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By R. B. Moore

IN 1914 the writer and K. L. Kithil announced, through Bulletin 70 of the U. S. Bureau of Mines, that the United States possessed the largest deposits of radium-bearing ore in the world. At that time a small amount of this ore was being, mined and shipped to Europe, and one company was preparing to start the extraction and recovery of radium in this country. The importance of our carnotite deposits in Colorado and Utah up to the publication of Bulletin 70 was not, however, appreciated by anyone.. At that time the total supply of radium in the world was about 10 g., calculated as element, and when I estimated that at least 100 g. of radium element, and probably 200 g., would be produced from our carnotite deposits, there was a considerable amount of skepticism. Up to the present time almost 200 g. have come from this ore, and some writers put this figure a little higher. Uranium ore was found in the copper mines of the Union Minière du Haut-Katanga at Luiswishi in 1913. Later discoveries were made at other places in the Belgian Congo, and the Katanga ores now represent the largest supply of radium-bearing ore in the world. In fact, since 1923 very little radium has been produced from American carnotite, the Katanga ore representing practically the world's supply of. radium since that date.

Jan 1, 1930

By K. A. Gschneidner, O. D. McMasters, T. J. O’Keefe

X-ray diffraction, differential thermal, ad rnetallo-graphic methods were used to establish the Dy-Pb Phase diagram. Lead additions lower the 1377°C transformation temperature of dysprosium to 1360°C leading to an inverted peritectic reaction. The 327°C melting point of lead is lowered by dysprosium additions to about 326°C yielding a eutectic reaction. A second eutectic reaction occurs at 13.3 at. pct Pb and 1200°C. The dysprosium-richest intermetallic compound DysPb3 melts congruently at 1695°C and crystallizes in the hexagonal Mn5Si3 (D8,) type structure. The peritectic decomposition temperatures for the remaining compounds are Dy5Pb, at 1555C, DyPb2 at 955C, and DyPb3 at 880°C. A fifth compound near the DyPb stoichiometry exists over a 310°C temperature range decomposing at 1130°C by means of an inverted peritectic reaction and melting incongruently at 1440°C. The crystal structures of the compounds are discussed. A systematic study of the rare earth-lead alloy systems is underway in an effort to supply information concerning the alloying behavior of the rare earth metals. The Dy-Pb phase diagram is the fourth system to be investigated in this study. The Yb-Pb,1 Y-Pb,2 and Eu-Pb 3 diagrams have been published recently. Utilization of the rare earth series of metals as a research tool in this manner should yield a better understanding of alloy formation. EXPERIMENTAL PROCEDURE Materials. The lead used in this investigation was obtained from Cominco Products, Inc., and was specified to be 99.99 pct pure. The dysprosium was prepared in this Laboratory by the calcium reduction of the fluoride followed by distillation of the dysprosium. The major impurities in the dysprosium in ppm are: A1 (<40), Ca (400), Er (<50), Gd (<200), Ho (<200), Mg (<50), Si (30), Ta (400), Tb (<100), Y (<50), 0 (651, H (15), N (not detected), F (430), C (35). Alloy Preparation. Most of the alloys were prepared by melting weighed amounts of dysprosium and lead in sealed tantalum crucibles. The tantalum crucibles were sealed by are-welding in a He-Ar atmosphere welding chamber. Thus the alloys are in contact with He-Ar at about 1 atm pressure. Homogenization was achieved by holding them in the liquid state for about 1 hr, cooling, inverting the crucibles, remelting, and repeating the process at least twice. Since these alloys were prepared in sealed tantalum crucibles, chemical analysis for final composition was thought to be unnecessary. No detectable reaction of these alloys with the tantalum crucible was observed by metallographic examination. Metallographic evidence was also used to confirm the homogeneity of some of the alloys prepared in this manner. The compositions of a few alloys, which were prepared by nonconsum-able are-melting, were corrected for the small weight losses involved by assuming that the weight loss is due to vaporization of lead. The specimens obtained from the alloy samples were prepared under a dry-argon atmosphere because they were rapidly attacked by air and moisture. Thermal Analysis. Differential thermal analysis methods were used to determine the liquidus curves and reaction horizontals of the system. Both Pt vs Pt + 13 pct Rh and W + 5 pct Re vs W + 26 pct Re thermocouples were used to measure the temperature. An X- Y recorder was used to record the specimen temperature and differential electromotive force between the specimen and molybdenum standard. The arrest temperatures were measured potentiometri-cally. The accuracy limits (* values) associated with the reaction temperatures obtained by this method were estimated on the basis of both the reproducibility of the particular temperature value and the accuracy of the thermocouple at a given temperature. Liquidus temperatures were obtained from cooling arrest data while both heating and cooling arrest data were used to establish the horizontals of the diagram. Heat treatments during the thermal analyses of the alloys between 40 and 70 at. pct Pb were necessary in order to approach equilibrium conditions. The samples were held at temperatures between the various peritectic horizontals for l to 2 hr before the thermal analyses were continued. The entire range of compositions was investigated at the expense of a minimum amount of materials by adding appropriate amounts of lead to master alloys. More than sixty alloys were analyzed by this differential thermal method and for each alloy the results given herein are taken from two or three heating and cooling cycles. X-Ray and Metallographic Methods. Slice specimens for metallography and powder specimens for X-ray diffraction were prepared from rod-shaped samples which had been melted in sealed 0.62 5-cm-diam tantalum crucibles. The specimens were heat-treated in sealed tantalum crucibles which were protected by sealing them in argon-filled quartz ampules. Quenching was accomplished by breaking the ampules in ice water after heat treatment. X-ray powder specimens were sealed in 0.3-mm-diam glass capillaries under a dry-argon atmosphere. Copper, iron, and chromium radiation were used to obtain the powder patterns for these alloys. More than 150 powder patterns were obtained for specimens of various compositions and heat treatments. Included in these were several patterns for specimens which had purposely been oxidized. Patterns from specimens which had been accidentally exposed

