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Sec 1 This Institute is incorporated under the Membership Corporations Law of the State of New York, its corporate name is American Institute of Mining, Metallurgical, and Petroleum Engineers, Incorporated, and its abbreviated title AIME Sec 2 The objects of this Institute, as set forth in its certificate of incorporation, are "To promote the arts and sciences connected with the economic production of the useful minerals and metals, and the welfare of those employed in these industries by all lawful means, to hold meetings for social intercourse and the reading and discussion of professional papers, and to circulate by means of publications among its members the information thus obtained, and to establish and maintain a place for meeting of its members, and a hall for the reading of papers and delivery of addresses, and a library of books relating to subjects cognate to the sciences and arts of mining and metallurgy "

Jan 1, 1960

The amendment to Article III, Section 1, of the Constitution, raising the dues from $10 to $12 a year commencing in 1917, which was submitted to the members for letter ballot, was passed by a vote of 1,039 to 401.

Jan 4, 1916

Some 800 analyses for Cu, Zn, Pb, Ag, Fe, S, CaO, MgO, C02, Si02 and Al203 were used to examine the elemental constitution of the Mount Isa ores, and, for comparison, those of the Northern Leases area.Lead: zinc relations show little order from sample to sample, but means of large zones (Le. individual orebodies) show a highly linear inverse relationship. Lead: silver ratios are fairly constant from sample to sample, and means show a highly linear direct relationship.Sulphide sulphur appears to be related to carbonate and alumina abundance, sulphur decreasing systematically with increase in the ratio CO2:AI203. In all graphs involving members of the group CO2-A203-S the lead-zinc and copper ores conform to the same simple trends, confirming that the two ore types and their host rocks are genetically closely related. The earlier deduction of a volcanic-sedimentary origin for the sulphides of the Black Star orebodies appears compatible with present findings.On the basis of such an origin, the pattern of sulphide occurrence and mineralogy in the Mount Isa district is possibly related to the configuration f the upper surface of the "Eastern Creek Volcanics", the facies patterns of the overlying sedimentary rocks, and to the oldest faults affecting the two.INTRODUCTIONStanton (1962) described and interpreted the chemical constitution of the Black Star No. 2 lead-zinc and the 650750 copper orebodies at Mount Isa. The investigation has now been extended to the other principal orebodies of the area.The earlier study showed that the host rocks of the two adjacent Black Star orebodies were constitutionally closely related, and suggested that this was probably due to their original formation in contiguous zones of sedimentation. On this basis the silica-dolomite of the copper ore appeared to represent the slightly more shoreward, "limey" member, probably a fairly massive algal calcareous sediment laid...

Jan 1, 1963

By J. M. Rotter, R. A. Lohnes, J. Y. Ooi

Reliable prediction of wall pressures in a coal silo demands careful characterization of the stress-strain properties of the coal. This paper describes the results of triaxial compression tests on a typical fine-ground coal used in power generation and describes the fitting of several different constitutive models to the data. The fitted parameters for the constitutive models are well suited to application infinite element analyses. The different models are used to give the reader flexibility in choosing a model and to show the shortcomings of some models. To explore the quality of the fitting process, the fitted results are compared with finite element predictions of the triaxial test itself.

Jan 1, 1998

By S. Feinstein, I. Reznik, Y. H. Hatzor, O. Shitrit, Y. O. Rosenberg, H. J. Vinegar, S. Nguyen

