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By John T. Eash

During the course of the war the supply of tin in this country has steadily decreased and a continued effort has been made since the beginning of the emergency to use alloys that are either tin free or contain appreciably lower quantities of it. Even with the use of emergency materials, the tin supply has become particularly acute, and it is estimated that after recovery of the far eastern mines probably one to two years will be required before our normal annual supply is attained. One of the major tin-consuming fields is the bronze industry. Recommendations were made several years ago by the War Production Board toward the use of bronzes of lower tin content and manganese bronze and yellow brass for the customarily used tin-bronze alloys. The substitution of 88-5-5-2 copper-nickel-tin-zinc for composition G received considerable attention before the war, and consequently this alloy and modifications of it have been adopted for some current structural and pressure castings requiring high mechanical properties. Another of the commonly used bronzes is the 10 per cent tin 10 per cent lead copper alloy for heavy-duty bearings. It has been known for some time that about 2.5 per cent nickel could be substituted for an equal amount of tin with retention of equivalent casting and tensile properties; however, it was desirable that a still greater conservation of tin be made. Consequently the present paper deals with the development of an alloy of low tin content that could be used for bearings. Selection or Composition While the final criterion of the suitability of an alloy for a bearing must of course be determined by practical applications where all the ramifications of service conditions are encountered, certain short cuts in the initial selection of compositions are necessary to avoid an extensive and time-consuming investigation. The present program has been to match certain metallurgical properties of the 80-10-10 bronze that have been shown by previous investigations to be related to good bearing qualities. Such properties as microstruc-ture, hardness, compressive yield strength, ductility, and castability have been considered, and certain simple wear tests have been made. Of the readily assessed characteristics desired in a good bearing alloy, one can be certain that a uniform distribution of lead is required in a leaded bronze. It has generally been found that segregation of lead is one of the chief difficulties with this type of material. Tin has a powerful effect on lead distribution, and with 10 per cent present produces a random distribution of chunky particles, such as are shown in Fig' 1. I' In the tin bronze contains the hard alpha + delta

Jan 1, 1946

Duval Corporation, the copper -producing subsidiary of Pennzoil United, Incorporated, operates two plants south of Tucson, Arizona: Esperanza and Sierrita. The newest, the Sierrita mine and mill, came into full production in 1971 at a rate of 72,000 tpd (now 87,000 tpd) which at that time made it the largest designed copper plant in the world. In 1968, the largest milling units then available were selected; these included fourteen 16$ ft by 19 ft long ball mills (5 m by 5. 9 m), each equipped with a 3,000 hp motor. The primary crushing plant contains two 60 in. by 89 in. Allis- Chalmers gyratories; each is constructed for two-side dumping by 120- and 150-ton trucks. Access to the maintenance areas to service the drive motors, eccentrics, apron feeders, and speedup conveyor is through hatches between the crushers. This area also contains a good- sized elevator and a stairway. The rather shallow dump pocket, the 300 -ton crushed ore pocket, and the overall design make for a fine, compact arrangement for two 60 in. crushers installed in a total height of 89 ft (27. 13 m) from dump level to top-of-concrete at the lowest level beneath the collecting or speedup belt. This efficient layout is reflected in the short changeout time for mantle and concaves (8 and 30 hours respectively) and the low crushing cost of $0.0416 per ton (first half of 1975). The crushers are installed on 64 ft (19.5 m) centers (Bougainville and Palabora, both with 54 in. crushers, are 55 and 43 ft or 16.76 and 13. 1 m respectively). The switchgear room is located between the truck dumping positions on one side of the installation and the hydroset units on the other side, with dust collecting equipment mounted over the latter. Protection of these units from rocks falling from the trucks is provided by concrete walls. Truck dumping is controlled from a central 25 ft high (7.6 m) control room, which permits fine observation of operations. Removal of oversize boulders from the dump pockets is done with a hydraulic grapple or a rock hook or sling attached to the auxiliary crane hook. Spare mantle stands and a maintenance room with a zinc melting furnace are located at one end of the surface facilities. The maintenance room is equipped with an access door for the crane, which is equipped with 80- and 20-ton hoists. The two 84 in. (2. 135 m) apron feeders receiving the 6 in. (152 mm) crushed rock discharge onto a 72 in. (1. 83 m) collecting and wear belt operated at a speed compatible with the 60 in. (1.524 m) takeaway belt running at a speed of 630 fpm (192 mpm). This belt feeds a 54 in.

