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By J. T. Ryan

SiR Humphry Davy's epoch-making treatise delivered on Nov. 9, 1815, before the Philosophical Society of London, first announced and demonstrated a flame safety lamp for detecting methane in mine air. Coal mining was known in England as early as 1180, more than 600 years before the advent of a safe lighting and gas-detecting appliance. Mining, no doubt, took place first in outcrop coal, and danger from gas did not become apparent until the workings became deeper and more extensive. Davy's lamp did much to increase the safety of miners, and a modified Davy lamp is still widely used for detecting the presence and approximate amount of combustible gas in mine air. No factor in coal mining has improved more rapidly during the past 15 years than safety measures. This progress came about not only as a result of the workmen's compensation laws now in force in most of the states of the Union, but because the research work of the U. S. Bureau of Mines had demonstrated some dangers incident to coal mining which were not considered to be serious matters prior to that time. Electric power and explosives are necessary but potential hazards in American coal mines. A general mechanization program is now under way, while the mines are growing deeper and more extensive. Consequently ventilation problems are more complex and a demand has been created for positive detecting devices for carbon monoxide as well as methane. Scientists and manufacturers, working hand in hand during the past 10 years, have invented and developed appliances which make systematic and effective examination of mine air possible. The object of this paper is to describe briefly the approved testing devices in use today, with the hope that they may lead to the saving of life. Two Groups of Dangerous Coal Mine Gases Dangerous coal mine gases may be placed in two groups: (I) those that are dangerous because of their explosive properties, and (2) those that are dangerous because of their toxic properties. Methane (CH4) is the important gas of the first group, and in so-called gaseous mines constantly exudes from the coal and rock; it may accumulate and form an explosive mixture. The explosive range of methane-air mixtures is well defined. The low limit is 5 per cent., the high limit 15 per cent., and the most explosive

Jan 1, 1927

By Oswald J. Heinrich

At the Baltimore meeting of the Institute in February, 1879, Mr. Eckley B. Coxe, then president of the Institute, called attention in his address to the subject of Secondary Technical Education, and

Jan 1, 1881

By Oswald J. Heinrich

AT the Baltimore meeting of the Institute in February, 1879, Mr. Eckley B. Coxe, then president of the Institute, called attention in his address to the subject of Secondary Technical Education, and

Jan 1, 1881

By B. H. Kear, H. G. F. Wilsdorf

After a few percent strain, dislocations in disordered Cu3Au are arranged in groups in ulell-defined slip planes, which contrasts with the more or less random distribution of dislocations in the corresponding ordered alloy. Unusually long and straigkt dislocation segments observed in the ordered alloy are interpreted in terms of cross slip of screw dislocations from their slip planes into cube planes. It is proposed that cross slip, and also dislocation reactions of the Lomer type, play an important role in work hardening of the ordered alloy. ACCORDING to to dislocations in the disordered lattice are only partial dislocations in the corresponding super lattice, so that in the super-lattice dislocations must occur as pairs, with each pair connected by a ribbon of antiphase boundary. Recently, striking confirmation of the theory was obtained by Marcinkowski, Brown, and Fisher3j4 in a study of dislocation configurations in thin foils of the Cu3Au alloy by transmission electron microscopy. In the present investigation, the observations made by these workers have largely been confirmed. Moreover, it has been found that dislocations in the plastically deformed ordered alloys are often unusually long and straight. This paper is mainly concerned with these new observations. A tentative explanation is offered in terms of cross slip of dislocations in the ordered alloys. PREPARATION OF SPECIMENS An alloy close to the stoichiometric composition Cu3Au was prepared by melting together high-purity copper and gold (99.995 pct) under vacuum in a graphite mold. The resulting ingot was rolled down and cut into strips 1 1/2 by 1/8 by 0.005 in. in thickness. All strips were annealed at 900°C for 3 hr and cooled to 500'~. A few of the strips were quenched from ~OO'C, while the remainder were slowly cooled in the temperature range 400" to 150°C at intervals of a few degrees every 24 hr to produce a high degree of order. After this annealing treatment the average grain size of the material was about 0.1 mm. Material in different states of order was obtained by reannealing fully ordered material at temperatures below the critical temperature: followed by quenching. The equilibrium degree of order attained at a particular reannealing temperature was obtained from the paper by Keating and warren.' A simple hand-operated tensile jig was used to give the strips the required amounts of plastic strain. Following straining, a number of specimens were cut from each strip by means of an electrolytic cutting device.= Finally, each specimen was thinned electrolytically using the polishing procedure described by Bakish and Robertson,7 and these thin sections were examined in the Philips EM-100B electron microscope operating at 100 kv. EXPERIMENTAL OBSERVATIONS Antiphase boundaries. The average domain size in the ordered alloys was found by direct observation to be about 0.25 p. Also, it was observed that the grown-in configuration of antiphase boundaries was changed little by plastic deformation, at least up to 6 pct strain. Fig. 1 shows an electron micrograph of a thin foil of partially ordered alloy (S - 0.8) after 3 pct strain. The micrograph shows a few dislocation pairs and a somewhat crystallo-graphic network of antiphase boundaries. The normal to this foil was found by diffraction analysis to be approximately 11121, with principal directions in

