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By Ö. Valla, A. Muan, K. Grjotheim

The solid miscibility gap between CaO and "FeO" in the system CaO-"FeO" in contact with metallic iron gradually closes as MnO is added as a third component. On the liquidus surface of the system CaO-"FeO"-MnO, this situation gives rise to a boundary curve which originates in the bounding system CaO-''FeO", proceeds a short distance into the composition triangle of the ternary system, and then vanishes. A new approach has been used to follow the course of this boundary curve. It is shown that the curve vanishes at a much lower temperature and at much lower MnO contents than has been inferred previously. The oxides CaO, "FeO", and MnO together with SiO2 are the major components of the slag phase formed in the basic oxygen process. Hence, the system CaO-"Fe0"-MnO is of considerable technological interest. In addition, the system is theoretically interesting. The three oxides above all have sodium chloride structure. The size differences among the cations are such that in the bounding "binary"* systems CaO-MnO and "FeO"-MnO complete solid-solution series are formed,1,2 whereas in the system CaO-"FeO" there is only limited mutual solubility in the solid state,3 see Fig. 1. Hence, a solid miscibility gap exists in a part of the "ternary"* system Ca0-"Fe0"-Mn0, and a liquidus boundary curve originates at the peritectic point, point a in the insert in Fig. 1, in the system CaO-"FeO", extends a certain distance into the ternary system, and then vanishes. Very little quantitative information on the system Ca0-"Fe0"-Mn0 is available in previous literature. Fischer and Feischer4 sketched isothermal sections at 1600° and 1700°C, and estimated that the vanishing boundary curve extends to approximately 1650°C and a composition of 42 wt pct MnO. The work to be described in the following was aimed at locating experimentally the "vanishing" liquidus boundary curve, determining its end point, and deriving representative isothermal sections below and above the temperature of the end point of this curve. EXPERIMENTAL METHODS The phase relations were determined by the well-known quenching method. (For a description of this method as applied to oxide systems, see the original paper of Shepherd, Rankin, and wright.5) Prereacted oxide mixtures of desired compositions were contained in iron crucibles and heated in a vertical tube furnace in an atmosphere of purified nitrogen. After equilibrium was judged to have been attained, the samples were quenched to room temperature, the phases present were identified by microscopic and X-ray examination, and total iron oxide contents were determined by chemical analysis. "Fisher Certified" reagents (MnCO3, CaCO3, Fe2O3, and Fe) served as starting materials. The calcium carbonate and ferric oxide were heated in air at 400° and 800°C, respectively, for 24 hr prior to use. Manganese oxide was prepared by decomposing the carbonate in a graphite crucible in hydrogen atmosphere at 1000°C. The product was ground under acetone to -100 mesh. The iron powder was used directly without any preparation. Dicalcium ferrite was used as a starting material for some mixtures. It was made by mixing some of the above prepared CaCO3 and Fe2O3 in proportions corresponding to the formula 2CaO. Fe2O3. This mixture was then fired in air for

Jan 1, 1969

By E. H. Edwards, J. Washburn, E. R. Parker

The mechanism of strain hardening was discussed in connection with some recent observations on the stress-induced motion of dislocation boundaries and on the simple shear deformation of zinc, cadmium, and copper crystals. THE nature of work hardening accompanying plastic deformation of metals has long been the subject of speculation. As early as 1925, Becker' suggested a thermodynamic process to explain slip based on the statistical probability for the occurrence of local glide steps. The theory that plastic glide was due to the presence of lattice defects known as dislocations was first introduced in 1934 by Taylor: and Orowan, and Polanyi.4 Taylor presented a rather complete analysis of the behavior and interaction of dislocations. He assumed a certain type of stress field to be associated with a single dislocation, and postulated that work hardening would result from plastic flow in two ways. First, dislocations following one another across a slip plane might encounter a barrier which would impede their movement, thereby gradually diminishing and eventually stopping plastic flow on that plane. The second mechanism involved the concept of dislocations traveling in opposite directions on nearby planes attracting one another to form a meta-stable lattice. The interaction of dislocations in such an array was shown to be large enough to require an increase in stress before plastic flow would continue. Kochendorfer5,6 extended the concept of work hardening involving a back stress due to an accumulation of like dislocations on a slip plane. He suggested that the formation of a new dislocation is hindered by those dislocations already present. The hardening resulting from the interaction of newly forming dislocations with bound dislocations was designated as "formative hardening," and indicated to be the only type of hardening necessarily connected with the slip process. In another detailed picture of work hardening, Mott and Nabarro7 assumed that sufficient dislocations are primarily present in the crystal as a consequence of growth irregularities (mosaic structure) to initiate slip. Work hardening resulted from the interaction of migrating dislocations with localized internal stresses produced during deformation. More recently, however, Mott8 as modified his treatment to include the generation of dislocations at Frank-Read sources and he accounts for the hardening by the formation of sessile dislocations. It has been widely accepted that metals of hexagonal structure, usually possessing only one set of glide planes, exhibit mechanical behavior markedly different from those having cubic structures. However, the differences in behavior of the two classes may be less fundamental than suggested. Rohm and Kochendorfer9 have shown that a single crystal of aluminum subjected to approximately simple shear gives the type of stress-strain curve associated with the hexagonal metals. Similar behavior has been demonstrated for gold and silver crystals by Andrade and Henderson.'" Thus when glide in cubic metals is restricted to one plane, the strain hardening characteristics closely resemble those of the hexagonal metals. It would appear, therefore, that an investigation of some of the strain hardening properties of single crystals sheared in simple glide might provide a more complete understanding of the phenomenon of work hardening. This report relates a number of recent observations made on single crystals of copper, zinc, and cadmium tested in simple shear. In addition, pertinent observations of effects accompanying stress-induced movement of dislocation boundaries are reported. Experimental Procedure and Results Details of the production and advantages of the type of single-crystal shear specimen employed for the following tests have been presented in a previous publication.n The method of testing makes

