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By C. K. Blankenship

The competitive forces now driving the electric utility industry and the constraints of regulation at all levels are changing the way we will all do business now and into the next century. Utilities are using every technological tool at their disposal to find optimum solutions to challenges in every aspect of operations. In fuel planning and procurement, these tools include models that predict which dispatch line-up will comply with environmental regulations at minimum cost. This paper presents how those tools are being used for Duquesne Light Company.

Jan 1, 1995

By J. W. Spretnak, G. W. Powell, J. H. Bucher

Tlze room-temperature tensile fracture oj smooth, round specitnens of three ultrnhigh- strength steels tempered to a wide range of strength levels was studied by means by light and electron-microscopic examination of the fracture surfaces. The fracture of AISI 4340 and 300 M at all the strength levels studied, and H-11, except after tempering at 1200° and 1300°F, occurs in three stages. The initiation of fracture is internal (except in some lightly tcmpeved specimers in which fracture is initiated at surface flaws), and is nucleated largely by separation at metal-second phase intevjaces. TIze voids grow and, coalesce to form a crack. When the crack has reached a sufficienl size, rapid propngutio~z ensues. Failure in this stage of fracture usually occurs by dimpled rupture of inicroshear stefis. In the case of H-11 tempered in the 1125° to 1300°F range, fracture in the shear steps is predominantly by concentrated deformation without void formation. The termination of fracture is usually occomplished by the formation of a shear lib in which fracture occurs by shear dimpled rupture. In the case of H-11 tempered at 1200° and 1300°F, no shear lip was obserued, and the radial elelments extend to the surface—a true termination slage does not exist. ThE tensile fracture of several metals and alloys has been investigated.2-4 In the case of polycrystal-line materials, cup-cone fracture usually results. The mechanism of cup-cone fracture may be summarized as follows.5 Cavities are formed in the necked region of the specimen. They usually are initiated by inclusions or second-phase particles. The cavities extend outwards by means of internal necking, and a crack lying about perpendicular to the length of the specimen is formed in the necked region. Subsequent crack growth occurs by the spread of bands of concentrated plastic deformation inclined at an angle of 30 to 40 deg to the tensile axis. Cavities are formed in the bands of concentrated deformation. The deformation bands zigzag across the bar with the net result that mac-roscopically the crack extends about perpendicular to the specimen axis. The final separation, or cone formation, appears to occur by continued crack propagation along one of the deformation bands out to the surface of the specimen. The micromechanics of the tensile fracture of ultrahigh-strength steels have not been thoroughly investigated. Larson and carr6,7 studied the tensile-fracture surfaces of AISI 4340 with a low-power microscope and reported that three stages of fracture could be observed in general. A centrally located region characterized by circumferential ridges, an annular region characterized by radial surface striations, and a peripheral shear lip were found. It was first pointed out by 1rwin8 that the central region is very probably one of fracture initiation and slow growth, and that the annular, radially striated region is one of rapid crack growth. Presumably the crack grows slowly, assuming roughly a lenticular shape, until it is large enough for the initiation of rapid propagation. In this investigation, it was attempted to determine the fine-scale aspects of the room-temperature tensile fracture of some ultrahigh-strength steels, and to relate the variation in fracture mode with microstructure. The steels studied were AISI 4340, 300M, and H-11 tempered to a wide range of strength levels. I) EXPERIMENTAL PROCEDURE The compositions of the steels studied are given in Table I. The steel was received in the form of hot-rolled bar stock 5/8 to 1 in. in diameter from which oversized specimens were machined and heat-treated. The heat treatments employed are given in Table 11. Subsequent to heat treatment, the specimens were ground to the final dimensions and stress-relieved by heating for 1 hr at 350°F (with the exception of the as-quenched steel). Standard smooth round specimens of 0.252-in. diameter and 1-in. gage length were tested in a Tinius Olsen Universal Testing Machine using a cross-head speed of 0.025 in. per min. The relatively coarse aspects of the fracture topography were determined by light-microscopic examination of sections through the fracture surface of nickel-plated specimens. A direct carbon-replication technique9 was used in the electron-microscopic study of the fracture surfaces. The replicas were examined in the electron microscope, and stereo pairs of electron micrographs were taken. The stereo pairs were then examined using a Wild ST4 Mirror Stereoscope. Carbide and inclusion particles extracted in the replicas were analyzed by selected-area electron diffraction. II) EXPERIMENTAL RESULTS The mechanical testing data are summarized in Table 111. The values reported are the average of

Jan 1, 1965

By T. Richard Davies

This paper primarily deals with the application of wearfacing materials and their role within a well-established preventive maintenance welding program. Some examples are shown, as well as a synopsis of possible benefits and a compendium of methods and materials.