Jan 1, 1962

By H. W. Mossman

Three aspects of magnetite smelting are discussed. The first is the working out of equilibrium conditions for eliminating sulfur. The second is the influence of magnetite solubility on the difficulty of tapping the reverb matte. The third is an approximation of the equilibrium conditions in the reverb gases which govern whether magnetite is mode or reduced in the reverb slag by these gases and by any iron sulfide in the slag. MAGNETITE has had a varied history in the Hurley Smelter since its start in 1939. Magnetite determinations on the smelter products are made regularly only on the monthly composite samples. Variations on the monthly averages are shown in Table I. Magnetite which drops from the slag and matte in the reverb has some slight bottom buildup which comes and goes, but no substantial accumulation from this source has been found at the end of a normal nine months' furnace campaign. However, there has been some low grade magnetite bearing material mixed with considerable A1,0,, which has slid down from the bottom of the sloping flue between the reverberatory furnace and the waste heat boilers. This accretion has required drilling and blasting near the skimming end of the furnace. The magnetite has interfered with tapping at times. When the smelter was first started, tapping trouble from magnetite was extremely severe. Increasing the reverberatory furnace temperature by putting in an air preheater and a Dutch oven has helped greatly, although there still is occasional tapping trouble. When the present series of physical chemistry articles on copper smelting started coming out in 1950, they were read with interest, but no immediate application was seen for them. Results of some laboratory work in 1952 aroused a much stronger interest in this physical chemistry. A series of melts was made on some converter slags, which had magnetite in very large grain sizes, with the object of reducing the grain sizes in the slag, as it was known that it was easier to handle in the reverb in that condition. Anything done in the tests greatly reduced the grain sizes—even in the controls, where nothing was done except melt the slag and cast it. There was more magnetite in the slags after the tests than before, and with wide variations. There were no obvious reasons lor much of what happened in these tests. Much of the base material published in English in this field was made available for study. Recalculations were made on many of the type problems, and part of the data was reduced to local temperatures and compositions. Explanations were found for what happened in the 1952 series of tests on converter slags, and the same principles turned out to be a description of much of what magnetite does in the reverb. This article is to present the results of that study, from the viewpoint of applying the technical material in definite numerical form to the operating conditions in both the converters and the reverberatory furnaces at the Hurley smelter. Table I. Magnetite Variations on Monthly Averages, 1939 to 1955 Pet Magnetite Lowest Highest Average Converter slag 13.6 43.3 25.4 Roverb slaa 2.7 20.9 8.7 Matte 28 15.9 98 In general it was found that magnetite is made or reduced in both the converters and the rever-beratory furnace, depending on variations of temperature, matte composition, and reverb gas composition occurring in ordinary plant operation. Within reasonable limits, the field conditions for formation or reduction can be predicted, and probably can be set up and maintained. Converter conditions affecting magnetite formation can be put into numerical values better than for the reverb from purely technical calculations. The converter can be operated so as to keep the magnetite in the slag down to between 12 and 14 pct and still give satisfactory life for the converter brick. This depends upon having converter flux available which will make a slag with a good separation without raising the temperature too high. In the Hurley reverb and others with similar conditions, it is likely that a compromise of conditions will give a reasonably good control of combustion and still keep the magnetite from building up on the bottom. This discussion consists of three main parts. The first is the working out of the equilibrium conditions in the converter for determining in which direction the reaction 3 Fe,,O, (s) + FeS (1) F? 10 FeO (1) + SO, will go under actual converter operating conditions. The second deals with the influence of the solubility of magnetite in the slag and matte in the reverb on the difficulty of tapping matte. The third is an approximation of the equilibrium conditions in the

Jan 1, 1957

By T. R. Ingraham

The thermal decomposition of Nickel (II) sulfate was examined by determining the total pressure of SO3, SO2, and O2 developed over a sample when it was heated in an evacuated system fitted with a Pyrex-bellows mercury manometer. The equilibrium pressures were combined with previously published data on the oxide and sulfide systems to establish a predominance volume diagranz for the Ni-S-O system over the normal range of roasting temperature and gas composition used in the treatment of nickel minerals. The protective coating of NiSO4 which forms on the oxide during roasting, and which causes the sulfation reaction to almost cease, may be disengaged readily by temperature cycling. When nickel sulfide is roasted in air, either nickel sulfate or nickel oxide may be formed by altering the experimental conditions. Because of the differing stabilities of various metal sulfates and oxides, possibilities exist for separating some of the constituents of mineral mixtures by the technique of selective sulfation and oxidation roasting. The experimental information required to make reasonable predictions of the possibilities of this technique is incomplete and, in some cases, in conflict. As part of a general program of providing some of the required information, a study EXPERIMENTAL Materials, Apparatus, and Procedure. Fisher certified reagent-grade nickel sulfate hexahydrate was used in the source material for both nickel sulfate and nickel oxide. The manufacturer's analysis indicated: 0.001 pct Fe, 0.005 pct Cu, 0.003 pct Co, 0.001 pct C1, other 0.15 pct. The material was dehydrated, before use, by heating at 420°C for 5 hr. The oxide was prepared by roasting the sulfate in an open tray in a muffle furnace for 10 hr at950°C. Preliminary experiments which were made by thermogravimetric analysis and X-ray diffraction analysis established that only one crystal form of nickel oxide and one of nickel sulfate were involved in the equilibrium shown in Eq. [I]. The apparatus used in the experiments was described previously.1,3 The main feature is a corrosion-resistant, flexible Pyrex bellows which is used to protect the mercury in a u-tube manometer. Before a run was begun, the apparatus was washed thoroughly with hot hydrochloric acid and water and dried by heating under vacuum. The mixed sulfate -oxide sample was contained in a platinum bat rest-

