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By R. V. Norris

The anthracite coal trade, with a shipment averaging about 70,000,000 tons per year, differs essentially from other coal business, iii the fact that the larger sizes, comprising about 65 per cent. of the total, are used almost exclusively for domestic purposes, principally during the winter months; and that for proper combustion close-sizing is imperative, so that eight sizes are made; broken, egg, stove, and nut, known as prepared sizes, and pea, buckwheat, rice, and barley, known as steam sizes. These sizes are made in the regular course of preparation, and but little variation in the natural percentage of each is practicable. Thus it is necessary either to work the collieries intermittently, or to dispose of all the varying sizes in their fixed proportions. Unfortunately, the market does not at all times absorb the coal in proper proportions, and rather than interfere with the regular operation of the mines it has been found economical to provide storage for the excess. While the cellars of' the consumers are the great storage-plants of the country, this capacity is not under the control of the trade, except so far as reduced prices during the spring and summer months may tempt the use of this reserve. The same remark applies but with less force to the retail yards, which, though usually small, store in the aggregate a large amount of anthracite. While every effort is made to utilize to the utmost the individual storage-capacity of the country, there still remains the necessity for taking care of the irregularities of the market by the construction of storage-plants under the direct control of the producers. Such plants are usually the property of the railroads, and are situated at points most convenient from a traffic stand-point.

Jan 1, 1912

By Shiou-Chuan Sun

DISCUSSION shiou-chuan sun (author's correction) - Too late for revision, the author realized that the data on stability of froth presented in this paper were not based on the same height of froth column. It is recommended herein that the stability of froth is better calculated by dividing the time elapsed until the froth disappeared by the volume (in cu cm per sq cm) or height of the froth column, and is expressed in sec per cm. The stability data in Tables I, II, and III of this paper were recalculated, in accordance with this recommendation, and the results are shown respectively in the accompanying Tables VII, VIII, and IX. Compared with Tables I, II, and III of the paper, the stability data of Tables VII, VIII, and IX are more reliable in disclosing the persistence of froth of various frothers having different frothabilities. For example, the conclusion given in the paper that the stability of froth is generally proportional to the volume of froth becomes invalid when the stability of froth is expressed in sec per cm, as shown in Tables VII, VIII, and IX. Table VII shows that both the stability and volume of froth increase with the addition of small amounts of n-hexyl alcohol, and that further additions of n-hexyl alcohol decrease the stability but increase the volume of froth. Table VIII indicates that the stability of different types of frothers is governed by their chemical composition and is difficult to predict from their frothability. Table IX shows that the stability of different frothers with similar nonpolar groups tends to increase with a decrease of the hydro- philic property of their polar groups, and consequently with the decrease of their solubilities. This is further substantiated by the fact that the stability of frothers, as listed in Table IX, is generally lowered by the solvent effect of absolute ethyl alcohol. Table IX shows also that the stability of froth tends to increase with the decrease of the volume of froth* provided the frothers used have similar nonpolar groups and dissimilar polar groups.

Jan 1, 1953

By Shiou-Chuan Sun

DISCUSSION shiou-chuan sun (author's correction) - Too late for revision, the author realized that the data on stability of froth presented in this paper were not based on the same height of froth column. It is recommended herein that the stability of froth is better calculated by dividing the time elapsed until the froth disappeared by the volume (in cu cm per sq cm) or height of the froth column, and is expressed in sec per cm. The stability data in Tables I, II, and III of this paper were recalculated, in accordance with this recommendation, and the results are shown respectively in the accompanying Tables VII, VIII, and IX. Compared with Tables I, II, and III of the paper, the stability data of Tables VII, VIII, and IX are more reliable in disclosing the persistence of froth of various frothers having different frothabilities. For example, the conclusion given in the paper that the stability of froth is generally proportional to the volume of froth becomes invalid when the stability of froth is expressed in sec per cm, as shown in Tables VII, VIII, and IX. Table VII shows that both the stability and volume of froth increase with the addition of small amounts of n-hexyl alcohol, and that further additions of n-hexyl alcohol decrease the stability but increase the volume of froth. Table VIII indicates that the stability of different types of frothers is governed by their chemical composition and is difficult to predict from their frothability. Table IX shows that the stability of different frothers with similar nonpolar groups tends to increase with a decrease of the hydro- philic property of their polar groups, and consequently with the decrease of their solubilities. This is further substantiated by the fact that the stability of frothers, as listed in Table IX, is generally lowered by the solvent effect of absolute ethyl alcohol. Table IX shows also that the stability of froth tends to increase with the decrease of the volume of froth* provided the frothers used have similar nonpolar groups and dissimilar polar groups.