Jan 1, 1954

By George D. Dub

Owens Lake is at present a source of important nonmetallic minerals, sodium carbonate (soda ash, Na2CO3); sodium sesquicarbonate (trona, Na2CO3.NaHCO3.-2H2O) and borax, (Na2B4O7.10H2O). Owens Lake is a closed basin in the southern part of Inyo County, California, at the southern end of Owens Valley, east of the Sierra Nevada Mountains and west of the Coso and Inyo Mountains. Broadly considered, it is in the Great Basin area, but at no time was it a part of Lake Lahontan. Closed Basins Closed basins are phenomena of arid or semiarid regions where outflow is considerably less than inflow and where accordingly soluble salts concentrate in residual liquid. When inflowing waters originate in areas where the rocks are predominantly marine sediments, the residual basin liquid is primarily a chloride, brine; if the rocks are largely igneous, the residual brine tends to be alkaline containing carbonates, and sometimes borates as well. Since normally, both types of rocks occur in regions contiguous to closed basins, residual liquids do not often fit into the two broad divisions mentioned. No two residual brines are exactly alike, just as no two ore deposits are exactly alike. Even out of the same closed basin, it is possible to get widely different analyses of liquids since underground and surface flows, as well as local evaporation rates and other conditions might have a marked influence on residual-liquid compositions. At Searles Lake, The American Potash and Chemical Corporation is building a plant to process a lower-level brine which is considerably higher in sodium carbonate and borax, and lower in potassium chloride, than that company has processed for many years. F. W. Clarke1 has classified waters of closed basins as follows: I. Chloride type; largely sodium chloride (NaC1) and of oceanic type, such as Great Salt Lake. Related is -the Dead Sea, a bittern residue of magnesium, potassium and sodium chlorides. a. Sulphate type; largely sodium sulphate (Na2SO4) with considerable sodium chloride such as Sevier Lake, Utah; Laramie Lakes, Wyoming; Dale Lake, California. 3. Carbonate type; high in carbonates and fairly high in sulphates, such as Moses Lake, Eastern Washington; and the Nebraska Potash Lakes. 4. Carbonate—chloride type; lower in carbonates than type 3 and about equally rich in sulphates, such as Pyramid Lake, Nevada. 5. Sulphate—Carbonate type; quite high in sulphates and carbonates, such as Pelican Lake, Oregon. 6. Triple type; considerable quantities of carbonates, sulphates and chlorides

Jan 1, 1948

By George D. Dub

Owens Lake is at present a source of important nonmetallic minerals, sodium carbonate (soda ash, Na2CO3); sodium sesquicarbonate (trona, Na2CO3.NaHCO3.-2H2O) and borax, (Na2B4O7.10H2O). Owens Lake is a closed basin in the southern part of Inyo County, California, at the southern end of Owens Valley, east of the Sierra Nevada Mountains and west of the Coso and Inyo Mountains. Broadly considered, it is in the Great Basin area, but at no time was it a part of Lake Lahontan. Closed Basins Closed basins are phenomena of arid or semiarid regions where outflow is considerably less than inflow and where accordingly soluble salts concentrate in residual liquid. When inflowing waters originate in areas where the rocks are predominantly marine sediments, the residual basin liquid is primarily a chloride, brine; if the rocks are largely igneous, the residual brine tends to be alkaline containing carbonates, and sometimes borates as well. Since normally, both types of rocks occur in regions contiguous to closed basins, residual liquids do not often fit into the two broad divisions mentioned. No two residual brines are exactly alike, just as no two ore deposits are exactly alike. Even out of the same closed basin, it is possible to get widely different analyses of liquids since underground and surface flows, as well as local evaporation rates and other conditions might have a marked influence on residual-liquid compositions. At Searles Lake, The American Potash and Chemical Corporation is building a plant to process a lower-level brine which is considerably higher in sodium carbonate and borax, and lower in potassium chloride, than that company has processed for many years. F. W. Clarke1 has classified waters of closed basins as follows: I. Chloride type; largely sodium chloride (NaC1) and of oceanic type, such as Great Salt Lake. Related is -the Dead Sea, a bittern residue of magnesium, potassium and sodium chlorides. a. Sulphate type; largely sodium sulphate (Na2SO4) with considerable sodium chloride such as Sevier Lake, Utah; Laramie Lakes, Wyoming; Dale Lake, California. 3. Carbonate type; high in carbonates and fairly high in sulphates, such as Moses Lake, Eastern Washington; and the Nebraska Potash Lakes. 4. Carbonate—chloride type; lower in carbonates than type 3 and about equally rich in sulphates, such as Pyramid Lake, Nevada. 5. Sulphate—Carbonate type; quite high in sulphates and carbonates, such as Pelican Lake, Oregon. 6. Triple type; considerable quantities of carbonates, sulphates and chlorides