Jan 1, 1967

By D. H. Baldwin, R. E. Reed-Hill

Six compositions of polycrystalline ZY-0 alloys, containing up to 4.2 at. pct 0, were tested in tension between 77° and 600° K. The data obtained from each of the compositions corresponded closely to a rela-ion between yield stress and absolute temperature of the form In s/so = BT, where oo is the yield stress extrapolated to zero degrees and B is a constant. In agreement with others who have observed this relationship, it is shown that the activation energy may be expressed as Ho In so/s. In the present specimens Ho is approximately 18,000 cal per mole and is apparently independent of temperature and composition inside the limits of the investigation. It is also demonstrated that this form of activation energy cowesponds to a strain rate sensitivity parameter RT/Ho. Oxygen was also noted to have an effect upon the operative deformation mechanisms. With increasing oxygen concentration there was an increased tendency to observe both nonbasal slip and cross-slip phenomena. Oxygen does not seriously inhibit twinning more than it does slip. Twins were observed in all specimens tested. It is becoming increasingly evident that interstitial atoms in solid solution are able to interact strongly with mobile dislocations. Stein, Low, and seybolt,&apos; have shown that, if the carbon concentration in bcc iron is lowered below the solubility limit, its flow stress temperature dependence is markedly reduced. This suggests that carbon atoms in interstitial solid solution may be responsible for the pronounced temperature dependence of the flow stress normally observed in iron. This view has recently been challenged by Leslie and sober2 who observed a strong flow stress temperature dependence in iron to which a trace of titanium had been added in order to remove carbon atoms from solution. Since the interstitial concentration must be reduced below approximately 1 ppm in order to produce a pronounced effect on the flow stress temperature dependence,&apos; studies of the effect of interstitials on the flow stress in iron necessarily involve serious experimental difficulties in alloy preparation. There are other metals, however, in which strong effects of interstitial solutes upon both the flow stress and its temperature dependence are observed. Of particular significance is zirconium which, according to Domagala and Mcpherson, 3 is capable of dissolving 28.6 at. pct O. The O-Zr alloy system is an almost ideal system for studying the interaction of interstitial atoms with deformation modes since it is possible to form alloys capable of study over an extensive range of compositions. Mills has made such a study using single crystals oriented primarily for single prismatic slip4 and has found an effect of oxygen concentration on the flow stress temperature dependence analogous to that observed in iron due to carbon by Stein, Low, and Seybolt. The present paper is specifically concerned with the effect of oxygen on deformation in polycrystalline zirconium. Although plastic flow in this type of specimen is much more complex than that reported for the single-crystal work, and involves both slip (on several different types of planes) and mechanical twinning, the results of this investigation are in general agreement with the single-crystal observations concerning the effect of oxygen on the temperature dependence of the flow stress. In addition, they also demonstrate that oxygen affects the acting deformation systems. This is in contrast to single-crystal results4 that showed only single slip on a prism plane. EXPERIMENTAL PROCEDURE Material. High-purity hot-rolled zirconium strip, 0.2 in. thick by 4 in. wide, of 0.10-mm average grain diameter, was used for forming alloys. It was obtained from the Carborundum Metals Co., Akron, N.Y., whose analysis indicated the major impurities were, in wt ppm: Hf, 540, C, 145; Fe, 100; and 0, <80. The plate texture was similar to a wire texture, with basal planes generally parallel to the rolling direction and basal poles randomly distributed about the rolling direction. The heat treatments described below did not appreciably alter the basic texture. specimen Preparation. Small threaded-end tensile specimens were machined from the plate with axes perpendicular to the rolling direction. These transverse specimens had gage sections 1 in. long by 0.060 in. in diam. The small gage section diameter was dictated by the fact that the alloys were formed by dif-

