdenum characterizing this system. Some preliminary experiments indicated a difference in reversion behavior in specimens aged above and below 850°F. This suggested that knowledge of the reversion behavior might shed some light on these two questions. Hardy and Heal11 have defined reversion as the loss of hardness, electrical conductivity, or other properties caused by heating for a very short time at temperatures above the aging temperature. The precipitate reversion referred to in this paper is entirely distinct from "austenite reversion" often referred to in maraging steel literature. This latter phenomenon could more accurately be called the diffusion-controlled formation of austenite. Much has been written about reversion in several of the aluminum-base binary and ternary alloys.&apos;2-16 Two explanations for reversion have been advanced. Based on the classical nucleation theory of Volmer-Weber-Becker-Doring,17 Konobeevski18 suggested that reversion was due to the size of the precipitate particles

Jan 1, 1968

By Ray W. Carpenter, Ralph Hultgren, Raymond L. Orr

Heats of formation of Au-Zn alloys of compositions Au3Zn and AuZn were rneasured at several temperatures by liquid tin solution calorimetry. The data for Au3Zn show that much smaller heat and entropy effects are associated with the a2- a1 transformation than for the a1 - a transformation. This result is consistent with reported X-ray diffraction studies which indicate that the a1 and (a2 phases have ordered structures which only slightly differ, whereas the a phase is disordered. The heat of formation of AuZn(ß&apos;) does not change significantly between 322o and 8000 K, confirming that the ß&apos; alloy remains highly ordered to high temperatures. Not many measurements have been made of the energies of order-disorder transformations in alloy phases. In many cases these energies may be conveniently determined by measuring heats of formation as a function of temperature. Many alloy phases approximately obey Kopp&apos;s law of additivity of heat capacities; as a consequence their heats of formation do not change with temperature. In ordered alloys, however, a decrease in degree of order with increasing temperature will contribute an endothermic component to the heat of formation. Thus, for ordered phases where other anomalous contributions to heat capacity in the alloy or its pure components are absent, a measured change in the heat of formation may be attributed to the energy of disordering. A portion of the Au-Zn phase diagram1" is shown in Fig. 1. The composition AuZn(ß&apos;) has a well-ordered superlattice structure of the CsCl type at room temperature. The alloy is considered to remain ordered to its melting point,1 although, from X-ray diffraction data and electromotive force measurements of zinc activities, Terpilowski3 concluded that considerable disordering began at -673°K. The Au3Zn composition is a disordered fee solid solution (a) at high temperatures. Between 693o and 543oK, the ailoy orders to form the a1 phase, having a structure based on the L12(AuCu3) type, but with a regular distribution of ordering faults in the c direction, making it tetragonal. Below 543oK, the alloy orders further by means of slight atomic shifts, without diffusion, yielding a complex long-period orthorhom-bic superlattice structure, a2.4-7 Heats of formation of Au3Zn(a2) and AuZn(ß&apos;) were measured at 363°K by Biltz et al.8 by aqueous solution calorimetry. Heat capacities of several alloys between 14.9 and 38.4 at. pet Zn in the range from 373" to 723°K were measured by Iwasaki ef al.6 Stoichiometric Au3Zn, however, was not included in the alloys measured, the nearest composition being 26.1 at, pet Zn. Also, the heating rate used for the measurements, 2 deg per min, may not have permitted equilibrium to be reached in the neighborhood of the transformations. Heat contents of solid and liquid AuZn were measured between 488o and 1138°K by Kubaschewski.9 It was decided to determine the ordering energies of Au3Zn and to check on the possible disordering of AuZn by measuring the heats of formation of these alloys at several temperatures. The present paper reports measurements by liquid tin solution calorimetry of the heats of formation of the three phases of Au3Zn at temperatures within their regions of stability, and of AuZn at a low (322°K) and a high (800°K) temperature. EXPERIMENTAL Alloy Preparation and X-Ray Examination. Ten-gram ingots of Au3Zn and AuZn were prepared by melting weighed amounts of gold (99.95 pct Au) and zinc (99.99 pet Zn) together in sealed evacuated Vycor tubes at 1060°K, followed by rapid quenching. The resulting ingots were homogenized at 900°K for 10 days. The Au3Zn alloy lost no weight in preparation, the loss from the AuZn alloy was enough to introduce an uncertainty of only 0.1 at. pet in its composition. Filings for X-ray examination were taken from various parts of each ingot and strain-annealed at 723°K for 20 min in vacuo. The Au3Zn filings were quenched rapidly to retain the disordered phase. Diffraction patterns were determined with an X-ray dif-fractometer using copper Ka radiation.

