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By James W. Halley

"Fine-grained" steels have been standard products for many years. This paper describes an investigation of the effects of some of the more common grain-growth inhibitors used to produce these steels. The properties, response to heat-treatment, and the deoxidation practice necessary to produce "fine-grained" steels have been the subjects of numerous papers. It has been shown that a substantial quantity of aluminum is necessary to produce this type of steel, and it has been indicated that titanium and zirconium produce a similar result. For some time there was a lively argument between those who felt that aluminum refined the grain and those who felt that aluminum oxide was the effective agent. This argument subsided with an apparent victory for those promoting aluminum oxide. Most of the work on grain size has divided steels into two broad classifications based upon McQuaid and Ehn's original carburizing test at 1700°F or some similar test at this temperature. If the steel coarsened at this temperature, it was coarse grained, while if it did not it was he grained. Quantitative data concerning how much a certain amount of aluminum raised the coarsening temperature and affected the properties, all other factors being constant, are very meager. Similar information regarding titanium and zirconium is practically nonexistent. The published information on aluminum yields some quantitative information. Epstein, Nead and Washburn' gave the first quantitative information on the amount of aluminum necessary to prevent grain growth at 1700°F. They showed that the addition of a pound to a pound and a half of aluminum per ton (0.05 to 0.075 per cent) would assure steel that would not coarsen during a McQuaid-Ehn test at 1700°F. Herty, McBride and Houghg reported that a steel containing 0.020 per cent acid-soluble aluminum coarsened between 1742° and 1832°F., while steels containing 0.016 per cent and 0.017 per cent acid-soluble aluminum coarsened between 1652° and 1742°F. compared with similar aluminum-free steels that started to coarsen at 1580°F. Herty3 also showed a marked improvebent in low-temperature notched impact resistance with increasing acid-soluble aluminum up to 0.030 per cent. McQuaid4 compared the coarsening of steels containing 0.011 to 0.204 per cent total aluminum and found a maximum coarsening temperature at 0.098 per cent aluminum. He had no samples between 0.015 and 0.098 per cent aluminum. Houdrement6 found a marked decrease in coarsening temperature with aluminum contents above 0.50 per cent. Kinzel, Crafts and Egan6 found that aluminum and zirconium improved the low-temperature impact resistance of a low-alloy steel, while titanium was detrimental. Sims7 shows that sulphur increases the coarsening temperature and that the maximum coarsening temperature is between 0.04 and 0.06 per cent added aluminum. The fact that the coarsening temperature passes through a maximum with increasing aluminum content appears to be well established. The aluminum content that

Jan 1, 1947

By R. W. Gurry

In the course of a series of measurements of the rate of diffusion of carbon in austenite at about 960°C. (1760°F.) and III0°C. (2030°F.), it became necessary to determine carbon concentration when austenite is saturated with graphite, because of indications that the hitherto accepted value of the solubility of graphite in this temperature range is too low. Specimens of carbonyl iron were saturated at temperature in a hydrogen-toluene atmosphere which was in equilibrium with graphite. The carbon conteDt of the steel was measured either by combustion or by the loss of weight after substantially complete decarburization in an atmosphere of hydrogen saturated at room temperature with water vapor. The absence of carbide or graphite as a phase at temperature was demonstrated by microscopic examination of saturated specimens that had been drastically quenched, The mean result of concordant ments-—taeats-is-^that gamma iron saturated with respect to graphite contains 1.39 per cent carbon at 957°C .(i755°F.) and 1.89 per cent at III0°C. (2030°F.). These two values, plotted along with what appears from the literature to be the best value (0.69 per cent) at the iron-graphite eutec-toid temperature (738°C. or I360°F.), yield a curve that is almost linear, and which, continued upward through a short interval, leads to a limiting solubility of 1.98 per cent at the iron-graphite eutectic temperature II35°C. (2075°F.) instead of the commonly accepted value of about 1.7 per cent. Experimental Procedure The samples of carbonyl iron, with the exception of those for microscopic examination, were cylinders 3.8 cm. long, 0.475 cm. in diameter. They were placed in a graphite holder and suspended in a vertical furnace, the temperature of which was controlled precisely by means of the resistance of the platinum winding.' In the empty furnace the range of temperature over a distance of 4 cm. at the heat center was 9°C. (I6°F.) but at a given location the temperature remained within +o.5°C. (o.9°F.); it was measured by inserting a platinum-rhodium couple calibrated at the melting point of palladium, 1555°C. Since the presence of the sample probably improved the uniformity of temperature, it was decided to consider the average temperature of the sample to be 3°C. lower than the maximum measured in the empty furnace with the same setting of the controller. Data were obtained in the vicinity of two temperatures, 960°C. (I760°F.) and III0°C. (2030°F.). The samples for microscopic examination, 2.2 cm. long, 0.475 cm. in diameter, were suspended individually on pure iron wire; for these no temperature correction was applied. The carburizing gas passed into the furnace was dry, oxygen-free hydrogen saturated at room temperature with toluene; this quickly produced on the