Jan 1, 1969

By C. W. Chen

MARTENSITE transformations in ß Cu-Al alloys have been studied by Greninger1 and other investigators. According to Greninger, the parent phase ß1 an ordered body-centered-cubic structure obtained from ß phase by suppressing the eutectoid decomposition, transforms into an ordered hexagonal-close-packed phase in composition containing 12.9 to 14.7 pct Al. The M, temperature decreases with increasing aluminum content; for the alloy containing 14.5 pct Al, for example, the ß1??&apos; transformation occurs below room temperature. More recently, Kurjumow2 studied the transformation in ß Cu-Al alloys with the addition of nickel. His report stimulated new interest in the subject due to the observation of completely reversible transformation without hysteresis in the transformation temperature ranging from 10° to —10°C. In the present paper some characteristics are described of the transformation of Cu-Al-Ni alloys that were partly studied by Kurjumow. Experimental Procedure High purity copper (99.999 pct) and aluminum (99.99 pct) and electrolytic nickel were used in the preparation, by the Bridgman technique, of single crystal specimens which contained aluminum and nickel of 14.5 and 0.5 to 3.0 pct, respectively. Polished surfaces were prepared mechanically. Specimens were then chemically etched to remove distorted material, homogenized at 1000°C for several hours, and quenched drastically to room temperature in a 10 pct NaOH bath to produce the parent phase ß1. The transformation was studied under a microscope and, in some cases, recorded by means of motion pictures. A device similar to that designed by Greninger and Mooradian3 was used to cool and reheat the specimens. Results and Discussion When the specimens were cooled below room temperature, the ß1 to ?&apos; transformation began at 10°C with the appearance of ?&apos; crystals In relief, Fig. la. As the specimen temperature dropped further, the transformation continued, either by the growth of the ?&apos; crystals, with the ß1 — ?&apos; interface moving into the ß1 phase, Fig. 1c and 1d, or by the formation of new ?&apos; crystals, Fig. 1b. As a consequence of the former process, banded structure is observed as a common feature of the low temperature phase. According to the theory of the formation of martensite by Wechsler, Lieberman, and Read,&apos; the bands of ?&apos; phase are probably twin-related, as is the case in the diffusionless phase change of In-T1 alloys,5 but this was not revealed by X-ray tech- niques. New ?&apos; crystals, in needle form, often emerged suddenly across the ?&apos; bands during the transformation. These acicular crystals then grew, both in length and in width, see Fig. 2a through 2d. The transformation on cooling is completed at about -35°C. Upon heating, the reverse transformation started at —10° C, in a manner nearly opposite to the transformation on cooling, and completed at 35°C. There was no noticeable change in the transformation temperature when the nickel content was varied within the limits previously mentioned. Through control of the specimen temperature, the transformation can be started, stopped, or reversed at will. This phenomenon has frequently been observed in the martensite transformation of many nonferrous alloy systems. Other systems are Au-Cd6 and In-Tl.5 ow-- ever, in the latter systems, the transformation is accomplished by single interface motion if the specimen composition is homogeneous and the temperature gradient in the specimen is uniform and sharp, whereas in the Cu-Al-Ni specimens, only multiple interface transformation is observed. The speed of the interface motion appears to be a functionof the rate of temperature change and the temperature gradient across the specimen length. In one case, in which the temperature increased at the rate of 10°C per min and there was no temperature gradient along the specimen axis, the speed of the disappearance of a ?&apos; plate was determined, by the study of the motion pictures made, to be 26 µ per sec. Quench markings were observed on the polished surfaces of specimens. The markings were grouped into one or more sets of different orientations, and were parallel in each set. The ?&apos; plates formed in subsequent transformation were parallel to the markings, indicating that the ?&apos; plates and the quench markings had the same geometric relation-ship to the ß1 matrix. The quench markings on two intersecting surfaces of a specimen were therefore used in the determination of the habit plane of transformation, by the trace method suggested by Barrett.&apos; Results obtained from five sets of markings in three specimens indicate that the habit plane is an irrational plane about 2" from one of the {221} planes. This is very close to the habit plane (3" from 221 planes) of ß Cu-Al alloys containing more than 13.0 pct Al.1 The martensite transformation of Cu-Al-Ni alloys is reproducible. No sluggishness was found between consecutive transformation cycles, although a slight difference in the distribution pattern of the ?&apos; plates was observed, compare Figs. Id and 2d. The transformation can be strain-induced. This characteristic has been tested by a simple method. When a specimen was elastically strained slowly in a vise, ?&apos; plates were gradually produced in the same fashion as during transformation on cooling, This test was done at room temperature, and thus above the M,

Jan 1, 1958

By T. P. Floridis, J. H. Young

The need for close control of the phosphorus content of steels has led to numerous investigations on the equilibria of the dephosphorization reactions. Winkler and chipman1 have established the general conditions for effective dephosphorization of steel. They are high slag basicity, high oxygen potential, and low temperature. Other investigators have made additional contributions to the understanding of the dephosphorization process. The current status of the understanding of the dephosphorization of steel is concisely presented by Bodsworth2 and by Ward. The investigation reported in this communication was undertaken with the purpose of establishing the effect of additions of barium oxide and calcium fluoride on the dephosphorizing capacity of slags. Esin and Gel&apos;d4 proposed that barium oxide, being more basic than calcium oxide, should cause an increase in dephosphorizing capacity when added to steelmaking slags, or when used as a substitute for calcium oxide. Derge5 has also proposed that in the future conventional slags might be replaced by BaO-Al2O3 slags. There is, however, no experimental evidence confirming the effect of barium oxide on the dephosphorizing capacity of slags. The effect of calcium fluoride on the dephosphorization of steel is not clearly understood. It is generally recognized that additions of calcium fluoride are beneficial. It is not clear, however, whether calcium fluoride affects the equilibrium of the dephosphorization reaction, or whether it simply causes an increase in the fluidity of the slag and, consequently, faster approach to equilibrium. The experimental procedure consisted in equilibrating synthetic molten slags with liquid copper at 1550°C under a gas stream containing argon, hydrogen, and water vapor. In all experiments the argon to hydrogen ratio was approximately 4:1, and the hydrogen to water ratio was 5.42:l. Molybdenum crucibles were used as containers for the slag and metal. Under the above-described conditions of temperature and composition of atmosphere, there was no observable attack of the crucibles by the metal, slag, or atmosphere. Copper was used instead of iron, because iron attacks molybdenum. The equilibration was made in a tubular furnace consisting of a recrystallized alumina tube. The alumina tube was heated by electrical resistance. A Pt-40 pct Rh wire winding was used for most runs. A silicon carbide tubular resistor was also used for some runs. Temperatures were measured with a Pt-Pt-Rh (10 pct Rh) thermocouple and kept constant within ±5°C. Equilibrium was approached from both sides, i.e., by adding the phosphorus either as oxide in the slag or as a phosphorus-rich alloy of phosphorus and copper. The holding time at the equilibrium temperature was 6 hr. At the end of each run the crucibles were rapidly cooled and removed from the furnace. The slag and metal were separated and analyzed. The experimental results are shown in Table I. The phosphorus content of the slag is expressed both as percent phosphorus pentoxide and as percent phosphorus. The basicity ratio is computed by dividing the number of moles of basic oxides-oxides of barium, calcium, magnesium, and sodium—by the number of moles of acidic oxides- oxides of aluminum, phosphorus, and silicon. Calcium fluoride is not included in the computation of the basicity ratio; i.e., calcium fluoride is assumed to be neither basic nor acidic. The distribution ratio of phosphorus—percentage of phosphorus in the slag divided by the percentage of phosphorus in the metal- is plotted in Fig. 1 against the basicity ratio. The results indicate that slags containing barium oxide have greater dephosphorizing capacity than slags containing calcium oxide. The high dephosphorizing capacity of slags containing sodium oxide and the low dephosphorizing capacity of magnesia-containing slags which already have been reported in the literature2 are confirmed by the results of this investigation. It appears that calcium fluoride has a beneficial effect on the distribution of phosphorus between slag and metal in acid slags only. Although the obtained distribution ratios between the phosphorus contents of slag and copper are not directly applicable to the dephosphorization of steel, they are sufficient for evaluating the effect of slag additions on the dephosphorizing capacity of slags in general. An increase in the ratio of distribution of phosphorus between slag and metal indicates lowering