Jan 1, 1969

By Walter A. Backofen, Donald H. Avery

Intense primary and cross-slip traces were observed in easy glide on Cu: 6 pct-A1 single crystals deformed in tension. A mechanism of cooperative source operation is developed which recognizes that both the applied stress and that resolved from dislocation arrays contribute to the stress on a source. By considering the stress from arrays, the activation of secondary systems in easy glide and the difference in appearance of slip markings on pure and solid-solution crystals are explained. Furthermore the cross-slip system is predicted to be second favored during easy glide for a large range of orientations. FROM many studies of easy glide in fcc metal crystals, a pattern of observations has developed which still requires full explanation. In pure copper, aluminum, and silver, the easy-glide, slip traces are fine and uniform with steps 50 to 100A on the primary* system.1-4 Yet it has not been at all clear why slip in the pure materials should occur in this way. On the contrary, one might expect, with simple slip and a low hardening rate, that the weakest sources on the primary system would operate extensively to produce strong slip traces. In a brass, however, intense slip steps up to 0.5 are found on both primary and cross-slip system4,6-9 seeger10 has pointed out that in a material of low stacking-fault energy such as a brass, extensive easy-glide cross-slip could not be thermally activated, but rather must originate on the cross-slip plane. In further explanation, the wilsdorfs4 and seeger10 noted that only the cross-slip system would not harden differently than the primary, whatever the latent hardening tendencies of other systems; therefore, in brass, which exhibits high latent hardening, the cross-slip is the only secondary system which could be expected to operate extensively, once it became active. The explanation of how the cross-slip system, on which the applied stress is low, becomes activated has been related, in a nonspecific way, to pileups by Hirsch11 and Honeycombe.l2 In the work to be described, deformation traces on solid-solution Cu: 6 pct A1 single crvstals were studied with conventional light microscopy and a mechanism proposed to explain both the observed activation of secondary systems in easy glide and the difference in appearance of slip markings on pure and solid-solution crystals. EXPERIMENTAL In preparing the Cu-A1 single crystals, rectangular bars approximately 0.4 in. by 0.4 in. by 6 in. were first machined from ingots made by vacuum melting copper of 99.999 pct and aluminum of 99.99 pct purity. They were then packed in powdered graphite and the crystals grown in a modified Bridgman apparatus at a rate of 5 mm per hr under an atmosphere of purified nitrogen. Crystal perfection as indicated by Laue back-reflection spots was good except for the last inch to freeze, which was discarded. Orientations so determined are shown in Fig. 1. Tensile specimens with a reduced gage section were obtained by masking the ends with electrical tape and chemically polishing in a solution of 20 pct HNO3, 24 pct HAc, 52 pct H3PO4, 2 pct Hcl, and 2 pct H20 at room temperature.13 Finally, the specimens were electropolished in a 30 pct ortho-phosphoric acid bath at 12 v. Tensile deformation was performed in an Instron testing machine at a strain rate of approximately 1 pct per min. During easy glide, extensive cross-slip was observed with all crystals, in spite of the low Schmid factor on this system, and in no case was any other secondary trace encountered. A certain sequence of slip-line development was associated with the extension. At yielding, intense primary traces passing completely through the crystal developed at some point on the gage section, generally near the grips; weak cross-slip traces also appeared, extending from the primary traces into the adjacent undeformed crystal, Fig. 2. Further extension led to primary traces adjacent to the first-formed and connecting with those of the cross-

Jan 1, 1963

THE illustration below shows the design of the face of the clock to be erected as a memorial to the American engineers who gave their lives overseas in the World War. It will be placed in the tower of the library at Louvain, which has been rebuilt through contributions from America. Large clocks seldom have figures on their faces, and each hour on this one will be marked by a gold star, the four faces containing 48 stars, one for each State of the Union. Contributions toward the memorial have been received from every one of the States, and from Alaska and Hawaii as well, but additional subscriptions are invited by the Com-mittee which is arranging for the memorial.

Jan 3, 1928

By Graham John Mitchell

In May, 1919, a crosscut on the 1400-ft. level of the Briggs mine, a Calumet, & Arizona property at Bisbee, Ariz., penetrated a deposit of pyrite and chaleopyrite that had replaced quartzite and limestone and was cut, especially in the quartzite, by numerous fractures. As the drift advanced through these fractured portions, they were found to be free from copper-sulfate minerals although &apos;(sooty" chalcocite was present in the fractures and also as coatings and replacements of pyritc. Howcver, sulfate soon made its appearance in certain fractures, filled them, and began to form stalactites. Several centers of incipient growth were noted for subsequent study, observations in the drift at frequent intervals during the past 17 months affording the opportunity of determining the rate of formation of the stalactites. The one to attain the greatest growth during the period of observation measured 27 in. long and averaged 1 in. in diameter. A close study showed it to be composed of small interlocked crystals of chalcanthite radiating from a small central tube-like opening. The trielinic form of the crystals was evident with prism and pinacoid faces recognixable. Assuming that the supply of copper-sulfate solution was constant (luring the 17 months—-and such an assumption is supported by the fact that the number of drops of excess solution falling, per minute, from the tip of the stalactite at the times of observation checked closely—thc stalactite, it has been estimated, grew at the rate of 1.58+ in. per month. DISCUSSION James I?. Kemp," New Yolk City.—J. S. Curtiss, in the monogr:tpll on Eureka, Ncv., say:: that a stalactite formed in limestonc grew 5/8 in. in three weeks and continued to grow as long as it had a drop of water or solution at the lower end. Another grew at the rate of 3/8 in. in the same time. In the case cited by Mr. Mitchell, the growth was about 136 in. in three weeks. A year ago we had an instance of the formation of cubes of galena on a railroad spike. A mine in northeastern Oklahoma had been idle about two years but on account of pumping of neighboring properties, the circulation of water was active throughout the drift. A cube of galena,

Jan 1, 1922

United Engineering Trustees, Inc. W. D. B. HOTTER, JR., &apos;44 A. L. QUENEAU, &apos;45 ALBERT ROBERTS, &apos;43 The Engineering Foundation GEORGE D. BARRON, &apos;44 F. F. COLCORD, &apos;46 A. L. QUENEAU, &apos;45