Jan 1, 1942

By R. W. Gurry

In the course of a series of measurements of the rate of diffusion of carbon in austenite at about 960°C. (1760°F.) and III0°C. (2030°F.), it became necessary to determine carbon concentration when austenite is saturated with graphite, because of indications that the hitherto accepted value of the solubility of graphite in this temperature range is too low. Specimens of carbonyl iron were saturated at temperature in a hydrogen-toluene atmosphere which was in equilibrium with graphite. The carbon conteDt of the steel was measured either by combustion or by the loss of weight after substantially complete decarburization in an atmosphere of hydrogen saturated at room temperature with water vapor. The absence of carbide or graphite as a phase at temperature was demonstrated by microscopic examination of saturated specimens that had been drastically quenched, The mean result of concordant ments-—taeats-is-^that gamma iron saturated with respect to graphite contains 1.39 per cent carbon at 957°C .(i755°F.) and 1.89 per cent at III0°C. (2030°F.). These two values, plotted along with what appears from the literature to be the best value (0.69 per cent) at the iron-graphite eutec-toid temperature (738°C. or I360°F.), yield a curve that is almost linear, and which, continued upward through a short interval, leads to a limiting solubility of 1.98 per cent at the iron-graphite eutectic temperature II35°C. (2075°F.) instead of the commonly accepted value of about 1.7 per cent. Experimental Procedure The samples of carbonyl iron, with the exception of those for microscopic examination, were cylinders 3.8 cm. long, 0.475 cm. in diameter. They were placed in a graphite holder and suspended in a vertical furnace, the temperature of which was controlled precisely by means of the resistance of the platinum winding.' In the empty furnace the range of temperature over a distance of 4 cm. at the heat center was 9°C. (I6°F.) but at a given location the temperature remained within +o.5°C. (o.9°F.); it was measured by inserting a platinum-rhodium couple calibrated at the melting point of palladium, 1555°C. Since the presence of the sample probably improved the uniformity of temperature, it was decided to consider the average temperature of the sample to be 3°C. lower than the maximum measured in the empty furnace with the same setting of the controller. Data were obtained in the vicinity of two temperatures, 960°C. (I760°F.) and III0°C. (2030°F.). The samples for microscopic examination, 2.2 cm. long, 0.475 cm. in diameter, were suspended individually on pure iron wire; for these no temperature correction was applied. The carburizing gas passed into the furnace was dry, oxygen-free hydrogen saturated at room temperature with toluene; this quickly produced on the

Jan 1, 1942

By P. R. Sperry

Aluminum alloy 3S was examined to determine the relationship between the applicable phase diagram and the microstructures produced under conditions tending toward non-equilibrium as well as equilibrium. The only phases present in the alloy were aluminum, (Mn,Fe)A16, aAI(Mn,Fe)Si, and silicon. It was found that the alloy behaved essentially as though it belonged to the Al-Mn-Si system. The effects of various solidification rates and of various homogenization treatments were studied. Particular attention was given to a nonuniform precipitate distribution which reflected the solute distribution in the cast aluminum solid solution. Some possible explanations for the solute distribution based on characteristics of the Al-Mn-Si or AI-Mn systems were discussed. ALUMINUM alloy 3s (aluminum plus 1.25 pct Mn) is one of the oldest and the most widely used of the wrought aluminum alloys. In spite of this fact, there are some aspects of the metallography of 3s which have not been revealed or which have not been brought to the attention of the metallurgist working with the alloy. The purpose of this investigation was to demonstrate in general how some of the fundamentals of equilibrium phase diagrams may be applied to the study of an alloy and to show specifically the origin of certain metal-lographic features of aluminum alloy 3s. It is hoped that this information can serve as a basis for solving specific problems relating microstructure to such things as mechanical properties, recrystallization, grain growth, and surface finish. Literature Survey Considering that alloy 3s contains only one alloying addition, manganese, it seems reasonable that it would have characteristics pertaining to the A1-Mn system. However, those characteristics might well be altered by the impurities in the commercial alloy, especially iron and silicon. Therefore, not only the simple binary but also more complex systems had to be reviewed. One characteristic of the A1-Mn system worthy of mention is its sluggishness, which leads to conflicting solubility data1,'2 and makes it relatively easy to supersaturate the solid solution and shift the eutectic composition during rapid solidification.3'4 Another characteristic is the strong influence which small amounts of impurities, especially iron, exert in lowering the solid solubility of manganese. Undissolved manganese is ordinarily present in an intermetallic phase, MnAl,, having an orthorhombic crystal structure. A metastable precipitating phase designated as "G" has been reported but small amounts of iron and silicon suppress this phase entirely." Raynor has established that MnA1, can dissolve a considerable amount of iron by a substitution of iron for manganese atoms up to 50 atomic pct.'" Consequently, in the Al-Mn-Fe liquidus diagram, the eutectic valley separating the primary aluminum and (Mn,Fe)Al, fields is fairly close to a constant Mn + Fe content." This is an important consideration in regulating the composition of 3s and also other aluminum alloys with high manganese additions."

Jan 1, 1956

By J. M. De La Garza Cárdenas

The total production of oil in Mexico in 1942 was 34,715,547 bbl.; that is, 8,179,254 bbl. less than the amount produced in 1941. The daily average for 1942 was 95,111 bbl., while that for 1941 was I17,520 barrels. The reduction of 22,409 bbl. per day in production was due to a decrease in exports because of the scarcity of tankers. This reduction has affected principally the Northern and Southern districts of the Tampico area; in December the Northern district (Panuco) was producing less than zoo0 bbl. per day with 445 wells shut down and the Southern district (GolJen Lane) less than $000 bbl. per day with 140 wells shut down. Drilling.—During the year 14 wells were drilled in Mexico, of which five were productive. Seven dry holes were drilled in Ebano, Tampico Northern district, in search for new producing fractures; plans are being made to explore this area by means of resistivity measurements. Three producing wells were completed in Poza Rica. Drilling in this important field will continue during 1943 to determine its extension to the west. In the Isthmus of Tehuantepec no wells were completed. Two gas wells were completed in northeastern Mexico. Wildcats.-—Two wildcat wells were completed during the year, to test the extension of the Golden Lane to the south; both were unproductive. In the Isthmus a deeper sand was found in the El Plan field* This discovery has made an important increase in the reserves in this field and will allow an increase in the production of the area during 1943. In December, three exploration wells were being drilled, two in the Tampico area and one in the Isthmus. Operating Technology.— Satisfactory progress has been made in Mexico in the application of petroleum engineering to the development and operation of the oil fields. Mud control is now an established practice and has lessened the number of fishing jobs and has allowed more open hole to be carried than formerly was possible. During 1942 the first effective use was made of determinations of bottom-hole pressure, and other modern production practices such as gun-perforating and squeeze-cementing are being introduced. Plans for 1943 The oil production during 1943 is expected to be slightly larger than that of 1942, especially if tankers are available for exporting crude and refined products. The Government proposes to increase exploration and drilling activities to increase the reserves and potential production of the country; about 50 wells will be drilled during the year, including wildcat wells in northeastern Mexico, Tampico and Isthmus areas and exploitation wells in the Ebano-Panuco, Poza Rica, El Plan and Cuichapa fields. Acknowledgment Mr. A. Barnetche, of the Production Department, "Petr6leos Mexicanos," assisted the writer in the preparation of the statistical data.