Jan 1, 1979

By D. Weinstein

The high temperature mechanical properties of stainless steels after fast neutron irradiation are discussed in the light of effects caused by lattice dattmage and effects caused by helium generated from n,a transmutations. Embrittlement at high temperatures is due to helium accumulation at grain boundaries and to cavity formation and proPagation along grain boundaries. Following from the embrittlement mechanism, it is suggested that when deformation occurs by mechanisms associated with super plasticity, helium ac-curnulation at boundaries should be attenuated and cavities, if formed, should be nonpropagating. As the mean free Path between interphase boundaries of a two-phase Fe-Cr-Ni alloy was decreased, the degree of superplastic deforrnation at 870°C increased, as vneaszired by total elongation and by the expottent m = a log 'a/a log 'i. This alloy and type 304 stainless steel were cyclotron irradiated in an a-particle beam to a helium concentration of -1 x 10 atom He per atom. The stainless steel specimen was embrittled, but the ductility of irradiated two-phase Fe-Cr-Ni alloys correlated with the values of. m during 'defor-malion. The .finest grained, helium-injected specimens that deforrned with highest m values exhibited the largest elongations to ,fracture. These results could be correlated with metallographic observations of cavity behavior: the propensity for intergranular propagation was lessened as the m value increased. It is concluded that superplastic deformation is ef-fectizle in attenuating helium embrittlement at elevated temperatures. One of the principal problems associated with development of fast breeder reactors is application of alloys such that suitable fuel cladding results. Stainless steels and other Fe-Cr-Ni alloys, because of highly acceptable nuclear characteristics, represent the primary materials for this component, and an exhaustive research and development effort is being conducted. The main deficiency of these materials has been a severe loss of ductility at high temperatures after fast neutron irradiation. An extensive body of mechanical property data and microstructural observations has provided an adequate phenomenological description of embrittlement; in conjunction with transmission electron microscopy studies, a reasonably acceptable embrittlement mechanism has been obtained. Following from this mechanism, it is suggested in the present work that ductility would be enhanced if deformation could occur by mechanisms associated with the phenomenon of superplasticity. Experiments to test this hypothesis have been conducted, and the results are presented and discussed in this paper. IRRADIATION EMBRITTLEMENT AT HIGH TEMPERATURE Austenitic stainless steels have been irradiated to accumulated fast neutron fluences of 1020 to 1022 nvt at temperatures between 60" and 600°C. Specimens that have been exposed to these conditions and subsequently tensile tested at temperatures between 600" and about 900°C exhibit approximately 5 pct total elongation to fracture.'-3 For unirradiated specimens receiving a nearly identical thermal exposure, total elongation at these test temperatures is about 45 pct. Examination of irradiated specimens has shown that fracture propagation is entirely intergranular. These phenomenological aspects of irradiation embrittle-ment at elevated temperatures are well known and are not generally disputed. Although the explanation of this phenomenon has been controversial, a mechanism for ernbrittlement has emerged that accounts reasonably well for the observed mechanical behavior. The controversy resulted primarily from an indeterminate role of neutron-in-duced lattice damage, if any, and a presumed, but experimentally unverified, contribution to embrittle-ment from helium generated by n,a transmutations. Recently, Holmes and coworkers4 have conducted experiments that separate these effects, and the results are instructive in formulation of the ernbrittlement mechanism. Holmes el al.4 irradiated type 304 stainless steel in EBR-I1 to a fluence of 1.4 x 1022 nvt (E > 0.18 mev); the irradiation temperature was 538" * 48°C or, in terms of absolute melting point, 0.49 * 0.03 T,. After irradiation, tensile tests were conducted at temperatures of 21" to 870.C, the specimens first being annealed for 30 min at each test temperature. In addition, thin sections of irradiated specimens were annealed for 1 hr at identical temperatures, electro and examined by transmission electron microscopy. Thus, for a given temperature, it was possible to correlate mechanical properties with the defect structure. At room temperature, the yield stress of irradiated specimens was a factor of 2.5 higher than unirradi-ated specimens exposed to an equivalent thermal history. Electron microscopic examination of the irradiated specimen revealed a high density of lattice damage in the form of Frank sessile dislocation loops and polyhedral voids. Holmes et al.4 concluded that the presence of this defect substructure caused the increase in yield stress and that after irradiation in a fast neutron flux at 0.49 Tm, substantial lattice dam-age persists. Annealing at progressively higher tem-

Jan 1, 1970

By R. I. Jaffee, R. M. Evans

IN recent years, the interest in liquid metals as heat-transfer media for power plants has been very great. The possibility of the development of nuclear power plants has increased this interest and served as the impetus behind much research on low melting metals and alloys for such purposes. The principal reasons for consideration of liquid metals as heat-transfer media lie in their high thermal conductivity and consequent high heat-transfer coefficients, stability at high temperatures, and the high ranges of temperature possible. The element gallium possesses some of the requisite properties for a heat-transfer liquid. It is a unique material, having a low melting point and a high boiling point. Pure gallium melts at 29.78oC, and suitable alloying will produce a metal which melts below room temperature. The boiling point is about 2000°C. As it is a liquid metal, the heat-transfer characteristics would be good. Gallium is not now readily available, due in part to a lack of uses for the metal. Nevertheless, it is not a rare element, and a sufficient supply of gallium exists to permit its consideration for this use. Since gallium has some promise as a heat-transfer liquid, owing to its unique properties, research on the subject was carried on at Battelle Memorial Institute at the request of the Bureau of Ships, U.S.N. The research had as its objectives the determination of the effect of alloying on the melting point of gallium, and the study of the corrosion of possible container materials. In this research, alloys were found which had significantly lower melting points than pure gallium, but none which simultaneously fulfilled other additional requirements, chiefly the corrosion problem. Neither was it found possible to reduce the melting point of certain otherwise suitable alloys appreciably by small additions of gallium or gallium alloys. The results gave little hope that gallium alloys can be developed which enhance the good properties and minimize the undesirable characteristics of elemental gallium. Thus, gallium now appears less promising than other metallic heat-transfer media. The experimental thermal-analysis techniques used in this work have been described.' Experimental Results As a first approximation, the development of low melting gallium alloys was based on alloying elements suitable for use in a nuclear power plant, which also lowered the melting point of gallium. Information from the literature, summarized in Table I, indicates that. tin, aluminum, and zinc are the only suitable elements which cause a lowering of the melting point of gallium. Indium and silver also lower the melting point of gallium, but are of little interest for use in nuclear power plants. Of the elements reported not to lower the melting point of gallium, there is some ambiguity on the behavior of copper. Weibke3 obtained solidus arrest temperatures of 29°C for Cu-Ga alloys from 60 to 90 pct Ga, 0.8C lower than the generally accepted melting point. This may be the effect of a eutectic close to gallium, or, more likely, the result of impurities, or experimental error. The seven elements listed in Table I whose effects were not known were of potential interest if they lowered the melting point of gallium. Their effects were determined experimentally for this reason. Binary alloys containing nominally 2 pct of each of these elements were prepared in the form of 2-g melts by placing the components in a graphite crucible and holding them in an argon atmosphere at 370°C for 5 hr. These melts were then subjected to thermal analysis. In all cases. the solidus temperature was the melting point of gallium. Since these elements (As, Ca, Ce, Mg. Sb, Si, and T1) did not lower the melting point of gallium, they were not considered further as components of a eutectic-type alloy. Ga-Sn-Zn Alloys Preliminary considerations of this system for low-melting alloys were encouraging. All three binary systems were of the simple eutectic type. The composition and melting points of the eutectics were as follows: Sn-9 pct Zn (199°C), Ga-8 pct Sn (20°C), and Ga-5 pct Zn (25°C). Therefore, the probability of a ternary eutectic was high. For reasons to be discussed later, aluminum could not be used as an alloying constituent, leaving the Ga-Sn-Zn system as the only one of interest for low-melting gallium-