Jan 1, 1942

By H. T. Sumsion, A. U. Seybolt

The results of an investigation of vanadium-rich V-O solid solutions are presented, indicating the structure and lattice parameters of two solutions, a and ß, and their approximate temperature-composition existence. The a solution is the terminal body-centered cubic one, and contains up to 3.2 atomic pct 0. The ß solution has an ordered body-centered tetragonal structure, is formed at 1270°C, and exists from about 15 to 22 atomic pct 0. From the evidence available, the various phase boundaries have no appreciable temperature dependence. Evidence has been found for a polymorphic transformation in pure vanadium at 1550°C. IN an earlier investigation&apos; dealing with the preparation of pure vanadium by calcium reduction of the oxide, it was found that small amounts of oxygen drastically reduced the ductility of the metal. Because this effect was so marked, it was decided to make a study of the solubility of oxygen in solid vanadium. This report deals principally with this solubility and the nature of the phase relationships in the vanadium-rich region, particularly at temperatures below 1300 °C. However, during the investigation enough data on the V-0 system were obtained to make it appear worthwhile to present a tentative phase diagram up to the composition VO. The only significant prior work found on this system are the contributions of Klemm and Grimm,&apos; and Mathewson et al.3 Klemm and Grimm prepared a wide range of V-O compositions by powder techniques including the compositions VO.l, VO.2, VO.3 and VO4 (9.1, 16.8, 23, and 28.6 atomic pct 0, respectively). The first three compositions were found to consist of a body-centered tetragonal solid solution, while the last also showed lines of VO (NaCl structure). They found that the parameter c, increased and the parameter a, decreased with increasing concentration of oxygen. For their composition VO.27, or about 16.8 atomic pct O, they cite the values a, = 2.948A, c, = 3.53A, and c/a = 1.2. Klemm and Grimm made no attempt to determine the solid solubility limit nor to construct a phase diagram. They did, however, give some data on the homogeneity range of VO, and they proposed a structure for the body-centered tetragonal solid solution; these points will be taken up later. Materials and Preparation of Samples The vanadium used in this investigation was prepared in the laboratory by the method previously mentioned.&apos; A typical analysis is as follows: Fe, 0.007 pct; Si, 0.02; Ca, 0.06; C, 0.224; O2, 0.044; N,, 0.0017; H2, 0.003; and V, 99.34 ±0.3 (assay). The vanadium assay is probably low by about the error given. The impurities total about 0.36 which, if subtracted from 100, gives a purity of about 99.6. At the time material was being prepared for this work no suitable technique was available for melting vanadium without appreciable contamination. The procedure adopted therefore was to cut the calcium-reduced regulii into slices which were then rolled to strip about 0.025 in. thick for oxygen diffusion. Pieces of rolled vanadium of approximately 0.025xY4xl in. and weighing about 0.3 g were suspended in a vertical fused-silica tube which was part of an ordinary gas absorption apparatus. The silica tube was heated by an electric resistance-tube furnace which could be raised around the silica tube or lowered away from it as desired. This apparatus had no novel features which require detailed description. Other than the silica tube and furnace, it consisted of a glass system evacuated by a liquid nitrogen trapped mercury diffusion pump, a mercury-operated gas burette, a McLeod and a Pirani gage, and a mercury manometer. It was also equipped with suitably located stopcocks for isolating various parts of the system; the vacuum ordinarily attained was between 10" and 10-6 mm of mercury. Oxygen generated by decomposing MnO2 was passed through anhydrous magnesium perchlo-rate before introducing it into the gas burette and thence to the absorption chamber.

Jan 1, 1954

H. M. BOYLSTON E. B. YOUNG SECRETARY-TREASURER OF SECRETARY-TREASURER OF BOSTON SECTION MONTANA SECTION R. M. KEENEY F. G. MOSES SECRETARY-TREASURER OF SECRETARY-TREASURER OF COLORADO SECTION UTAH SECTION

Jan 7, 1919

By J. M. Steininger

The mechanism of flotation of sphalerite with sulthydryl compounds of different acidities has been interpreted in terms of chemisorption of unionized collector molecules on hydroxylated zinc surface sites. At a constant collector concentration of 1.5 x 10 -4 M, the upper pH limit of flotation has been found to be directly and closely related to the negative logarithm of the collector acidity constant, pKa The experimental results indicate that flotation is prevented when about 95% of the collector is ionized. At low pH, collector adsorption decreases because of the reduction in number of hydroxylated zinc surface sites. These two effects explain why sphalerite does not float with the more acidic sulfhydryl compounds. Typical anionic collectors for sulfide minerals are short chain xanthates or dithiophosphates which are nearly completely ionized in the generally alka- line flotation circuits.1,2 Flotation mechanisms have therefore been usually interpreted in terms of ionic collector adsorption and competition with ionic depressants of the same sign as the collector. A few attempts have been made, however, to consider adsorption of neutral collector molecules. Contact angle measurements obtained by Wark and cox3 have been interpreted by Cook and Nixon4 as representing adsorption of the free acid molecules on sulfide minerals such as pyrite, galena, chalcopyrite, and sphalerite. Cook and Nixon tried to show that the free acid concentration of the collector was substantially constant along the curves relating collector addition and minimum pH for bubble adherence. However, in a review of the subject,5 Sutherland presented new data on the flotation of pyrite with mercaptobenzothiazole which did not support the hypothesis of constant concentration of either collector ions or collector free molecules. He further showed that, with the assumption of constant occupation of surface sites, both theories lead to relations between pH and collector addition having similar mathematical expressions. In fact, with the more likely assumption of variable occupation of surface sites which results in more complex relationships, it appears impossible to