Jan 1, 1943

By J. M. De La Garza Cárdenas

The total production of oil in Mexico in 1942 was 34,715,547 bbl.; that is, 8,179,254 bbl. less than the amount produced in 1941. The daily average for 1942 was 95,111 bbl., while that for 1941 was I17,520 barrels. The reduction of 22,409 bbl. per day in production was due to a decrease in exports because of the scarcity of tankers. This reduction has affected principally the Northern and Southern districts of the Tampico area; in December the Northern district (Panuco) was producing less than zoo0 bbl. per day with 445 wells shut down and the Southern district (GolJen Lane) less than $000 bbl. per day with 140 wells shut down. Drilling.—During the year 14 wells were drilled in Mexico, of which five were productive. Seven dry holes were drilled in Ebano, Tampico Northern district, in search for new producing fractures; plans are being made to explore this area by means of resistivity measurements. Three producing wells were completed in Poza Rica. Drilling in this important field will continue during 1943 to determine its extension to the west. In the Isthmus of Tehuantepec no wells were completed. Two gas wells were completed in northeastern Mexico. Wildcats.-—Two wildcat wells were completed during the year, to test the extension of the Golden Lane to the south; both were unproductive. In the Isthmus a deeper sand was found in the El Plan field* This discovery has made an important increase in the reserves in this field and will allow an increase in the production of the area during 1943. In December, three exploration wells were being drilled, two in the Tampico area and one in the Isthmus. Operating Technology.— Satisfactory progress has been made in Mexico in the application of petroleum engineering to the development and operation of the oil fields. Mud control is now an established practice and has lessened the number of fishing jobs and has allowed more open hole to be carried than formerly was possible. During 1942 the first effective use was made of determinations of bottom-hole pressure, and other modern production practices such as gun-perforating and squeeze-cementing are being introduced. Plans for 1943 The oil production during 1943 is expected to be slightly larger than that of 1942, especially if tankers are available for exporting crude and refined products. The Government proposes to increase exploration and drilling activities to increase the reserves and potential production of the country; about 50 wells will be drilled during the year, including wildcat wells in northeastern Mexico, Tampico and Isthmus areas and exploitation wells in the Ebano-Panuco, Poza Rica, El Plan and Cuichapa fields. Acknowledgment Mr. A. Barnetche, of the Production Department, "Petr6leos Mexicanos," assisted the writer in the preparation of the statistical data.

Jan 1, 1943

By R. F. Mehl, M. L. Swanson, J. P. Hirth, G. M. Pound

The sorption of zinc in a-brass was investigated. By analyzing the elliPtically polarized light reflected from the surfaces of a-brass solid-vapor couples, their optical constants and surface compositions were determined. It was shown that the surface compositions approached an equilibrium value exponentially with time. The bulk diffusion coefficient, surface diffusion coefficient and condensation coefficient for a-brass solid-vapor couples were calculated near 350°C from the experimental results. It has been assumed during various diffusion experiments with bimetallic solid-vapor couples (Dunn,1 Seith and Kraus,2 Jenkins,3 Balluffi and Seigle,4 and others) that the surface composition of such couples reaches its equilibrium value immediately. Although this assumption is probably valid at elevated temperatures, it is certainly not expected to be confirmed at lower temperatures.* consequently, it was decided that a fundamental experiment, designed to solve this problem as well as to examine all aspects of bimetallic sorption, would be the investigation of the rate of approach to equilibrium of the surface composition of solid-vapor couples. The sorption (i.e., adsorption, absorption, and desorption) of metals is obviously dependent on the state of the metallic surface and substrate. Much recent theoretical8"9 and experimental10-12 work on condensation and evaporation of pure metals has shown that the surface contour in atomic dimensions is a critical factor in sorption, so that any deformation or contamination13,14 of the surface exerts a profound influence. It is also known that absorption or diffusion occurs much more rapidly along grain were and in other deformed regions. Thus an effective sorption experiment must utilize a strain-free and clean surface, or at least a clearly defined and reproducible surface state. A mechanically polished surface is in a highly deformed17-l9 state. Because the degree of this deformation is uncertain, it is clear that for most experimental purposes a strain-free (i.e., etched, chemically polished, or electropolished) surface is desirable. Although adsorbed gas molecules on metallic surfaces are unavoidable even in the highest vacuum,20 oxide films can in some cases be removed by suitable vacuum annealing. Electrolytically polished and annealed surfaces, being relatively without strain, smooth and free of bulk oxides, were considered adequate for the present sorption experiments. Measurement of the surface composition of a solid-vapor couple presents a unique problem. Since experiments to determine the approach to equilibrium of the surface composition must be performed at temperatures where diffusion is very slow, solute penetration is so small that conventional analytical methods are unsuitable. An optical reflectivity method utilizing a polarization spectroscope was employed because of the relatively small penetration (~200A) of visible light in metals. By this method, elliptically polarized light reflected from the metal surface was analyzed and compared with that reflected from samples of known composition. The a phase of the copper-zinc system was