Jan 1, 1978

By P. Nyul, S. Caplan, M. F. Lamorte, T. Gonda

Therrnal resistance is measured on GaAs laser diodes in the temperature range 77" to 300°K. These data show that typically the thermal resistance increases fifteen times from 77 to 300°K. The increase is attributed to the decrease in thermal conductivity of Gds. Nornzalized thermal-resistance experimental and calculated from published thermal-conductivity data) values are compared and good agreement shown. In many practical applications, it is desirable to obtain high average power output from laser diodes. Average power output is determined in large measure by the external efficiency and thermal resistance of the diode.'-3 The average input power with which a laser diode may be pumped is dependent on the thermal resistance. Thermal resistance has been studied over the temperature range from 77" to 300°K which is usually of interest in practical applications. This correspondence presents data on thermal resistance for GaAs laser diodes and compares them to values calculated from thermal-conductivity data. Junction heating arises from three distinct components: 1) joule heating, 2) nonradiative recombinations, and 3) reabsorption of radiation. In high-threshold lasers, joule heating predominates.4 Typically, the nonradiative recombination is not negligible and contributes significantly to junction heating. The loss caused by reabsorption may be of the order of the joule-heating loss.' The fabrication and the geometry of laser diodes have been described in other publications.5 The data presented were obtained from diodes which were fabricated from the same region of a wafer to minimize differences. Fig. 1 shows four types of heat sinks employed in this study. In types 1 and 2, the laser cavities are alloyed; pressure contacts are employed in types 3 and 4, and indium is plated on the inside faces to improve the intimacy of contact between the diode and the heat sink. The thermal resistance R is defined by the following relationship: where AT is the junction-temperature rise above the heat-sink temperature for the diode power dissipation Pd. For measurement of thermal resistance, a direct current is passed through the junction, in the range from 1 to 50 pct of threshold. The diode is placed in the laser mode by superimposing a current pulse of low-duty cycle on the direct current.3 Junction-temperature rise is obtained from the time-resolved laser spectrum at the leading edge of the current pulse.' Typically, the laser line shift is 2.5A per OK. The resolution employed in these experiments is 0.5A; therefore, the accuracy in the temperature measurement is 0.2K. Typically, the temperature rise is 2" to 20°K. The power dissipation is the difference between the dc power input, and the radiation power exiting the diode. The measurement accuracy is less than 1

Jan 1, 1968

By Shiou-Chuan Sun

THERE are several postulations for the mechanism of slime-coating. Ince1 proposed the electrostatic hypothesis, del Giudice2 suggested the chemical theory; Bankoff3 reported that slime-coating is inhibited whenever the surfaces of the slime particles are ionized sufficiently for Brownian movement to occur, and is facilitated when they are not. These postulations were examined by experiments involving the Burton tube, the cataphoretic cell, and slime-coating. The techniques involved will be described elsewhere in this paper. The electrostatic hypothesis, supported by migration experiments in a Burton tube, was abandoned by del Giudice.2 Experiments of mineral migration in a Burton tube, after the idea of Ince with modification, indicated that this apparatus was inefficient for accurate observation. A cataphoretic cell was constructed. The experimental data, as shown in this paper, enable the correlation and advancing of this hypothesis. Del Giudice's chemical hypothesis considered that slime particles are attached to the galena surfaces by a cement formed as the result of a metathetical reaction between the sulphide and the slime particles. Applying this hypothesis, the calcite slime particles were cemented to the galena surface by the formation of a cementing lead carbonate. In a similar way, lead silicate was considered as the cement of quartz slime coating on galena. Experiments verifying this hypothesis were carried out according to the following procedure: Fresh mineral particles selected from the central part of the broken lumps and examined first under a microscope for ascertaining a clean surface were placed in a small beaker (1.5-in. dia. by 2.5 in. deep) half filled with one kind of slime pulp. The dilution of all slime pulp was 4: 1 and the particle size of slime was minus 400-mesh. After agitation for 5 min. the specimen was removed by a sharp-end tong and washed gently three times in three separate distilled water baths. It was dried before a small electrical fan and examined immediately under a microscope. A photograph was made of the observation. With reference to the published solubilities,4 several of these experiments, which contradict the chemical hypothesis, are listed as following: I. The coating of fluorite slime on the galena surface was light but the coating of galena slime on the fluorite surface was heavy. In a similar way, the coating of quartz slime on the galena was fairly light, but the coating of galena slime on the quartz particle was heavy. According to the chemical hypothesis, using the same constituents and regardless of which is the slime, the density of the coating should be expected to be the same. This, however, as shown above, is not true. Moreover, fluorite (CaF2), having a solubility of 16 mg.[4a] per liter of water, should not be