Jan 1, 1968

By Amitava Gangulee, Michael B. Bever

The effects of changes in short-range order on some thermodynamic, electrical, and mechanical properties of the silver-rich Ag-Mg solid solutions have been investigated. The heats of formation at 273°K of several alloys after two different thermal treatments were measured by liquid metal solution calorimetry. Their electrical resistivities were measured at temperatures ranging from 4° to 773°k. Tensile tests and microhard-ness measurements were carried out at room temperature. An effective short-range order parameter and an effective interaction energy are defined. The measured changes in the properties are interpreted in terms of these quantities. The short-range order parameter is evaluated separately from calorimetric and X-ray data; the results are in fair agreement. Increasing short-range order lowers the resistivity of- the silver-rich Ag-Mg alloys but does not measurably affect the temperature dependence of the resistivity. Short-range order increases the yield stress of these alloys but does not affect the ultimate tensile stress. Changes in the effective short-range order parameter independently obtained from measured changes in the heat of formation, resistivity, and yield stress are in fair agreement. THE thermodynamic, electrical, and mechanical properties of solid solutions have been investigated as functions of composition in many alloy systems, but the effects of configurational variables, such as short-range or long-range order, on these properties have received much less attention. The silver-rich terminal solid solutions in the system Ag-Mg are well-suited for the investigation of the effects of order. The composition of these solid solutions may be varied over a range extending to about 27 at. pct Mg.1 Appreciable short-range ordering is likely to take place in these solid solutions; the degree of short-range order can be varied by suitable thermal treatments. In the alloy Ag3Mg long-range ordering is possible.2 This paper is primarily concerned with the effects of short-range order on some thermodynamic, electrical, and mechanical properties of the silver-rich Ag-Mg solid solutions.-The long-range ordering transition is the subject of a concurrent paper.3 1) EXPERIMENTAL PROCEDURES 1.1) Preparation of Specimens. Ag-Mg alloys containing up to 26.4 at. pct Mg were prepared by melting 99.99 pct pure Ag (obtained from Baker Chemical Co.) and 99.99 pct pure Mg (obtained from Eastman Chemi- cal Co. or Johnson Mathey & Co.) in graphite crucibles under molten potassium chloride. The solidified ingots were homogenized at 823°K for 10 days in evacuated Vycor capsules. A surface layer 6 in. thick was then removed by machining. The ingots were swaged into rods (2.5 mm diam) which were drawn into wires (1.0 mm diam). The wire specimens were annealed at 773°K for 24 hr in vacuum and either quenched into iced brine or cooled to room temperature over a period of 15 days. The average grain diameter (obtained from the linear intercepts) was about 0.1 mm. 1.2) Calorimetry. The heats of formation of the alloys were measured in a tin solution calorimeter as the difference in the heat effects of alternate additions of samples of an alloy and the mechanical mixture of its component elements. The additions were made from 273°K to the bath at about 623°K. The heat effect on addition of the alloys was approximately 5.5 kcal per g-atom. The procedure and the method of calculation have been described.4 1.3) X-Ray Diffraction. Short-range order parameters of two alloys were calculated from diffuse scattering intensities,&apos; obtained with briquettes made from powdered specimens. A GE-XRD 5 diffractometer with filtered copper radiation was used. The absolute intensities were based on calibrations with paraffin. 1.4) Resistivity Measurements. The electrical resistivities were measured by a potentiometric method in which the potential drop across the specimen was compared with that across a standard resistance. For measurements below room temperature, the specimens were immersed in liquid helium (4°K), liquid nitrogen (78°K), dry ice and trichloroethylene (195°K), or ice and water (273°K). Measurements above room temperature were made on specimens held in a furnace under helium or in vacuum. The resistivity measurements were reproducible to ±0.5 pct or less. 1.5) Mechanical Tests. Tensile tests were carried out at room temperature with a Tinius-Olsen XY elec-tromatic Universal testing machine. Wire specimens (nominal gage length 1 in.) were used. The strain was measured with an extensometer, which had a sensitivity of 4 x 10-5; the strain rate was 10-3 min-1. Microhard-ness measurements were made with a square-pyramid indenter and a 100-g load.11 2) RESULTS AND DISCUSSION All measurements were made on two parallel sets of specimens, one quenched and the other slowly cooled from 773°K. The reported values are the averages of at least three measurements. In the range of magnesium concentration substantially below that of the composition Ag3Mg, quenched alloys have less short-range order than slowly cooled alloys. As the magnesium concentration approaches that of Ag3Mg, slow cooling develops long-range order. Quenching sup-