Jan 1, 1953

The importance, validity, and acceptability of the Paley Report has become the great issue of the minerals industry. A 12-man panel made up of representatives of various facets of the industry sat down at the 1952 Mining Congress Show in Denver in an attempt to clarify the many opinions currently held concerning the report. Howard Young, deputy administrator of the Defense Materials Procurement Administration, noted that the report had caused more discussion and serious thinking than any other government report, He did not subscribe to the "have not" principle, however, but called it a "need more" principle. Horace M. Allbright, president of U. S. Potash Corp., believes industry should continue to develop the report. He cited the editorial appearing in the September, 1952 edition of Bining Engineering,,, "Don't Let It Die," as expressing his feelings. Edward Snyder, president of Combined Metals Reduction Corp., says the statistics used in the report were used to promote a false international policy. He sees the report as a master plan for the development of the superstate. He also stated that the-ore reserves given in the Pales, report represent only 10 pct of what is actually available.

Jan 1, 1952

By C. E. Sims, G. A. Lillieqvist

The following presentation is more a summation of observations made and conclusions drawn over a period of three to four years than a report of experiments made with a definite objective. Tests were made, of course, to verify hypotheses, but the large mass of data was collected incidental to other work. Lack of space forbids the presentation of more than a few of those data. On the other hand, it is unfortunate that some of the most illustrative examples are not available. The work is admittedly incomplete but is presented in the hope that it will stimulate further work along a similar line. The answers to many of the questions left unanswered by this paper will be found only after long and painstaking study, such as is scarcely possible in a commercial laboratory where problems of immediate importance consume most of the time. It is also hoped that this study will in some measure throw light on some of the dark spots of steelmaking, and help to clarify some of the annoyingly vague references to "good practice." Only cast steel was studied and the conclusions apply to cast steel, but it is the authors' opinion that the quality of rolled or forged steel is largely dependent on the quality of the ingot, which is a casting. Although this discussion deals almost entirely with inclusions, it is important that the reader understand that only those natural inclusions which owe their origin to the solubility of certain materials in steel and to the chemical reactions of steelmaking are considered. Accidental inclusions, such as entrapped particles of slag, nozzle, runner, or mold arc outside its province and are not considered. Such accidental inclusions are always dangerous and should be carefully avoided. Relation between Ductility and Nonmetallic Inclusions In September, 1928, an acid electric furnace was put into operation at the Indiana Harbor plant of the American Steel Foundries. Following the usual practice, aluminum additions were made in order to obtain sound castings. Immediately the problem of low ductility arose. Previous to that time, only open-hearth steel was produced, and, aluminum being used rarely in that steel, there had been practically no trouble with low ductility in the lower carbon range.

Jan 1, 1932

By C. E. Mills

STUDIES of ground movement and subsidence resulting from mining operations cover a broad field. It is also a very important consideration and one that eventually affects nearly every mining operation of any extent. The damage to surface buildings, removal of mine surface plants and replacements of main hoisting shafts must reach an astounding expenditure were these cost figures available over a period of years. Were the subject one of mathematical formula and dependent only on the angle of draw of the walls or country rocks, the problem would be relatively simple. Consideration must be given to depth and extent of workings; structural conditions of the orebodies and enclosing country rock; physical properties of these same rocks, including schistosity, bedding, jointing, slips, dikes, faults, and other structural weaknesses. The method of mining, extent of openings, supporting pillars, and amount of moisture or water underground are other pertinent factors. With such a large number of variable factors, it is apparent that the subject of subsidence is one that must be based largely on experience and results at other mines where conditions are somewhat comparable. In presenting this paper, it is the intent to describe the progressive results of past ground movements at the United Verde, to bring out physical conditions largely responsible for this movement, but to refrain from theorizing on the ultimate extent of subsidence resulting from future mining operations.