"We report the effects of porosity and total organic carbon (TOC) content on the mechanical behavior of organic-rich chalks, based on laboratory measurements. The organic-rich chalks are from the Senonian Ghareb and Mishash formations. This deposit is an immature oil-prone source rock, about 300 meters thick, situated at shallow depths (several hundreds of meters) in the Shefela basin in central Israel. The study is based on core data collected from the Zoharim borehole.Three main phases comprise the bulk volume of this rock: 1) carbonate mineralogical phase (predominantly calcite); 2) pores (mostly saturated with brine); and 3) solid organic phase (type IIS kerogen). The mechanical effects of porosity (23-45%) and TOC (4-19%) are evaluated based on bulk incompressibilities vs. volumetric fractions analysis. Volumetric fractions of mineralogical and organic phases were obtained by porosity, density and TOC measurements. Saturated rock bulk moduli were measured through ultrasonic velocity measurements. Saturated rock bulk moduli correlate best with the sum of kerogen and pore volume fractions, rather than each of them separately. The volumetric fractions were also used to define the range of bulk modulus of the organic-rich chalk, using the Hashin-Shtrikman upper and lower bounds. Based on Marion's Bounding Average Method (BAM), the normalized stiffness factor of the chalk averaged at w=0.197, indicating a soft pore geometry. Using the average normalized stiffness factor, the BAM predictions showed excellent match with the saturated rock bulk modulus and the sum of kerogen and pore volume fractions. Based on our findings, it is inferred that kerogen grains do not support the solid skeleton but are distributed between the minerals grains. We also propose here a straightforward approach for calculating the Hashin-Shtrikman bounds, based on TOC and porosity measurements. The mechanical effects of porosity and TOC indicated here are very important for the understanding of the influence of immature kerogen presence in highly porous chalks, as the Ghareb and Mishash formations."

Jan 1, 2015

By M. C. Villeneuve

A numerical modelling approach was developed to explicitly simulate geomechanical characteristics of intact rock: mineralogy, grain size and fabric. The approach involved creating a representative constitutive model for each of three common rock-forming minerals: mica, quartz and feldspar. The constitutive models developed are valid within the low confinement realm of excavation boundaries, where tensile fracture processes dominate. The mineral types were assigned to numerical elements, which were associated with each other through an algorithm created in a finite difference model, FLAC 2D (Itasca 2007a), to simulate real crystal geometries and orientations. The numerical models were used in a parametric investigation of the geomechanical characteristics and compared with published observations of the rock yielding process in laboratory testing. This approach has allowed the explicit grain-scale investigation of the impact of geomechanical characteristics on rock yielding at low confinement, leading to an improved mechanistic understanding of excavation-scale rock yielding processes at excavation boundaries.

May 1, 2009

By Ali P. Gordon, James Williams

Joints gasketed with viscoelastic seals often receive an application of a secondary torque, i.e., retorque, in order to ensure joint tightness and proper sealing. The motivation of this study is to characterize and analytically model the load and deflection re-torque response of a single 25% glass-fiber reinforced polytetrafluorethylene (PTFE) gasket-bolted joint with serrated flange detail. The Burger-type viscoelastic modeling constants of the material are obtained through isolating the gasket from the bolt by performing a gasket creep test via a MTS electromechanical test frame. The re-load creep response is also investigated by re-loading the gasket after a period of initial creep to observe the response. The modeling constants obtained from the creep tests are used with a Burger-type viscoelastic model to predict the re-torque response of a single boltgasket test fixture in order to validate the ability of the model to simulate the re-torque response under various loading conditions and flange detail.