Jan 1, 1962

The production of bituminous coal in Pennsylvania started a few years after that in Virginia. For forty-five years all the production came from the Pittsburgh bed, and since that time its total output in value has exceeded that of any other single mineral deposit in the world. It has been impossible to separate the outputs of the four counties producing it at that time, as none of it came from Greene County until many years later, and the figures in the first column of Table 49 represent all the coal coming from Allegheny, Fayette, Washington and Westmoreland Counties. The years prior to 1790 are grouped on account of space limitations, they were made in the same manner as the later ones. All outputs have been rounded into hundreds, or thousands of tons because their accuracy does not warrant further subdivision. All figures are net tons, and all represent production and not shipments; this distinction in the early days when transportation facilities were scarce meant a great deal. In the period 1820-1835, more coal was used in salt-making in western Pennsylvania than for any other single purpose, but not a ton of such coal was handled by any transport agency. Later the iron industry used large quantities of local coal without shipment, and other instances can be cited. Prior to 1832, the only means of shipment of coal, excepting by wagons, was by the rivers, and this tonnage was larger than is generally believed. In Table 23, all of the available data about production from the Pittsburgh bed have been shown. After 1844 the river shipments above Pittsburgh are accurately reported by the Monongahela Navigation Company, but its figures do not include tonnage loaded on the river below Dam No. 1, nor that brought down Saw Mill Run. Prior to 1844 no tonnage records were kept although the amount moved was considerable. The information about local production, not shipped, whether used domestically or for manufacturing is necessarily incomplete. Column 7, shipments by railroad, river and canal, includes the figures in column 4, and until at least 1857 are quite evidently too low. The figures in column 9 are largely estimates based on the data shown and prior to 1870, and perhaps later, are much more likely to be lower, rather than higher, than the actual ones. This table was prepared over a period of years, and as additional data were found the revisions necessary were almost always upward. Clearfield County was the next producer in the state, shipments-in small amounts going down the Susquehanna River; some were made later via the West Branch Canal. Railroad facilities became available in 1862, and the growth in production increased rapidly and steadily. All available data and the estimated totals are in Table 24. Butler County began mining a few years after Clearfield, the early output

Jan 1, 1942

By O. A. E. Jackson

In all three principal dunite dikes, minute disseminated grains of pyrrhotite are to be seen in the fresh rock in amounts less than 1 pct. Traces of nickel have been reported from many such specimens. Nickeliferous pyrrhotite is the principal nickel-bearing mineral, although olivine contains very minor amounts of nickel in solid solution, substituting for magnesium and iron. Such occurences are of no economic significance. Nickel analyses of possible economic significance (up to 3 pct Ni) have been returned from samples taken near the margins of the two large western dunite dikes. These samples were from schistose zones along and within the dike walls. The mineralized shear zones, up to 30 ft wide, strike N 30 W and dip from 65" to 90" west. They consist of layers of hard, fresh, green-gray dunite, mineralized with finely disseminated pyrrhotite. These layers alternate with layers of soft, black, oxidized schistose dunite up to 1 ft wide, well mineralized with pyrrhotite and tabular-fractured pentlandite in crystals up to 1/4-in. diam. Sulfides identified in polished sections were chiefly pyrrhotite and pentlandite, with much smaller amounts of chalcopyrite. There are two varieties of pyrrhotite. One develops a fine polish and conforms to published descriptions. The other is slightly darker and browner, does not polish as well, and displays flamelike growth or intergrowth structure when etched with nitric acid fumes. Pyrrhotite and pentlandite crystallized contemporaneously in the form of veins in the olivine and as fillings of the interstices of euhedral olivine crystals. Chalcopyrite replaces the other sulfides. The most striking feature of the deposits is their occurrence in intimately fractured and schistose zones in the dunite where faulting has been localized along the contacts of the sediments and the dunite dikes. Though far too little detailed work has been done to be certain, there appears to be a higher concentration of sulfides in the shear zones along the margins of dike apophyses, sediment reentrants, and strike and dip irregularities. Much more study would be required in order to submit a plausible theory of formation of the nickel deposits, supported by laboratory and field evidence. Results of the present study, especially localization of the metals by fracturing, certainly indicate that the nickel concentrations were brought about by precipitation of nickel sulfides from fluids that used the fractures as a means of ingress. Studies of nickel sulfide deposits throughout the world have demonstrated that such sulfides were deposited at considerable depths and at temperatures of several hundred degrees. This would rule out the possibility of groundwater as the transporting medium. The two possible sources of hot nickel-bearing fluids are quartz diorite and gabbro. The first possibility is that large volumes of fluids, expelled from the quartz diorite magma during its crystallization, obtained a supply of nickel by altering the nickeliferous olivine of the dunites. The likelihood of this process is lessened by the recognition that, although the dunites show some alteration to serpentine, talc, and trem-olite, they are for the most part remarkably fresh outside the schistose zones. The worldwide association of nickel with basic rocks, together with the presence of pyrrhotite in thin films on joint planes of the gabbro on Jumbro Mountain, favors the gabbro source. The writer extends thanks to C. Phillips Purdy, Jr., project chief of American Metal Climax Inc., for giving him the opportunity to study these deposits. He is also indebted to the officers of American Metal Climax Inc. and Discovery Mines Inc. for permission to publish this article.

Jan 1, 1961

By E. P. Valby

A method is presented in which the hydrocarbon weight fractional analyses of the well effluents from a true gas-tip-portion well and a true dark-oil-ring well furnish the basic data for determining the properties of any mixture of production from the two sources. The effect of injected gas on these properties can be computed. Comparison of similar properties of a given well effluent to the properties of mixtures of the basic well effluents will give the relative amounts of production from the two sources and injected gas. Presentation of Problem Oil fields having an area of gas-cap, or light-oil production, over the dark-oil area present a problem of increasing gas-oil ratios during the production of an oil well. The use of pressure maintenance in a field such as above, or a cycling project, wherein residue gas from an absorption plant is injected into the reservoir, adds to the problem, because a part of the increased gas-oil ratio may be due to the production of injected gas. The usefulness of a means to determine the amount of injected gas in a well effluent led to the development of a method in which well-effluent analyses of a true gas-cap-portion well and a true dark-oil-ring well furnish the basic data needed to determine whether injected gas of a known composition is present in a given well effluent. Development op Method The method presented herein is based upon the principle that if two fluids, or well effluents, having different properties are mixed in known proportions by weight, • the properties of their mixtures can be computed. Conversely, if these basic data are known, the properties of a mixed well effauent can be used to compute the relative amounts of each type of well effluent in the mixture. A convenient way of making these computations is to have the fractional analyses of the base well effluents and mixtures expressed in weight fractions The mathematical equations expressing the properties of mixtures will be developed for a two-component system consisting of the well effluents from the dark-oil ring and the gas-cap area. The added effect of injected gas, a third component, will also be developed. The following assumptions have been made in developing the method: 1. The gas-cap and dark-oil fluids exist initially in separate portions of the reservoir. 2. The fluids in each portion exist in a single phase at pressures above the saturation pressure for the dark-oil fluid. 3. The composition of each fluid is substantially constant throughout its respective portion of the reservoir. 4. Changes in reservoir pressure and temperature will not change the composition of the fluid entering the well bore. 5. The injected gas displaces original gas-cap fluid 100 per cent. •• 6. AS the injected gas travels through a