Jan 1, 1948

R. B. BRINSMADE, Puebla, Puebla, Mexico (communication to the Secretary*) :-As I have previously considered the general equations for mine-valuation,1 I will here discuss only that part of Mr. Finlay's paper criticized by E. E. White, along with some of the general principles of mining taxation, under the following six topics: 1. Rate of interest earned by sinking-fund. 2. Rate of interest earned by investment. 3. Place of royalty item in valuation. 4. Place of taxation item in valuation. 5. Calculation of future ore-prices. 6. Valuation of undeveloped mining land. 1. Rate of Interest Earned by Sinking-Fund.-Mr. Denny's quoted opinion of a 3 per cent. rate is evidently based on English conditions before the Boer war. The great destruction of liquid capital by the great wars and catastrophes since 1897, along with the flood of gold and consequent rise in the average index of commodity prices, has undoubtedly increased the rate of interest at least one-half per cent. In the United States sound savings banks now allow 4 per cent. on deposits, while safe bonds are quoted at prices which yield even 5 per cent.. 2. Rate of interest Earned by Investment.-Gold-mines with uncertain and scanty ore reserves, the objects of Mr. Denny's quoted remarks, are quite different from iron,. copper, or coal mines with definite mineral reserves blocked out ahead by drilling or driving. While 10 per cent is often all inadequate yield from gold-mines, the latter class of mines are often, as safe as city real estate and resemble it as an investment in more ways than one.

Jan 10, 1913

By T. R. Graham, R. S. Dean. J. R. Long

Although considerable information has been published concerning manganese additions to the aluminum bronzes, these data refer principally to the two-phase alloys containing 8 to II pet aluminum, with some incidental information on the single-phase alpha solid solution alloys. Straussl gives an excellent review of the literature on aluminum bronzes up to 1927 and a bulletin2 of the Copper Development Association brings the bibliography up to 1938. The most extensive work on manganese addition is that of Rosenhain and Lantsberry,3 who investigated alloys ranging up to 10 pct aluminum and up to 10 pct manganese, and presented data on the physical properties of sand-cast, chill-cast, and several wrought alloys. The greater portion of the work, however, concerns alloys ranging from 7 to 10 pct aluminum in the cast condition and sheds little light on the properties of wrought single-phase alloys. Corson4 notes that alloys containing 10 pct manganese and 4 pct aluminum are not hardenable by heat-treatment, while one with 15 pct manganese and 5 pct aluminum can be hardened, but he gives no tensile data on it. In fact, no systematic study of the properties of single-phase wrought alloys containing substantial amounts of manganese has been reported. It should also be noted that most of the investigators worked with alloys containing significant amounts of iron and silicon, and that these impurities, whether intentionally added or accrued from the raw materials used in the preparation of the alloys, may be expected to have a direct bearing on the resultant properties. In the work on alloys made with electrolytic manganese, the Bureau of Mines has undertaken an investigation of the copper-manganese-aluminum system for the purpose of establishing the constitution and properties of high-purity alloys. In a recent paper6 the authors reported a metallographic investigation of the alpha solid solution area and the adjoining fields of the copper-manganese-aluminum system up to 50 pct manganese. This study established the alpha-phase boundaries at 50" temperature levels from 400' to 850°C. It was found that the addition of manganese to copper-aluminum alloys decreases the solub%ty of aluminum in the alpha phase (Fig I), curving the boundary toward the copper-manganese base line. The greatest change occurs in the first 20 to 25 pct manganese, which decreases the solubility of aluminum to approximately 4 pct. With larger amounts of manganese the alpha boundary tends to parallel the copper-manganese binary. At 13oo'F and higher,