Jan 1, 2015

By Carl A Zapffe, George A. Moore

In a previous publication from this laboratory1 the conclusion was drawn that the embrittling effect of occluded hydrogen on iron and steel must result from the precipitation of the gas within small openings through the crystal structure. A review of the literature then published indicated the real existence of crystal substructures of the mosaic type and gave presumptive evidence that the "rift openings" associated with hydrogen occlusion are the same as the L'disjunctions" of the mosaic theory. Inasmuch as no clear conception of the actual nature of the crystal fragments and openings inherent in these theories has been established, it was considered desirable to obtain direct micrographic evidence of their true configuration. In this direction the standard methods of metallographic attack have proved useless. Recourse was had, therefore, to the direct examination of the surfaces upon which failure by cleavage had resulted from the presence of hydrogen. It was expected that such a surface, unaffected by chemical attack and undis-torted by polishing, would reveal the fine openings penetrated by the gas and thus indicate the real nature of the substructures. Materials Studied In commercial steels, defects caused by hydrogen take various external forms according to such factors as structure and purity, and to the extent to which hydrogen either occludes or produces chemical changes. The present work relates only to the well-known transcrystalline hydrogen embrittlement—-"pickling brittleness "—in which elemental hydrogen occluded within the grain develops cleavage characteristics. For example, hydrogen embrittlement resulting from hydrogen reaction products is excluded from detailed study. It was also desirable, as far as possible, to have the study uncomplicated by other elements which themselves cause brittleness. Three types of iron were used. These included ingot iron, decarburized free-machining steel, and a specially prepared, aluminum-killed, electrolytic iron. This choice allowed a comprehensive selection of unavoidable minor impurities and inclusions, which could subsequently be shown to have no specific effects on the structures within the grains. Carbon, and any other impurities free to react with hydrogen, were first removed by a preliminary annealing treatment in hydrogen. Subsequent treatment in vacuo extracted much of this hydrogen and its gaseous reaction products. The resultant ductile iron was then embrittled either by a second anneal in hydrogen, followed by rapid cooling, or by cathodic electrolysis. Group 1: Prepared from Armco Ingot Iron As shown in Table 1, the principal impurity in this iron was oxygen, which survived the purifying treatment to a large extent. Two series were prepared. The first (Nos. SIO-24) was purified only one day in hydrogen and dehydrogenized one day in nitrogen. Containing both hydrogen and its

Jan 1, 1943

By Carl A. Zapffe, George A. Moore

In a previous publication from this laboratory1 the conclusion was drawn that the embrittling effect of occluded hydrogen on iron and steel must result from the precipitation of the gas within small openings through the crystal structure. A review of the literature then published indicated the real existence of crystal substructures of the mosaic type and gave presumptive evidence that the "rift openings" associated with hydrogen occlusion are the same as the L'disjunctions" of the mosaic theory. Inasmuch as no clear conception of the actual nature of the crystal fragments and openings inherent in these theories has been established, it was considered desirable to obtain direct micrographic evidence of their true configuration. In this direction the standard methods of metallographic attack have proved useless. Recourse was had, therefore, to the direct examination of the surfaces upon which failure by cleavage had resulted from the presence of hydrogen. It was expected that such a surface, unaffected by chemical attack and undis-torted by polishing, would reveal the fine openings penetrated by the gas and thus indicate the real nature of the substructures. Materials Studied In commercial steels, defects caused by hydrogen take various external forms according to such factors as structure and purity, and to the extent to which hydrogen either occludes or produces chemical changes. The present work relates only to the well-known transcrystalline hydrogen embrittlement—-"pickling brittleness "—in which elemental hydrogen occluded within the grain develops cleavage characteristics. For example, hydrogen embrittlement resulting from hydrogen reaction products is excluded from detailed study. It was also desirable, as far as possible, to have the study uncomplicated by other elements which themselves cause brittleness. Three types of iron were used. These included ingot iron, decarburized free-machining steel, and a specially prepared, aluminum-killed, electrolytic iron. This choice allowed a comprehensive selection of unavoidable minor impurities and inclusions, which could subsequently be shown to have no specific effects on the structures within the grains. Carbon, and any other impurities free to react with hydrogen, were first removed by a preliminary annealing treatment in hydrogen. Subsequent treatment in vacuo extracted much of this hydrogen and its gaseous reaction products. The resultant ductile iron was then embrittled either by a second anneal in hydrogen, followed by rapid cooling, or by cathodic electrolysis. Group 1: Prepared from Armco Ingot Iron As shown in Table 1, the principal impurity in this iron was oxygen, which survived the purifying treatment to a large extent. Two series were prepared. The first (Nos. SIO-24) was purified only one day in hydrogen and dehydrogenized one day in nitrogen. Containing both hydrogen and its