Jan 1, 1945

By John C. McDonald

This report is a continuation of an earlier one with a similar titlell to which the reader is referred for such details of procedure as do not appear here. A brief summary will be given of the objects and methods of the research at this time. It was pointed out in the previous paper that properties of magnesium alloys, as reported by different workers, often differ. This was thought to be due to differences in the preparation of the specimens tested. The ideal way to compare alloys would be to prepare each of them in so many ways that reasonable assurance would exist that the best properties of each alloy had been developed. The alloys would then be compared on the basis of these properties. The treatment actually given the alloys of these studies is necessarily a compromise with completeness. Nevertheless, enough treatments have been given so that a considerable degree of assurance exists that the properties obtained are representative of the alloy as such, and not merely of the manner in which it was worked and heat-treated. Procedure The alloys were cast in the form of cylindrical ingots, which were extruded into bar before being rolled. Some of the ingots were made in 2-in. diameter molds whereas all of the alloys of the previous paper were cast in 4.in. molds. For the larger ingots, the extruded bar was 2 by 0.140 in.; lor the 2-in. ingots, it was I by 0.100 in. Preliminary work indicated that strip rolled from both sizes of the bars would have substantially the same properties. The bar was rolled at some temperature in the range 400' to 800°F., for a total reduction of 50 per cent. The temperature and reduction per pass were so selected that no cold cracking occurred, but were such as to leave the metal in a semihard condition. The hot-rolled metal was annealed at several temperatures in the range 400' to 800°F.; it was also cold-rolled, and given solution and aging treatments. Two standard test bars corresponding to each treatment were pulled in tension, and the average was obtained. Twelve treatments, in all, were given each alloy. The properties of the annealed specimen having the highest percentage of elongation were selected to represent the alloy in the annealed condition. The properties of the cold-rolled specimen having the highest yield strength were selected to represent the alloy in the cold-rolled condition. These properties were then plotted against percentage composition by analysis. Smooth curves were drawn among the points, thus graphically averaging the results over the entire range studied. The percentage of elongation in the cold-rolled state was not plotted, since it was always low, and varied in an irregular manner with composition. Alloys Studied Magnesium-calcium Alloys.—Fig. I presents the results on the alloys of the rnag-

Jan 1, 1941

By John C. McDonald

This report is a continuation of an earlier one with a similar titlell to which the reader is referred for such details of procedure as do not appear here. A brief summary will be given of the objects and methods of the research at this time. It was pointed out in the previous paper that properties of magnesium alloys, as reported by different workers, often differ. This was thought to be due to differences in the preparation of the specimens tested. The ideal way to compare alloys would be to prepare each of them in so many ways that reasonable assurance would exist that the best properties of each alloy had been developed. The alloys would then be compared on the basis of these properties. The treatment actually given the alloys of these studies is necessarily a compromise with completeness. Nevertheless, enough treatments have been given so that a considerable degree of assurance exists that the properties obtained are representative of the alloy as such, and not merely of the manner in which it was worked and heat-treated. Procedure The alloys were cast in the form of cylindrical ingots, which were extruded into bar before being rolled. Some of the ingots were made in 2-in. diameter molds whereas all of the alloys of the previous paper were cast in 4.in. molds. For the larger ingots, the extruded bar was 2 by 0.140 in.; lor the 2-in. ingots, it was I by 0.100 in. Preliminary work indicated that strip rolled from both sizes of the bars would have substantially the same properties. The bar was rolled at some temperature in the range 400' to 800°F., for a total reduction of 50 per cent. The temperature and reduction per pass were so selected that no cold cracking occurred, but were such as to leave the metal in a semihard condition. The hot-rolled metal was annealed at several temperatures in the range 400' to 800°F.; it was also cold-rolled, and given solution and aging treatments. Two standard test bars corresponding to each treatment were pulled in tension, and the average was obtained. Twelve treatments, in all, were given each alloy. The properties of the annealed specimen having the highest percentage of elongation were selected to represent the alloy in the annealed condition. The properties of the cold-rolled specimen having the highest yield strength were selected to represent the alloy in the cold-rolled condition. These properties were then plotted against percentage composition by analysis. Smooth curves were drawn among the points, thus graphically averaging the results over the entire range studied. The percentage of elongation in the cold-rolled state was not plotted, since it was always low, and varied in an irregular manner with composition. Alloys Studied Magnesium-calcium Alloys.—Fig. I presents the results on the alloys of the rnag-