Jan 1, 1953

By H. T. Green, R. M. Brick

True stress-strain data on alloys of pure iron with up to 2.4 pct Al were obtained in the temperature range +100° to —185°C. Alumi-num was found to reduce yield and flow stresses of iron at low temperatures but to have little or no effect on ductility. The effects of temperature and composition on strain hardening are discussed. SEVERAL independent studies of the behavior of high purity iron binary alloys at low temperatures are now in progress in attempts to evaluate systematically the variables affecting the low temperature brittleness of ferritic steels. This paper reports the results of one such investigation in which the tensile properties of aluminum and aluminum plus silicon ferrites were measured from 100" to —192°C. True stress-natural strain data have been obtained in order to evaluate as many as possible of the parameters which describe the behavior of the materials involved. In comparable studies at the National Physical Laboratory in England, iron and iron alloys of high purity have been produced' and tested at subat-mospheric temperatures.' True stress-natural strain curves were obtained there also. The purest iron contained 0.0025 pct C and 0.001 pct O and N. Even this, as normalized at 950°C following hot rolling, showed little ductility at -196°C. The grain size was ASTM No. 3, and the room-temperature yield strength was 17,800 psi (which seems too high for pure iron). Some of the NPL irons contained considerably more oxygen and demonstrated intergran-ular fracture at —196°C. The authors2 carefully differentiated between intergranular fractures associated with excessive oxygen content and transcrys-talline cleavage with little ductility encountered at —196°C in the purer material. The cleavage stress was half again as great as that associated with inter-granular fracture. Test Material, Preparation, and Procedures Of a number of Fe-A1 alloys produced, eight were considered to be sufficiently pure for testing. Partial chemical analyses (Table I), low observed yield points, and high ductilities indicate these alloys to be comparatively pure for vacuum-melted irons of sizable ingots, 5 Ib or more. To produce the binary Fe-A1 alloys, electrolytic iron was melted in air, cast into slabs, and rolled to strips 0.010 in. thick. These strips, joined into a continuous ribbon and wound into 2 1/2 in. diameter spools, were subjected for four weeks to a moving atmosphere of purified dry hydrogen in a stainless-steel tube at 1050" to 1150°C. Charges of these spools were melted in beryllia crucibles under good vacuums (1 micron), and aluminum (99.97 pct Al) was added to the melts. Compositions of these alloys are recorded in Table I. The ingots were hot forged and then cold rolled at least 65 pct to 3/8 in. rods which were vacuum annealed to the desired grain size, approximately ASTM No. 4, prior to machining into tensile test bars. All tensile specimens had gage sections 1 in. long, with a fillet of 1.5 in. radius to the shoulder. Gage diameters were 0.250 in, except for a few rods where additional cold work required use of a 0.200 in. gage section. After machining, 0.002 in. was removed from the gage diameter using 240, 400, and 600-grit metallo-graphic papers. The final polish with 600 grit left the fine scratches running in the longitudinal direction. By this means, surface metal strained during machining was removed. A few specimens heat treated after machining were similarly reduced 0.004 in. to remove any material affected chemically by the atmosphere during heat treatments, as is discussed in a later section. Tensile tests of the eight alloys at constant temperatures from +100° to —185°C were performed in apparatus which has been described." The essentials include a double-walled insulated metal vessel which contained the liquid heat-transfer medium surrounding the test specimen. A constant temperature was maintained by means of a pyrometer which regulated the pressure of dry air driving liquid air through a copper coil. Temperature variation was less than ±2°C during a specific test. For axial straining, two lengths of case-hardened chain, terminating in simple shackles, loaded the specimen through threaded grips. The lower grip bar passed through a hole in the bottom of the test vessel to which it was joined by a thin-walled

Jan 1, 1955

By S. Fróes Abreu

The importation of liquid fuels in Brazil amounts to about 1,300,000 tons; 30 per cent of this total being gasoline for automobiles and airplanes. Statistics show an increasing consumption of gasoline notwithstanding the use, daily greater, of fuel alcohol, a domestic fuel of low cost in the producing areas. The Brazilian market for petroleum products has shown a tendency to grow as the country has steadily developed— expansion of the industries in the areas already developed, and also the recent occupation and population of new zones. The war has brought many difficulties to Brazil. It has limited the importation of petroleum products, which came mostly from the United States, Venezuela and Mexico. During the first days of the present crisis automobile and truck traffic was vastly curtailed; but the national activity in the production of more fuel alcohol, and principally in the development of producer gas, has done much to provide a solution for the transportation crisis threatened at the beginning of the gasoline shortage. Many thousands of cars and trucks in Brazil are now burning producer gas made from charcoal. Producer gas for automobiles has been improved by Brazilian engineers and many new types have appeared. - Production There is only one producing oil field in Brazil—the one in the State of Baia, near the Baia de Todos os Santos. The first producing oil well was drilled in Lobato in 1939 by the Federal Government following studies carried on in that area by Oscar Cordeiro, S. Froes Abreu, G. Paiva and J. Amarai. The field and laboratory work done by these men was published in 1936 in a paper entitled " ContribuicBes para a Geologiado Petroleo no Reconcavo (Baia)," which indicated the potentialities of the area as an oil producer. After the discovery of the first producing well, the Brazilian Government issued a decree establishing an area, 60 km. in radius, as a National Reservation. This first well was drilled in Lobato. In no other locality in Brazil is any petroleum produced. Although many wells have been drilled in the southern part of Brazil, on the northeastern coast and in the Amazon Basin, they yield nothing. Actual production is below 1000 bbl. daily. The producing area covers Itaparica Island, Candeias and Arath, Joanes Island and Lobato. Since very little geological information was available, especially about subsurface structure, it was necessary to drill several wildcats in order to collect the data essential in the choosing of good locations. The United Geophysical Co. was contracted to perform a geophysical survey to clear up the major problems concerning the oil deposits in the Reconcavo beds. The best wells in Reconcavo are on Itaparica Island, where the daily flow amounts to about zoo bbl.; in other sectors production is considerably lower and most of the wells are pumping. In the Aratfi sector, near an area of very light oil, there is an important gas field, with wells producing about I0,000,000 cu. ft. daily.