Jan 1, 1943

By Walter Crafts, C. M. Offenhauer

In a previous study of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is 2 more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about I to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, I per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to I-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudo-martensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously1 that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By Walter Crafts, C. M. Offenhauer

In a previous study of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is 2 more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about I to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, I per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to I-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudo-martensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously1 that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By Eric R. Jette, Earl S. Greiner

The constitution of the iron-rich iron-nickel-silicon alloys at 600ºC. has been determined by X-ray crystal structure methods. If the results are represented isothermally by a conventional equilateral triangular plot, the single-phase solid solution (body-centered cubic structure) of the iron-rich alloys at 600°C. is bounded by four two-phase areas. Three three-phase areas separate the respective two-phase areas. Three of the phases in the polyphase areas have the same structures as the corresponding phases in the binary iron-nickel or iron-silicon alloys. There are two phases, however, that have not been reported previously and whose structures have been tentatively identified in the present investigation. The solid solubility limit of the iron-rich solid solution at 600°C. was determined by comparing the lattice constants of this body-centered cubic phase in the single and polyphase areas. It has been shown that the lattice constants of ternary alloys may be represented by surfaces whose distances above the ternary plot (representing atomic composition) are proportional to the lattice constants of the respective alloys. Then the intersections of the surface for the single-phase alloys with those for the corresponding phase in the polyphase alloys, at the respective temperatures, are the solid solubility limits of the single-phase alloys at these temperatures. The data of the present investigation were evaluated by this method, although by a system of orthographic projection it was possible to represent the three-dimensional model in two-dimensional diagrams. This investigation has shown that the solid solubility of nickel in the body-centered cubic phase of the iron-rich iron-nickel-silicon alloys at 600°C. changes from 3.7 weight per cent (Owen and Sulley) in the binary iron-nickel alloys to 5.0 per cent in the ternary alloys containing 4.8 per cent silicon. Further additions of silicon increase the solid solubility limit of nickel to 16.0 per cent at 12.6 per cent silicon. From this point the limit decreases to 9.5 per cent nickel at 15.5 per cent silicon and is, of course, zero at 15.3 per cent silicon in the binary iron-silicon alloys.

Jan 1, 1943

By Earl S. Greiner, Eric R. Jette

The constitution of the iron-rich iron-nickel-silicon alloys at 600ºC. has been determined by X-ray crystal structure methods. If the results are represented isothermally by a conventional equilateral triangular plot, the single-phase solid solution (body-centered cubic structure) of the iron-rich alloys at 600°C. is bounded by four two-phase areas. Three three-phase areas separate the respective two-phase areas. Three of the phases in the polyphase areas have the same structures as the corresponding phases in the binary iron-nickel or iron-silicon alloys. There are two phases, however, that have not been reported previously and whose structures have been tentatively identified in the present investigation. The solid solubility limit of the iron-rich solid solution at 600°C. was determined by comparing the lattice constants of this body-centered cubic phase in the single and polyphase areas. It has been shown that the lattice constants of ternary alloys may be represented by surfaces whose distances above the ternary plot (representing atomic composition) are proportional to the lattice constants of the respective alloys. Then the intersections of the surface for the single-phase alloys with those for the corresponding phase in the polyphase alloys, at the respective temperatures, are the solid solubility limits of the single-phase alloys at these temperatures. The data of the present investigation were evaluated by this method, although by a system of orthographic projection it was possible to represent the three-dimensional model in two-dimensional diagrams. This investigation has shown that the solid solubility of nickel in the body-centered cubic phase of the iron-rich iron-nickel-silicon alloys at 600°C. changes from 3.7 weight per cent (Owen and Sulley) in the binary iron-nickel alloys to 5.0 per cent in the ternary alloys containing 4.8 per cent silicon. Further additions of silicon increase the solid solubility limit of nickel to 16.0 per cent at 12.6 per cent silicon. From this point the limit decreases to 9.5 per cent nickel at 15.5 per cent silicon and is, of course, zero at 15.3 per cent silicon in the binary iron-silicon alloys.