Jan 1, 1934

By A. H. Daane, A. S. Wilson

The U-Cr system is of the simple eutectic type with some solid solubility of chromium in r and ß uranium. The eutectic occurs at 20 atomic pet Cr and melts at 859°C. The maximum solubility of chromium in y uranium is 4 atomic pet at the eutectic temperature, and in ß uranium the solubility is estimated to be 1 atomic pet. y-ß and ß-a transformations were found to occur at 737°estimatedt° and 612°C respectively. DURING 1944 and 1945, the U-Cr constitution diagram was studied in this laboratory as a part of a research program on uranium metallurgy in the Manhattan Project, and the work was described in a Manhattan Project report issued in December 1945. This paper is based on that report, which has been declassified. Prior to this study, it had been shown by other Manhattan Project workers that the low Cr-U alloys could be quenched to retain the form of uranium. Experimental The uranium used in this work was massive metal prepared in this laboratory and contained less than 0.1 pct of other elements. The chromium was 200 mesh powder obtained from the A. D. McKay Co. and was found on analysis to be 99.5 pct Cr with 0.3 pct Fe the major impurity. Alloys, weighing 400 to 600 g, were prepared by induction heating the components to 1700°C in slip-cast ZrO, crucibles in a vacuum of 3x103 mm Hg. TO prevent too violent agitation of the melt by the induction field with subsequent crucible breakage and sample loss, the ZrO2 crucible was placed in a graphite crucible, which was surrounded by a layer of powdered carbon insulation 2 to 3 cm thick. Polished vertical sections of the alloys were examined microscopically to confirm their homogeneity. Heating and cooling curves were taken on the alloys by reheating them in ZrO, crucibles to 1200°C and inserting a mullite-protected chromel-alumel thermocouple into the melt by means of a slip seal in the vacuum head of the furnace. A recording potentiometer traced the curves which had a slope of from 3" to 6" per min. Samples of the alloys were prepared for metallo-graphic examination by conventional mechanical polishing techniques followed by an electrolytic polish in an ethylene glycol-phosphoric acid-ethyl alcohol bath. The structure of the alloys was brought out clearly by this procedure so that no further etching was required. Samples for chemical analysis were taken from drillings from the top, center, and bottom sections of the alloys. The uranium was determined by titra-tion with Ce(SO1)2, while the chromium was titrated with FeSO,; the uranium and chromium totaled at least 99.6 pct in all of the alloys prepared. X-ray samples were prepared by filing bulk specimens in a helium-filled glove box and annealing the resulting powder in a zirconium-gettered helium atmosphere. A 114.6 mm diam Debye-Scherrer camera and a Weyland nonsymmetrical self-focusing camera were used with filtered copper radiation to obtain the powder X-ray diffraction data. Results The data obtained in this study have been combined to construct the constitution diagram of the U-Cr system shown in Fig. 1 where the arrests observed in cooling curves are indicated by dots. The liquidus arrest was quite distinct in thermal data taken on alloys in the range 0 to 20 atomic pct Cr. The eutectic arrest was not observed in studies on the 2.5 and 4.5 pct Cr samples but appeared in the 7.5 pct samples, which suggested some solubility of chromium in y uranium. On quenching from 859 °C, the 2.5 pct sample showed but one phase while the 4.5 pct sample contained a small amount of the eutectic along the grain boundaries; see Figs. 2 and 3. From this the maximum solubility of chromium in r uranium has been set at 4 pct. X-ray studies on these samples showed that the r phase was not retained at room temperature by quenching, but in each case a pattern was observed .which has been identified with the ß phase of uranium. Thermal data show the y-ß transformation of uranium lowered to 737°C as a consequence of this solubility. On quenching from the ß range (660°C), precipitation of chromium in the primary uranium is observed in the 2.5 and 4.5 pct Cr samples (see Figs. 4 and 5),