Jan 1, 1968

By O. P. Arora, M. Metzger, G. R. Ramagopal

Single-phase aluminum containing 0.0001 to 0.06 pct Cu was studied in strong acid, mainly through observations of hydrogen evolution. The strong influence of copper was exerted almost entirely through the imposition after a certain delay time of an auto-catalytic localized-corrosiott reaction. Additions of cupric ion to the acid produced lower accelerations. The significance of the quantity and distribution of copper was discussed, and the implications for intergranular corrosion and neutral chloride pitting were indicated. AN investigation of intergranular corrosion in single-phase high purity aluminum exposed to hydrochloric acid indicated the copper content of the metal to have an influence on corrosion at lower levels than previously suspected.&apos; The work reported here was a closer examination of the action of copper but dealt with general corrosion to gain the advantage of having a continuous measure of corrosion through the volume of hydrogen evolved, the reduction of hydrogen ion to hydrogen gas being the principal or only cathode reaction in strong hydrochloric acid. Previous work on the hydrochloric acid corrosion of aluminum was sometimes insufficiently structure-conscious and the need for care in evaluating it arises from the low solubility of the iron impurity,&apos; and of some alloying elements, and the known or possible presence in many of the compositions studied of second phases leading to greatly increased corrosion rates.3 These increases are attributed to the presence of low hydrogen-overvoltage cathodes provided by the second phase.3&apos;4 For the present single-phase work, a few studies which used high-purity base material and small copper additions5-&apos; provide the essential information most unambiguously. The corrosion rate was shown to be increased markedly by the introduction into the acid of small quantities of the ions of copper (and of certain other metals) which cement on the aluminum and provide cathodes of low overvoltage.5 When there was sufficient copper in the aluminum, the same result was produced during the course of corrosion leading to a rate which increased with time as the reaction was stimulated by one of its products (autocatalytic reaction). In 2N (7pct) HC1, an accelerating rate was observed at 0.1 pct Cu but not at 0.01 pct.5,7 The present work dealt with corrosion rate and morphology and their correlation with the quantity and distribution of copper catalyst for copper contents from 0.0001 to 0.06 pct. PROCEDURE A lot of high-purity aluminum containing 0.0021 pct Cu, 0.001 pct Fe and 0.003 pct Si (Alloy A) was alloyed with copper to yield aluminum containing 0.014 pct Cu (B) and 0.06 pct Cu (C). Later it was found necessary to include the lower copper Alloy K which contained 0.0001 pct Cu, 0.0004 pct Fe and 0.0004 pct Si. The upper limit for any other element can be confidently estimated as 0.0005 pct. No element other than copper appears to be present in quantities sufficient to have an effect on general corrosion as great as the observed effect of the copper in A, B, and C. The only other heavy metal detected by spectrographic examination was silver (< 0.0001 pct). The acid was made up from a selected lot of 37 1/2 pct CP hydrochloric acid containing 0.1 ppm heavy metals (mainly Pb), 0.05 ppm Fe, and < 0.008 ppm As and from water distilled from 1 megohm-cm demineralized water and believed to have contained negligible quantities of heavy metals influencing corrosion. Acid strength was adjusted to within 0.05 pct HCl of the stated value by using precision specific gravity measurements. Test blanks 10 by 41 mm were sheared from 1.65-mm cold-rolled sheet. Edges were finished by filing. The blanks were annealed in air at 645°C for 24 hr in alundum boats and rapidly water quenched. The anneal is thought to have produced a substantially homogeneous solid solution—for iron, copper, or silicon, for example, the annealing temperature was 200°C or more above the solvus-and the quench is considered to have preserved the high-temperature structure except for the condensation of lattice vacancies into dislocation loops.&apos; The 0.06 pct Cu alloy did not appear unstable in respect to slow precipitation reactions at room temperature since two pairs of tests failed to show significant differences between specimens heat treated 3 1/2 years earlier and specimens heat treated 1 or 2 days before.

