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By Ellsworth Daggett

Jan 1, 1879

By R. Gibson, F. S. Turneaure

The tin deposit of Monserrat, Bolivia, consists of one major vein 1600 m in length. The ore is unusual because of the notable quantity of teallite, even though cassiterite is the principal tin mineral. The deposit, formed at shallow depth under a wide range in temperature, may be classed as xenothermal. Polished sections reveal a complex history of replacement, with low-temperature minerals deposited before high-temperature minerals. THE Monserrat mine1'2 of the Compania Minera Monserrat is located 11 km east of Callipampa, a station on the Antofagasta and Bolivia Railway some 55 km south of Oruro. The mine lies in a group of low hills on the east slope of a broad valley of northerly trend. The valley is separated from the altiplano to the west by a prominent ridge that rises several hundred meters above the surrounding country (fig. 1). Callipampa is situated a short distance west of the ridge, near the east margin of the altiplano. The elevation at Callipampa is 3700 m and at the mine 4100 m. Monserrat is near the northeast limit of what may be called the Poopo-Pazna district, an area of northwesterly trend about 25 km in length and 15 km in width. The district includes the tin-silver prospects of Poopo to the north, the zinc-tin deposits of Salvador to the southwest, and the tin-tourmaline veins of Avicaya to the south. Along the prominent ridge already mentioned, which forms the western limit of the district, there are several tin and antimony prospects including the Trinacria mine, which is famous for its fine specimens of cylindrite. Although the principal ore mineral is cassiterite, the tin deposit of Monserrat contains an unusual amount of teallite. Cassiterite occurs in part as a finely granular mixture with sulphides of iron and zinc and in part as needle tin of late hypogene age. The regional structure is dominated by broad folds typical of the Andes of Central Bolivia and similar to those of the Llallagua district. The area adjacent to the Monserrat mine is underlain by thin-bedded shale somewhat variable in color but predominantly dark grey to black. The beds strike northwest and dip at low angles southwest, forming part of the northeast limb of a major syncline. A resistant sandstone formation, which underlies the black shale, crops out along a ridge east of the mine and also accounts for the high ridge to the west that was previously mentioned. No igneous rock of any kind is found at Monserrat. The nearest intrusive bodies known are the irregular dikes and masses of quartz porphyry at Avicaya about 10 km to the south. Mine Workings The Monserrat property includes one main vein, which can be followed on the surface for 1600 m, and two or three branch veins of distinctly minor importance. The principal haulage tunnel, referred to as the San Carlos adit or level III, extends northeasterly for 1700 m, following the main vein throughout most of this distance. Above the adit, there are

Jan 1, 1951

By F. S. Turneaure, R. Gibson

The tin deposit of Monserrat, Bolivia, consists of one major vein 1600 m in length. The ore is unusual because of the notable quantity of teallite, even though cassiterite is the principal tin mineral. The deposit, formed at shallow depth under a wide range in temperature, may be classed as xenothermal. Polished sections reveal a complex history of replacement, with low-temperature minerals deposited before high-temperature minerals. THE Monserrat mine1'2 of the Compania Minera Monserrat is located 11 km east of Callipampa, a station on the Antofagasta and Bolivia Railway some 55 km south of Oruro. The mine lies in a group of low hills on the east slope of a broad valley of northerly trend. The valley is separated from the altiplano to the west by a prominent ridge that rises several hundred meters above the surrounding country (fig. 1). Callipampa is situated a short distance west of the ridge, near the east margin of the altiplano. The elevation at Callipampa is 3700 m and at the mine 4100 m. Monserrat is near the northeast limit of what may be called the Poopo-Pazna district, an area of northwesterly trend about 25 km in length and 15 km in width. The district includes the tin-silver prospects of Poopo to the north, the zinc-tin deposits of Salvador to the southwest, and the tin-tourmaline veins of Avicaya to the south. Along the prominent ridge already mentioned, which forms the western limit of the district, there are several tin and antimony prospects including the Trinacria mine, which is famous for its fine specimens of cylindrite. Although the principal ore mineral is cassiterite, the tin deposit of Monserrat contains an unusual amount of teallite. Cassiterite occurs in part as a finely granular mixture with sulphides of iron and zinc and in part as needle tin of late hypogene age. The regional structure is dominated by broad folds typical of the Andes of Central Bolivia and similar to those of the Llallagua district. The area adjacent to the Monserrat mine is underlain by thin-bedded shale somewhat variable in color but predominantly dark grey to black. The beds strike northwest and dip at low angles southwest, forming part of the northeast limb of a major syncline. A resistant sandstone formation, which underlies the black shale, crops out along a ridge east of the mine and also accounts for the high ridge to the west that was previously mentioned. No igneous rock of any kind is found at Monserrat. The nearest intrusive bodies known are the irregular dikes and masses of quartz porphyry at Avicaya about 10 km to the south. Mine Workings The Monserrat property includes one main vein, which can be followed on the surface for 1600 m, and two or three branch veins of distinctly minor importance. The principal haulage tunnel, referred to as the San Carlos adit or level III, extends northeasterly for 1700 m, following the main vein throughout most of this distance. Above the adit, there are