Jan 1, 1947

By R. S. Dean. J. R. Long, T. R. Graham

Although considerable information has been published concerning manganese additions to the aluminum bronzes, these data refer principally to the two-phase alloys containing 8 to II pet aluminum, with some incidental information on the single-phase alpha solid solution alloys. Straussl gives an excellent review of the literature on aluminum bronzes up to 1927 and a bulletin2 of the Copper Development Association brings the bibliography up to 1938. The most extensive work on manganese addition is that of Rosenhain and Lantsberry,3 who investigated alloys ranging up to 10 pct aluminum and up to 10 pct manganese, and presented data on the physical properties of sand-cast, chill-cast, and several wrought alloys. The greater portion of the work, however, concerns alloys ranging from 7 to 10 pct aluminum in the cast condition and sheds little light on the properties of wrought single-phase alloys. Corson4 notes that alloys containing 10 pct manganese and 4 pct aluminum are not hardenable by heat-treatment, while one with 15 pct manganese and 5 pct aluminum can be hardened, but he gives no tensile data on it. In fact, no systematic study of the properties of single-phase wrought alloys containing substantial amounts of manganese has been reported. It should also be noted that most of the investigators worked with alloys containing significant amounts of iron and silicon, and that these impurities, whether intentionally added or accrued from the raw materials used in the preparation of the alloys, may be expected to have a direct bearing on the resultant properties. In the work on alloys made with electrolytic manganese, the Bureau of Mines has undertaken an investigation of the copper-manganese-aluminum system for the purpose of establishing the constitution and properties of high-purity alloys. In a recent paper6 the authors reported a metallographic investigation of the alpha solid solution area and the adjoining fields of the copper-manganese-aluminum system up to 50 pct manganese. This study established the alpha-phase boundaries at 50" temperature levels from 400' to 850°C. It was found that the addition of manganese to copper-aluminum alloys decreases the solub%ty of aluminum in the alpha phase (Fig I), curving the boundary toward the copper-manganese base line. The greatest change occurs in the first 20 to 25 pct manganese, which decreases the solubility of aluminum to approximately 4 pct. With larger amounts of manganese the alpha boundary tends to parallel the copper-manganese binary. At 13oo'F and higher,

Jan 1, 1947

By M. G. F. Söhnlein

In a Bolivian tin concentrator an appliance was needed to furnish a suitable product for fine jigging from a pulp of the following composition: Mesh Per Cent. + 20 8.0 40 36.5 + 60 9.0 + 80 10.5 + loo 5.5 4-150 1 4.5 -150 25.0 99.0 Loss 1.0 In the jigging process, particles of cassiterite as small as 0.1 mm. can be recovered, provided they are not present in excess, that is if the material treated on the jigs contains sufficient coarser sands to keep the interstices between the grains open. Removing the fine material from the pulp by screening, to prepare the jig feed, is impractical, because that means the use of an 80-mesh screen or finer, and is not as effective as hydraulic classification. A screen would eliminate mineral grains from the feed which can be recovered by jigging, and in addition would throw more gangue and low-grade middling on to the jig. Formerly a one-spigot Richards vortex classifier had been used to accomplish this separation, but the tangential openings through which the hydraulic water enters became clogged by fine vegetal matter with which the water was contaminated, and which it was impossible to remove from the water. Consequently the work of the classifier was imperfect and a good deal of slime was sent to the jig. When designing the plant referred to in this paper it was decided to use a hydraulic classifier as shown in Fig. 1. The essential features had been copied from the Anaconda classifier.' This remodeled classifier consists of a truncated pyramidal wooden hopper A with an attachment of cast iron

Jan 1, 1917

By L. F. Mondolfo

An aluminum-copper 4 pct Cu alloy aged at room temperature for times increasing up to 78,000 hr was annealed at 170°C and the hardness and electrolytic potential determined during retrogression and subsequent aging. It was found that there is a change of behavior with increasing time of aging at room temperature which indicates that, given sufficient time, retrogression disappears and the room-temperature aging tends to progress directly into the high-temperature aging. WHEN an age-hardenable alloy is aged at low temperature after proper quench, changes of properties and structure take place. If this material is then heated to a temperature higher than the previous one, but still well below the solution-treatment temperature, the properties and structure changes produced by the low-temperature aging disappear and the material apparently reverts to the as-quenched condition, ready to undergo aging from the start. This phenomenon is termed reversion or retrogression. First discovered by Cayler3l it has been extensively investigated, as shown by the number of publications listed in the bibliography. From all the information available the following facts emerge: 1) All aeg hardenable alloys can be retrogressed.9 13, 19) 24, 33, 41, 42, 44, 47, 53, 54, 61, 63, 64, 66, 73, 78 2) All the properties changes that take place during aging at low temperature can be retrogressed to a smaller or larger extent.4,5,7,12,19,32,49,50,57,58,60,67,75,76 67, 75, 76 3) Retrogression requires that the temperature be well above the preliminary aging temperature. If the difference between the two temperatures is small, Little or no retrogression takes place.7,20,32,33,41,49 57. 65 4) The larger the difference between the preliminary aging and the retrogression temperature, the faster and more complete is the retrogression.20, 33, 41, 48, 65, 67 5) Retrogression can be produced several times in the same alloy.23, 35,41,57,59 6) The amount of retrogression that can be obtained depends on the aging temperature. Alloys aged at the lower temperatures can be retrogressed the most. The higher the preliminary aging temperature the less retrogression can be produced.20,25,33,37,44,49,54,58 7) Appreciable retrogression after aging at elevated temperatures is possible only if the preliminary aging has lasted a short time. After a certain time of aging at high temerature retrogression does not take place any more,:5,6,9,10,11,21,40,58 8) During retrogression the Guinier-Preston zones redissolve and disappear. The other precipitate

Jan 1, 1960

The work of Americanism is getting under way slowly but surely. The purpose of the Auxiliary is to build for all time and it therefore wishes to lay its foundations securely, feeling that the true understanding and adoption of the ideals of Americanism are the safeguards of our country from the perils of anarchy, and the bond of union between labor, brains, and capital. We hope to have reports from several directors to print next month and here append the report of the Director for the New York Section, Mrs. W. Y. Westervelt.