Jan 1, 1944

By S. Fróes Abreu

The importation of liquid fuels in Brazil amounts to about 1,300,000 tons; 30 per cent of this total being gasoline for automobiles and airplanes. Statistics show an increasing consumption of gasoline notwithstanding the use, daily greater, of fuel alcohol, a domestic fuel of low cost in the producing areas. The Brazilian market for petroleum products has shown a tendency to grow as the country has steadily developed— expansion of the industries in the areas already developed, and also the recent occupation and population of new zones. The war has brought many difficulties to Brazil. It has limited the importation of petroleum products, which came mostly from the United States, Venezuela and Mexico. During the first days of the present crisis automobile and truck traffic was vastly curtailed; but the national activity in the production of more fuel alcohol, and principally in the development of producer gas, has done much to provide a solution for the transportation crisis threatened at the beginning of the gasoline shortage. Many thousands of cars and trucks in Brazil are now burning producer gas made from charcoal. Producer gas for automobiles has been improved by Brazilian engineers and many new types have appeared. - Production There is only one producing oil field in Brazil—the one in the State of Baia, near the Baia de Todos os Santos. The first producing oil well was drilled in Lobato in 1939 by the Federal Government following studies carried on in that area by Oscar Cordeiro, S. Froes Abreu, G. Paiva and J. Amarai. The field and laboratory work done by these men was published in 1936 in a paper entitled " ContribuicBes para a Geologiado Petroleo no Reconcavo (Baia)," which indicated the potentialities of the area as an oil producer. After the discovery of the first producing well, the Brazilian Government issued a decree establishing an area, 60 km. in radius, as a National Reservation. This first well was drilled in Lobato. In no other locality in Brazil is any petroleum produced. Although many wells have been drilled in the southern part of Brazil, on the northeastern coast and in the Amazon Basin, they yield nothing. Actual production is below 1000 bbl. daily. The producing area covers Itaparica Island, Candeias and Arath, Joanes Island and Lobato. Since very little geological information was available, especially about subsurface structure, it was necessary to drill several wildcats in order to collect the data essential in the choosing of good locations. The United Geophysical Co. was contracted to perform a geophysical survey to clear up the major problems concerning the oil deposits in the Reconcavo beds. The best wells in Reconcavo are on Itaparica Island, where the daily flow amounts to about zoo bbl.; in other sectors production is considerably lower and most of the wells are pumping. In the Aratfi sector, near an area of very light oil, there is an important gas field, with wells producing about I0,000,000 cu. ft. daily.

Jan 1, 1944

By W. F. Stowasser

downdraft traveling grate process to agglomerate pelletized iron ore concentrates has been successfully demonstrated in a pilot plant at Carrollville, Wis. Work there followed several years of development in the Allis-Chalmers Mfg. Co. laboratories, and the pilot plant phase was carried out in cooperation with Arthur G. McKee & Co., consultants and engineers to the iron and steel industry. End result of the process is conversion of iron ore concentrates into a form which can easily be transported and smelted in the blast furnace. Process Description The first of two process steps incorporates the art of balling and prepares the concentrates for burning. The second step consists of burning the green balls on the grate machine to the hardness required for shipping and handling purposes and for reduction in blast furnaces, see Fig. 1. Facilities are provided at the pilot plant to receive carload quantities of concentrate. The concentrates are loaded into a 50-ton bin direct from railroad cars. Because of the variable moisture content of the concentrates after shipment in an open railroad car it is necessary to repulp and refilter the concentrates to maintain a uniform and proper moisture content for the balling operation. Concentrates are conveyed to slurry tanks, and the slurry, at 50 to 60 pct solids, is pumped to a 4x4-ft drum filter. The filter provides feed of uniform moisture to the plant. Magnetite concentrates are normally filtered to produce a cake containing about 10 pct moisture, a necessary requirement for the following balling operation. The filtered concentrate is conveyed to a rotary bin table feeder which acts as a surge bin for the filter cake and delivers a steady flow of concentrates to the balling drum. It is often desirable to make additions to the concentrates as they are fed to the balling drum. These additives, such as bentonite, increase the strength of the finished green pellet and improve ballability of the concentrate. A vibrating feeder supplies additive to the feed belt, and the additive is mixed with the concentrate in the balling drum. The balling drum, shown in Fig. 3, is 8x3-ft diam. An oscillating cutting bar maintains the lining in the drum by trimming off the buildup of excess concentrate as it forms. The drum is operated in closed circuit with a lx4-ft rod-deck vibrating screen. Undersize pellets or seed pellets from the screen are returned to the balling drum until they grow to the desired size. Size of pellets is controlled by the opening in the screen deck. The formation of pellets in the balling drum is affected by many variables. Some of these are: the size distribution of the feed, the particle shape of the concentrate, the feed rate to the drum, the moisture in the concentrate, the speed of rotation of the drum, the slope of the drum, and the type of trimming obtained with the cutting bar. In this process, attempts are made to control the pellet size within the limits of % to 5/8 in. diam. The screened oversize pellets are conveyed under a coal feeder where sufficient powdered coal is added to the belt to produce desired results in the burning process. The top size of the coal successfully used has been 20 mesh, and anthracite was used in the test program. Fig. 4 illustrates the vibrating screen and the coal feeder. The pellets and free coal are conveyed together to the 5x3-ft diam- reroll drum that rolls the coal onto the surface of the pellets. This drum is also equipped with a cutting bar. The prepared pellets, containing bentonite, water, and surface coal, are elevated to the traveling grate, which consists of a continuous strand of 37 pallets. Each pallet, with a grate bar area 2 ft wide by 1 1/2 ft long, has 14-in. high side plates, Fig. 5. Feeding and distribution of the green balls to the grate is handled by a short conveyor which oscillates back and forth across the 2-ft width of the grate. An adjustable vertical plate located several inches in front of the head pulley of the oscillating conveyor controls the height of the bed and levels the moving bed of pellets. This method of feeding prevents segregation of various size pellets as well as fines and produces a uniform, permeable bed. The pallet train moves under the furnace and across four windboxes, located beneath the pallet frames, see Fig. 2. As the green pellets are deposited on the grate, partial drying of the pellets begins over a 2-ft long updraft windbox. The low temperature air reduces the moisture in the pellets in the lower level of the bed and this operation is essential to prevent sagging of the bed during later stages of the Process. The air used for this drying is recuperated from cooling the pellets on the grate, and supplemental heat, required for starting the Process, is obtained from an auxiliary burner. The pellets are then moved by the grate into the furnace and over an 8-ft windbox, designated as the downdraft waste windbox. Products of combustion are exhausted from this windbox to atmosphere. The furnace, shown in Fig. 6, is constructed with three chambers to provide downdraft drying, preheating, and ignition, respectively, to the pellet bed as it passes through. Overall length of the furnace is 5.57 ft; however, the exterior wall ends may be moved to reduce the length and also adjusted to Obtain the bed height desired, The drying, preheating, and ignition sections of the furnace are supplied with medium temperature