Jan 1, 1943

The Carol concentrator of the Iron Ore Company of Canada was completed in 1962 and through expansions .has become one of the largest iron ore beneficiation plants in the world with a new feed capacity of 147,000 It (149,350 mt) per day. The Carol ore contains specular hematite and magnetite having a bulk density of 166 lb/cu ft (2660 kg/m3) It is crushed to 7 in. (178 mm) for feeding to autogenous mills for subsequent concentration by Humphrey’s s spirals. Mine ore is loaded into 150 It (152 mt) trucks by 10 cu yd (7.65 m3) shovels and then, at a transfer station, is loaded into side-dump rail cars of 100 It (101.6 mt) capacity. Each train is composed of twenty cars and operates completely automatically as it hauls the ore to the primary crushing station, dumps its loads, and returns to the loading station, a 11.5 mile (18.5 km) round trip. The primary crushing station is composed of two 60 in. by 89 in. gyratories installed in 1962 and 1972. Neither crusher has a grizzly ahead of it. Rail cars are dumped into 300 It (305 mt) pockets which contain splitters for spider protection. What few oversize boulders which do not go through the crushers are drilled and blasted. Liners in the dump pocket and the crushed ore pokcet under the crusher are manganese steel. Both crushers are direct driven by 700 hp (5 15 kw) motors, and the hardness of the ore causes them to pull almost their full power. Crushed ore drops into a 160 It (163 mt) pocket equipped with radioactive level control units for low-intermediate, and high level indications. The low level unit stops and starts the apron feeder, the intermediate unit controls the feeder speed, and a high-level reading stops the car dumping. Apron feeders are 72 in. (1. 82 m) wide and are variable speed from 12 to 52 fpm (3. 7 to 15. 8 mpm). They discharge to 60 in. (1.52 m) wide takeaway belts. Crushing capacity, nominal and maximum for each crusher, is 4, 000 ltph (4. 064 mtph) and 4,600 ltph (4,674 mtph). The crusher eccentric is removed by a conventional cart which travels on removable rails. Mainshaft removal and maintenance is done by 65 and 75/10 ton cranes. Mainshaft replacement requires eight hours, a usual time, but concave replacement requires only sixteen hours, a result of a fine maintenance organization. Mantle relining is done at grade next to a repair track. Zinc is used for backing both mantles and concaves. Both the electrical rooms and the dust collectors are located below grade,

Jan 1, 1978

By Shiou-Chuan Sun

There are several postulations for the mechanism of slime-coating. Incel proposed the electrostatic hypothesis, del Giudice2 suggested the chemical theory; Bankoff3 reported that slime-coating is inhibited whenever the surfaces of the slime particles are ionized sufficiently for Brownian movement to occur, and is facilitated when they are not. These postulations were examined by experiments involving the Burton tube, the cataphoretic cell, and slime-coating. The techniques involved will be described elsewhere in this paper. The electrostatic hypothesis, supported by migration experiments in a Burton tube, was abandoned by del Giudice.2 Experiments of mineral migration in a Burton tube. after the idea of Ince with modification, indicated that this apparatus was inefficient for accurate observation. A cataphoretic cell was constructed. The experimental data, as shown in this paper, enable the correlation and advancing of this hypothesis. Del Giudice's chemical hypothesis considered that slime particles are attached to the galena surfaces by a cement formed as the result of a metathetical reaction between the sulphide and the slime particles. Applying this hypothesis, the calcite slime particles were cemented to the galena surface by the formation of a cementing lead carbonate. In a similar way, lead silicate was considered as the cement of quartz slime coating on galena. Experiments verifying this hypothesis were carried out according to the following procedure: Fresh mineral particles selected from the central part of the broken lumps and examined first under a microscope for ascertaining a clean surface were placed in a small beaker (1.5-in. dia. by 2.5 in. deep) half filled with one kind of slime pulp. The dilution of all slime pulp was 4:I and the particle size of slime was minus 400-mesh. After agitation for 5 min. the specimen was removed by a sharp-end tong and washed gently three times in three separate distilled water baths. It was dried before a small electrical fan and examined immediately under a microscope. A photograph was made of the observation. With reference to the published solubilities,' several of these experiments, which contradict the chemical hypothesis, are listed as following: I. The coating of fluorite slime on the galena surface was light but the coating of galena slime on the fluorite surface was heavy. In a similar way, the coating of quartz slime on the galena was fairly light, but the coating of galena slime on the quartz particle was heavy. According to the chemical hypothesis, using the same constituents and regardless of which is the slime, the density of the coating should be expected to be the same. This, however, as shown above, is not true. Moreover, fluorite (CaF2), having a solubility of 16 mg.4a per liter of water, should not be

Jan 1, 1943

By Shiou-Chuan Sun

There are several postulations for the mechanism of slime-coating. Incel proposed the electrostatic hypothesis, del Giudice2 suggested the chemical theory; Bankoff3 reported that slime-coating is inhibited whenever the surfaces of the slime particles are ionized sufficiently for Brownian movement to occur, and is facilitated when they are not. These postulations were examined by experiments involving the Burton tube, the cataphoretic cell, and slime-coating. The techniques involved will be described elsewhere in this paper. The electrostatic hypothesis, supported by migration experiments in a Burton tube, was abandoned by del Giudice.2 Experiments of mineral migration in a Burton tube. after the idea of Ince with modification, indicated that this apparatus was inefficient for accurate observation. A cataphoretic cell was constructed. The experimental data, as shown in this paper, enable the correlation and advancing of this hypothesis. Del Giudice's chemical hypothesis considered that slime particles are attached to the galena surfaces by a cement formed as the result of a metathetical reaction between the sulphide and the slime particles. Applying this hypothesis, the calcite slime particles were cemented to the galena surface by the formation of a cementing lead carbonate. In a similar way, lead silicate was considered as the cement of quartz slime coating on galena. Experiments verifying this hypothesis were carried out according to the following procedure: Fresh mineral particles selected from the central part of the broken lumps and examined first under a microscope for ascertaining a clean surface were placed in a small beaker (1.5-in. dia. by 2.5 in. deep) half filled with one kind of slime pulp. The dilution of all slime pulp was 4:I and the particle size of slime was minus 400-mesh. After agitation for 5 min. the specimen was removed by a sharp-end tong and washed gently three times in three separate distilled water baths. It was dried before a small electrical fan and examined immediately under a microscope. A photograph was made of the observation. With reference to the published solubilities,' several of these experiments, which contradict the chemical hypothesis, are listed as following: I. The coating of fluorite slime on the galena surface was light but the coating of galena slime on the fluorite surface was heavy. In a similar way, the coating of quartz slime on the galena was fairly light, but the coating of galena slime on the quartz particle was heavy. According to the chemical hypothesis, using the same constituents and regardless of which is the slime, the density of the coating should be expected to be the same. This, however, as shown above, is not true. Moreover, fluorite (CaF2), having a solubility of 16 mg.4a per liter of water, should not be