Jan 1, 1951

By J. M. Hodge, H. M. Reichhold, R. D. Manning

The work reported in this paper represents the first of a series of investigations of the factors governing notch toughness in ferritic materials. This paper is concerned with two of these factors, namely, the effect of ferrite grain size, and the effect of an alloy content of 3.64 pct nickel. Every effort has been made to hold other variables constant in order to isolate the effects of these two. The compositions chosen for study were 0.02 pct carbon in order to eliminate the variable of carbide distribution, one steel being a plain carbon steel and the other alloyed with 3.64 pct nickel. Both steels were fully killed with silicon, without an aluminum addition, so that the variable of deoxidation practice was held constant. Finally, all samples were water quenched from 1200°F and tempered 24 hr at 300°F in order to hold the quench aging variable constant and to decrease the tendency for grain boundary carbide precipitation. The ferrite grain size was varied by heat treatment over the range of A.S.T.M. grain size numbers of about 2 to 6 in each steel so that the effect could be evaluated individually and the notch toughness of each steel could be compared at comparable ferrite grain sizes. Materials and Experimental Procedure MATERIALS The materials for this investigation were obtained from Battelle Memorial Institute as 250 Ib induction melts which were forged into 134 in. square bars and then rolled into 0.6 in. square bars. Both steels were deoxidized with silicon, the aluminum content being held to a minimum. Check analyses of these materials are shown in Table 1. HEAT TREATMENT FOR GRAIN SIZE The grain size was varied by varying both the cooling rate and the austeni- tizing temperatures. The heat treatments used and the corresponding grain size values are tabulated in Table 2. A semicontinuous network of grain boundary carbide was observed in the microstructure of both the plain carbon and nickel steels in the as rolled condition and an additional heat treatment was necessary to eliminate this variable. Therefore, all samples were water quenched from 1200°F and tempered 24 hr at 300°F as a final treatment before testing. This heat treatment served to eliminate to a great extent the grain boundary carbide network and to hold the quench aging effect constant. IMPACT TESTING Standard keyhole Charpy impact tests were machmed after heat treatment. A bath of acetone and dry ice was used for refrigeration of the samples down to temperatures of —94°F (—70°C) and a bath of either methyl-cyclohexane or isopentane, cooled by

Jan 1, 1950

By Edward Saibel

The object of this study is to investigate the effect of prior tensile strain on the fracture stress of a metal. This is done in a theoretical manner starting from the point of view developed by the author in his thermodynamic theory of the fracture of metals. The results are compared with experimental findings and a rational interpretation is given to work that shows the variation of fracture stress with prior tensile strain. Theory op Fracture In the thermodynamic theory of the fracture of metals developed earlier,' it was shown that a critical strain energy per unit volume exists. This is characteristic of the material and may be calculated from basic thermodynamic data. When the area under the flow curve reaches this characteristic value, fracture occurs. The general criterion for fracture is expressed in the form « ± u. = JLJW/V [r] where L, is the latent heat of melting of the unit volume, AV/V is the ratio of the change in volume of the unit volume to the original unit volume on passing from the solid to the liquid state, uo is an initial energy density, and J is a conversion factor, which changes energy from heat to mechanical units. As a first approximation, it will he assumed that the stress-strain curve is the straight line that extends from the point of maximum load onward, extrapolated back to the axis of zero strain, for example, line AC of Fig. I. If the point of intersection of the straight line with the axis of zero strain is represented by v., and the slope of the straight line is p, the analytic form of the flow curve is a = (To + pd [2] and the strain energy per unit volume evaluated from Jud6 becomes (u,6 + pa2/2). When this latter term reaches the critical value JLbVIV (designated hereafter by M), it is assumed that fracture occurs. This leads to the following expression for the fracture stress: a; = V«r.2 + iMp [3] The strain at fracture may be calculated from Eqs 2 and 3. It is assumed in this paper that the initial energy density is either zero or that it may be neglected. Furthermore, the complicated states of stress and of strain existing in the necked region of the tensile specimen after "necking" has occurred will also be neglected for the present, so that average conditions will be assumed to prevail. Flow and Fracture Curves LudwigZ attributed to materials two fundamental properties: (I) a resistance to flow and (2) a resistance to fracture. He reasoned that at each instant the