Jan 1, 1956

By Rodney P. Smith

The diffusivity of carbon (0.1 wt pct C) in Fe-Nz alloys (0 to 100 pct Ni) has been determined for the temperature range 860° to 1100°C. As a function of nickel content, the diffusivity has a maximum near 60 pct Ni (the maximum diffusivity being about 1.3 times that in the absence of nickel); the activation energy has a maximum between 40 and 50 pct Ni and a maximum between 80 and 90 pct Ni. The difference between the minimum activation energy and that in iron is about 3000 cal pev g-atom; Do has a minimum between 40 and 50 pct Ni and a maximum between 80 and 90 pct Ni. The results cannot be rationalized by an approximate thermodynamic treatment. THE diffusivity of carbon has been determined in a number of iron alloys over a limited concentration range. It seemed desirable to investigate a system which allows an extended range of alloy composition within a single-phase region. The Fe-Ni system is ideal in this respect, in that all alloys from 100 pct Fe to 100 pct Ni are fee in a convenient temperature range.&apos; The carbon diffusivity was determined by a decar-burization method. The experimental procedure was identical with that used to determine the diffusivity of carbon in y Fe-Co alloys.2 The experimental data are given in Table I. A small correction (order of a few percent) has been made to the measured carbon loss to correct for the carbon lost from the ends of the cylinders.&apos; Since the diffusivity of carbon varies with carbon content the measured diffusivity is an average value for a carbon content between zero (surface) and that at the center of the sample at the end of the decarburization periods. In making the correction in D to 0.1 wt pct C it is assumed that the measured D corresponds to the arithmetical mean of the carbon content at the surface and at the center of the sample at the end of the decarburization period.3 Since this correction is small (<4 pct in D) and since for our decarburization times the changes in carbon content at the center of the sample was small the mean carbon content could have been taken as half the initial value. It is further assumed that the change in D with carbon content for the alloys is the same as that for the diffusion of carbon in iron. From the data of Wells, Batz, and Mehl4 and of smith5 the correction of D from the mean carbon content to 0.1 wt pct C is 0.3 (0.1 - mean wt pct C). The results for iron are given in Ref. 2. Within the experimental error log Do.l%C for each alloy is a linear function of 1/T; the constants for the equation determined by the method of least squares are given in Table I. The deviations of the experimental points from the least-squares line are of the order of 2 pct in D. A comparison of our results for the diffusivity of carbon in nickel with those of other investigations is shown in Fig. 1. The lower curve in Fig. 1 is a linear extrapolation of values calculated* from the equation of Diamond6 for the relaxation time (temperature range 100° to 500°C). The results indicate a small increase in the activation energy over the temperature range 100° to 1400°C; however, it is difficult to say whether the change in Q is real or experimental error. Certainly the change in Q is less than the variation of 5 kcal per g-atom in the diffusivity of carbon in a iron.6 The experimental data for all the alloys are plotted in Fig. 2. As a function of nickel content the diffusivity has a maximum near 60 wt pct Ni at all temperatures investigated and possibly a minimum between 80 and 90 wt pct Ni for temperatures below 1000°C. The activation energy, Q, and log Do are plotted as a function of the nickel content in Fig. 3. Due to the limited temperature range of our experiments neither Q nor Do can be determined precisely; the activation energies appear to be consistent to ±0.3 kcal per g-atom; however the deviation from the absolute values may be considerably larger, see Table II. The Do values probably have little significance. The solid line for Do in Fig. 3 represents the values required to reproduce the experimental values for D when Q has values represented by the upper solid line The diffusivity of carbon may be expressed in terms of the mobility B22, the activity coefficient r2,

Jan 1, 1967

By R. M. Waterstrat

The phases occurring in the V-Mn system were studied by means of X-yay diffraction and metallo-paphic techniques, using are-melted alloy specimens annealed in the temperature range 800° to 1150°C and quenched. The bcc solid solution extends at 1250°C all the way from vanadium to 6-manganese. Below 1050°C the a-phase is formed, and the terminal a-manganese phase is stabilized up to about 900°C by vanadium in solid solution. IN the only previous general survey of the V-Mn system Cornelius, Bungardt and Schiedtl reported the existence of three intermediate phases corresponding to the approximate compositions VMn,, VMn, and V5Mn. The phase VMn8 has recently been identified as a o phase2 but the alloy VMn was found to have a bcc structure2 corresponding apparently to the vanadium solid solution rather than to the large cubic unit cell reported by Cornelius et al. 1 Subsequent work by Rostoker and Yamamoto3 has shown that the vanadium-base bcc solid solution extends to at least 15 pct Mn at 900°C. An alloy corresponding to the composition VMn, was examined by Elliott,4 who reported that the as-cast sample as well as samples annealed at 1200o and 1300°C had bcc structures, but that annealing at 1000°, 800") and 600°C produced two phases. One of these phases was apparently the bcc solid solution and the other resembled the o phase structure. Hellawell and Hume-Rothery5 established the phase relationships in manganese-rich alloys above 1000°C, and showed that the o phase in this system is replaced by the 6 Mn (bcc) solid solution at temperatures above 1050°C. These results suggest that a continuous bcc solid solution may exist above 1050°C between vanadium and 6 Mn. The present investigation was undertaken in order to develop more complete information in regard to this system. EXPERIMENTAL METHODS The alloys used in the present work were prepared by arc-melting electrolytic manganese having a minimum purity of 99.9 pct and vanadium lumps with a purity of 99.7 pct. The major impurities present in these metals were carbon, nitrogen, and oxygen and this would account for the small percentage of nonmetallic inclusions observed metal-lographically. The arc-melting was at first performed under a helium atmosphere and it was necessary to keep the melting times as short as possible in order to minimize the loss of manganese by vaporization. It was later found that the evaporation of manganese was considerably reduced when the melting was done under argon atmosphere. The final composition of each alloy was calculated by assuming that the total weight loss during melting was due to evaporation of manganese. Compositions which were calculated in this manner agreed reasonably well with the results of chemical analysis, as shown in Table I. Spectrographic analysis revealed the presence of contamination by tungsten, but in no case was the percentage of tungsten greater then 0.4 at. pct. The specimens were in each case broken in half and the fractured section was examined visually and microscopically for evidence of inhomogeneity. Each specimen was homogenized at temperatures near l100°C, as shown in Table I. After this treatment most specimens consisted of large columnar grains of the bcc vanadium solid solution. The etchant used in most of the metallographic work consisted of 20 pct nitric acid, 20 pct hydro-flouric acid, and 60 pct glycerine. It was found that this etchant would clearly delineate the phases present in these alloys although it does not produce any striking contrast between the phases. For certain manganese-rich alloys, a 1 pct aqueous solution of nitric acid was used. This etchant gave a brown color to the a-manganese phase, whereas the o phase was virtually unattacked and appeared very light as shown in Fig. 1. The etchants used by Cornelius et a1.l were found to produce spurious effects in some of these alloys. In particular, the vanadium-rich alloys etched in hot sulfuric acid often appeared to consist of two phases when both X-ray diffraction and etching with the glycerine-acid mixture indicated the presence of single phase bcc solid solution. A few percent of what appears to be an oxide or nitride phase was found at the grain boundaries and in the interior of the grains, especially in the vanadium-rich alloys. All alloys were annealed in sealed silica tubes containing 1 atm of pure argon and these tubes were then quenched in cold water. Although some manganese loss occurred during annealing, the loss seemed to be confined to the surface of the speci-