Jan 1, 1943

By A. H. Daane, B. J. Beaudry

The uranium-antimony system has been investigated by metal-lographic, therma1, and X-ray methods. There are four intermediate phases present: USbz and U,Sb, which undergo peritectic decomposition at 1355°and 1695OC, respectively, and USb and UPSbs which melt congruently at 1850"and 1800°C, respectively. There is a eutectic between U4Sbs and USb which melts at 1770°C. USb,, UsSb4, and USb are line compounds while U4Sbs has a solubility range of approximately 1.5 at. pct at 1770°C. Single crystal studies of the cmpo,und U.,SbS show it to have a hexagonal unit cell with a, = 9.268A and c, = 6.201b;. The solubility of uranium in liquid antimony increases from 0.1 at. pct at 650°C to 1.5 at. pct at 900°C. The solubility of antimony in a, P, and y uranium is approximately 0.02, 0.1, and 0.05 at. pct respectively. IN nuclear reactors, uranium has been used mainly in the form of solid-fuel elements of the pure-metal or uranium-rich alloys. However, in attempts to overcome some of the difficulties inherent in solid-fueled reactors, reactor concepts hive been proposed in which the fuel would be a liquid metal, such as a solution of uranium in another metal. The liquid-metal fueled reactor (LMFR) designed by the Brook-haven National Laboratory of the U.S. Atomic Energy Commission is an example. Fuel for this reactor is a solution of approximately 800 ppm enriched uranium in bismuth.' To design such liquid-fueled reactors, a background of information on the solubility of uranium in various metals is obviously necessary. The present study was undertaken to pro- vide information for such purposes, and in addition, to study the uranium-antimony system as a whole. A preliminary literature search showed that some work had been done previously on uranium-antimony alloys, but no attempts had been made to make a complete study of this system. The solubility of uranium in antimony was studied by Cammack and Bridger' up to 950°C and by Hayes and Gordon up to 900'-C. Their data are plotted in Fig. 1. Ferro, in an X-ray structure study of some uranium-antimony compounds, found USg to be tetragonal with a, = 4.2724 and c, = 8.471A, U3S4 to be b.c.c. with aO = 9.095A and USb to be f.c.c. with a,= 6.191A. Colani reported the compound UsSba; however, work by Ferro and by this Laboratory showed that this compound does not exist. foote' reported several intermediate phases in the uranium-antimony system, but did not identify them. Greninger' re-ported finding an intermediate phase in a 5 at. pct Sb

Jan 1, 1960

By Arthur F. Taggart

THIS paper proposes two units (believed to be new) for designating, respectively, capacity and efficiency for primary and intermediate crushers. CAPACITY Operators know that the tonnage of rock that can be put through a particular crusher in a given time depends primarily upon the crusher setting (open setting for jaw and primary gyratory types; closed setting for cones and rolls), and secondarily on the lithological character of the feed, its maximum size, the amount of fines that it carries, its moisture content, and the method of feeding. The proposed capacity formula states a relationship between all of these factors in one empirical equation, as follows: TR = TR80KK'K" [I] in which TR, called "reduction tons per hour" is the new capacity unit; T = tons of feed per hour requiring crushing; R50 = "80-per cent reduction ratio"; and K, K', and K" are constants that approximate quantitatively lithological character, moisture content and method of feeding, respectively. T is the hourly weight of material in the feed that is coarser than the coarsest discharge of the crusher. (For marked departures from the normal specific-gravity range-2.5 to 2.7-T should be adjusted to corresponding volumes.) In the crushers under consideration, discharge from the crushing zone is limited by the minimum dimension of a particle. Passage of the particle through a square-mesh screen is limited by its intermediate dimension. Sheppard' has shown that the ratio of maximum: intermediate: minimum average dimensions of broken rock fragments may be taken as about I.5: I: 0.6 for nonslabby rocks such as are normally crushed, while for slabby rocks, such as thin-bedded sedimentaries, foliated schists and gneisses, and the blocky metamorphics, the ratio is of the average order I.3:I:0.3, with the possibility of an even smaller relative value for the thickness. Gaudin2 finds the corresponding ratio for -150 + 200-mesh pyrite to be I.4: I : 0.6. The hourly tonnage of crusher-plant feed, therefore, should be diminished by the amount thereof that would pass a square-mesh screen of which the aperture is S times the crusher setting, where S is a shape factor with the value I.7 for rocks that break exceptionally cubic; 2 for ores and rocks that exhibit no appreciable tendency to slabbiness; and 2.5 to 3, or, in exceptional cases, even larger, when the tendency to slabbiness is marked. R80 is the quotient of the theoretical square-mesh aperture that would pass 8o per cent of the feed, divided by the aperture that would pass S0 per cent of the product. It is best determined from screen tests.