Jan 1, 1948

By Edward Saibel

The object of this study is to investigate the effect of prior tensile strain on the fracture stress of a metal. This is done in a theoretical manner starting from the point of view developed by the author in his thermodynamic theory of the fracture of metals. The results are compared with experimental findings and a rational interpretation is given to work that shows the variation of fracture stress with prior tensile strain. Theory op Fracture In the thermodynamic theory of the fracture of metals developed earlier,' it was shown that a critical strain energy per unit volume exists. This is characteristic of the material and may be calculated from basic thermodynamic data. When the area under the flow curve reaches this characteristic value, fracture occurs. The general criterion for fracture is expressed in the form « ± u. = JLJW/V [r] where L, is the latent heat of melting of the unit volume, AV/V is the ratio of the change in volume of the unit volume to the original unit volume on passing from the solid to the liquid state, uo is an initial energy density, and J is a conversion factor, which changes energy from heat to mechanical units. As a first approximation, it will he assumed that the stress-strain curve is the straight line that extends from the point of maximum load onward, extrapolated back to the axis of zero strain, for example, line AC of Fig. I. If the point of intersection of the straight line with the axis of zero strain is represented by v., and the slope of the straight line is p, the analytic form of the flow curve is a = (To + pd [2] and the strain energy per unit volume evaluated from Jud6 becomes (u,6 + pa2/2). When this latter term reaches the critical value JLbVIV (designated hereafter by M), it is assumed that fracture occurs. This leads to the following expression for the fracture stress: a; = V«r.2 + iMp [3] The strain at fracture may be calculated from Eqs 2 and 3. It is assumed in this paper that the initial energy density is either zero or that it may be neglected. Furthermore, the complicated states of stress and of strain existing in the necked region of the tensile specimen after "necking" has occurred will also be neglected for the present, so that average conditions will be assumed to prevail. Flow and Fracture Curves LudwigZ attributed to materials two fundamental properties: (I) a resistance to flow and (2) a resistance to fracture. He reasoned that at each instant the

Jan 1, 1948

By G. P. Airey, R. F. Mehl, T. A. Hughes

The coarsening of cementite particles in a ferrite matrix has been studied in a series of steels with 0.15 pct C only and 0.15 pct C plus 1 pct Ni, Mn, and Cr, respectively. Two initial states were employed: quenched nartensite, and quenched and cold-rolled martensite. A series of tempering temperatures between 500' and 700" and tempering times of up to 190 hr were used. The structures were studied by replica and transmission electron microscopy. Particle size distribution curves were determined. From the average size value coarsening curves were obtained. These were plotted in accordance with the Wagner analysis assuming diffusion control. A discussion of the significance of the results is given. L HE reactions occuring upon the tempering of martensite have long engaged the attention of metallurgists. The latter stages, when cementite particles coarsen in a ferrite matrix, have been studied both qualitatively and quantitatively. Studies of such coarsening processes have recently been spurred by the publication of the Lifshitz-Wagner theory1, and the extension of this to the a Fe-Fe3C system by Oriani3 and by Li, Blakely, and einold. Following Wagner the coarsening process is often designated as "Ostwald Ripening". The only quantitative data on the rate of coarsening, except for the work of Hyam and uttin,' in the a Fe-Fe3C system are those of Bannyh, Modin, and odin' for a commercial eutectoid steel and those of Heckel and ereorio" using a pure eutectoid steel. The data of Bannyh, Modin, and Modin have been employed by 0riani3 to derive the a Fe-FeE interface energy. The reaction is one of the most important ones in steel and is worthy of detailed study. This is the purpose of the present study. Laboratory heats were prepared; these were steels with approximately 0.15 pct C, selected so that the num ber of carbide particles would be relatively small and thus so that the overlapping diffusion fluxes would be minimized, presumably a desirable circumstance.'-3 In addition to Fe-C alloys, comparable heats containing 1 pct of Ni, Mn, and Cr, respectively, were included with a view of appraising the effect of alloying elements. This report includes an account of the micro-structures observed, primarily with the electron microscope, and of kinetic data and their interpretation. MATERIALS AND TREATMENT The alloys were prepared from electrolytic iron ("Plastiron") and high-purity graphite; these were melted in a zirconia crucible using a vacuum furnace. The alloy steels were made by adding electrolytic nickel, electrolytic manganese, and "vacuum grade" chromium, respectively, under a partial pressure of argon. Each melt was poured into a mold within the vacuum furnace and cooled in the mold. The ingots were 2 in. in diam. and 8 to 10 in. long. The analysis of the alloys is given in Table I. These ingots were hot-rolled to strip 0.1 in. thick, then cold-rolled to 0.05 in. and each alloy split into two batches. One batch was austenitized at 1200 for 1 min, quenched in cold brine, then cold-rolled to 0.02 in.; samples given this treatment are hereinafter designated as "worked". The other batch was cold-rolled to 0.025 in., austenitized at 1200" for 1 min, and quenched in cold brine; such samples are hereinafter designated as "quenched". These two batches were then tempered, as below. The purpose of the treatment given the first batch was to provide an initial structure of cold-worked martensite, with the expectation that the additional defect structure created by cold work would encourage a higher rate of nucleation of cementite on tempering and hence a more uniform distribution of cementite particles. Individual specimens were sealed in evacuated quartz or Pyrex tubes, then tempered in a muffle furnace. The temperature control was better than 3'C at 700. Tempering treatments wer: performed at 400°, 500°, 550°, 600°, 65o°, and 700C for time periods between 15 min and 190 hr. PREPARATION OF SPECIMENS Specimens for optical and replicalelectron microscopy were mounted, polished conventionally, and etched with 2 pct nital. For electron microscopy, single-stage "formvar" replicas were made, dry-stripped and rotary-shadowed with chromium at an angle of 30 deg. Carbide extraction replicas were prepared from electropolished specimens usirig the method described by Smith and uttin.' Thin foils for electron transmission microscopy were prepared by chemical thinning in an H202-HF bath prior to electropolishing in a chromium trioxide-acetic acid solution. The most