Jan 1, 1969

By M. E. Nicholson

IN order to understand how alloying elements influence hardenability through their effect on the rate of pearlite nucleation, it is advantageous to use a model to describe the mechanism of pearlite nucleation. The model which currently is most widely accepted has been described, among others, by Mehl et a1.1-3 and by Hultgren.4 Pearlite nucleation involves both the nucleation of ferrite and cementite.4,5 Cementite nucleates first, followed rapidly by the formation of ferrite at the cementite-austenite interface. The conclusion that cementite is the first phase to form during the process of pearlite nucleation is based, in part, on the observation that proeutectoid cementite is continuous with pearlitic cementite and that the orientation relationship between proeutectoid ferrite and austenite is not the same as that between pearlitic: ferrite and austenite.6,7 Recently, Smith" has suggested that the dissimilarity in orientation relations, instead of indicating the pearlite nucleation sequence, may be evidence that pearlite is nucleated in a grain adjacent to the one in which it is growing. Also, Modin9 as shown that ferrite in pearlite may be continuous with proeutectoid ferrite, and he has also shown that proeutectoid cementite is not always continuous with pearlitic cementite. Thus, it appears that the conclusion that cementite always initiates the pearlite nucleation process is open to serious question. The author believes that important evidence on the question of how pearlite is nucleated exists in the rates of pearlite nucleation of different steels. The following is a review and analysis of this evidence. If the nucleation of pearlite is considered as requiring the nucleation of both cementite and ferrite, then the time to nucleate pearlite consists of the time to nucleate the first phase to form, plus the time to nucleate the second phase at the advancing interface of the first. In the absence of a visible proeutectoid constituent, the second phase must be nucleated very soon after the first. If cementite is the first phase to form, the time to nucleate pearlite will be approximately equal to the time to nucleate cementite. Since the nucleation rate of a phase precipitating from solid solution increases with the degree of supersaturation, it would be expected that the rate of pearlite nucleation should increase with increasing carbon content. Digges10 has shown that the reverse is true. He determined the quenching velocity necessary to suppress pearlite formation in plain carbon steels as a function of carbon content. His results showed that pearlite formation could be suppressed with lower quenching velocities as the carbon content of the steel increased, at least up to 1.25 pct C, the highest carbon content he studied. Thus, it would appear that in plain carbon steels, in which pearlite is formed at the nose of the C-curve, the hypothesis that cementite nucleates first does not predict the observed change in nucleation rate. The validity of the current model of pearlite nucleation may be tested also by determining whether or not it predicts the relative influence of alloying elements on the nucleation of ferrite and pearlite. This can be done best by a consideration of the changes in a TTT diagram produced by a change in alloy content. Hultgren4 suggests that the general shape of a TTT diagram, as well as what constituents will form from austenite, can be predicted from the relative nucleation rates of ferrite and cementite. This is illustrated in the schematic diagram for a hypoeutectoid low alloy or plain carbon steel shown in Fig. 1. The C-curves for ferrite nucleation and for cementite nucleation are represented by AA&apos; and CC&apos;, respectively, and the start of ferrite formation and pearlite formation are represented by AX and BC respectively. Adopting the current theory, Hultgren suggests that between X and C pearlite forms directly from austenite, since cementite nucleates more rapidly than ferrite in this temperature range. Conversely, at temperatures above Tx, ferrite is the first to nucleate and proeutectoid ferrite is formed. As proeutectoid ferrite grows, austenite becomes enriched in carbon, so that the cementite nucleation rate is increased. As a result, pearlite nucleation above temperature Tx is accelerated, beginning along BX and not along C&apos;X. Temperature T, represents the temperature at which the nucleation rate of ferrite equals that of cementite. By extending this type of reasoning, using the principle that the nucleation rates of ferrite and cementite increase with increasing supersaturation of austenite with respect to these phases, it should be possible to predict the influence of carbon content on the shape of the isothermal transformation diagram. For a slightly higher carbon content than that of Fig. 1, the curve A A&apos; representing the beginning of ferrite formation should be moved to the right, whereas the curve CC&apos; representing the beginning of cementite formation should be moved to the left. For hypoeutectoid plain carbon and low alloy steels, this prediction does not agree with existing data. Instead, as the carbon content is increased, the pearlite curve moves to the right along with the ferrite curve and always appears to join it tangen-tially. This behavior is demonstrated by the TTT-curves of a series of manganese steels" shown in Fig. 2. In this figure, TTT-curves for four alloys containing 0.20, 0.40, 0.60, and 1.20 pct C are superimposed. The curves are identified in the figure legend. The letters F, P, and C are used to indicate the start of ferrite, pearlite, and cementite formation, respectively. In these steels, it appears that the nucleation of ferrite and of pearlite are related processes. As a result of the foregoing evidence, it is proposed that the formation of pearlite be considered as follows: Either ferrite or cementite may nucleate from austenite, then grow until the remaining phase is nucleated at the interface between the growing phase and austenite. (The growth of the first phase to form is assumed to be accompanied by a composition change in the adjoining austenite.) According to this model there are two modes of nucleating pearlite: 1—where cementite initiates the succession

Jan 1, 1955

By J. Handin, S. J. Bauer, M. Friedman

Deformation mechanisms in room-dry and water-saturated specimens of Charcoal Granodiorite, shortened at 10-4s-1, at effective pressures (Pe) to 100 MPa and temperatures to partial melting (?1050°C) are documented with a view toward providing criteria to recognize and characterize the deformation for geological and engineering applications. Above 800°C strength decreases dramatically at effective pressures ? 50 MPa and water-weakening reduces strength an additional 30 to 40% at Pe = 100 MPa. Strains at failure are only 0.1 to 2.2 percent with macroscopic ductility (within this range) increasing as the effective pressures are increased and in wet versus dry tests. Shattering (multiple faulting) gives way to faulting along a single zone to failure without macroscopic faulting as ductility increases. Microscopically, cataclasis (extension microfracturing and thermal cracking with rigid-body motions) predominates at all conditions. Dislocation gliding contributes little to the strain. Precursive extension microfractures coalesce to produce the throughgoing faults with gouge zones exhibiting possible Riedel shears. Incipient melting, particularly in wet tests, produces a distinctive texture along feldspar grain boundaries that suggests a grain-boundary-softening effect contributes to the weakening. In addition, it is demonstrated that the presence of water does not lead to more microfractures, but to a reduction in the stresses required to initiate and propagate them.