Jan 1, 1962

J. Chipman (Massachusetts Institute of Technology, Cambridge, Mass.)—The fact that the experimental work has been applied to copper rather than iron and that the paper is presented to the Iron and Steel Division, I regard as rather significant. It shows the unity of metallurgy and the fallacy of trying to cut it up by metals. This result for the solubility of sulphur in molten copper correlates with Professor Schuhmann&apos;s finding that the published data on the other side of the copper-sulphur miscibility gap are also in error. I should like to ask the author to say a little bit more about the sulphur capacity of the slag. T. Rosenqvist (author&apos;s reply)-—I hope that Dr. Chipman will find the derivation of the expression for sulphur capacity more clearly explained in the printed version of the paper than in the oral discussion at the meeting. I feel that this quantity, which actually is the ratio of two activities, can be measured more easily than the individual activities. Even if the ratio CaO/CaS is chosen as the standard state, the expression can be used for any slag, even for slags completely free of lime, and it represents a way to put the desulphurizing power of all slag constituents into one bag. Some doubt has been expressed as to whether oxygen ions really exist in calcium oxide and in molten slags. From a thermodynamic view point that question is of minor importance. The term oxygen ion activity, or any activity for that matter, is defined rigorously by the equation: activity = exp u/RT, where u is the change in free energy connected with the transfer of one mol of ions from the standard state into the slag. Whatever happens to the ion in the slag is of no concern to the thermodynamicist. Regardless of whether the ion is "free" in the slag or not, or whether it is present in a very small amount, its activity can always be expressed, and for a thermo- dynamic calculation that is all we need. However, ionic activities will only be of some real value if they are simple functions of the slag composition, or can be measured easily. Concerning the real nature of the oxygen in the slag, my feeling is that the oxygen atom has a rather multiplex nature depending on how strongly it is tied by covalent forces or polarized by the other atoms or ions present. The oxides of iron, cobalt, and nickel differ from calcium oxide and blast furnace slags as to the amount of free electrons that can give rise to electronic conductivity. In slags we know that the conductivity is mostly ionic. The fact that reversible emf&apos;s can be obtained with oxygen electrodes in certain salt melts, indicates a significant amount of oxygen ions in these melts. But extended work, e.g. polaragraphic studies and measurements of transference number; are needed to obtain quantitative information about the real structure of the slags. D. E. Babcock (Republic Steel Corp., Youngstown, Ohio)—-With reference to the ion, it might be well to remember Dr. Moses Gomberg. All of his life he had no use for the ionization theory and he contributed greatly to the field of chemistry on the assumption there was no such thing as ions. I do not think we have to worry about whether the oxygen is ionic or not. I think one thing specifically should be brought to your attention and this I think is one of the important contributions of Dr. Rosenqvist. He pointed out what we know as oxygen potentials or what is described as oxygen potentials. I have used this concept for a long period of time and I want to state that if this concept is properly applied, it vitiates much of what we have in the literature, or makes our usual ideas regarding oxidation seem primitive. That one thing is more valuable than almost all the rest of the discussion as a fundamental basis on which to build a reasonable in-

Jan 1, 1952

By J. David Lowell

On a low desert ridge six miles south of Cairo lie the remains of the world&apos;s earliest known mining town, Maadi of predynastic Egypt. Copper artifacts, including ingots and an ax head which was spoiled in casting, indicate that copper ore was smelted in Maadi, and radio- carbon dating has placed the activity at least as early as 3300 B.C. This is the oldest documented date for the use of smelted copper, the first industrial metal used by man. Copper was mined and smelted almost as early in Asia Minor and on the island of Cyprus, but the greatest development of the budding copper industry was in Egypt, where specific historical records show that about 10,000 tons of copper metal were produced in the period from 1850-1580 B.C. Contrast this production of 35 tons per year with the present world production of about 7 million tpy.

Jan 1, 1970

By Terkel Rosenqvist

J. Chipman (Massachusetts Institute of Technology, Cambridge, Mass.)—The fact that the experimental work has been applied to copper rather than iron and that the paper is presented to the Iron and Steel Division, I regard as rather significant. It shows the unity of metallurgy and the fallacy of trying to cut it up by metals. This result for the solubility of sulphur in molten copper correlates with Professor Schuhmann&apos;s finding that the published data on the other side of the copper-sulphur miscibility gap are also in error. I should like to ask the author to say a little bit more about the sulphur capacity of the slag. T. Rosenqvist (author&apos;s reply)-—I hope that Dr. Chipman will find the derivation of the expression for sulphur capacity more clearly explained in the printed version of the paper than in the oral discussion at the meeting. I feel that this quantity, which actually is the ratio of two activities, can be measured more easily than the individual activities. Even if the ratio CaO/CaS is chosen as the standard state, the expression can be used for any slag, even for slags completely free of lime, and it represents a way to put the desulphurizing power of all slag constituents into one bag. Some doubt has been expressed as to whether oxygen ions really exist in calcium oxide and in molten slags. From a thermodynamic view point that question is of minor importance. The term oxygen ion activity, or any activity for that matter, is defined rigorously by the equation: activity = exp p/RT, where p is the change in free energy connected with the transfer of one mol of ions from the standard state into the slag. Whatever happens to the ion in the slag is of no concern to the thermodynamicist. Regardless of whether the ion is "free" in the slag or not, or whether it is present in a very small amount, its activity can always be expressed, and for a thermo- dynamic calculation that is all we need. However, ionic activities will only be of some real value if they are simple functions of the slag composition, or can be measured easily. Concerning the real nature of the oxygen in the slag, my feeling is that the oxygen atom has a rather multiplex nature depending on how strongly it is tied by covalent forces or polarized by the other atoms or ions present. The oxides of iron, cobalt, and nickel differ from calcium oxide and blast furnace slags as to the amount of free electrons that can give rise to electronic conductivity. In slags we know that the conductivity is mostly ionic. The fact that reversible emf&apos;s can be obtained with oxygen electrodes in certain salt melts, indicates a significant amount of oxygen ions in these melts. But extended work, e.g. polaragraphic studies and measurements of transference number; are needed to obtain quantitative information about the real structure of the slags. D. E. Babcock (Republic Steel Corp., Youngstown, Ohio)—-With reference to the ion, it might be well to remember Dr. Moses Gomberg. All of his life he had no use for the ionization theory and he contributed greatly to the field of chemistry on the assumption there was no such thing as ions. I do not think we have to worry about whether the oxygen is ionic or not. I think one thing specifically should be brought to your attention and this I think is one of the important contributions of Dr. Rosenqvist. He pointed out what we know as oxygen potentials or what is described as oxygen potentials. I have used this concept for a long period of time and I want to state that if this concept is properly applied, it vitiates much of what we have in the literature, or makes our usual ideas regarding oxidation seem primitive. That one thing is more valuable than almost all the rest of the discussion as a fundamental basis on which to build a reasonable in-