Jan 1, 1969

By Henry A. Wriedt, John Chipman

Equilibrium in the reaction of hydrogen gas with oxygen in liquid nickel, iron, and their alloys has been studied at temperatures of 1500° to 1700°C. The equilibrium con^stant, 0/p, [% O], is greater in nickel then in iron by a factor of 45 at 1594 C. In the alloy steel range, the activity coefficient of oxygen is slightly increased by the presence of nickel. NICKEL is more resistant to oxidation than is iron, in both the solid and the liquid state. The solubility of oxygen,' however, is greater in nickel than in iron. The two metals and their liquid alloys constitute an interesting series of solvents in which the chemical behavior of oxygen can be studied by methods similar to those used for liquid iron.' Experimental Method The furnace assembly, including the tube for preheating the entrant gas, was the same as that described by Gokcen and Chipman. " controlled mixture of water vapor, hydrogen, and argon preheated to approximately the melt temperature impinged upon the surface of the liquid metal. The charge, weighing about 35 g, was held in an alumina crucible at constant temperature in the gas stream for a prolonged period, then quenched in a stream of cold helium, sampled, and analyzed for oxygen by the vacuum fusion method. The temperature of the liquid metal was measured by a Leeds and Northrup optical pyrometer which had been checked against a similar instrument certified by the National Bureau of Standards. The freezing point of electrolytic iron in hydrogen, taken as 153j°C, was used for frequent checks. This point and the emissivity determined by Dastur and Gokcen" established the transmissivity of the optical system. The emissivities of Fe-Ni alloys at 1535" are known from the experiments of Smith and Chipman,h nd it was assumed that the variation with temperature is parallel with that of iron. Recorded temperatures are probably accurate within k5"C. Preparation of Gas Mixtures—Since the ratio H,O:H, had to be much higher for nickel than for iron, it was found convenient to construct two separate systems for gas preparation. In both cases, the ratio of argon to the sum of the other two gases was at least 4, the proportion of hydrogen and water vapor being varied to produce H,O:H, ratios from 0.2 to 70. The primary purpose of the argon was to diminish the errors caused by thermal separation' in the gas phase. In addition, its presence lowered the saturation temperature required for a given H,O:H, ratio and minimized the difficulty of preventing condensation. Both systems were modifications of the one used by Dastur and Chipman.' For H,O:HI ratios greater than 2, welding-grade argon purified by passage over magnesium chips at 620°C was mixed with hydrogen produced in the electrolytic cell shown in Fig. 1. The design was based on a cell used by Bodenstein and Pohl," the important features of which area nickel wire cathode and sodium hydroxide solution as the electrolyte. The rate of hydrogen evolution was controlled by a calibrated ammeter and a rheostat in the circuit, the current efficiency being taken as 100 pct. Precise control of the argon flow rate at about 300 ml per min was achieved by means of a capillary flowmeter having a controlled pressure drop. The molar flow rate was measured volumetrically before each run, and a constant flowmeter setting was maintained thereafter. The A-H, mixture passed through a water saturator of the type previously described," the saturation temperature being maintained automatically with a fluctuation of k0.04"C. The tube between saturator and furnace was heated by a resistance winding. The system was of all Pyrex construction with the exception of the stainless steel tube containing magnesium chips. This was connected with the glass tubing at the upstream end by

Jan 1, 1957

By W. H. HOWARD

Discussion of the paper of W. H. HOWARD, presented at the Salt Lake meeting, August, 1914, and printed in Bulletin No. 92, August, 1914, pp. 2029 to 2046. EDGAR M. DUNN, Anaconda, Mont.-Mr. Howard's contribution to the literature on the Cottrell, method. of electrostatic precipitation of flue dust and fume at Garfield has been read by us with peculiar interest for the reason that we have been experimenting with the same process at Anaconda for the past eight or nine months, with a view to adapting it, under the direction of Dr. Cottrell and the supervision of J. 0. Elton. and D. R. Kellogg, representing the Anaconda Smelter Commission, to the treatment of very large volumes of gases, using large pipe treaters and various forms of electrodes, with about 150,000 volts rectified current. For the amount of gas to be treated here the 5-in. pipe means a cost for installation that is practically prohibitive, clue to the tremendous number of pipes required, and we have been trying to cut installation costs by increasing the size of the pipes. All of our work is still in the experimental stage, but we have obtained good clearance on our mixed roaster and converter gases with a 3-ft. pipe treater 20 ft. long, and with 5 ft. velocity therein; and with a discharge electrode of a single No. 29 nichrome wire set in the middle of the 3-ft. pipe, with. a consequent electrode spacing of 18 in. The current used has been 150,000 to -160,000 volts, .and of course there have been troubles galore, relating particularly to insulation and rectifying. Blast-furnace gases did not give so good clearance, probably due to relative poverty in moisture and S03 content of blast gases to roaster-converter gases; making the blast gases the poorer in conductivity. Our experience has been so satisfactory that in our next experimental plant-built on a somewhat larger scale-we shall attempt a 4-ft. treater with greater electrode spacing, and a somewhat higher voltage-up to 220,000. As to fractional precipitation, mentioned by Mr. Howard, this idea was tried out at Anaconda last year (1913). Formerly our furnaces (Bruntons) roasting flue dust for the recovery of white arsenic, passed so great an, amount of dust to the settling kitchens that refining of this arsenical product was necessary to eliminate the 5 to 20 per cent. dust contained therein. Consequently the manufacture of As203 was effected in two stages, with a reverberatory roast of the first .product to obtain the final 99.5 per cent. As203. -