Jan 1, 1962

By W. M. Baldwin, J. T. Brown

The effect of hydrogen on the ductility, c, of SAE 1020 steel at strain rates, i, from 0.05 in. per in. per rnin to 19,000 in. per in. per rnin and at temperature, T, from +150° to —320°F was determined. The ductility surface of the embrittled steel reveals two domains: one in which and the other in which The usual "explanations" of hydrogen embrittlement are in accord with the first of these domains only. THE purpose of this investigation was a fuller A characterization of this of the investigation effects of varying temperature and strain rate on the fracture strain of hydrogen-charged steel. To be sure, it is known that low and high temperatures remove the embrittlement that hydrogen confers upon steels at room temperature,1 * see Fig. la and b, and that high strain rates have a similar effect,&apos;-&apos; see Fig. 2a, b, and c. However, the general effect of these two testing conditions on the fracture ductility of hydrogen-charged steels is not known, i.e., the three-dimensional graphical representation of fracture ductility as a function of temperature and strain rate is not known—only two traverses of the graph are available. The need for such a graph is not pedantic. To demonstrate this point, Fig. 3a, b, and c shows three of many three-dimensional graphs, all possible on the basis of the two traverses at hand. The important point (as will be developed in the Discussion) is that each of them would indicate a different basic mechanism for hydrogen embrittlement. It will be noted that the four types of ductility surfaces in Fig. 3a, b, and c may be characterized as follows: Material and Procedure Tensile tests were made at various temperatures and strain rates on a commercial grade of % in. round SAE 1020 steel in both a virgin state and as charged with hydrogen. The steel was spheroidized at 1250°F for 168 hr to give the unembrittled steel the lowest possible transition temperature. The steel was charged cathodically with hydrogen as follows: The specimen was attached to a 6 in. steel wire, degreased for 5 min in trichlorethylene, rinsed with water, and fixed in a plastic top in the center of a cylindrical platinum mesh anode. The assembly was placed in a 1000 milliliter beaker containing an electrolyte of 900 milliliters of 4 pct sulphuric acid and 10 milliliters of poison (2 grams of yellow phosphorous dissolved in 40 milliliters of carbon disulphide). A current density of 1 amp per sq in. was used which developed a 4 v drop across the two electrodes. All electrolysis was carried on at room temperature. Temperatures for tensile tests were obtained by immersing the specimens in baths of water (+70° to + 150°F), mixtures of liquid nitrogen and isopen-tane (+70° to —24O°F), and boiling nitrogen (-240" to-320°F). Specimens were tested in tension at strain rates of 0.05, 10, 100, 5000, and 19,000 in. per in. per min. The 0.05 and 10 in. per in. per rnin strain rates were obtained on a 10,000 lb Riehle tensile testing machine, the 100 in. per in. per rnin rate on a hydraulic-type draw bench with a special fixture, and the 500 and 19,000 in. per in. per rnin rates on a drop hammer. The fracture ductility of hydrogen-charged steel at room temperature and normal testing strain rates (-0.05 in. per in. per min) is a function of electro-lyzing time, dropping to a value that remains constant after a critical time.&apos;* Under the conditions of • The hydrogen content of the steel continues to increase with charging time even after the ductility has leveled off to its saturated value.&apos; this research the saturated loss in ductility occurred at approximately 30 min, see Fig. 4, and a 60 min charging time was taken as standard for all subsequent tests. After charging the steel with hydrogen, the surface was covered with blisters. These have been described by Seabrook, Grant, and Carney.&apos; The original diameter of the specimen was not reduced by acid attack, even after 91 hr. Results The ductility of both uncharged and charged specimens is given as a function of strain rate in Fig. 5, and as a function of temperature at four different strain rates in Fig. 6. These results are assembled into a three-dimensional graph in Fig. 7. It is seen that the locus of the minima in the ductility curves of the charged steels divides the ductility surface into two domains. At temperatures below the minima,