Jan 1, 1952

United Engineering Trustees, Inc. W. D. B. MOTTER, JR., &apos;40 A. L. QUENEAU, &apos;38 ALBERT ROBERTS, &apos;39 The Engineering Foundation GEORGE D. BARRON, &apos;40 F. F. COLCORD, &apos;38 A. L. QUENEAU, &apos;38

Jan 1, 1937

By R. R. Judd, H. W. Paxton

The rate of dissolution of cementite was studied in three low-carbon materials: a zone-refined Fe-C alloy, an Fe-0.5pct Mn-C alloy, and a commercial low-carbon steel. The materials were spheroidized, ad then held isothermally at temperatures above the Al. The isothermal anneal was interrupted periodically by a water quench and the specimens were analyzed by quantitative metallography for the amount of aus-tenite formed during the anneal. The results of this study were compared with an analytical model for the process, which assumes that carbon diffusion in aus-tenite is the rate-controlling step for the cementite dissolution process. The correlation between the model and the experimental data is excellent for the zone-refined Fe-C alloys; however, the Fe-0.5 pct Mn-C alloys and the commercial steel deviate from the calculated model. This deviation is thought to be a result of manganese segregation between the carbide and the matrix. The rate of nucleation of austenite at carbide interfaces was reduced by the manganese addition and enhanced by the presence of ferrite-ferrite grain boundaries. PREVIOUS investigations of the nucleation and growth of austenite from ferrite-carbide aggregates are not entirely satisfying for at least one of several reasons. The most prevalent of these is a lack of quantitative data. Engineering studies have been run on many steels with little control over important parameters such as composition and initial aggregate structure. The data obtained are valid only for material with identical chemistry and thermal history. A more informative approach to the problem of aus-tenitization would be to determine the mechanism that controls the rate of solution of carbide in austenite and how it is modified by alloying elements. This information could then be used to calculate an austeniti-zation rate for any material, provided its composition and structure are known. The object of the present work is to establish the rate-controlling step for cementite dissolution in Fe-C austenite and to investigate the modification of this rate by small manganese additions. The composition and structure of the material used were carefully controlled and all measurements were designed to allow a quantitative analysis of the kinetic process that controls the austenitization rate. A MODEL FOR DISSOLUTION OF CEMENTITE Cementite dissolution has been analyzed mathematically by a model that approximates the material used in the experiments. This model postulates a regular ar-array of identical cementite spheroids with 4 C( diam, embedded in a grain boundary- free ferrite matrix. The analysis provides a detailed description of the dissolution of one carbide spheroid and a generalization of the solution by summation over all the carbides in the material. The carbides may be isolated by defining identical, space-filling cells of ferrite around them. If the cell dimensions are greater than the diameter of the austenite sphere resulting from complete dissolution of the carbide, and no interaction (through diffusion in ferrite) takes place between cells during the dissolution process, the model need concern only one cell, since the solution in each cell is identical. In the experimental material, the dimensions of the cell, the carbide, and the final austenite sphere are approximately 24, 4, and 8 p, respectively; use of the single cell is therefore justified. The experimental observations are made on the austenite nodules that form around each carbide during the dissolution process. The model concerns the growth of these austenite nodules. The attendant shrinking of the carbide can be obtained from the same analysis by an extension of the calculations. Several a priori assumptions are necessary to make the analysis of the growth problem tractable. They are: 1) carbon diffusion through the austenite nodule is the rate-controlling process; 2) local equilibrium exists at all interfaces, 3) the austenite nucleus that forms on each carbide instantaneously envelops the carbide; 4) during the austenite growth process, the diffusion flux of carbon in ferrite is insignificant; 5) a quasi-steady state exists in the austenite concentration field; that is, at any instant during the dissolution process, the austenite carbon concentration gradient closely approximates that for a steady-state solution; and 6) the effects of capillarity on the dissolution rate of the carbides can be neglected. Referring to Fig. 1, a mass balance at the y-a interface for an infinitesimal boundary movement gives: Where rb is the outer radius of the austenite shell, C1 and C are carbon concentrations at the interface in austenite and ferrite, respectively, see Fig. 2, is the diffusion coefficient of carbon in austenite for the concentration of carbon at the interface, and t is time. The fifth assumption permits the austenite carbon concentration to be approximated by the Laplace solution for the spherical case. Therefore, where C(Y) is the carbon concentration at r, and A and B are constants. Local interfacial equilibrium fixes the boundary conditions for the diffusion problem. They are:

Jan 1, 1969