Jan 11, 1914

By G. L. F. Powell, L. M. Hogan

Large degrees of undercooling have been produced in bulk samples, 400 g, of copper and Cu-O alloys by melting in a slag of commercial soda-lime glass. The maximum degrees of undercooling obtained for copper, hypoeutectic and hypereutectic Cu-O alloy samples were 208°, 218°, and 97°C, respectively. No grain refinement as a function of undercooling was observed for the pure copper samples although in Cu-O alloys there was a marked decrease in grain size at degrees of undercooling greater than 150°C. The grain size change is the result of recrystallization during or immediately following the freezing process. THE first report of a large degree of undercooling in a bulk sample of metal was that by Bardenheuer and Bleckmann1 who produced 258°C undercooling in a 150-g sample of iron by melting in a glass slag. Garbeck2 extended this technique to nickel, cobalt, and copper and reported maximum undercoolings of 220°, 230°, and 60°C, respectively, while Fehling and schei13 undercooled a large number of metals as 2 to 20-g samples slagged in glass. Subsequently, one of the authors4 applied the glass slag technique to bulk samples, -500 g, of silver and obtained a maximum undercooling of 250°C. It was concluded that impurities which normally catalyze nucleation at small degrees of undercooling could be removed by oxidation and solution in a glass slag. Large undercooling of bulk samples of iron, nickel, cobalt, and their alloys has also been reported by other authors.8'10'11 Of the five metals mentioned above, copper was the only one which had not been undercooled by at least 200°C in the form of a large bulk sample. Since copper exhibits high liquid solubility for oxygen, as do iron, nickel, cobalt, and silver, it was considered that a much larger degree of undercooling than that observed by Garbeck for copper should be possible. This was obtained by modification of the technique applied previously to silver and the results of undercooling experiments with 400-g samples of copper form part of this paper. The undercooling behavior of Cu-O alloys was also studied, and the influence of a small oxygen content on the grain structure of undercooled copper was observed metallographically. 1) EXPERIMENTAL The copper was obtained as oxygen-free high-conductivity copper, the major impurities of which were Fe 0.01 pct, Zn 0.015 pct, and Si 0.02 pct. Melting was carried out in an open-ended vertical cylindrical fur- nace wound with Kanthal wire. A schematic diagram of the experimental setup is shown in Fig. 1. The undercooling technique previously used for silver4 involved preoxidation of the samples by melting the silver in contact with the atmosphere. The oxygen content was subsequently reduced by freezing the samples under a glass slag, whereby the oxygen was evolved as gas and the continuous glass cover over the surface prevented re-solution of oxygen when the sample was remelted. Since oxygen is not released as gas when a Cu-O alloy is frozen. the technique had to be modified for use with copper. Also, copper is oxidized during fire refining so that an in situ preoxidation treatment of the samples was not considered to be a prerequisite to large undercooling. Experimentation proved that this supposition was correct. Instead, it was found that care was necessary during melting of the copper under glass to minimize oxygen pickup. Samples of copper weighing 400 g were prepared by adding 50-g pieces of copper to a vitreous silica crucible partly filled with soda lime glass at a temperature of approximately 1000°C. Each piece was quickly immersed in the glass slag and held at this temperature until the oxide coating on the surface decomposed. This change was easily observed and coincided with the formation of gas bubbles in the glass

Jan 1, 1969

By J. A. Retty

HOLLINGER CONSOLIDATED GOLD MINES LTD., through two subsidiary companies, has the exclusive right to prospect in two contiguous areas in central Labrador. This paper presents the results of the mineral exploration that has been conducted in these areas to date. One of the subsidiaries of Hollinger. the Labrador Mining and Exploration Co. Ltd., under an agreement with the Newfoundland Commission of Government, holds an area of approximately 20,000 sq.mi., which comprises the western portion of Newfoundland-Labrador and includes the upper basins of the Hamilton and Naskaupi Rivers. (The main area of the deposits so far discovered may be located on a map at about 67° west longitude and 55° north latitude.) The other subsidiary of Hollinger, Hollinger North Shore Exploration Co. Ltd., has licensed from the Province of Quebec an area of approximately 3900 sq.mi. lying directly north of the Newfoundland-controlled area at the headwaters of the George, Whale, and Swampy Bay Rivers. The M. A. Hanna Co., large iron-ore operators of the Lake Superior region, are participating in the project. With the exception of the years 1940 and 1941, field work has been conducted in the Newfoundland area since 1936, and in the Quebec area since 1939. Hollinger assumed the management in 1942. The field base is at Hollinger Lake, some 300 miles north of the tidewater base at Seven Islands, Quebec, which, in turn, is just west of the mouth of the Moisie River, on the Gulf of St. Lawrence. Air transportation is used exclusively to service the operation.