Jan 1, 1941

By J. H. Frye, J. W. Caum

One of the most important methods of increasing the hardness of metals is alloying. In spite of the widespread use of alloys, the fundamental mechanism of alloy hardening is little understood. This is true even of the simplest form of alloy, the primary substitutional solid solution in which solute atoms substitute for solvent atoms at random in the crystal lattice of the latter. A fundamental understanding of the mechanism of solid-solution hardening would be important, both because of the widespread use of alloys of this type and because of the light that might be thrown on the mechanism of hardening in more complicated alloys. The present investigation covers the hardness of primary substitutional solid solutions of zinc, gallium, germanium, and arsenic in copper. All of these are B subgroup elements from the third period of the periodic table. Cu, Zn, Ga, Ge, and As are in groups I, 11, 111, IV, and V, respectively, and lie just to the right of the transition elements. Because of their cost, it is unlikely that Ga-Cu and Ge-Cu solutions will ever be of commercial importance, but the intent of this study has been to discover some of the fundamentals of solid-sdstion hardening and it has been demonstrated that such a study is far more likely to be fruitful if alloys are made up on the basis of position of the constituent elements in the periodic table than if commercially important alloys are used. Many investigations have been made of the hardness of substitutional solid solutions, but lack of attention to purity of the metals, infrequent grain-size control, and the use of varying testing techniques have virtually invalidated any comparisons. The purpose of most investigations has been to solve immediate problems and not to solve the primary problem: Why do solid solutions usually have a greater hardness than pure metals? Previous Work Theories have been advanced to account for the hardness of solid solutions, but relatively little experimental work has been done. Perhaps the first experimental work of merit was that of Norbury1 on copper alloys. He showed that for several alloys there was a correspondence between the hardness of one atomic per cent solutions and atomic volume of the solutions. More recently Brick, Martin, and Anger2 have investigated the effect of various solute elements on the hardness of copper. They concluded that the hardening effect of these solutes depended "primarily on the difference in atomic volume of solute and solvent or, relatedly, upon the parameter change of the solvent or the extent of solid solubility." They also observed the effect of these elements on hardness and texture developed by rolling and by annealing of rolled material. Austin3 examined the hardness and response to heat-treatment of various iron solid solutions. Frye and Hume-Rothery4 measured the hardness of a number of silver solid solutions containing approximately 1 References are at the end of the paper.

Jan 1, 1943

By J. H. Frye, J. W. Caum

One of the most important methods of increasing the hardness of metals is alloying. In spite of the widespread use of alloys, the fundamental mechanism of alloy hardening is little understood. This is true even of the simplest form of alloy, the primary substitutional solid solution in which solute atoms substitute for solvent atoms at random in the crystal lattice of the latter. A fundamental understanding of the mechanism of solid-solution hardening would be important, both because of the widespread use of alloys of this type and because of the light that might be thrown on the mechanism of hardening in more complicated alloys. The present investigation covers the hardness of primary substitutional solid solutions of zinc, gallium, germanium, and arsenic in copper. All of these are B subgroup elements from the third period of the periodic table. Cu, Zn, Ga, Ge, and As are in groups I, 11, 111, IV, and V, respectively, and lie just to the right of the transition elements. Because of their cost, it is unlikely that Ga-Cu and Ge-Cu solutions will ever be of commercial importance, but the intent of this study has been to discover some of the fundamentals of solid-sdstion hardening and it has been demonstrated that such a study is far more likely to be fruitful if alloys are made up on the basis of position of the constituent elements in the periodic table than if commercially important alloys are used. Many investigations have been made of the hardness of substitutional solid solutions, but lack of attention to purity of the metals, infrequent grain-size control, and the use of varying testing techniques have virtually invalidated any comparisons. The purpose of most investigations has been to solve immediate problems and not to solve the primary problem: Why do solid solutions usually have a greater hardness than pure metals? Previous Work Theories have been advanced to account for the hardness of solid solutions, but relatively little experimental work has been done. Perhaps the first experimental work of merit was that of Norbury1 on copper alloys. He showed that for several alloys there was a correspondence between the hardness of one atomic per cent solutions and atomic volume of the solutions. More recently Brick, Martin, and Anger2 have investigated the effect of various solute elements on the hardness of copper. They concluded that the hardening effect of these solutes depended "primarily on the difference in atomic volume of solute and solvent or, relatedly, upon the parameter change of the solvent or the extent of solid solubility." They also observed the effect of these elements on hardness and texture developed by rolling and by annealing of rolled material. Austin3 examined the hardness and response to heat-treatment of various iron solid solutions. Frye and Hume-Rothery4 measured the hardness of a number of silver solid solutions containing approximately 1 References are at the end of the paper.

Jan 1, 1943

By G. R. Countryman, G. H. Thomas, I. Fatt

Procedures are described for scaling up turbulent friction pressure drops of inelastic fluids and laminar friction pressure drops of inelastic/elastic fluids in circular ducts. The laminar scale-up calculation method does not require the use of arbitrarily defined flow models, and can be performed directly from laminar capillary viscometry data, or by means of a numerical solution of the integral equation for Poiseuille flow, inputing true shear rate-shearing stress data from other simple shearing experiments. e.g. Couette and plane Poiseuille flows. The turbulent scale up procedure, based on the Dodge-Metzner correlation for inelastic fluids, requires evaluating the characteristic rheological parameters at the existing wall stress under laminar flow conditions. Since the wall stress is the quantity sought in a pipe flow scale-up problem, a trial-and-err or solution is indicated, and therefore a calculation method has been developed and programmed for a medium-size magnetic-drum-memory digital computer. Laminar flow data from other somces may also be adapted to this turbulent scale-upmetbod for inelastic fluids. Comparisons between predicted and experimentally observed friction losses for generalized Reynolds numbers ranging from 3 — 32.000 indicate an error band width of less than 10 per cent in circular tubes ranging in size from 0.5 to 1.89 in. in diameter. INTRODUCTION Considerable effort has been devoted to the problem of establishing an accurate procedure for predicting the laminar and turbulent flow behavior of non-Newtonian materials in pipes from data obtained with small scale laboratory or pilot plant equipment, as evidenced by the variety of reviews on the subject.1-3 Although there does not appear to be a universally accepted scale-up method for laminar flow, the various methods currently employed can be divided into two principal categories depending upon whether a functional relationship between shear rate and shearing stress is assumed, or whether a generalized correlation is adopted. The methods of Hedstrom4 and of Weltmann are typical of the first category, while the generalized correlation of Metzner and Reed6 illustrates the second category. Under appropriate conditions either approach yields a reasonably satisfactory description of laminar flow behavior. In contrast, the prediction of non-Newtonian transitional and turbulent friction losses is much more complex. For example, certain types of materials have been observed to sustain laminar flow at much higher flow rates than Newtonian liquids. When turbulence does develop, the level of agitation is damped to such an extent that the friction factor is lower than that of the original low viscosity Newtonian medium. Extremely broad transition regions and "diameter effects" have also been noted in the turbulent flow of certain chemically complex materials.' Aside from the latter complications, which are attributed in part to viscoelasticity, the existence of a variety of scale-up procedures for predicting transitional and turbulent friction losses is due principally to the chioce of viscosity parameter used in the definition of the Reynolds number. In contrast, Dodge and Metzner have demonstrated that the friction factorReynolds number relationship for inelastic fluids can be satisfactorily generalized provided the fluid property parameters are evaluated at the existing wall shearing stress.* However, the quantity usually sought in a pipe flow scale-up problem is the wall stress, and hence a trial-and-error procedure is required when dealing with non-power law materials. In this report we discuss two calculation methods: (1) for determining the wall stress, and hence the friction pressure drop, for inelastic fluids under conditions of turbulent flow by means of an iterative procedure, and (2) for predicting laminar friction pressure drops without recourse to arbitrarily defined flow models.