Jan 1, 1968

By G. W. Thomas

A mathematical tnodel of forward combustion in an oil reservoir is treated in this paper. The model describes a radial system having a vertical section of essentially infinite thickness, all of which is permeable to gas flow. Combustion, however, is presumed initiated over a limited thickness of the total vertical section. In the interval supporting cotnbustion, the mechanisms of radial conduction, convection and heat generation are taken into account. Above and below the burning interval, heat transport in the radial direction is by cottduction and convection. Vertical heat losses from the ignited interval are accounted for by conduction alone. A general solution is presented for the temperature distribution caused by radial movement of the combustion front. The results show that no feedback of heat occurs into the ignited interval when convection and conduction are acting in the bounding media. Peak temperatures are also 5 to 10 per cent higher than in the case where heat transport in the bounding media is by conduction alone. We arbitrarily define vertical coverage to be that fraction of the total ignited interval which is at 600F above atnbient, or greater, at any given time. The radial distance at which the vertical coverage becomes zero is the propagation range of the combustion front. It was found that an increase in vertical coverage results when the oxygen concentration, fuel concentration or gas-injection rate is increased. Moreover, the combustion front can be propagated 10 to 15 per cent further than in the case where only conduction is acting above and below the ignited interval. INTRODUCTION In the theoretical treatment of forward combustion in a radial system, one of the problems encountered is the determination of the transient temperature distributions caused by an expanding cylindrical heat source. Bailey and Larkin&apos; and Ramey&apos; simultaneously presented analytical solutions to the problem assuming heat transport by conduction alone. In a subsequent publication, Bailey and Larkin3 included the effects of both conduction and convection while treating linear and radial models. In this latter work, however, vertical heat losses were largely neglected. Selig and Couch&apos; dealt with a radial model in which both conduction and convection were acting. Only a limiting case involving vertical heat losses was considered, however. Namely, temperatures on the boundary of the bed of interest were set equal to zero. Solutions thus obtained were representative of a system having a maximum vertical heat flux. Chu5 recently treated a more general case in which a permeable bed was considered bounded by impermeable media. Conduction and convection took place within the bed, and only conduction outside of the bed. The effects of vertical heat losses were included in his study. Solutions were obtained by numerical techniques. This paper is an extension of the theoretical work of other authors pertaining to forward combustion in a radial system. In particular, a mathematical model of the process is treated in which heat generation occurs over a small vertical interval of a larger permeable section. In the interval supporting heat generation, and above and below this interval, the mechanisms of radial conduction and convection are also presumed acting. Heat losses from the ignited interval are accounted for by vertical conduction. An analytical solution for the temperature distribution caused by radial movement of the burning front is presented. The effects of certain process variables are indicated and comparisons with Chu&apos;s results are made. THEORY To render the mechanism of forward combustion tractable to mathematical treatment, we idealize the problem to the extent of assuming continuous reservoir media possessing homogeneous and isotropic properties. The following additional assumptions are implicit in this analysis. 1. The thermal parameters, i.e., heat capacities, thermal conductivities and thermal diffusivities are invariant with temperature and pressure. Moreover, the bounding media possess the same thermal properties as the bed of interest. 2. The temperatures of the porous media and its contained fluids at any point and at any time are equal. 3. The reaction rate between the oxidant gas and the fuel is infinite. This assumption implies that the incoming oxygen concentration instantaneously goes to zero within an infinitesimal distance, i.e., the width of the combustion zone is negligible. 4. The rate of gas injection is constant and corresponds to the average rate throughout the lifetime of the project. 5. The fuel concentration is constant throughout the volume of rock swept out by the burning zone. 6. There is complete burnoff of fuel. This assumption demands that the rate of propagation of the burning front equals the rate of fuel burnoff. In a radial system, with a

By A. T. Yu

On March 18, 1964, the 52,000 dwt ore carrier M/S Houn Maru called at the port of Coquimbo, Chile, on its maiden voyage to take iron ore to Japan. This marked the first time a vessel of this size ever arrived at Coquimbo, and the event promises to drastically accelerate henceforth iron ore exports from Coquimbo province. This was made possible by the use of a versatile shiploading system recently installed by Cia. Minera Santa Fe. Coquimbo has no mechanized loading facilities other than six conventional 5-ton cargo cranes. Limited space on shore further complicates the stocking problems. To make the best of what is available, a meticulously organized semi-hand-loading system has been developed through the years.

Jan 5, 1964

By W. C. Ellis

For the purpose of developing combinations of higher strength and conductivity than are obtainable by heat treatment alone in the age-hardening copper alloys, an investigation has been made of the effects of combinations of heat treatment and strain hardening on the properties in question. The investigation was limited to a study of the alloys containing nickel and silicon; nickel, silicon and cadmium; and cobalt and silicon. The experiments made in the course of the study are described in this paper in three groups: 1. The effect of hard drawing subsequent to heat treatment on the mechanical and electrical properties of the alloys. 2. The effect of low-temperature aging on the properties of heat-treated and hard-drawn alloys. 3. The effect of strain on the rate of aging of quenched nickel-silicon-copper alloys.

Jan 1, 1931