Jan 1, 1943

By George A. Roberts, Robert F. Mehl

When austenite first forms from aggregates of cementite and ferrite, it is not homogeneous.' This inhomogeneity, consisting of both undissolved carbide and carbon concentration gradients, has a profound effect upon the characteristics of the austenite-pearlite reaction. The rate of formation of pearlite is increased by the presence of inhomogeneities2,3 and because of this (when the inhomogeneity is sufficiently pronounced) the hardenability of the steel is reduced.4,5 In addition to this kinetic effect, inhomogeneities will cause the direct formation of spheroidite from austenite at temperatures just below the Ael.6-11 The present report attempts quantitatively to evaluate the effect of inhomogeneity upon the rate of the austenite-pearlite reaction, presents a method of determining the austenitizing cycle necessary to obtain "practical homogenization," and considers the effect of inhomogeneity on hardenability. Effect of Inhomogeneities on Reactions Near Ae1 Few quantitative data are available representing the effect of undissolved carbide or of carbon concentration gradients upon the rate of the austenite-pearlite reaction. In procuring such data it is necessary to separate the effect of undissolved carbide from that of grain size, and this is not an easy task. The solution of carbide is always accompanied by an increase in the grain size, both effects contributing to the retardation of the reaction rate. It has been possible to make this separation and to obtain data that are independent of the grain size by the use of steels that differ only with respect to their tendency toward grain-coarsening. Steels C and D in Table I are such steels. These steels have been studied previously ;l,2,12 they differ only with respect to deoxidation practice; C was deoxidized with silicon while D, teemed from a portion of the same heat, was deoxidized with additional auantities of aluminum in the mold. C therefore is what is normally termed a "coarse-grained steel" and D is a, "fine-grained steel." These steels can be compared at the same grain size by heating c to a low temperature and heating D to an elevated temperature, producing a structure of austenite plus residual carbide in the former while forming homogeneous austenite in the latter; in such

Jan 1, 1943

By George A. Roberts, Robert F. Mehl

When austenite first forms from aggregates of cementite and ferrite, it is not homogeneous.' This inhomogeneity, consisting of both undissolved carbide and carbon concentration gradients, has a profound effect upon the characteristics of the austenite-pearlite reaction. The rate of formation of pearlite is increased by the presence of inhomogeneities2,3 and because of this (when the inhomogeneity is sufficiently pronounced) the hardenability of the steel is reduced.4,5 In addition to this kinetic effect, inhomogeneities will cause the direct formation of spheroidite from austenite at temperatures just below the Ael.6-11 The present report attempts quantitatively to evaluate the effect of inhomogeneity upon the rate of the austenite-pearlite reaction, presents a method of determining the austenitizing cycle necessary to obtain "practical homogenization," and considers the effect of inhomogeneity on hardenability. Effect of Inhomogeneities on Reactions Near Ae1 Few quantitative data are available representing the effect of undissolved carbide or of carbon concentration gradients upon the rate of the austenite-pearlite reaction. In procuring such data it is necessary to separate the effect of undissolved carbide from that of grain size, and this is not an easy task. The solution of carbide is always accompanied by an increase in the grain size, both effects contributing to the retardation of the reaction rate. It has been possible to make this separation and to obtain data that are independent of the grain size by the use of steels that differ only with respect to their tendency toward grain-coarsening. Steels C and D in Table I are such steels. These steels have been studied previously ;l,2,12 they differ only with respect to deoxidation practice; C was deoxidized with silicon while D, teemed from a portion of the same heat, was deoxidized with additional auantities of aluminum in the mold. C therefore is what is normally termed a "coarse-grained steel" and D is a, "fine-grained steel." These steels can be compared at the same grain size by heating c to a low temperature and heating D to an elevated temperature, producing a structure of austenite plus residual carbide in the former while forming homogeneous austenite in the latter; in such