Jan 1, 1941

By O. Schaaber

THE existence of isothermal martensite without the presence of the athermal form thus far has been established only in a 6 pct Mn, 0.6 pct C steel' and a 23 pct Ni, 3.5 pct Mn-Fe alloy.' " Theoretical calculations, moreover, predict this phenomenon to be observable only within narrow composition limits.' In all other steels, it is generally considered that martensite inevitably forms upon cooling below the M, temperature. Even the possibility of isothermal martensite transformation has been questioned for the common steels, although Averbach and Cohen" showed that martensite formation does not stop when the cooling is arrested and that as much as 5 pct retained austenite may transform isothermally to martensite at room temperature. This note presents some observations made on a number of low alloy steels by dilatometric and inductive measurement of isothermal austenite transformation at various temperatures close to the Ms. In the dilatometer studies reported previously, austenitized specimens were immersed in a salt bath maintained at proper temperature and immediately inserted into the dilatometera8' already at thermal equilibrium with the bath. The dilation was read on a dial gage divided into 0.001 mm intervals. Induction measurements were made with a separate excitation source and a test coil, sealed together in a salt-tight case. The increase in the induced voltage caused by the transformation products was recorded at a paper speed of 4 mm per sec. The steels studied are listed in Table I. Specimens prepared from steels 1 and 2 were in the form of hollow cylinders 60 mm long, 10 mm ID, 14 to 16 mm OD, while specimens of steels 3 and 4 were in the form of notched-bar impact test specimens, DVM type. Specimens of steel 5 were in the form of strip, 7x0.3 mm. In the bainitic range, curves representing the amount of transformation vs log time have a single step appearance, as is shown by the curve of isothermal transformation at 320°C in Fig. 1.Y ransformation proceeds in accordance with the Austin-Rickett equationo P -----------= k t" 100-P where P is the percentage of transformed austenite, t is time, and k and n are constants. Near the M, temperature, however, complex curves having two steps are observed. As shown in a previous papera for steel 3 isothermally transformed at 245"C, such complex curves can be separated into two separate transformations or reactions each following the Austin-Rickett equation. The values of n in the equation for a number of such two-step transformations are listed in Table 11 for steel 1. At transformation temperatures of 250" and 235"C, evidence of a second transformation is very

Jan 1, 1956

By William Klement

THIS note reports the results of some attempts to metastably extend the primary solid solubilities of tin, antimony, and silicon in silver by rapidly quenching these binary alloys from the melt. The procedures employed for alloy preparation, quenching, and Debye-Scherrer X-ray diffraction measurements have been described elsewhere. The lattice parameters obtained for the fcc structures in Ag-Sn and Ag-Sb alloys are presented in Fig. 1, together with the results of Owen and Roberts.4 Alloy compositions are believed to be within ± 0.2 at. pct on the basis of negligible weight losses during preparation and reproducibility of the lattice spacings. Diffraction lines from hcp phases were also detected: the relative visual intensities of the fcc (200) and hcp (10.2) lines were estimated as shown in Table I. The intensities were m for the (200) and w for the (10.2) in a quenched 13.6 at. pct Sn alloy: a lattice spacing for the fcc structure could not be reliably obtained, however. As seen from Fig. 1, the lattice parameters vary linearly with composition up to 8., at. pct Sb and 13." at. pct Sn, respectively, which may be compared with the maximum equilibrium solid solubilities5 of 7.2 at. pct Sb and 11.5 at. pct Sn, respectively. The presence of the hcp phases, which were found together with the fcc phases, is a consequence of the competition1,3 in the nucleation and growth processes during solidification. Both the fcc and hcp phases must be of the same composition for those alloys in the range where lattice parameters increase linearly with solute concentration. Despite much effort with quenched alloys of several silicon concentrations, it has not been possible to obtain lattice spacings of sufficient precision to establish the variation of lattice parameter with composition or the maximum metastable solid solubility: the equilibrium data5 suggest little primary solid solubility of silicon in silver. For alloys containing less than 16 at. pct Si, the Si (111) diffraction line was not observed. For alloys containing 10 to 25 at. pct Si, faint low-angle lines were detected and these could be indexed as the (10.0), (00.2). and (10.1) reflections of an hcp structure with a = 2.87 + 3A. c = 4.52 ± 3A, and c/a = 1.57 ± 2. The 16 at. pct Si alloy was quenched to, and examined at. liquid nitrogen temperatures without any change observed in the proportion of the hcp phase. These experiments were jointly supported by the Atomic Energy Commission and Office of Naval Research under a program directed by Professor P. Duwez. The assistance of H.-L. Luo and A. Abu-Shumays is acknowledged. as is the support afforded