Jan 1, 1982

By D. A. Elgillani, M. C. Fuerstenau

With either sulfonate or oleate as collector, quartz responds to flotation with moderate additions of calcium only at moderately high pH, where some portion of the activator has hydrolyzed to caOH+ . Calculations of the concentrations of various ionic and precipitated species of calcium and collectors suggest that the products of [(CaOH+) (RSO3)] and [(CaOH+)(01-)] determine whether flotation is obtained under specific conditions. Ion products on the order of 10-12 were calculated for both the sulfonate and oleate systems. The activating effect of calcium ion in nonmetallic flotation systems is of considerable interest because of the normal presence of calcium in natural water. As a result, this phenomenon has received quite some attention in the past. Kraeber and Boppel1 showed that quartz could be activated by calcium above pH 10 with sulfonate as collector. The feasibility of selectively separating quartz from hematite with calcium activation at relatively high pH was demonstrated by Clemmer, Clemmons, Rampacek, Williams, and stacy.2 Cooke and Digre3 showed with a bubble pick-up method that the minimum quantity of calcium ion required as activator for complete pick-up of particles occurs at pH 11.5 for an addition of 20 mg per liter sodium oleate. They also showed that larger additions of calcium (10-fold increase per unit decrease of pH) must be added for complete bubble pick-up as the pH is reduced. Schuhmann and Prakash,4 using a vacuum flotation technique, found that quartz could be floated with moderate additions of calcium chloride and oleic acid at neutral pH, providing the metal ion was present in stoichiometric excess over the quantity needed to form the normal soap with oleic acid. They also reported that calcium will function as an activator only in basic media. More recently, Eigeles and volova5 have shown that essentially complete flotation of quartz is obtained with 6 x 10-4 mole per liter calcium chloride and 1.7 x 10-5 mole per liter sodium oleate at pH 11.6. while no flotation is obtained at about pH 10.9 and below. The importance of adsorption of activator and collector at the air-liquid interface is also demonstrated in these systems. The important role that metal ion hydrolysis assumes in quartz activation systems was also demonstrated recently.6-8 A detailed investigation of metal activation in sulfonate flotation of quartz was undertaken in one system7 and yielded a number of interesting and important observations. Quantification of the data of this system7 to the extent desired was not possible, though, because certain species could neither be ignored nor accounted for accurately. These difficulties can be circumvented when calcium is involved as activator. This detailed analysis was undertaken to obtain a more quantitative explanation of calcium and metal ion activation in quartz flotation. EXPERIMENTAL MATERIALS AND METHODS Sodium alkyl aryl sulfonate9 mol wt 450, and pure potassium oleate were used as collectors. All other reagents used were reagent grade in quality, i.e., n-amyl alcohol as frother, KOH for pH adjustment, and calcium chloride. Conductivity water, made by passing distilled water through an ion exchange column, was used in the investigation. Quartz was prepared by leaching the sized sample (48 x 150 mesh) with HC1 until no iron could be detected in the leach liquor. The experimental equipment and procedure were the same as that described previously.6,10 EXPERIMENTAL RESULTS As the presence of precipitates was noted in all of the systems to which ca++ and collector were added, experiments were undertaken to determine the solubility products of calcium sulfonate and calcium oleate using a nephelometer. With this technique, collector is titrated into a known solution, which in this case was 5 x 10-5 mole per liter CaCl2 at pH 5.5. Upon precipitation of the calcium-col lector salt, e.g., calcium oleate, light is scattered and detected

Jan 1, 1967

By Walter E. Mingramm

The construction of rescue-apparatus on the principle of furnishing the wearer with air from a tank containing it under high pressure was given up by inventors about 20 years ago. Such an apparatus must supply continuously, for not less than two hours, nearly 50 liters of pure air at atmospheric pressure per minute, or a total of 6,000 liters (210 cu. ft.); and a cylinder containing this amount of air, under 2,000 lb. pressure, would weigh 180 lb. or more. Although apparatus has been recently constructed utilizing liquefied air, no experiments have been made to determine the regularity with which the air is generated, nor has its chemical composition been determined. Moreover, the chief objection to such an apparatus is the difficulty and cost of keeping on hand a large quantity of liquefied air. The system of furnishing the wearer of a breathing-apparatus with air which has given the best results, and has found

Jan 1, 1909

[Officers and Committees of the Society of Petroleum Engineers 1 Officers and Staff of the AIME 2 Officers of AIME Sections Predominantly Petroleum in Membership 3 Society of Petroleum Engineers By-Laws 5 Alphabetical List of Members 9 Geographical List of Members 110 List of Members by Companies 127 Advertisers&apos; Index 137]

Jan 1, 1961

By Robert D. Pehlke, Robert G. Blossey

The solubility of nitrogen in liquid iron alloys containing chromium and nickel has been measured in the temperature range 1550° to 1700°C at nitrogen pressures to 1 ah. The solubility surface has been determined for the iron corner of the liquid metallic ternary to 21 pct Cr and 11 pct Ni. The nitrogen solubility increases markedly with increasing chromium concentration. The effect of nickel and the bilinear effect of chromium and nickel are less pronounced. The solubility data are presented, and interaction coefficients are determined from them. THE solubility of nitrogen has been determined in liquid Fe-Cr-Ni binary and ternary alloys. The work in this laboratory has been directed toward a relatively limited composition range, that is to 20 pct Cr and 10 pct Ni. Humbert and Elliott1 studied the entire ternary; but their investigation was, understandably, limited in this region. The region considered here is of the greatest commercial interest. Knowledge of the solubility in these alloys is important for accurate control of the nitrogen levels in various stainless alloys. Also, the solution behavior is of theoretical interest because of its deviations from ideal dilute solution rules at moderate chromium levels. The nitrogen solubility in the iron binaries has been reviewed extensively by Pehlke and Elliott,2 particularly with respect to the dilute solution interaction parameters. More recently, there has been some interest in accurate descriptions of solute interactions in non-dilute solutions. Most of the investigators have adopted a power series fit to accommodate any non-linearity of the data. The Taylor series expansion is adopted here, using Scimar&apos;s notation3 for the interaction coefficients. An alternate method, instead of requiring higher-order expansions, is to restrict the range of the composition variables. A planar solubility surface may be fitted to the data and its range of validity determined. If the axes of this region are displaced from the usual origin—the solvent pure iron—a change in standard state is required for the solution reaction. The standard state for nitrogen in iron is commonly taken as the hypothetical 1 wt pct solution of nitrogen in pure iron. However, for an alloy of appreciable solute concentration, e.g., 13 pct Cr or 18 pct Cr-8 pct Ni, it would be convenient to use that base composition as the solvent and redefine the solution reaction in terms of that particular solvent. A brief derivation follows, defining the free energy of solution in an alloy solvent. Following the treatment by Darken and Gurry,4 for a nitrogen pressure of 1 atm: N (pure component) = N (1 wt pct in solvent) [l] AG=RT ln100/ wt pct N in solvent Taking this reaction for an alloy solvent less the same reaction for a pure iron solvent the relation between nitrogen in pure iron and an alloy is: N (1 wt pct in Fe) = N (1 wt pct in alloy) [2] AG = RT[ln(wt pct N in Fe) - ln(wt pct N in alloy)] Adding ½ N2(g) to each side and subtracting N (1 wt pct in Fe) = ½N2(g), the total free energy for solution of nitrogen in an alloy, referred to nitrogen at infinite dilution, is given by: ½N2 (g) = N (wt pct in alloy) L3] AG°3 = -RT ln(wt pct N in alloy )PN =1 atm Interaction parameters have their usual meaning when ?G°/RT is differentiated with respect to the proper composition coordinate. If desired, the dilute alloy treatment may be utilized with the new reference composition as the solvent. This has. indeed, been done by Small and pehlke5 in their analysis of nitrogen solution in liquid 18-8 alloys. The Sieverts&apos; technique was used for this investigation on an apparatus previously described.6 The materials used in this program were: Ferrovac-E iron, 99.95 pct, Crucible Steel Co.: chromium, 99.95 pct, Union Carbide Metals Co.; nickel, 99.9 pct, International Nickel Co. Recrystallized alumina crucibles were used with no evidence of crucible-melt reaction. RESULTS The locations of the ternary compositions investigated are shown in Fig. 1. The data of Turnock7 and small8 have been alluded to in other papers.5&apos;6 There were no departures from Sieverts&apos; law observed for