Jan 1, 1945

By C. C. Martin, M. C. Fuerstenau, R. B. Bhappu

Experiments revealed that quartz could not be floated in conductivity water at any pH with a long-chained sulfonate as collector. Various cations, Fe+++, Al+++, Pb++, Mn++, Mg++, Ca++, are shown to function as activators when the pH is such that the cation hydrolyzes. As the solubility product of the various cation-sulfonates was exceeded at the concentrations involved, precipitated cation-hydroxy-sulfonates, e.g. Fe (RSO3)2OH must be functioning as the collector in these systems. A number of interesting and important phenomena were revealed during a study of the response of beryl to sulfonate flotation.' Perhaps the most important observation is the effect of various cations on flotation response. Experiments showed that in the vicinity of pH 3.0, about three ppm Fe+++ increased recovery of leached beryl from about 60 pct to 100 pct, whereas about four ppm Ca+ + reduced the recovery from 60 pct to 2 pct. With the exception of ferric iron, all of the other cations investigated reduced flotation recoveries drastically at this pH. As postulated in the paper Sulfonate Flotation of Beryl,' even with low concentrations of polyvalent cations and sulfonate (on the order of one ppm), cation-sulfonates are precipitated immediately in solution. That is, blueish-white clouds were noted to form after given additions of salts to known solutions of sulfonate. Furthermore, it was also shown that one-hundred-fold more ferric iron than necessary to form the precipitated cloud had to be added to effect complete flotation. In other words, some form of solid ferric sulfonate was found to function as the collector in these systems. The form of the compound was suggested to be Fe(RCO3)2OH and that bonding to the surface occurs through the hydrogen of the hy-droxyl. As the form of this compound was shown to be dependent on the extent of hydrolysis of ferric iron, it is to be expected that ferric iron will be unique in some pH region. Further, if the presence of hydroxyl in the neutral precipitated compound is of primary importance, then other cations should function similarly but at different values of pH, since they will hydrolyze at different concentrations of hydroxyl ion. To determine whether the premise of hydrolysis is correct, the study was extended to quartz systems in the presence of various cations at various values of pH. EXPERIMENTAL MATERIALS AND METHOD Sodium alkyl aryl sulfonate was chosen as the collector for the experimental work. The reagent, sup plied as a courtesy sample by the Shell Chemical Co., has the following physical properties: Physical form Solid — Finely ground Sulfonate content, pct wt 95 to 97 Molecular Weight 450 to 470 Number of carbon atoms in hydrocarbon chain 25 to 30 Melting Point, OF 250 to 260 Solubility Soluble in water with gel formation at concentrations higher than 25 pct by weight sulfonate This particular reagent was chosen because of its high active sulfonate content, long-chain length, and solid form. With the exception of the frother and sulfonate, all other chemicals were of reagent grade quality. Conductivity water was used in the experimental work. This water, made by passing distilled water through an ion exchange column, had an average measured conductivity of one micromho. Experiments were conducted at room temperature in a small glass flotation cell with the following procedure: 1) A redetermined amount of water and salt solution (e.g. CaC12 . 2H2O) were combined and the pH adjusted to a given value. 2) A given amount of sulfonate was added so that final solution volume was 130 cc. 3) One drop of Dow 250 frother was added. 4) Five g of quartz were added and the system conditioned for five minutes. 5) The pH of the system was measured (termed flotation pH).