Jan 1, 1965

By D. M. Buck

greatly restricted and every effort is being made to do away with it where possible, and to substitute a lead-base babbitt or a babbitt with 50 per cent,. tin. The amount of solder having the composition 50 lead, 50 tin, used by the Pennsylvania Lines East, during 1917, was approximately 100,000 lb., but there is reason to believe that a large portion of this can be replaced by a 60-lead, 40-tin solder with satisfactory results, and instructions have been issued to this effect. Inregard to the quantity of new tin used, it is not possible to give the amount, except approximately and from calculations based on the 1917 consumption of bearing metals by the Pennsylvania Lines East only, it is estimated that about 770,000 Ib. of new tin were required in a total of about 21,380,000 lb. of all kinds of bearing metals turned out by the foundry or purchased in the market. No change has been made in the specifications for bearing metals for some years, as the metals used have been satisfactory. A large proportion of the bearing metals are made up from old material re-melted and brought to standard composition by some addition of new metal, and every effort is being made to utilize old material to the best advantage and reduce the amount of new metal of all kinds purchased. For a number of years no tin has been used in the lining metal of either passenger or freight car journal-bearings, except such small amounts as come in from re-melting old linings. No change has been made in phosphor bronze used for rod bushings, as we should expect some trouble from bushings pounding out of shape if a phosphor bronze were used which contained less tin or more lead than the present specifications call for. In this, as well as in the case of all other bearing metals, we expect to use our utmost endeavors to economize and to substitute for tin wherever possible. The Tin-plate Industry D. M. Buck,* Pittsburg, Pa.—During the first 5 months of 1918, approximately 11,000,000 Ib. per month of pig tin were consumed in the United States. Solder, bearing metals, bronzes, etc, used about 5,500,000 lb.; collapsible tubes a little more than 250,000 lb.; tin-foil about 500,000 lb.; and the tin- and terne-plate industry somewhat less than 5,000,000 Ib. In an effort to reduce this consumption and thus conserve our tin supplies, several methods of procedure suggest themselves: 1. Salvage. The most careful and systematic collection and re-use of tin and tin-bearing materials is economically important, in that we thus

Jan 1, 1919

By Walter Crafts, C. M. Offenhauer

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

Jan 1, 1943

By Walter Crafts, C. M. Offenhauer

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

Jan 1, 1943

By Walter Crafts, C. M. Offenhauer

IN a previous study' of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is a more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about r1to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, 1 per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to 1-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudomartensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously' that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By A. T. Sweet, C. E. Plummer, H. W. St. Clair, S. F. Ravitz

The ammonium sulphate process for recovering alumina from clays was proposed by Rinman, Buchner, and others many years ago, and more recently various modifications have been investigated both here arid abroad.l-Is The process, as it is being investigated by the Bureau of Mines, involves the following essential steps: (I) sulphating the alumina in the clay by baking with ammonium sulphate at approximately 4m°C., (2) leaching with water to extract aluminum sulphate and ammonium sulphate, (3) crystallization of ammonium alum, (4) conversion of the alum to aluminum hydroxide by the ammonia gases given off during the baking, and (5) calcining the aluminum hydroxide to alumina. The over-all reaction between kaolinite (the most important alumina mineral in most clays) and ammonium sulphate during the baking step may be expressed as follows: Al2O8.2SiO2.2H2O + 4(NH4)2SO4 = (NH4)2SO4.A12(SO4)3 + 6NH3(g) + SH2O(g) + 2SiO2 [i] There is definite evidence. as is shown later. that the aluminum occurs in the baked product as anhydrous ammonium alum rather than as aluminum sulphate, provided the baking temperature is not much above 4o0°C. It is probable that the reaction takes place in the following steps: (NH4)2SO4 = NH3 -(- NH4HSO4 [2] Al2O3.2SiO2.2H2O + ;}NH4HSO4 = A12(SO4)3 + 3NH3 + 5H2O + 2SiO2 [3] A12(SO4)3 + (NH4)2SO4 = (NH4)2SO4.A12(SO4)3 [4] In onc modification of the process, the ammonium sulphate is first decomposed according to Eq. z, and the clay is treated with the resultant ammonium acid sulphate. Reaction I is highly endothermic. On the basis of the best available data, the following values have been estimated for the changes in heat content and free energy in calories per mol: -AH = 331,040 - 34T -AF = 331~040 - 78.3T log T - 671.8T The heat required' at 25OC is 320,900 cal. per mol, or 5650 B.t.u. per pound of alumina. The theoretical equilibcium pressure of NHa, based on the free-energy equation given above, increases very rapidly as the temperature rises above 3o0°C., and reaches one atmosphere at 43s°C. Ammonium alum solutions are almost ideal for crystallization. A saturated solution of ammonium alum contains 126. grams of AlzOa per 1000 grams of water at 90°C. and only 16.5 at 20 C. Hence most of the alum crystallizes out when the