Jan 1, 1969

By P. R. Sahm, T. V. Pruss

Mixtures of fine GaSb and Gds as well as preal-loyed GaSbl,As, powders were hot-pressed at 690°C and 25,000 psi. Dense alloys with compositional gradients of less than 5 pct were obtained from mixtures containing about 20 mol pct GaAs. For x < 0.2, there were increasing compositional gradients, and for x > 0.2 a GaAs-rich second phase appeared in the microstructure. Densification as well as alloying wlechanisms were enhanced by dissociation and, possibly, oxidation reactions of the powders. Densification of coarse prealloyed material, however, primarily depended on plastic flow phenomena and required temperatures just below solidus and pressures of 50,000 psi Thermal and electrical properties were measured. Although in no case was the figure of merit of melt growth materials approached closer than to within 25 pct, the better overall thermoelectric properties were found in coarse-grained, prealloyed materzals which had been compacted to near theoretical density and where grain regrowth had been induced. Similar results are believed to hold in other binary 111-V compound systems if processed under similar conditions. The densification of unalloyed GaSb powders during pressure sintering (= hot pressing) was shown to depend strongly on powder particle size.&apos; Fine powders displayed a "liquid skin effect" that enhanced compaction through the presence of liquid gallium, whereas coarse powders compacted predominantly by plastic flow. The liquid skin effect, in particular, appeared attractive for alloying GaSb with other, higher-melting, III-V compounds, and to densify these in a one-step operation. This is of special interest in the case of III-V compound alloys as the conventional techniques of melt growth or long-time annealing of powders2 is very time-consuming, especially in cases where a large separation of liquidus and solidus can be expected, such as in GaSbl alloys.2 It was felt that experimentation with this system, a particularly unfavorable example, would allow us to extrapolate to several other, more favorable, cases. Uses of Ill-V alloys are most evident in thermoelectric energy conversion devices.&apos; For this reason certain thermal and electrical properties were measured and compared to those of melt-grown material and to hot-pressed prealloyed powders. EXPERIMENTAL PROCEDURE The pressure-sintering apparatus has been previously described.&apos; Using this equipment, both powder mixtures of GaSb with GaAs and prealloyed GaSbl-,As, powders were hot-pressed. The mixtures were pre- pared from cast GaSb and melt-grown GaAs. The prealloyed material was obtained from stoichiometric melts, initially heated to 1200°C in an evacuated quartz ampoule, and then annealed in the solidus-liquidus interval. A typical annealing cycle consisted of a heating to 850°C (1 hr), extended successive annealing at 750°C (65 hr), and slow stepwise cooling (25 hr) to below solidus. The GaSb, GaAs, and GaSbl-,As, materials were ground in a vibrational mill to particle sizes below 500 . A jet mill reduced these further where necessary. Fine powders were analyzed by Coulter counting for their size distribution. Average sizes by volume were calculated from the data. Mixing of the powders, where necessary, was carried out through rapid vibrational motion. The hot-pressing operation consisted of a degassing period of 15 hr, in most cases at 690°C, followed by several hours of compression, normally 25,000 psi at 690°C for powder mixtures and 50,000 psi at 710°C for prealloyed powders. The chemical compositions were confirmed by X-ray fluorescence analysis. To estimate the degree of solid solution achieved, lattice parameters were determined and interpreted according to Vegard&apos;s rule. In addition, optical microscopy helped to correlate the relative amounts of the phases present as well as the degree of porosity to the measured density. To reveal grain boundaries, polished surfaces were etched4 with H 2 O:H 2 O:HCl = 2:l:l. In several cases microscopic concentration gradients were monitored by electron-probe analysis-. RESULTS AND DISCUSSION Alloying and Densification of GaSn-GaAs Powder Mixtures. After preliminary experiments showed that no appreciable alloying took place in mixtures of coarse GaSb (22.5 p) and fine GaAs (2.5 p) powders, hot pressing of mixtures was confined to fine powders only (3.1 and 2.4 p, respectively). Alloying and densification apparently occurred simultaneously with a grain regrowth mechanism which depended on the presence of a liquid phase.&apos; The liquid phase was provided by the dissociation of GasbS particles into liquid gallium and antimony above 555°C. This not only enhanced grain regrowth and densification, as in unalloyed GaSb,&apos; but took on additional importance here for the alloying process. Alloying was speeded up by the resulting liquid-solid interaction as compared to the very slow solid-state diffusion process normally expected at these temperatures.&apos; The results obtained with a series of powder mixtures have been compiled in Table I. It is seen that for mixtures with less than 20 mol pct GaAs, x < 0.2, the compositional ranges increased and for x > 0.2 a GaAs-rich phase and free antimony appeared in the microstructure in sizable amounts.