Jan 1, 1955

By David Turnbull

IN isothermal recrystallization processes, new crystals generally grow into the matrix until they impinge upon other new crystals or an external surface, at constant linear rates G. Before impingement the perceptible course of growth can be described by the equation: 1 = G(t-7) C1I where, G = dl/dt, 1 is a crystal dimension measured in a constant direction, t is the time, and 7, the nucleation period, is a positive intercept on the time axis. Fig. 1 is a schematic representation of I as a function of time for a recrystallizing grain. G is dependent upon temperature, driving energy (strain or surface energy), relative grain and boundary orientations, but is generally independent of time. The frequency of nucleation, fi, (time" volume") can be defined by the equation: N = 1/fV [2] where ? is the mean nucleation period and V is the volume of the specimen that has not recrystallized. The kinetics of primary and secondary recrystallization generally can be described satisfactorily in terms of the parameters N and G only.&apos;-" After recrystallization is complete the average grain size 7 increases with time by "normal grain growth;" didt, the average rate of grain growth, is strongly time dependent and has not yet been precisely related to G for the motion of the individual grain boundaries constituting the system. It has been suggested4* " that the elementary act in grain boundary migration is closely related to the elementary act in grain boundary self-diffusion. Although the distance of atom movement in the two processes may be somewhat different, there is reason to expect that the activated states may be very similar, so that the free energy of activation for grain boundary migration should be of the same order of magnitude as for grain boundary self-diffusion. Therefore, it is desirable to develop a satisfactory basis for comparing data on self-diffusion and grain boundary migration and to make such comparisons where possible. Theory The formal theory of grain boundary migration rates is analogous to the theory for the rate of growth of crystals into supercooled liquids reviewed elsewhere 6-8. Boreliuss has shown that the latter theory describes, within the theoretical uncertainty, the growth of selenium crystals into supercooled liquid selenium. Motto and more recently Smolu-chowski" have derived expressions for the rate of boundary migration in recrystallization. The treatment to be presented is similar to Mott&apos;s excepting that the formalism of the absolute reaction rate theory will be used. The atomic mobility, M, in grain boundary migration is defined by: G = -M6p/6x where p is the chemical potential per atom and x is the coordinate measured in the direction of grain boundary movement. Let AF be the free energy difference per gram atom on the two sides of the boundary and k the thickness of the boundary. For RT>>AF the potential gradient across the boundary (6p/6x) is essentially linear and it follows that: SF/8x = - aF/N\ [4] where N is Avogadro&apos;s number. According to the Nernst-Einstein equation, M is related to a diffusion coefficient, Do, for matter transport in grain boundary migration by the equation: M = Da/kT [5] Substituting eqs 4 and 5 into eq 3 gives the basic relation between Do and G: G = DoaF/\RT [6] Do values may be calculated from experimental values of G from eq 6 and directly compared with the coefficient of self-diffusion within the crystal, DL, or the grain boundary self-diffusion coefficient D,. However, a more convenient, though equivalent, basis for comparing atomic mobility in grain boundary migration and self-diffusion is through the constants of the absolute reaction rate theory. According to this theory diffusion coefficients may be written:" D = k2(kT/h) exp [-AF,/RT] 171 aFa, the free energy of activation, is related to the measured energy of activation, Q, by the equation: AFA = Q - T aSx - RT [8] where aSa is the entropy of activation. Substituting eqs 8 and 7 into eq 6 gives: G = ek(kT/h) (aF/RT) exp [(AS,,)C/R] exp C-Qc/RTI C91 where the subscript G refers to boundary migration. The relationship between the driving free energy and the free energy of activation in boundary migration is indicated schematically in Fig. 2. Experience indicates that the variation of G with temperature can be described by: G= Go exp [- Qc/RT] [10] where Go and Qc are generally temperature independent over wide ranges of temperature. Comparison of eq 9 with eq 10 gives:

Jan 1, 1952

By Y. Rama Murthy, Amit Ranjan, Pankaj Satija, Sunil K. Tripathy

"Chromite ore is one of the essential raw materials for ferroalloy and steel industry. Boula-Nuasahi belt of Sukinda region, Odisha, is the only resource of chrome ore in India. The Indian deposits are typically characterized as ferruginous and siliceous type ores. Chrome ore is generally mined by open cast mechanized mining method in these areas. The run-off-mine (ROM) ore from the various mines is sent to beneficiation plant via stack yard for further up gradation. The chrome ore beneficiation plant (COB) has two major circuits: comminution (feed preparation) circuit and beneficiation circuit. In the comminution circuit, the ore is reduced to <1 mm size ( for complete liberation of chromite) from -220 mm size in stage wise crushing (primary and secondary crusher) followed by screening and single stage grinding in a ball mill. Fraction having size <1 mm is further upgraded in the beneficiation circuit to produce three different concentrates having specific chemical composition and Cr: Fe ratio. The beneficiation plant is a complex circuit comprising of combination of unit operations (spiral concentrator, Floatex density separator, hydrocyclone, wet shaking tables). In the recent years the overall performance efficiency (grade and recovery) of the beneficiation circuit got affected with numerous changes in the run-off-mine ore quality (physical and chemical), mineralogical composition, etc.In this article, attempts made for improving the performance of chrome ore beneficiation circuit is discussed. The approach adopted and practical problems encountered while the process were also highlighted. The results of the trials has shown an improvement in the performance of the cyclone (d50 : 56 µm – 40 µm) and chromite losses into tailing to <12%.INTRODUCTIONChromite has a wide range of application, from refractories to metallurgical industries, and from chemical to non-ferrous alloy industries. Chromium ore (Fe2+, Mg) [Cr, Al, Fe3+]2O4) occurs exclusively in rocks formed by the intrusion and solidification of molten lava or magma which is very rich in the heavy, iron containing minerals such as pyroxene and olivine. With the richness of silica and iron, the chrome ores are broadly classified as siliceous type and ferruginous type. Pure chromite with composition FeCr2O4 is rare because magnesium usually substitutes for ferrous ion and aluminum and ferrous ion substitutes for chromium. Generally the associated gangue minerals are talc, quartz, hematite, goethite, limonite, gabro, serpentine, anorthosite, dunite, pyroxinite, etc. Around 96% of the chromite ore production is consumed for metallurgical applications of which 94% is converted into different grades of ferrochrome i.e., high carbon ferrochrome and charge chrome as shown in Figure 1. Stainless steel being the leading consumer of ferrochrome, any change in the dynamics of the stainless steel industry impacts overall chrome ore demands significantly."