Jan 1, 1943

By Frank S. Gardner, Morris Cohen, Dara P. Anita

There is a conspicuous lack of information in the literature on the precise role played by residual quantities of austenite in heat-treated steels. While retained austenite may be expected to have significant effects on such important properties as the impact resistance of tools and carburized parts and the dimensional stability of dies and gauges, attempts to study these problems systematically have suffered from the absence of a reliable technique for ascertaining absolute amounts of austenite. With such a technique, not only could the relation between residual austenite and service behavior be investigated, but the chemical and heat-treatment factors that control the quantity of the austenite could also be established. Thus, the retained austenite content could then be regulated in accordance with the results desired. The presence of retained austenite may be inferred by observing the changes in properties (hardness, specific volume, electrical resistance, etc.) due to its decomposition on subatmospheric cooling1 or on tempering,2 but these indirect measurements are not readily translated into percentages of austenite. Quantitative determinations have been achieved in special cases by means of dilatometric3 and magnetic4,5 measurements; but the former procedure requires that observations be conducted during the hardening and involves difficult calibrations that limit its applicability, while the magnetic procedure is very sensitive to extraneous influences and depends upon assumptions that can only be considered as approximate in an absolute sense. Furthermore, both the dilatometric and magnetic effects are volume properties, and cannot be conveniently utilized to explore variations in retained austenite from point to point in a suecimen. Some microstructures lend themselves to accurate evaluation of retained austenite contents by metallo-graphic methods,3 but such instances are not too common, and at best, the procedure is tedious. The characteristic X-ray diffraction lines of austenite permit positive identification of its presence; and in the present paper the intensities of these lines, or the corresponding densities* on photographic film, are employed to measure the amount of the retained austenite. In order to cancel the effect of uncontrollable variations in the X - ray exposure and film - development conditions, a standard reference line is Superimposed on each film by maintaining a foil of aluminum in the path of the incident X-ray beam during the entire exposure. With different samples of steel and the same foil, the ratio of the density of some one austenite line to the density of

Jan 1, 1943

By Morris Cohen, Dara P. Anita, Frank S. Gardner

There is a conspicuous lack of information in the literature on the precise role played by residual quantities of austenite in heat-treated steels. While retained austenite may be expected to have significant effects on such important properties as the impact resistance of tools and carburized parts and the dimensional stability of dies and gauges, attempts to study these problems systematically have suffered from the absence of a reliable technique for ascertaining absolute amounts of austenite. With such a technique, not only could the relation between residual austenite and service behavior be investigated, but the chemical and heat-treatment factors that control the quantity of the austenite could also be established. Thus, the retained austenite content could then be regulated in accordance with the results desired. The presence of retained austenite may be inferred by observing the changes in properties (hardness, specific volume, electrical resistance, etc.) due to its decomposition on subatmospheric cooling1 or on tempering,2 but these indirect measurements are not readily translated into percentages of austenite. Quantitative determinations have been achieved in special cases by means of dilatometric3 and magnetic4,5 measurements; but the former procedure requires that observations be conducted during the hardening and involves difficult calibrations that limit its applicability, while the magnetic procedure is very sensitive to extraneous influences and depends upon assumptions that can only be considered as approximate in an absolute sense. Furthermore, both the dilatometric and magnetic effects are volume properties, and cannot be conveniently utilized to explore variations in retained austenite from point to point in a suecimen. Some microstructures lend themselves to accurate evaluation of retained austenite contents by metallo-graphic methods,3 but such instances are not too common, and at best, the procedure is tedious. The characteristic X-ray diffraction lines of austenite permit positive identification of its presence; and in the present paper the intensities of these lines, or the corresponding densities* on photographic film, are employed to measure the amount of the retained austenite. In order to cancel the effect of uncontrollable variations in the X - ray exposure and film - development conditions, a standard reference line is Superimposed on each film by maintaining a foil of aluminum in the path of the incident X-ray beam during the entire exposure. With different samples of steel and the same foil, the ratio of the density of some one austenite line to the density of

Jan 1, 1943

By A. G. Journel

The indicator random function formalism allows a common coding of hard numerical data, constraint intervals, and soft qualitative information as prior probability distributions. The soft kriging algorithm allows interpolating these priors into posterior probability distributions of the unknown. The curse of ?negative kriged grades? is solved. The consideration of additional soft information, whenever deemed necessary, becomes possible within the framework of well-proven kriging techniques.