Jan 1, 1968

By Thomas T., Read

IT will be recalled that the first professor of metallurgy in the United States, appointed in 1855, never really gave any instruction in metallurgy and gradually turned into a professor of mineralogy. Metallurgy was included with. mining in the early curricula, two students who took the B.S. in mining engineering from the Polytechnic College of Pennsylvania in 1862 wrote their graduation theses on the metallurgy of copper and zinc, respectively. Even after Columbia School of Mines established, in 1868, a separate curriculum leading to the Ph.B. degree in metallurgy, very few students elected it,* preferring to take the mining degree, since they received in its curriculum nearly all the metallurgical instruction they would have in the metallurgical curriculum. For many years a large proportion of the men who attained prominence in metallurgical enterprises had taken mining rather than metallurgical degrees. METALLURGY Clifton Sorby&apos;s study of thin sections of rocks under the micro¬scope, beginning in 1850, and its subsequent development by Zirkel and others, had made microscopic petrography a recog¬nized department of geological research as early as 1873. . It was inevitable that the idea of studying polished metal surfaces-¬under the microscope by reflected instead of transmitted light would occur to investigators. Sorby himself did it in 1864, and Adolph Martens and Floris Osmond may -be especially named among the many European workers who, by the late eighties, had so developed the, necessary techniques that they were able to photograph polished metal surfaces at magnifications as high

Jan 1, 1941

By E. A. von Hahn, T. R. lngraham

The rates of cementation of pd11 on electropolislzed copper cylinders were studied in aqueous perchloric acid solutions at pdII concentrations from 0.02 to 0.1 mM. The cylinders were rotated at various high speeds to reduce the diffusion layer to a minimum thickness. In 0.1 M HClO4 solutions, the cemented palladium is dense, adherent, and shiny. The rate data indicate that there are two stages in the deposition. In the initial stage, the rate of cementation is first order with respect to the pd11 concentration. The second stage is much slower. The first stage is consistent with rate control by the diffusion of pd11 ions to the copper surface, and/or chemical reaction at the surface, and the second stage is consistent with rate control by the diffusion of copper ions from the cop-per surface through the deposit of cemented palladium, out to the main body of solution. In 0.001 M HClO4 solution, only the first stage is evident and the rate is more rapid than at higher acidities. This rate enhancement is attributed to PdOH+ ions that predominate at low acidity and aye more reactive than unhydrolyzed pdII. The deposit is porous and loose. All of the cemented deposits are Pd-Cu alloys rather than pure palladium. Activation energies are 9.5 kcal per mole in 0.1 M HClO4 and 7.4 kcal per mole in 0.001 M HClO4. CEMENTATION or displacement reactions occur between aqueous solutions of metal salts and immersed metals according to the general equation: where M1 is electrochemically the more noble metal. Although these reactions find considerable application in metals recovery processes (e.g., the cementation of copper192 or gold3), in electrorefining processes (e.g., solution purification before electrolysis4,5), and in plating and metal finishing processes,6 few studies have been made of the kinetics of such reactions.7-9 The rates of cementation reactions will depend on one or more of the following factors: a) chemical reactions at the metal-solution interface; b) ionic transport to or away from the interface;&apos;-&apos; and c) the character and adherence of the cemented deposit on the substrate metal,6a,10 because the deposit will inhibit the rate of transport of dissolving ions (M2m+) into the solution. In this investigation, the kinetics of the early stages of cementation were studied to obtain an understanding of the reaction mechanism under conditions in which the deposit was sufficiently thin that its inhibiting effect could be disregarded. To achieve this, very dilute solutions in Mn+ were used. Cementation rates are often fast in the initial stages and transport-contr01led.9 To find conditions under which transport control would shift to chemical control, very high stirring velocities were used to minimize the diffusion-layer thickness. Of the many possible cementation reaction systems, the palladous perchlorate-copper system was chosen for this investigation because it was believed to involve simple ion-for-ion exchange. In addition, there are no interfering side reactions, such as the reduction of hydrogen ions, and anion effects are usually absent in perchlorate solutions. EXPERIMENTAL Materials. Reagent-grade chemicals and redistilled water were used in all experiments. Palladous perchlorate stock solutions (0.02 M) were prepared by dissolving 99.5 pct Pd sponge (Johnson Matthey and Mallory Ltd.) in concentrated nitric acid, adding concentrated perchloric acid, evaporating twice slowly (to prevent PdO precipitation) to a small volume, and diluting in a volumetric flask. To prevent gradual hydrolysis of PdII, the stock solutions were made 0.4 M in HC104. Tests for pdIV (Ref. 11) and C1- were negative. Copper strips were cut accurately from 0.025-in. electrolytic tough pitch sheet, Sample 1 (Mines Branch stock or Canadian Copper Refineries Ltd., Montreal), or from oxygen-free, high-conductivity sheet, Sample 2 (American Metal Climax O.F.H.C. brand, sold by Utility Brass and Copper Corp., Brooklyn). The strips were annealed for 1 hr at 335°C in a stream of purified nitrogen (Union Carbide, Linde Division) which had

Jan 1, 1967

By Oliver Halston

THIS paper reports the results of the use of salt in some research work carried on during the past 3 years at the Salt Lake City Station of the Bureau of Mines, which is quartered in the University of Utah, whose Metallurgical Research Department cooperates with the Bureau. Only the work involving the use of salt is described. The authors wish equal credit given to J. F. Cullen, C. L. Larson, C. Y. Proutz, C. E. Sims, H. C. Neeld, R. W. Johnson, F. G. Moses, N. C. Christensen, and G. H. Wigton, who have all in some way contributed toward the success of this work. Lead Metallturgy Popularly Supposed to be Without Problems In bead metallurgy it is difficult at first to find metallurgical problems, as the prevailing conception has been that bead is so easily concentrated in its various ores and so easily reduced to its metallic form that the only progress possible is that involved in cheapening methods of mining and of reduction. Most of the problems usually considered are of minor nature and do not involve radical metallurgical changes. During the 3 years of work of the Salt Lake Station, the following classification of problems in the metallurgy of bend has been made: Those involving Oxidized Ores. 1. Ores containing only bead. s. Ores containing bead and silver or gold. 3. Ores containing bead and zinc, with or without silver. Sulphide Ores. 1. Ores of pyrite carrying some bead. 2. Ores of complex zinc-bead-iron sulphides. 3. Complex sulphides containing in addition precious metals.