Jan 5, 1919

By R. M. Brick

SEGREGATION can occur only in cast alloys that solidify over a range of temperatures with a difference in composition of liquid and solid phases within this range (ignoring monotectic systems and chemical reactions as in rimming. steels). Thus a pure binary eutectic alloy should show no segregation unless undercooling occurs for one of the solid constituents of the eutectic; e.g., chill-cast Al-12 per cent Si alloys. The phasial condition implicitly described above-solidification over a range of temperatures with a compositional difference between solid. and liquid-is found in nearly all solid solutions or combinations of solid solutions and eutectics. Metastable conditions of solidification of these types of alloys cause coring; i.e., a difference in composition between the center of a dendrite and its periphery. This composition gradient can be detected usually only by microscopic examination of etched structures; it is a "micro" effect. The gradient of more immediate industrial concern is that between the center of a cast section and the surface layers. This gross (macro) effect is usually determinable only by sampling of the casting at different points, with chemical analyses of each sample. If the compositional difference is of the same type as in coring, the first parts to freeze being enriched in higher-melting-point constituents, and vice versa, the effect is called normal segregation. When what is apparently the first part of the casting to solidify shows a higher concentration of low-melting-point constituents, the effect is called inverse segregation. These macro effects are quite separate from the micro effect, coring, yet the latter must be. described first because it is an essential part of the picture. CORING Solid Solution, Ni-Cu The nickel-copper phase diagram is reproduced in Fig. Ia, in which the solid lines represent the equilibrium liquidus 1, 3, 7 and solidus 2, 5, 8, and where the vertical 1, 4, 8 represents a 55 per cent Cu alloy. Upon moderately rapid solidification, the first dendritic axes of composition 2 (see Fig. 1b) start to form at a temperature somewhat under I. Solidification continues by growth of the dendritic nuclei, but at temperature 4 the increment solidifying is of composition 5. The rapidity of this crystal growth prevents diffusion of nickel and copper atoms from adjusting the composition of the entire dendrite at 5, which it would be under equilibrium conditions. The average composition of a specific dendrite, or the entire solid phase, at temperature 4 might be that indicated

Jan 1, 1944

By G. F. Comstock

A serious disagreement as to the effect of titanium on the hardenability of steel exists in published references to this subject. Kramer, Hafner and Toleman reported' that acid-soluble titanium decreased creased the hardenability, with multiplying factors less than unity according to [ ] Grossmann's system for calculating hardenability from the composition. Crafts and Lamont, on the other hand, showed2 how titanium increased the hardenability, with multiplying factors up to nearly 1.7 for about o.' per cent titanium, though reducing the hardenability with higher titanium. Their data, however, were obtained almost exclusively from steels containing more than 2.4 per cent manganese, and thus should be considered as applying mainly to manganese steels. Likewise, the results of Kramer, Hafner and, Toleman regarding titanium are subject to the limitation that although they eliminated the effect of the insoluble titanium compounds from their calculations, they did not make a similar correction for the carbon that was combined with titanium and therefore occurred in a similarly ineffective form. It is suggested by the literature referred to that titanium may have different effects on hardenability, depending on the mode of its occurrence in the steel. The different effects on strength, ductility, etc., have already been discussed by the author.3 In the present discussion some comparisons, on the basis of hardenability as revealed by the Jominy end-quench test, are given between three titanium and two non titanium steels heat-treated in the same ways, and between other titanium steels heat-treated in several different ways. COMPARISON OF PLAIN-CARBON, VANADIUM AND TITANIUM STEELS The five steels used for these experiments were commercially melted and rolled products received in the form of 1.5-in. rounds from three different mills. Their analyses are given in Table I. Steels 2, 3, and 4 are from the same basic electric heat, containing 0.009 per cent of phosphorus and the same amount of sulphur. They were made by, adding aluminum and low-carbon ferrotitanium to different ingots. No. I is a plain-carbon coarse-grained S.A.E. 1040 steel made by the basic open-hearth -process. No. 5

Jan 1, 1945

By W. C. Hagel, H. J. Beattle

DURING an earlier examination of high-temperature alloy, A-286, the presence of an unknown intermetallic compound was verified by X-ray diffraction. Owing to its prominent appearance at grain boundaries, the arbitrary name of G phase was assigned.' Low-silicon, vacuum-melted heats were found to contain no G phase, and this investigation was initiated to determine the influence of silicon on G-phase formation; additional work was undertaken to extend existing data on other intermetallic compounds and the aging reactions in this group of alloys. Taylor and Floyd2-4 observed that the Ni3Al phase (y') and the Ni8Ti phase (7)) are the only equilibrium intermetailic compounds in nickel ternary alloys containing up to 25 pct Cr and 10 pct Ti or Al. They found that y' can dissolve considerable nickel, chromium, and titanium, although the stoichio-metric composition of 7 remains fixed. Nordheim and Grant5 confirmed that age hardening of 80 pct Ni-20 pct Cr alloys of the Nimonic-80 type is due to precipitation of y', where aluminum is partially replaced by titanium. They suggest that optimum hardness and ductility result from increasing the titanium to aluminum ratio toward the end of the y' solubility range and from increasing total titanium and aluminum content to the maximum limit permitted by processing. Such information is necessary for those seeking to improve the service-temperature range and long-time structural stability of high-temperature components. Experimental Methods Compositions of materiais studied. are listed in Table I. Alloys D, E, and F possess representative A-286 analyses; silicon increases from less than 0.02 to 1.75 pct in alloys A to H; in alloys I to K, titanium decreases from 2.24 to 0.39 pct, and aluminum increases from 0.1 6 to 240 pct. Alloys A and D were solution heat treated for 4 hr at 1700°F, oil quenched and aged for 0.5, 2, 6, 12, 30, 70, 200, and 1000 hr at 1200°, 1300°, 1400°, and 1500°F. Therefore, a total of 33 specimens was required for each