Jan 1, 1986

By D. E. Calisaya, É. Poulin, A. Desbiens, R. del Villar, A. Riquelme

"A 15 cm diameter by 732 cm height, fully automated pilot column was installed in the Agnico-Eagle’s Laronde concentrator (Québec), for evaluating the performance of newly developed sensors and advanced process control algorithms. For eventual comparison purposes, the pilot column is continuously fed with the same slurry feeding the copper-circuit third cleaner column. It is equipped with electrical conductivity-based sensors for estimating the three variables used in the multi-variable control algorithm under evaluation: froth depth, fraction of wash water underneath the interface and gas hold-up. Froth depth is regulated using a PI controller while a predictive controller is formulated as the minimization of the tracking errors of gas hold-up and bias rate, keeping several operating constraints between their practical upper and lower limits. In particular, an unreachable high set point of gas hold-up can be selected in order to maximize the bubble surface area available for particle collection, while maintaining bias rate above a minimum value required for froth cleaning. The proposed strategy may indirectly help optimizing flotation column operation by using these secondary variables, without the implementation of a real-time optimization layer.INTRODUCTION Flotation columns have already been used for more than 30 years, mainly in cleaning stages, as a result of their froth washing capabilities. Despite their widespread usage, the measurement of some pertaining variables and their use for automatic control is rather scarce. Among the most important variables directly related to the process metallurgy (concentrate grade and recovery), froth depth (hf) is the only one normally monitored and controlled. Even though commercial sensors are available, gas hold-up (eg), a variable directly related to concentrate recovery, is seldom measured and never controlled. Another important variable, bias rate (Jb) responsible for concentrate cleaning, is simply not measured at all, as there is no commercial sensor available. Therefore, besides froth depth, continuous monitoring of process variables is limited to flow rates (gas, feed, tailings, wash water), i.e. to manipulated variables. Some new available sensors have been industrially tested, e.g. the Sonar-TracTM for eg (O'Keefe, Viega & Fernald, 2007) and the McGill hold-up sensor (Tavera & Escudero, 2002). Others, for bias rate and frother concentration, have been successfully validated at laboratory scale by Laval’s LOOP researchers (Maldonado, Desbiens, del Villar & Chirinos, 2008b; Maldonado, Desbiens, del Villar & Aguilera, 2010a). Some devices have been extensively used for monitoring bubble size and space distri-bution, e.g. McGill bubble viewer (Gomez & Finch, 2007) and others."

Jan 1, 2012

A wide range of models are available to computer process simulators that require a large number of parameters which often are physical measurements, coefficients, calibration factors and functional forms. This makes it intuitively difficult in the majority of cases, to tune the parameters in order to simulate a circuit in agreement with experimental results. Unit calibrators operating on the principle of least squares minimisation of the difference between predicted and observed plant data, are relatively successful when the parameters are unconstrained and independent of each other. In the case of constrained parameters, a set of carefully chosen transforms is presented which map the unconstrained variables being minimised into the constrained parameters required for the model. This technique has been used successfully with the Powell minimisation algorithm and it may also be applied to other unconstrained minimisation methods. Necessary model characteristics are discussed to enable these transforms to be applied to interdependent model parameters such as points of a graph. Examples of the calibration of several unit models using these transforms are presented. KEYWORDS: Calibration, optimisation, transforms, process simulation.

Jan 1, 1993