Jan 1, 1944

By Walter Crafts, C. M. Offenhauer

IN the course of study of the heat-treatment of low-alloy steels, the behavior of alloy carbides at subcritical temperatures was found to vary from that indicated by published investigations. In order to establish the ranges of formation of carbides in low-chromium steels, a series of steels containing up to 0.60 per cent carbon and 7.5 per cent chromium has been investigated with respect to the type of carbide formed after quenching and tempering at sub-critical temperatures. The types of carbide formed as products of austenite transformation to pearlitic and pseudomartensitic structures have also been determined. There have been numerous investigations of the carbides in chromium steels and diagrams of the iron-chromium-carbon system.' Westgren, Phragrnén and Ne-gresco2 identified the following carbides in chromium steels: Cementitc............ Fe3C (containing up to about 15 per cent Cr) Trigonal carbide...... Cr7C3 (containing up to about 55 per cent Fe) Cubic carbide........ Cr4C (or Cr23C6) (con- . taining up to about 25 per cent Fe) Orthorhombic carbide , cr3c2 (containing only a few per cent Fe) The most complete information on the iron-chromium-carbon system was supplied by Tofaute, Sponheuer and Bennek3 and by Tofaute, Küttner and Büttinghaus,4 By means of dilatometric and microscopic stzdies, sections of the diagram were established. This diagram showed the presence of the special carbide CI7C3 in steels containing about 2.5 per cent to 12 per cent chromium in low-carbon steels and from about 5 to 20 per cent chromium in I per cent carbon alloys from A1 to room temperature. The present investigation has shown that a carbide inversion from Fe3C to Cr7C3 occurs in chromium steels containing more than about I per cent chromium after hardening and tempering at subcritical temperatures above about 500°C. The type of carbide observed to be more stable in quenched and tempered specimens was also found to be produced by direct austenite transformation in corresponding ranges of temperature and composition. Experimental Procedure The steels used in these experiments were made in magnesia-lined crucibles in high-frequency furnaces and were poured into 2-in. ingots. The ingots were forged into I-in. round bars, and all heat-treatments were conducted on specimens of this section size. Sections of the bar stock 8 in. long were hardened in air, oil or water, depending on the composition, after heating to a temperature sufficiently high to dissolve the carbides. The low-carbon and low-chromium steels were not entirely martensitic after quenching, and contained some ferrite and pseudomartensite. No pearlite was observed in the microstructures after quenching. The quenched bars were cut into

Jan 1, 1942

By Walter Crafts, C. M. Offenhauer

IN the course of study of the heat-treatment of low-alloy steels, the behavior of alloy carbides at subcritical temperatures was found to vary from that indicated by published investigations. In order to establish the ranges of formation of carbides in low-chromium steels, a series of steels containing up to 0.60 per cent carbon and 7.5 per cent chromium has been investigated with respect to the type of carbide formed after quenching and tempering at sub-critical temperatures. The types of carbide formed as products of austenite transformation to pearlitic and pseudomartensitic structures have also been determined. There have been numerous investigations of the carbides in chromium steels and diagrams of the iron-chromium-carbon system.' Westgren, Phragrnén and Ne-gresco2 identified the following carbides in chromium steels: Cementitc............ Fe3C (containing up to about 15 per cent Cr) Trigonal carbide...... Cr7C3 (containing up to about 55 per cent Fe) Cubic carbide........ Cr4C (or Cr23C6) (con- . taining up to about 25 per cent Fe) Orthorhombic carbide , cr3c2 (containing only a few per cent Fe) The most complete information on the iron-chromium-carbon system was supplied by Tofaute, Sponheuer and Bennek3 and by Tofaute, Küttner and Büttinghaus,4 By means of dilatometric and microscopic stzdies, sections of the diagram were established. This diagram showed the presence of the special carbide CI7C3 in steels containing about 2.5 per cent to 12 per cent chromium in low-carbon steels and from about 5 to 20 per cent chromium in I per cent carbon alloys from A1 to room temperature. The present investigation has shown that a carbide inversion from Fe3C to Cr7C3 occurs in chromium steels containing more than about I per cent chromium after hardening and tempering at subcritical temperatures above about 500°C. The type of carbide observed to be more stable in quenched and tempered specimens was also found to be produced by direct austenite transformation in corresponding ranges of temperature and composition. Experimental Procedure The steels used in these experiments were made in magnesia-lined crucibles in high-frequency furnaces and were poured into 2-in. ingots. The ingots were forged into I-in. round bars, and all heat-treatments were conducted on specimens of this section size. Sections of the bar stock 8 in. long were hardened in air, oil or water, depending on the composition, after heating to a temperature sufficiently high to dissolve the carbides. The low-carbon and low-chromium steels were not entirely martensitic after quenching, and contained some ferrite and pseudomartensite. No pearlite was observed in the microstructures after quenching. The quenched bars were cut into