Jan 1, 1958

By E. Luttgen

The substance referred to in the title is the artificially prepared basic carbonate of magnesia, a compound of the carbonate with the hydroxide. It is the "block-magnesia " of commerce, the magnesia alba of the pharmacist. Its composition, though somewhat indefinite, may usually be expressed by the formula : 4MgCO3,Mg(OH)2, 5H2O; it is found to vary slightly with the details of the method of preparation. This substance has been employed recently with much reported success, as a non-conductor of heat. It is moulded to form coverings suitable for steam-pipes and their fittings and sectional jackets for boilers and cylinders; it is furnished also in forms suitable for lining refrigerators, walls and roofs of buildings, fire-proof safes, etc. It was with a view of ascertaining the efficiency of this material as a non-conductor, or, rather, its relative efficiency among the various non-conducting materials in use, that the writer made the experiments about to be described. A block of magnesium carbonate or magnesia alba exhibits some interesting physical qualities. It is a smooth, white, close-grained solid, in outward appearance resembling a block of Paris plaster. It possesses the lightness of cork, the porosity of sponge, and withal a degree of firmness and strength that, in view of its levity, is quite remarkable. To examine more closely the properties of this substance, a number of 1-inch cubes were sawed from the commercial block-carbonate, also some bricks, that is, blocks measuring accurately 2 x 4 x 8 inches, the dimensions of an ordinary brick. The cubes weighed from 2.2 to 2.7 grammes, the bricks weighed 140 to 175 grammes (5 to 6 ounces). The bricks were carefully measured and weighed, and placed in vessels containing distilled water, in which they became gradually submerged, owing to the displacement by water of the air inclosed in the structure of the magnesium carbonate. After twenty-four hours the blocks were removed from the water, dried superficially by contact with filter-paper and weighed. From the increase in weight, the volume of the water absorbed, and consequently that of the air displaced by it, were obtained. For example: a block which when dry weighed 155

Jan 1, 1887

By R. M. Waterstrat

The phase VPt has been reported to have a crystal structure of the B19 (AuCd) type at low temperatures.''' A recent investigation3 has indicated that this phase forms by an ordering reaction from the disordered fcc Pt(V) solid solution which is stable at high temperatures. olander4 has shown that there exists a close structural relationship between the orthorhombic B19 structure and a fcc structure possessing the L1 , (CuAu) type of atomic ordering. One may derive the B19 structure by a shear-like displacement of the atoms in the (002) plane of the ordered fcc lattice along the (010) direction through a distance equal to one eighth of the unit cell dimension, see Fig. 1, followed by an orthorhombic distor; tion of the cube axes such that a = 2.693A. h = 4.413A. and c =4.767A (estimated uncertainty 0.005A). Filings of this phase were produced by grinding a VPt alloy having the B19 structure with a diamond dental grinding disc. When these filings were subjected to X-ray diffraction examination, a pattern of a severely cold-worked fcc structure was obtained. higher solution temperature used (1200 C) or to a much lower oxygen pressure than he reported, due to transfer through many bends and tubes in his complicated apparatus. An interesting observation was made when an oxygen addition was attempted in a cylindrical sample of columbium in the cold worked condition rather than in the annealed condition. The interior of the sample re-crystallized during the oxygen addition at 900°C. The periphery, however. (1 mm) did not recrystallize. A similar sample without oxygen recrystallized throughout the entire cross section. It appears that the re-crystallization is hindered by interstitial oxygen. Presumably, the interior recrystallized before appreciable oxygen could diffuse into the core. This work was supported by the United States Atomic Energy Commission. When the filings were annealed in a high vacuum furnace for 15 min at 110O0C, however. they returned to the B19 structure perhaps as a result of reordering. Subsequent less-severe cold-working of these filings by grinding in a hardened steel mortar and pestle at room temperature produced a transformation into a cold-worked L1, structure which was slightly tetragonal. Thus it appears that the degree of deformation may influence the final structure. Considerable line-broadening was observed in X-ray diffraction patterns obtained from the cold-worked filings and therefore no attempt was made to obtain accurate lattice parameters. Vacuum annealing treatments at 1100'C intended