Jan 1, 1963

By Pol Duwez

The effect of the rate of cooling on titanium, zirconium, and thallium has been measured. For titanium-molybdenum alloys, it has been shown that for molybdenum concentration up to 8 pct the rate of cooling has no effect on the transformation temperature, but this temperature decreases progressively with increasing molybdenum content. HIGH strength to weight ratio and remarkable corrosion resistance have often been advanced as the main reasons for the present widespread interest in titanium. For the physical metallurgist, the existence of an allotropic transformation in titanium is another important factor. As in the case of iron, an allotropic transformation may be the fundamental property upon which the response of Ti-base alloys to heat treatment will depend. It is therefore important to study the mechanism of the allotropic transformation in pure titanium and the effect of alloying elements on this transformation. The existence of an allotropic transformation in titanium was reported in 1931 by Schulze.' Later, De Boer et al.' determined the transformation temperature from a discontinuity in the electrical resistivity at 882"42O°C. According to recent work," the transformation temperature is 882.5"&1°C. The crystal structure of the room temperature a phase is hexagonal close-packed and the high temperature fi phase is body-centered cubic.' It is well known that the high temperature /3 phase of pure titanium cannot be retained by quenching to room temperature. Several investigators'-' have shown that in alloys of titanium with metals which are soluble in the @ phase (for example, nickel, molybdenum, chromium, or manganese), the /3 structure can be retained, providing the concentration of the alloying element is greater than a certain critical value. Below this critical value, the alloys transform during quenching, at least partially, into a supersaturated solid solution having the a. structure. This behavior is very similar to that found in some iron binary alloys (Fe-Ni and Fe-Mh, for example) and is an indication that the transformation might occur by means of a shear mechanism analogous to that observed in the austenite-martensite reaction. The present investigation is mostly concerned with the effect of the rate of cooling on the /3 to a transformation. In the first portion of this paper, the results obtained on pure titanium are described. Two other pure metals, zirconium and thallium, were also investigated, because they both have an allotropic transformation like titanium, from an hexagonal close-packed to a body-centered cubic structure. For purposes of comparison, the effect of rate of cooling of the y to a transformation of iron was included in the experimental program. The second part of this paper deals with Ti-Mo alloys. The equilibrium phase boundaries involving the solid state reactions on the Ti-rich side of the system were determined and a study was made of the effect of the rate of cooling on the transformation of the high temperature /3 solid solution. Experimental Procedure Since the main object of the present investigation was the determination of the effect of the rate of cooling on a transformation temperature, the best technique available was that used by Greninger" in his study of the martensite arrest in carbon steels. This method consists of heating a very small sample of metal to which fine thermocouple wires are .attached, cooling the sample rapidly by means of a blast of helium, and recording the temperature vs. time curve on an oscillograph. The apparatus used in the present experiments was very similar to that described by Greninger, except that the heating coil was made of molybdenum wire (0.040 in. in diam) and the assembly was placed in a bell jar under vacuum of 10" mm Hg or less. The specimen con-

Jan 1, 1952

THE Boston meeting, August 29-31, was in many ways one of the pleasantest the Institute has enjoyed in years. Much hard work had been done by the committee, and with excellent results. The program had been purposely restricted with the idea that a smaller, more homogeneous audience formed the best basis for a regional meeting, and the event fully justified the plan. The one hundred and sixty-three members and guests attending were dominantly interested in geophysical methods or their application, and, despite the highly technical character of the papers and the unusual heat, attendance and interest in the technical sessions were well sustained. The exhibition of geophysical instruments and aerial maps added greatly to the meeting.

Jan 1, 1928

By William P. Blake

The literature of the occurrence of quicksilver-ore in the United States does not contain, so far as the writer is aware, any mention of the locality herein described. In the preliminary report * upon the resources of the transPecos region of Texas, von Streerwitz gives a long list of minerals observed as occurring there, but cinnabar is not mentioned. In the second report upon the same region † mention is made of the reported existence of cinnabarite in one of the mountain ranges north of the Sierra Carrizo and the Bofecillos, but the author adds: " In spite of my careful examination of the float, I have not yet found any traces of this metal (quicksilver) up to the present time." Prof. Dumble, also, in * Geol. Surv. of Texas, E. T. Dumble, 1889, p. 225 (first report). † Geol. Surv. of Texas, E. T. Dumble, 1890, p. 713 (second report)

Jan 1, 1896

By R. B. Adamson

The fatigue behauior of irradiated copper single crystals is compared to that of unirradiated crystals. Proton or electron irradiation did not substantially change the fatigue-lzj.e, fatigue-hardening-, or slip-line characteristics, while neutron irradiation did appear to alter these parameters. Postfatigue transrrzission electron microscopy did not reveal any marked structural differazces between irradiated and unirradiated specimens. Possible dislocatio~a mechanisnzs to explain the observed behavior are discussed. A LTHOUGH considerable work has been done to ascertain the effects of radiation on the tensile properties of metals (e.g., Refs. 1 and 2), little work of a basic nature has been reported on the effect of radiation on fatigue properties. Since it is known that fatigue properties are sensitive to conditions at the surface of the metal, it was felt that charged particle irradiation, whose penetration into metals is very small, might influence the fatigue behavior. To investigate this surface effect, irradiations with 10 mev protons and 1.5 mev electrons were conducted. The range of 10 mev protons in copper is about 0.236 mm; which is very small compared with the thickness of the specimens used (1.53 mm). The range of 1.5 mev electrons is somewhat larger, on the order of 0.665 mm.4 On the other hand, irradiation with reactor neutrons produces damage throughout the entire specimen, so neutron irradiation can be used to study the effects of bulk damage. The damage introduced by proton and neutron irradiation is expected to con!ist of large interstitial dislocation loops,5 tiny (<25A) vacancy loops," and/or various vacancy-rich regions called depleted zones.&apos; In the case of electron irradiation at about -173°C the damage is expected to consist of isolated Frenkel pairs.&apos; The effect of radiation damage on the fatigue properties of some engineering materials has been reported,g,10 but the results are not clear-cut or consistent. SPECIMEN PREPARATION Zone-refined copper was used to grow single crystals by a modified Bridgeman technique in an inert helium atmosphere. Using a seeding method, single crystals were obtained as bars about 6 in. long with a cross section of 3 by $ in. Immediately upon removal from the split-mold graphite boat in which they were grown, the very soft bars were embedded in a room-temperature setting plastic. They could then be cut into specimen lengths without introducing any long-range damage. Generally, four or five specimens plus one seed were obtained from each bar. The plastic material was dissolved away in acetone and the orientation of each sample was checked by the standard Laue X-ray