By L. K. Jetter, J. C. Ogle, C. L. McHargue

The preferred orientation developed in extruded aluminum rods has been studied as a function of extrusion temperature, extrusion speed, and position in the rod. Duplex <111> - <001> textures were developed for slow extrusion speeds at temperatures of 75° to 850° F and for fast extrusion speeds at temperatures of 75° to 450°F. The <001> texture component was shown to be primarily due to re crystallization occurring during fabrication. At fast speeds and billet temperatures of 650° to 850°F, complete recrystallization occurred and <115>or<118> textures were found. The nature of the structure has proved to be an important variable in the production and use of extruded products. Data on the flow of aluminum and aluminum alloys summarized by Smith,&apos; Hardy,&apos; and Locke3 show that the extruded product may possess a deformed, a partially recrystallized, or a completely recrystallized structure, and there may be important variations in the structure from surface to center and from front end to back end of an extrusion. This inhomogeneity in grain structure and preferred orientation affects both the mechanical properties of the extruded section and its response to subsequent heat treatment. Among the fabrication variables causing these differences are temperature, extrusion ratio, and extrusion speed. Preferred orientation which develops during the extrusion of rods is of the fiber-type and, in general, has been assumed to be similar to that developed in rods or wires fabricated by drawing through dies, rolling, or swaging.4 Because of the nature of the flow of metal being extruded, this assumption may be valid for material near the center of the rod but probably is not valid for the outer layers. In most face-centered-cubic metals a duplex fiber texture has commonly been reported in which the grains have either <Ill> or <001> directions parallel to the longitudinal axis. Sachs and fichiebold5 reported cold-drawn aluminum (99.6 pct pure) to have a single <Ill> fiber texture, and this observation was confirmed by Schmid and Wassermann.6 Hibbard, on the other hand, observed a minor <001> and a major <111> texture in 99.994 pct A1 cold drawn to 98 pct reduction in diam.7 He concluded that the <001> component resulted from a small amount of room-temperature recrystallization. Dayal8 concluded that the <001> component was an intermediate component which disappeared at high reductions in 99.5 pct pure Al. Kostron and schippers9 found duplex <111> — <001> textures in extruded 99.5 pct Al, as did Gow and cahnl0 and GOW11 for extruded rods of commercial and superpure aluminum. Similar duplex fiber textures have been reported for aluminum alloys by Kostron,= Unckel,13 and Van Horn.14 The present paper is concerned with the variation in texture with fabricating conditions and with the nature of the texture components in extruded aluminum rod (99.99 pct pure). EXPERIMENTAL PROCEDURE Ingots were chill-cast in cast-iron molds and machined to 3-in. diam by 6-in. long extrusion billets. These billets were upset to 3.125 in. diam and extruded through a flat-face 0.690-in. diam die. Quenching of the rods was accomplished by a water spray at the die face. The combinations of fabricat-

Jan 1, 1960

By H. H. Waller, D. L. Monroe, D. P. R. Smyth

ORIGINALLY slusher drift development was conventional, advancing the drift full 10x13-ft size at 6 ft per round. This proved dangerous and costly because the weak fractured rock of the orebody cannot be left raw for any length of time without its reacting with air to cause spalling. With this method of drift development ground conditions were worsened by many blasts and the length of time it was necessary for the drift to remain open. The problem was practically solved by driving 5x7-ft pilot drifts, then enlarging to 10x13-ft raw size by slashing with 50 to 60-ft diamond drill holes. Calculated explosive charges are used to avoid wall and roof shattering. Still more recently jacklegs and sectional steel replaced the diamond drills, resulting in the same work for less than half the cost, see Figs. 1 and 2. Crosscuts at right angles to the main haulage drifts, which cross through the center of the ore-body, divide the ore into 200-ft blocks. Eight slusher drifts on 50-ft centers draw from one block. The slusher floor is 11 ft above the rail, inclined at 1 pct grade, and the top is 18 ft below the undercut. The finished slusher drift is 102.5 or 120 ft long, depending on whether it is north or south of the crosscut. Drawpoints are staggered on each side at 17.5-ft centers. The completed drift is concreted and has an inclined ramp and a vent at the back. When crosscuts and gangways have been excavated and concreted and slusher stations installed, slusher drift development is carried out in 10 clear-cut stages: 1—Pilot drift is driven full length, 5x7 ft. 2—Station is excavated at half-way point in pilot drift. 3—Pilot drift is enlarged to slusher drift size by drilling. 4—Lower part of slash is blasted and mucked for floor concrete. 5—Floor is poured. 6—Inside upper half of slusher drift is blasted, mucked, and concreted. 7—Outside upper half of slusher drift is blasted, mucked, and concreted. 8—Drawpoints are excavated. 9—Drawpoints are concreted. 10—Concrete ramp and sheave protection recess are constructed at back end of slusher drift. A 5x7-ft pilot drift is driven full length on the center line of the proposed slusher drift using a jackleg and detachable bits on 7/8-in. hexagon steel. A 10-hp air-driven slusher and 36-in. scraper pull muck directly into cars in the main haulage crosscut, see Fig. 2. Operations are on a three-shift basis. One man mucks the previous round, drills, and blasts his round in a shift. Average advance is 5 ft per shift. When the pilot drift is completed, 7 ft of drift in the center section is enlarged to full raw slusher drift size of 10x13 ft, for the set-up to drill long slash holes. The pilot drift is slashed to full raw size from this station with jacklegs and tungsten carbide chisel bits on 7/8-in. sectional steel. Crews line the holes without trouble, working from controls established by the engineers, and the pattern of holes for the slash is arranged to give the least shattering effect