Jan 1, 1970

By Walter R. Hibbard, George H. Eichelman, William P. Saunders

INTEREST in the various products of the austenite eutectoid transformation in iron-carbon alloys, particularly as produced by the isothermal sub-critical techniques introduced by Davenport and Bain,1 has resulted in the application of similar heat treatments to other eutectoid transformations. Studies2.3.4 of the beta copper-aluminum alloy eutectoid transformation revealed microstructures very similar to those found in iron-carbon alloys. Investigations5,6 of the kappa phase in the copper-silicon alloy system have established the eutectoid type transformation of this phase into terminal alpha and intermediate gamma phases. This transformation has been reported' as forming a eutectoid type structure in a very sluggish manner by a long time isothermal treatment. The purpose of the present investigation is to study the products of the kappa eutectoid transformation as affected by transformation temperatures by means of microscopic and hardness methods. PREPARATION OF ALLOYS The copper-silicon alloy equilibrium diagram is shown in Fig I. The kappa field converges to a minimum silicon solubility at the eutectoid composition and all other completely kappa alloys are hypereutectoid. Thus, kappa decomposition involves first the proeutectoid precipitation of gamma, followed by the eutectoid transformation of a lower silicon content kappa (5.2 pct). Smith5 has indicated that alloys containing between 5.4 and 5.7 pct silicon could be hot rolled, but that the latter composition was desirable in order to permit a wider kappa temperature range. The alloy was prepared by adding Baker's commercially pure silicon to an ingot of copper-silicon alloy containing 2.34 pct silicon remaining from the investigation of Brick, Martin and Angier.7 Following the casting techniques of Krynitsky, Saunders and Stern,8 the ingot was melted under a heavy layer of dried charcoal in a clay-graphite crucible in a gas-fired pot furnace. Small lumps of silicon were held under the melt surface with a graphite rod. After the silicon had dissolved, the alloy was chill-cast and then remelted without a cover, stirred thoroughly, skimmed and poured into a preheated cast iron mold. The cast microstructure is shown in Fig 2. The casting, 10 X 4 X ¾ in., was annealed at 750°C for 21 brand furnace cooled. About 1/16 in. was milled from all surfaces. The plate was cut in half longitudinally and the bottom half analyzed by A.S.T.M. methods.9 The composition by weight was as follows: copper 94.24 pct, silicon 5.56 pct, iron 0.11 pct.

Jan 1, 1948

By J. H. Frye, C. P. Sun

A Frimary substitutional solid solution is a solution that has the same crystalline structure as the solvent metal, and in which solute atoms have replaced, solvent atoms at random on the host lattice. This replacement usually results in an increase in hardness as measured by conventional penetration tests. The factors controlling this hardness increase have been investigated in some detail for copper and silver alloys.1-5 It has been shown that at least four factors are involved: (1) concentration of solute, (2) row of the periodic table from which the solute is taken, (3) increase in lattice parameter of the solvent produced by the solute, and (4) work-hardening produced by penetration of the hardness indenter.2,4 The last factor is of the utmost importance. An investigation of the available data2,4 leads to the conclusion that the solutes that produce large hardness increases in copper or silver do so principally because they increase the rate of work-hardening. Furthermore, it is doubtful whether any solute produces appreciable hardening in the absence of plastic deformation and hence work-hardening. Thus it is clear that an understanding of the effect of solutes on rate of work-hardening is essential to an understanding of the hardness of solid solutions. This effect of solutes is also important in forming operations such as rolling or deep drawing. The presence of elements that confer rapid rates of work-hardening will increase power consumption and may well decrease the amount of deformation possible between anneals. This paper is the report of an investigation of the factors that control rates of work-hardening in copper and silver solid solutions. The silver and copper alloys considered are those in which all solutes are from the B subgroups or second short period of the periodic table. Experience with hardness, with structure of alloys, and with electrical conductivity has proved that results of a fundamental nature are most likely to be obtained if alloys are made up on the basis of the positions of the constituent elements in the periodic table. In previous work, Meyer analysis data on high-purity copper and silver solid solutions were obtained through careful measurements with fully calibrated equipment and these data were corrected to correspond to a standard grain size. From these corrected Meyer analysis data rates of work-hardening have now been calculated. Calculation of Rates as Work- HARDENING The Meyer analyses, on which this study of rates of work-hardening is based, consisted in determining impression

Jan 1, 1944

By J. H. Frye, C. P. Sun

A Frimary substitutional solid solution is a solution that has the same crystalline structure as the solvent metal, and in which solute atoms have replaced, solvent atoms at random on the host lattice. This replacement usually results in an increase in hardness as measured by conventional penetration tests. The factors controlling this hardness increase have been investigated in some detail for copper and silver alloys.1-5 It has been shown that at least four factors are involved: (1) concentration of solute, (2) row of the periodic table from which the solute is taken, (3) increase in lattice parameter of the solvent produced by the solute, and (4) work-hardening produced by penetration of the hardness indenter.2,4 The last factor is of the utmost importance. An investigation of the available data2,4 leads to the conclusion that the solutes that produce large hardness increases in copper or silver do so principally because they increase the rate of work-hardening. Furthermore, it is doubtful whether any solute produces appreciable hardening in the absence of plastic deformation and hence work-hardening. Thus it is clear that an understanding of the effect of solutes on rate of work-hardening is essential to an understanding of the hardness of solid solutions. This effect of solutes is also important in forming operations such as rolling or deep drawing. The presence of elements that confer rapid rates of work-hardening will increase power consumption and may well decrease the amount of deformation possible between anneals. This paper is the report of an investigation of the factors that control rates of work-hardening in copper and silver solid solutions. The silver and copper alloys considered are those in which all solutes are from the B subgroups or second short period of the periodic table. Experience with hardness, with structure of alloys, and with electrical conductivity has proved that results of a fundamental nature are most likely to be obtained if alloys are made up on the basis of the positions of the constituent elements in the periodic table. In previous work, Meyer analysis data on high-purity copper and silver solid solutions were obtained through careful measurements with fully calibrated equipment and these data were corrected to correspond to a standard grain size. From these corrected Meyer analysis data rates of work-hardening have now been calculated. Calculation of Rates as Work- HARDENING The Meyer analyses, on which this study of rates of work-hardening is based, consisted in determining impression