Jan 1, 1968

By B. L. Averbach, M. Cohen

THE PROBLEM MANY hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or subatmospheric cooling.1,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1948

By J. L. Swedlow, M. L. Williams

Two standard techniques for assessing fracture initiatlon in sheet specimens involve measurements of 1) the applied load of fracture which, via linear fuacture mechanics, is used to find the fracture tough/less SC (or SIC), and 2) the velative displacement of fracture of the opposite free faces of the crack (COD). In the past, the quantitative association between these two parameters has not been determined in detail, owing to the mathematical difficulties of computing. the stress and strain fields in the combined elasto-plastic flow near the crack. Using a numerical technique Irecently developed and applied to a 7075- TO&apos; aluminum plane stress specimern, the two parameters noted have been determined and associated quantitatirely. CONSIDERABLE effort has been expended in selecting a rigorous yet practical manner for predicting fracture. The usual point of departure is the modified Griffith criterion in which the mechanism of energy dissipation is reflected in Sc (or SIC), the parameter of "fracture toughness". Sc. differs from the surface energy used by Griffith in perfectly brittle fracture in that it includes the effect of plastic flow in the vicinity of the crack tip. Because plastic flow alters the stresses and displacements predicted from the usual linear theory of fracture mechanics, and because analysis of the nonlinear stress-strain field has presented great obstacles: only estimates of several pertinent quantities have heretofore been available. Nevertheless, two parameters have often been measured experimentally althougll usually just at the point of fracture. These are the energy release rate, or fracture toughness 9, (or I,), and the crack-opening displacement (COD). Such observations have been used in connection with frarture criteria and have been most helpful in design, although their application requires care with respect to changes in geometrical configuration from that upon which the criteria were based 1,2 Recently, however, a reasonably accurate and comprehensive method for analyzing the elasto-plastic flow has been devised and applied to plane stress specimens having arbitrary octahedral stress-strain curves.3 It thus becomes feasible to compute, albeit numerically, both the fracture toughness and crack opening as functions of the applied tensile loading on a typical specimen geometry. As a result, it is now possible to associate these two parameters quantitatively, and thereby to associate data obtained by one method with data pertaining to the other. ELASTIC ANALYSIS For the case of linear elasticity, it is relatively easy to deduce the connection between various measures of crack deformation. Using, for example, the singularity analysis, it is found that the leading term of the Airy stress function x f°r symmetrically applied loadings on an isotropic medium is

Jan 1, 1968

By M. Cohen, B. L. Averbach

Many hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or sub-atmospheric coo1ing.l,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained . steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1949

By B. L. Averbach, M. Cohen

Many hardened steels contain significant quantities of retained austenite even in cases where the carbon and alloy contents are low. In fact austenite has been detected in plain carbon steels containing as little as 0.50 pct carbon.1 Although the austenite may not be visible metallographically it causes definite dilatometric and magnetic effects during subsequent tempering or sub-atmospheric coo1ing.l,2,3 Therefore if the reactions which hardened steels undergo during later treatment are to be understood in quantitative fashion it is most essential to have some kind of absolute analysis for the austenite content in hardened steels. Once this analysis has been made, the austenite transformations can be followed precisely by observing such sensitive phenomena as changes in length or in magnetization but the latter techniques can only determine differences in the amount of retained austenite and usually an independent method must be found to put these differences on an absolute basis. In certain instances the microscope can be used advantageously to determine the percentage of retained austenite. The quenched sample is tempered to "darken" the martensite with minimum decomposition of the coexisting austenite, and then is polished and etched in the usual fashion. If the austenite is clearly distinguishable from the tempered martensite and other possible constituents its volume percentage can be measured quantitatively by point counting or lineal analysis.4 Such microscopic methods, however, depend on the etch being sufficiently discriminatory. In fairly coarse structures, if the austenite content is greater than 10-15 pct, there is probably little smearing of the etch, and point counting or lineal analysis seems to be quite accurate. On the other hand if the structure is difficult to resolve, as in commercially treated fine-grained . steels which contain less than 10-15 pct retained austenite, the latter may be completely obliterated by the dark etching martensite or it may appear disproportionately low as the interfaces between the two constituents become difficult to distinguish. Another disadvantage of the metallographic method is the need for "darkening" the martensite and hence the determination cannot be applied to as-hardened steels. The required tempering may spoil the structure for other tests or may even remove a small part of the austenite which is being sought. X ray methods hold more promise for covering the range under 15 pct retained austenite. Although austenite and martensite in the hardened steel may be chemically identical, their crystal structures and lattice parameters differ, and the intensity of an austenite diffraction line is some continuous function of the percentage of re-

Jan 1, 1949