Jan 1, 1942

By P. S. Kotval

THE nominal composition of "hconel" alloy 718 is 53 pct Ni-19 pct Cr-19 pct Fe-3 pct Mo-5 pct Cb-0.8 pct Ti-0.6 pct A1-0.05 pct C. The nature of the fine strengthening phase in this alloy has been open to debate. iseelstein' has identified the strengthening phase in the matrix of "Inconel" alloy 718 as a colum-bium-rich fcc y' precipitate. Wagner and Hall2 have noted that the metastable y' precipitate overages to an orthorhombic Ni3Cb precipitate which manifests itself as an acicular phase. These authors thus clarify Kaufman and Palty's earlier suggestion3 that the strengthening phase is orthorhombic NisCb. Wagner and Hall's review also reports work by cometto4 which suggests that the fine strengthening phase is not fcc / but bct with a Ni3V type structure and based on the Ni3Cb composition. Kir man and warringtons have shown by analysis of electron diffraction patterns that the metastable strengthening phase in a 55 pct Ni-15 pct 0-25 pct Fe-5 pct Cb alloy is bct and based on the Ni3Cb composition. Raymond~ has provided electron diffraction evidence which indicates that "y'" in "Inconel" alloy 718 is bct in structure. This note presents unambiguous data on the structure of the strengthening phase in "Inconel" alloy 718* ihe samples of this alloy were provided by courtesy of the ASTM task group on phase identification in 'superalloys. after prolonged aging. The electron diffraction patterns presented show that strengthening phase is not y' but bct as was proposed by cometto,' supported by Raymond,~ and identified by Kirman and warringtons in a similar system. Fig. 1 shows a selected-area diffraction pattern taken from a thin foil of "Inconel" alloy 718. The sample had received a final aging treatment of 1118 hr at 1200°F to precipitate the strengthening phase in the matrix. Fig. 2 is a schematic representation of the features in the diffraction pattern. The matrix orientation is (0i2)~,, as is shown by the indexed zone in Fig. 3. Figs. 4, 5, and 6 illustrate that the pattern, Fig. 1, is composed of three bct zones, viz., (0i4), (2i0), and (i01), respectively. These three zones correspond, respectively, to the alignment of the c axis of the bct unit cell with the [001], [loo], and [OlO] directions of the fcc matrix and serve to illustrate the fully coherent relationship of the precipitate structure with the matrix. The streaking of certain reflections is also consistent with the above orientation relationship. The interpretation of the diffraction effects presented here shows the metastable strengthening precipitate in "Inconel" 718 alloy to be bct in structure and demonstrates that it is not an ordered fcc y' phase.

Jan 1, 1969

By C. P. Kempter, H. L. Brown

Recently Chang determined heat content values of tantalum monocarbide and tungsten monocarbide from 325" to985°Kand 326" to 912"K, respectively, and, using other published data, made certain solid-state calculations for these two compounds. Chang expressed his Cp values in units of cal per "g-atom alloy" OK. We consider this poor usage since l) tantalum monocarbide and tungsten monocarbide are not alloys, and 2) an atom of an alloy does not exist. However, one may convert his values to units of cal per mole OK by merely doubling them. For TaC he calculated a Nernst-Lindemann parameter of 9.07 x lo-' "g-atom per cal". We have refined the calculation of this parameter and have arrived at a much lower value. Rather than approximating the true coefficient of volume expansion, a,, as three times the true coefficient of linear expansion, a, we calculated u, from the TaCo.99, lattice expansion data of Fries and Wahman.38 At 25°C a, = 14.3 x 10s per OC. The molar volume was calculated to be 13.32 cu cm per mole at 25°C. From the lattice compression data of Champion and Drickamer~' for "TaC,.,," we determined the room-temperature compressibility to be 3.5 x 10-l3 sq cm per dyne by plotting l/Vo (aV/ap) vs applied pressure. (Since the lowest pressure at which Champion and Drickamer reported a relative volume was 20 kbars, one may calculate an approximate compressibility only. Chang calculated a compressibility of 3.175 x 10-l3 sq cm per dyne from the same data.) We also used twice Chang's value of Cp = 4.38 cal per "g-atom alloy" OK for the molar heat capacity at 25°C. From these four values we calculated a Nernst-Lindemann parameter of 5.8 x 10-l4 mole per erg or 2.4 x lo- ' mole per cal, and a Cdil value of 2.3 x 10' ergs per mole OK. In the case of tungsten monocarbide no isothermal compressibility has been reported; however, one may calculate an approximate Nernst-Lindemann parameter by using the adiabatic compressibility reported by Brown, Armstrong, and Kempter." Using the WC lattice expansion data of Mauer and Bolz~l we calculated a, to be 12., x l0-' per OC. The molar volume was calculated to be 12.5 cu cm per mole. We used twice Chang's value of Cp = 4.18 cal per "g-atom alloy" OK for the molar heat capacity at 25°C. From these four values we calculated an approximate Nernst-Lindemann parameter of 2 x lo-' mole per cal, which is much

Jan 1, 1969