Jan 1, 1944

By J. H. Frye, C. P. Sun

A PRIMARY substitutional solid solution is a solution that has the same crystalline structure as the solvent metal, and in which solute atoms have replaced solvent atoms at random on the host lattice. This replacement usually results in an increase in hardness as measured by conventional penetration tests. The factors controlling this hardness increase have been investigated in some detail for copper and silver aneys.1-5 It has been shown that at least, four factors are involved: (I) concentration of solute, (2) row of the periodic table from which the solute is taken, (3) increase in lattice parameter of the solvent produced by the solute, and (4) work-hardening produced by penetration of the hardness indenter.2,4 The last factor is of the utmost importance. An investigation of the available data2.4 leads to the conclusion that the solutes that produce large hardness increases in copper or silver do so principally because they increase the rate of work-hardening. Furthermore, it is doubtful whether any solute produces appreciable hardening in the absence of plastic deformation and hence work-hardening. Thus it is clear that an understanding of the effect of solutes on rate of work hardening is essential to an understanding of the hardness of solid solutions. This effect of solutes is also important in forming operations such as rolling or deep drawing. The presence of elements that confer rapid rates of work-hardening will increase power consumption and may well decrease the amount of deformation possible between anneals. This paper is the report of an investigation of the factors that control rates of work-hardening in copper and silver solid solutions. The silver and copper alloys considered are those in which all solutes are from the B subgroups or second short period of the periodic table. Experience with hardness, with structure of alloys, and with electrical conductivity has proved that results of a fundamental nature are most likely to be obtained if alloys are made up on the basis of the positions of the constituent elements in the periodic table. In previous work, Meyer analysis data on high-purity copper and silver solid solutions were obtained through careful measurements with fully calibrated equipment and these data were corrected to correspond to a standard grain size. From these corrected Meyer analysis data rates of work-hardening have now been calculated. CALCULATION OF RATES OF WORK- HARDENING The Meyer analyses, on which this study of rates of work-hardening is based, consisted in determining impression

Jan 1, 1944

By J. R. Long, T. R. Graham, R. S. Dean

ALTHOUGH considerable information has been published concerning manganese additions to the aluminum bronzes, these data refer principally to the two-phase alloys containing 8 to 11 pct aluminum, with some incidental information on the single-phase alpha solid solution alloys. Strauss1 gives an excellent review of the literature on aluminum bronzes up to 1927 and a bulletin2 of the Copper Development Association brings the bibliography up to 1938. The most extensive work on manganese addition is that of Rosenhain and Lantsberry,3 who investigated alloys ranging up to 10 pct aluminum and up to 10 pct manganese, and presented data on the physical properties of sand-cast, chill-cast, and several wrought alloys. The greater portion of the work, however, concerns alloys ranging from 7 to 10 pct aluminum in the cast condition and sheds little light on the properties of wrought single-phase alloys. Corson4 notes that alloys containing 10 pct manganese and 4 pct aluminum are not hardenable by heat-treatment, while one with 15 pct manganese and 5 pct aluminum can be hardened, but he gives no tensile data on it. In fact, no systematic study of the properties of single-phase wrought alloys containing substantial amounts of manganese has been reported. It should also be noted that most of the investigators worked with alloys containing significant amounts of iron and silicon, and that these impurities, whether intentionally added or accrued from the raw materials used in the preparation of the alloys, may be expected to have a direct bearing on the resultant properties. In the work on alloys made with electrolytic manganese, the Bureau of Mines has undertaken an investigation of the copper-manganese-aluminum system for the purpose of establishing the constitution and properties of high-purity alloys. In a recent paper5 the authors reported a metallographic investigation of the alpha solid solution area and the adjoining fields of the copper-manganese-aluminum system up to 50 pct manganese. This study established the alpha-phase boundaries at 50° temperature levels from 400° to 850°C. It was found that the addition of manganese to copper-aluminum alloys decreases the solubility of aluminum in the alpha phase (Fig I), curving the boundary toward the copper-manganese base line. The greatest change occurs in the first 20 to 25 pct manganese, which decreases the solubility of aluminum to approximately 4 pct. With larger amounts of manganese the alpha boundary tends to parallel the copper-manganese binary. At 1300°F and higher,

Jan 1, 1947