Jan 1, 1956

By Axel Sahlin

Jan 1, 1898

By Paul Pietrokowsky

THE formation of intermediate phases in the solid state reaction of titanium with silicon, germanium. and tin (of subgroup 4B in the periodic table) was the subject of a recent paper.' Further investigation of the system Ti-Sn revealed the presence of a new compound on the titanium-rich side of this phase diagram. Metallographic evidence indicated that this new phase was situated very close to 75 atomic pct Ti. The present work is concerned with the crystal structure of the intermediate phase Ti3Sn. Iodide-process titanium metal, furnished by the New Jersey Zinc Co., was used in this investigation. The metallic tin was received in powder form from Charles Hardy Inc., New York; its purity was 99.80 pct. The titanium rod was machined into small cups 1 cm in diam and about 1.5 cm in height. The tin powder was compacted, melted, and then placed in the titanium cups and alloyed by electric arc melting in a helium atmosphere. The alloys, weighing approximately 8 g, were sealed in evacuated quartz tubes and homogenized at 1030°C for 5 days. This isothermal heat treatment was terminated by a water quench from the soaking temperature. Filings of the alloyed pellets were sealed in evacuated quartz bulbs and quenched from 1030°C into liquid argon. A 14.32 cm Debye Scherrer X-ray camera was used for the powder diffraction study. The copper radiation was filtered. X-ray diffraction data for the compound Ti3Sn are given in Table I. The data were indexed on the basis of a hexagonal lattice. Good agreement between observed and calculated values of the square of the reciprocal interplanar spacings was obtained with the following expression: (1/d)2 = 0.03810 (h2 + hk + k2) + 0.04410 (l2) Unit cell parameters were calculated by the method of least squares from the reflections (20.3), (22.2), (40.1), and (42.1); ao = 5.916 & 0.004A, co = 4.764 ± 0.004A, and co/ao = 0.805. The relatively small size of the unit cell allowed the unambiguous indexing of most of the reflections from planes of large spacing. An inspection of the assigned indices (Table I) reveals the absence of (hk.l) reflections when 1 2n. A systematic absence of this type leads to the space group C 6/mmc or one of its subgroups. Assuming, for a first approximation, a linear increase in the density of titanium with the addition of tin, it follows that there are two stoichiometric molecules of Ti3Sn per unit cell. The density calculated from X-ray measurements is 5.194 g cm-'. At this point it became clear that Ti3Sn might be isomorphous with Mg3Cd. The 6 Ti and 2 Sn atoms were placed in the following positions of space group D4 6h, C 6/mmc: 6Tiin (h): X 2X 1/4; 22;; 1/4;Xx 1/4;x2x3/4; 2x x 3/4; x x 4; with x = 516 2 Sn in (c): 1/3 2/3 1/4; 2/3 1/3 3/4.