Jan 8, 1917

By John Mathews

IT is twenty years since the writer made his first molybdenum steels and others were making them commercially five years earlier but the prevailing opinion seems to be that molybdenum steels are new; from time to time the daily press speaks of important discoveries in Europe and intimates that American steel makers have much to learn in regard to alloy steels. The earlier experiments were largely confined to the use of relatively high percentages of molybdenum in tool steels and permanent magnet steels. Recent developments deal with the types of alloy structural steels used for airplane, automobile, and other engineering requirements; in these, the molybdenum content is usually less than 1 per cent. The earlier types generally have been unsuccessful commercially, but the newer types of steels are becoming of increasing importance. The war showed our dependence on foreign sources for alloying metals; molybdenum alone is widely distributed in the United States and seemingly abundant. Many deposits are not now of commercial interest, owing to location and relatively small amount of ore available, but other deposits are being developed and a stable source of supply seems well assured. In tool and magnet steels, it was early found that 1 part of molybdenum was equivalent to from 2 to 2.5 parts of tungsten. The permanent magnet steels investigated by Madame Curie,1 in 1898, contained from 3.36 to 4.05 per cent. molybdenum; in 1902, Dr. E. L. French and the writer thoroughly tested steels containing from 2:0 to 4.0 per cent. These steels gave satisfactory tests for permanence, but were low in residual density and much more sensitive to hardening than the corresponding tungsten products. We discontinued work with them, therefore, and have no advice that such steels have been used commercially here or abroad.

Jan 2, 1921

By JAMES J. PURDY

The great majority of stokers used in residential heating installations are of the clinkering type. Because of inherent characteristics of the under- feed combustion process as it occurs in these small stokers, and because of the wide range of operating conditions encountered in residential heating schedules, not all bituminous coals can be used successfully in residential underfeed units. The problem of providing consumer satisfaction has been attacked by the joint efforts of the stoker manufacturers and the coal industry; the stoker manufacturers have been concerned with improving their product to permit the use of a wider range of coals, and coal operators have made every effort to market selected and better prepared coals. In view of the wide range of physical characteristics and chemical properties of bituminous coals, the problems are real and are receiving increased attention.

Jan 1, 1949

By P. Ramachandrarao, T. R. Anantharaman

EVER since the technique of quenching metals and alloys from the melt (splat cooling) was perfected a decade ago, it has been recognized that the grain size of products solidified by this technique may be extremely small.&apos; The further observation that fcc metals quenched from the liquid state contain very few dislocations has led to the inference2 that metals are subjected to negligible or no stresses during the rapid solidification characteristic of the "gun" or "piston-and-anvil" technique. Evidence for the incidence of appreciable densities of stacking faults has, however, been obtained in case of some splat-cooled fcc and hcp alloys,3 although not for pure metals. In the light of these earlier observations it was considered desirable to study X-ray line-broadening effects, if any, in fcc metals rapidly cooled from the melt. In the present work pure silver (>99.99 pct), aluminum (>99.99 pct) and lead (>99.9 pct) were quenched from the liquid state from temperatures about 50°C above the melting point by the "gun technique" and the resulting foils were subjected to X-ray examination in a Philips Diffractometer. The quenched foils (up to -10 u thick) did not generally stick to the substrate surface and could be easily transferred to the Diffractometer without introducing any plastic deformation. The profiles of the first five reflections from the foils were recorded in each case with Cu Ka, radiation at the slowest available scanning speed of 1/8 deg per min. To correct for instrumental broadening, profiles were also recorded from the metals annealed in vacuo at suitable temperatures. The integral breadths of the X-ray reflections were arrived at by a procedure described earlier.&apos; There was a distinct suggestion of preferred orientation in the recorded intensities of reflections from aluminum and lead foils. Such an effect was not observed in case of silver. In addition, the integral breadths of X-ray reflections from splat-cooled aluminum and lead were not significantly different from those recorded for the annealed metals. The analysis was therefore continued only for silver where the X-ray line broadening was appreciable. The pure diffraction broadening, B, was evaluated for each X-ray reflection (hkl) from silver from the observed, B, and instrumental, b, breadths with the aid of each of the three equations due to Scherrer,5 Anantharaman and Christian,6 and Warren and Biscoe,7 respectively: Bs= B-b BAC=B- b2/B Table I gives the values of particle size, 71, the lattice strain, E, arrived at by the use of the following well-known relations and on the assumption that all observed diffraction broadening could be attributed to lattice strain or particle size, respectively: E = 1/4 cos ? n= B cos ? where A is the wavelength of X-radiation and 0 is the Bragg angle. As no significant peak shifts or asymmetry could be detected in the profiles from the foils, the possibility of any significant contribution due to twins or stacking faults was ruled out. The absence of faults is by no means surprising since pure silver is known to develop stacking faults only on severe deformation and the stacking fault densities recorded so far for even silver filings have been extremely low.&apos; The very low values for percentage mean deviation from the mean value for the particle size in Table I strongly suggest that all observed broadening in splat-cooled silver can be attributed only to small particle size. This conclusion receives further support from the lowest mean deviation recorded for data computed from the Scherrer equation based on Cauchy profiles that are considered characteristic of particle size broadening. Further analysis for separation of particle size and lattice strain effects was considered unnecessary in view of the very large mean deviations obtained for strain values and also the earlier results suggesting absence of even detectable strain in metals and alloys quenched from the melt. It is to be stressed in this connection that the particles are actually grains and not cells formed by walls of high dislocation density usually encountered in deformed samples. As such, the absence of strains is not surprising. The present results are probably the first to record X-ray line broadening due only to small particle

Jan 1, 1970

By C. E. Arnold

IN attacking the problems of mining and treating large disseminated copper orebodies such as those occurring in the Miami or the Ray district of Arizona, one of the vital questions to be decided is, "What is the value of the orebody per ton in place?", The importance of this question lies in the fact that its answer helps to reach a decision concerning the method of mining to be employed. Let it be assumed for the moment that milling and smelting practice, especially the latter, is more or less standardized; and has a cost incapable of much variation, while the cost of mining has a considerable range owing to the large number of mining methods by which a mine may be exploited. For instance, a method may be adopted whereby all the available ore can be extracted perfectly clean (i.e., unmixed with barren material) at a relatively high cost; or a method may be used involving the necessary dilution of a portion of the ore with waste, in order to recover that portion, thus giving a relatively poorer extraction but nevertheless a lower working cost than the former method. Naturally, between these two extremes there lie intermediate methods. In order to arrive at a solution of a mining method problem that is of any value, it is necessary that accurate and complete estimates of mining, milling, and smelting performances and costs shall be established. By way of illustration is given a, case of mining, by the system of undercutting and caving, a large orebody with an average content of 1.50 per cent. copper supposedly all in sulphide form, to which case has been applied those results in mining operations which one may fairly assume are to be expected when viewed in the light of G. R. Lehman&apos;s ore-drawing experiments, which form the subject of a paper to be presented at this meeting. It is assumed that the total. costs will be as shown in the following table:

Jan 9, 1916

Jan 1, 1965