Jan 1, 1955

By H. S. Price, D. A. T. Donohue

The system of equations describing displacement of a hydrocarbon liquid by a hydrocarbon vapor in a porous medium where mass transfer takes place between the phases is solved numerically for a variety of gas injection processes. Even though the method of solution is quite general, only systems with three hydrocarbon components are considered. Computer simulations of displacement processes wherein mass transfer between phases is both considered and neglected are compared, and it is shown that neglecting mass transfer can give pessimistic displacement efficiencies. INTRODUCTION The role of the gas displacement process in the recovery of petroleum has been subjected to a series of detailed analyses; as a result, a number of predictive models have been published in the literature. However, because of major simplifying assumptions, most of these models do not completely represent the physical system. As a result, the effect of making the simplifying assumptions is unknown. Therefore, a complete representation of this process — one without major simplifying assumptions —should lead to a fuller understanding of the process, and perhaps to methods of improving it. The general method of developing a model for two-phase fluid flow in a porous medium is to solve simultaneously the continuity equation, the energy equations and the equation-of-state for each phase under the prescribed initial and boundary conditions. For an isothermal system, the energy equations reduce to the momentum equation. Darcy&apos;s law. However, since natural gas is the vapor state of the reservoir liquid, interphase mass transfer may take place with concomitant changes in both the intensive and extensive thermodynamic properties of each phase. It is this phenomenon that has often been omitted in previous mathematical models. An additional relation, then, which accounts for mass transfer between the phases, must be included with the other equations to specify a complete model. Completely formulating the equations to be solved is not a difficult task1 but obtaining their solution has been intractable up to now. Availability of large-memory, high-speed digital computers now makes an attack on this formidable problem possible. This paper presents a preliminary study of the problem. Since this investigation is intended to be exploratory, it is restricted to the linear, horizontal, isothermal, two-phase viscous flow of oil and gas in an oil reservoir. In the early development of predictive models of this process, the reservoir system was considered as a unit and various forms of the material balance equation were proposed.2-4 Pressure and saturation gradients were then added in the Buckley-Leverett model.5 The Buckley-Leverett formulation considered the fluids to be incompressible; thus, the mathematical model reduces to a steady-state system. In the 1950&apos;s, studies incorporating numerical techniques were being published.6-9 These mathematical models differed in the efficiency of finite difference techniques, the inclusion or exclusion of capillarity or the number of space dimensions considered. To solve these nonlinear, partial differential equations, each phase was considered to be homogeneous with time; therefore, mass transfer between phases was neglected. The effect of mass transfer on the gas displacement process was first reported by Attra. 10 He simulated the one-dimension flow system by a series of cells in each of which the fluids were equilibrated during a time step. In addition, the pressure throughout the system during each time step was predetermined and constant-phase velocities were calculated according to the Buckley-Leverett incompressible fluid flow model. Welge et al.11 developed a model for the displacement of oil by an enriched gas where composition is considered to be a dependent variable. Fluid

By Mark J. Klein, A. H. Clauer

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

Jan 1, 1965

By A. E. Powers

YEARS of experience and research have shown that molybdenum, tungsten, and vanadium are among the most useful and effective elements in augmenting the high-temperature strength of heat-treatable, ferritic steels. Grün, for example, in conducting an early survey of the strengthening effects of various elements in annealed, low-carbon steel at 750° and 930°F (400° and 500°C), found molybdenum and vanadium to be the most effective.&apos; Holtmann comprehensively studied the strengthening effects of molybdenum and vanadium in quenched and tempered steels at 930°F as a function of microstructure and carbon content.&apos; He found creep strength to be related to the saturation of the austenite at the austenitizing temperature, such that strength is increased by alloying until the y loop is reached—the terminus of complete austenitization. Beyond this degree of alloying, the inability to heat treat will result in lower strength, at least for low and moderate creep temperatures. Tungsten, owing to higher cost than molybdenum, has been little used in ferritic high-temperature steels. As a result, very few investigations have been made of its high-temperature strengthening ability, and reliable data on its effectiveness are not available. Grün rated tungsten far less effective than molybdenum at 930°F (500°C) on a weight-percentage basis.&apos; However, Tammann has found that tungsten is as effective as molybdenum in raising the recrystallization temperature of iron. For this reason Smith, in a review of the subject, has advocated a re-examination of the influence of tungsten on the high-temperature strength of steel.&apos; In any study of the effects of alloying additions on high-temperature properties, recognition must be made of the many variables involved, some of which can be controlled but have been ignored in innumerable investigations of the past, and some of which are difficult or even impossible to control. One cannot, for example, assign a definite strengthening index to any one alloying element, for this will be dependent upon the microstructure (heat treatment and mechanical treatment), testing con- ditions (temperature, time, and stress), and the complete composition of the steel (accompanying alloying elements, impurities, etc.). One ever-appear ing variable, hardness or tensile strength at room temperature, may be eliminated by heat treating all of the alloys to a single hardness level. An attempt may be made to control the variable of primary quenched structure by oil or water quenching the test specimen. The control of such variables may not be possible in cases of widely varying compositions, but at least such deviations should be recognized. Material and Procedure The alloys, made as 30-lb ingots in an induction furnace, all had the same base composition of about 0.18 pct C, 0.85 pct Mn, and 0.48 pct Si. They were grouped as follows: Mo Group—Containing up to 5.2 pct Mo. W Group-Containing up to 6.0 pct W. V Group—Containing up to 3.3 pct V. Mo-V Group—Containing up to 2.7 pct Mo and up to 1.4 pct V. Mo-W Group—Containing up to 2.5 pct Mo and up to 2.5 pct W. The steels, listed in Table I, were forged to ¾ in. sq bars and the material for the high-temperature, rupture-test specimens was further swaged to %-in. round bars.

Jan 1, 1957