Jan 1, 2016

By A. E. Desrosiers, R. L. Kathren, D. L. Haggard, J. M. Selby

INTRODUCTION The radiological health significance of thorium-230 stems from its tendency to separate from the uranium238 parent, concentrate in bone tissues, and to subsequently irradiate the radiosensitive tissues lining the bone surfaces and the bone marrow. Indeed, thorium-230 may be the radionuclide which contributes the major dose following intake of natural uranium (Hartley and Pasternack 1979). This is reflected by the most recent recommendations of the International Commission on Radiological Protection, which specify the limits shown in Table I for the annual intake of radionuclides by occupationally exposed workers (ICRP 1979). TABLE 1. Occupational Annual Intake Limits (microcuries per year) for Selected Uranium Nuclides and Daughters (ICRP 79) [Radionuclide Ingestion Inhalation] [Uranium-238 200 0.05 Uranium-235 200 0.05 Uranium-234 200 0.03 Thorium-234 300 200 Thorium-230 3 0.02 Radium-226 2 0.5] Clearly, the relatively low annual limit of intake for thorium-230 shows it to be of greater radiological concern than its parent radionuclides. Because of the greater toxicity and different metabolism of thorium-230, monitoring only for uranium-238 does not satisfactorily identify the possible hazard from thorium-230 nor does it provide any real indication of the metabolism or biodynamics of these two radionuclides. Thorium-230 has a half-life of 80,000 years and can be detected by direct counting of the alpha particles or photons emitted during its transformation to radium-226. The 4.69 and 4.62 MeV alpha particles are distinctive and specific indicators of thorium-230 and are emitted with abundances of 76% and 24%, respectively. The principal photon, a 68 keV gamma ray, is emitted in only 0.37% of the transformations and is, therefore, not useful for low level measurements. The other photons emitted have even lower yields, or, in the case of radium L x-rays, are non-specific and, hence, useless for quantification. High sensitivity measurements of thorium-230 currently are usually accomplished by wet washing of the sample substrate, quantitative chemical separation of thorium atoms, and, finally, direct measurement of the alpha particles emitted from a massless deposition. This procedure is complicated, expensive, and time-consuming, and subject to interferences from uranium, other actinides, and other thorium isotopes. Recently, the feasibility of low-level measurement of thorium-230 by neutron activation analysis (NAA) was demonstrated (Kathren, Desrosiers and Church 1980). Two principal variations of the NAA method were used in this study: 1) instrumental NAA technique and 2) post-irradiation radiochemical separations (RCS). Instrumental NAA procedure is a nondestrucive technique which is preferred because of its simplicity. The procedure is as follows: after irradiation with a known neutron fluence, the samples are transferred to a clean container and quantitative gamma spectroscopy performed. With the radiochemical separations procedure, the sample is initially treated as in the instrumental technique. However, after irradiation, a known amount of "carrier" is added to the sample. The element(s) of interest are then separated from the rest of the matrix by distillation, precipitation and extraction techniques. The resulting sample, now free of interferring elements, is then ready for gamma-ray analysis. The use of a "carrier" is to determine the loss of element-of-interest during the chemical separations process. The neutron activation cross section of thorium-230 has an epicadmium resonance value of 1,010 barns (Mughahghab and Garber 1976) and a thermal neutron cross section of 23 barns. The 25.52 hr thorium-231 produced releases two photons of significance: an 84 keV complex, (6.5% yield) and 25.6 keV (15% yield) (Lederer and Shirley 1978). The 84 keV complex is particularly useful for quantification since neither natural uranium, thorium, their daughters, or activation products emit photons in this region. However, the higher yield of the 25.6 keV photon may result in increased sensitivity if there are no other photons of similar energy emitted by other radionuclides in the sample. PRELIMINARY STUDIES Thorium-230 standard stock solution was prepared from a pure sample of the oxide purchased from Oak Ridge National Laboratory. From this stock solution a series of samples were prepared for irradiation in the TRIGA Mark I reactor at Reed College. Various dilutions were prepared as well as thorium-230 spiked urine samples. Irradiation times varied from 1 to 54 minutes in a neutron fluence rate of 1.84 x 1012 n/ cu m-sec. The neutron spectrum was abundant in thermal neutrons, having a Cd ratio of approximately 10. Treated urine samples were also analyzed by the NAA instrumental method. Analysis of untreated urine samples was not possible due to the high background

Jan 1, 1981

By Spottiswood DJ, Kelly EG, Bull WR

An empirical model of a 254mm water-only cyclone treating minus 17001im coal has been developed. The model includes the prediction of by-passing, SG50 corrected, water split and the use of a single equation forthe reduced performance curve of particles in the range 1700-106pm. Particles finer than 106pm did not separate according to specific gravity. Simulations using the model showed the water-only cyclone performance to be dependent more upon feed size distribution than washability when the feed contains high-ash fines, but that, with a deslimed feed, the yield and ash content of the over-flow should be controllable by altering feed rate or feed pulp density.

Jan 1, 1985

By Bernard H. Coyle, Robert N. Anderson, Judith A. Koperski

Between 6 and 8 million cars are shredded annually in the U.S. This shredded material is magnetically separated leaving a mixture of nonmetallics and nonferrous metals which can be further separated into its metallic and nonmetallic constituents. The metallic portion is a mixture of metalr; and alloys containing predominately zinc, and some aluminum, copper, and residual iron. Hydrometallurgical strategies are proposed for recovering the aluminum, copper, and zinc. The thermodynamic and kinetic behavior of the most promising alternatives to separation are discussed.

Jan 1, 1978

By P. R. Hervey, J. W. Foise

Quartz is the second most widely used electronic material, with approximately 3,000 t (3,300 st) of single crystals used each year. Most of the quartz used in electronics is synthetic and is manufactured using a technique called hydrothermal growth. This paper reviews the basics of hydrothermal growth, the feedstocks used for these crystals and the challenges to improving the growth process. This paper also reviews the processing of these crystals, the markets and applications where they are used, and the global outlook for these markets.

Jan 1, 2002

By Brian Hawkins, David Bradley, Rachel Gooding, Lucas Rodriguez, Patrick Wilson, Kevin Wegman, Dhruv Mittal

To better protect the nation and inform difficult defense and preparedness decisions, the Department of Homeland Security (DHS) is performing a probabilistic risk assessment of explosive terrorism risk to the US homeland. After assessing existing models against requirements, well-developed mathematical descriptions and algorithms associated with explosives hazards were combined with probabilistic modeling techniques used in DHS chemical, biological, radiological, and nuclear (CBRN) terrorism risk assessments to create a suitable interior explosive consequence model for assessing the impact to human health.

Jan 1, 2018

By Betty M. Miller

INTRODUCTION In September 1974 the Resource Appraisal Group of the Branch of Oil and Gas Resources, U. S. Geological Survey, was asked to aid the Federal Energy Administration (FEA) in its legal responsibility to generate by June 1975 an independent appraisal of the undiscovered oil and gas resources of the United States, both onshore and offshore. As a result, the Resource Appraisal Group accelerated its ongoing efforts to develop sound appraisal methods and procedures that would permit the systematic collection and evaluation of basic data from petroleum provinces throughout the Nation. The results of this independent appraisal of the Nation&apos;s conventional oil and natural gas resources were published by the U. S. Geological Survey as Circular 725, June 1975 (1). The estimates in this report of undiscovered recoverable oil and gas resources for the United States were made: (1) by carefully re- viewing a large amount of geological and geophysical information gathered on more than 100 different provinces by over 70 geologists within the Survey; (2) by applying a variety of resource appraisal techniques to each potential petroleum province; and (3) by application of subjective probability procedures to group appraisals for each petroleum province. These methods provide a range of estimates which are summarized in terms of low, high and mean values for the various provinces and groups of provinces, or regions. The basic data and procedures used are documented and are being incorporated into a dynamic, data- intensive system that can be upgraded, updated and reevaluated periodically. These data are open to public inspection. In this study, the primary emphasis was placed on conventional crude oil and natural gas in the onshore provinces and the provinces on the continental shelf out to water depths of 200 metres; estimates

Jan 1, 1977