Jan 1, 1953

By D. Y. F. Lai, R. J. Borg

Tracer diffusion of Fe59 has been measured in the a-stabilized Fe-1.8 at. pet V alloy from 700° to 1500°C. The activation energies are obtained in both the presence and absence of magnetic order. Furthermore, it is established that diffusion in the alloy is identical to that in pure iron and consequently the values of Do and Q accurately represent the temperature dependence for self-diffusion. The purpose of this investigation is to obtain an accurate estimate of the temperature dependence for self-diffusion in bee iron in both the presence and absence of magnetic order, and, in so doing, to establish the temperature range of the magnetic effect.&apos;" As the temperature interval suitable for diffusion measurements is severely limited in both bee phases of pure iron because of the intervening fee ? phase, the experiments were performed on an a-stabilized alloy containing 1.8 at. pet V. This alloy is bee over the entire range from room temperature to the melting point. Although there have been several independent investigations of self-diffusion of iron in a, iron,1, 3-6 there still exists considerable disagreement regarding the values of Do and Q for the paramagnetic region. The two systematic studies of diffusion in 6 iron6, 7 previously reported are also only in fair agreement; but in view of the extremely small temperature range available for diffusion studies, i.e., 1390o to 1535oC, this is not surprising. It is comparatively easy to obtain accurate values of Do and Q for the a-stabilized alloy inasmuch as measurements can be made over the entire temperature range -700o to 1500°C. However, in order to assume that these same values apply to pure iron requires careful comparison of the data in the a, and 6 regions in both the alloy and pure iron. We have made several measurements in the appropriate temperature ranges and are unable to establish any systematic difference between the diffusion coefficients of iron in pure iron and in the alloy. We therefore conclude that the values obtained for the alloy are truly applicable to pure iron; the complete evidence favoring this conclusion will be discussed later in this paper. EXPERIMENTAL The experimental methods will be given here only in barest detail since they have been thoroughly de- scribed elsewhere.l, 7 The alloy was prepared by induction melting and chill casting under argon. Diffusion samples were machined from the ingot and annealed in hydrogen for several days at 900°C to give an average grain diameter of 1 to 2 mm. The penetration profiles of the tracer were established by a sectioning technique, the residual activity being counted after the removal of each section. The tracer used is Fe59 which emits two high-energy ? rays of 1.098 and 1.289 Mev, respectively; these were detected by a ? scintillation counter equipped with a pulse-height analyzer. For the measurements in the temperature range -700o to 1130°C the samples were vapor-plated with Fe59, encapsulated in quartz under vacuo, and annealed in resistance-heated furnaces which are controlled to ±1°C. The specimens diffused at higher temperatures are prepared as edge-welded couples, the two halves being separated by a thin washer of the alloy to prevent sintering. The diffusion anneal is then carried out by inductive heating under a dynamic vacuum. The temperature is monitored pyrometrically. RESULTS The diffusion coefficients are obtained from the penetration profiles in the usual way using the error-function complement relationship. The results over the entire temperature range are shown in Fig. 1. In the linear region, 900o < t > 1500°C the least mean squares (lms) values of the diffusion coefficients are given by D = 1.39 exp[-(56.5 ± 1) x 103/Rt] cm2/sec [l] The average departure of the measured diffusion coefficients from the values given by Eq. [1]In order to determine whether or not the slope is truly constant over the entire range from 900" to 1500°C, the data are arbitrarily divided into two groups, the first containing values between -900" and 1133°C and the second between -1133o and 1500°C. The lms values for the two groups are given by Eqs. [2] and [3]: D = 0.519 exp[-55.7 x 103 /RT](900° to 1133°C) cm2/sec [2] D = 1.45 exp[-56.7x103/RT] (1133° to 1500=C) cm2/sec [3] Thus, there is no significant difference between the high- and low-temperature segments of the linear region. This not only assures us of the consistency of the values obtained by induction heating as compared to those obtained from the resistance-heated

Jan 1, 1965

By D. C. Sisti

A technical evaluation of actual performance of the preparation facilities at Beaver Consolidated Coal Co.&apos;s Stinson plant is presented, with special emphasis on sulfur reduction in 1-1/2 x 3/8 in., 3/8 in. x 28 mesh, and 28 mesh x 0 size fractions. Evaluation of the heavy medium section of the plant and the treatment of the 28 mesh x 0 coal is included. Design criteria considered for sizing and selection of equipment required to process 1-3/4 million tpy of raw coal on coals of divergent washability characteristics are summarized.

Jan 1, 1976

By FREDERICK W. BRADLEY

MINING is one of the world&apos;s oldest industries and has pioneered the civilization of all new lands. Today, mining is not only one of the essential and basic industries of the world, but it is constantly growing in magnitude and becoming more important. We think that the most of our mining is done in the western portion of this country, but in the old State of Pennsylvania mines and products derived from minerals account for about two-thirds of the entire productive wealth of the State. As our forests are denuded, mining furnishes substitutes for timber in the way of concrete, gypsum board, etc. ; and as civilization increases the wants of mankind, mining is producing metals and minerals of every kind in increasing quantities. For example: during the past ten years the per capita consumption of copper in the United States has increased 49 per cent, to 14.56 Ib.; that of lead has increased 35 per cent, to 12.24 lb.; and that of zinc has increased 70 per cent, to 9.57 Ib. The consumption in European countries is about one-half of this, and there are other countries with large populations where there is virtually no consumption; so when all other countries catch up in consumption with our country, you can see what an enormous demand there is going to be for the metals, even to the extent of having to find substitutes for copper, lead, and zinc.

Jan 1, 1929

By C. S. Barrett, A. Smigelskas

There have been no publications of the deformation and recrystallization orientations of the metal beryllium, yet pronounced textures would certainly be anticipated since it is close-packed hexagonal in structure. Having an axial ratio approximately that of magnesium, beryllium probably deforms by nearly the same slip and twinning mechanisms that operate in magnesium, and the textures are likely to be similar or but slightly different from the magnesium textures. In the tests reported below this is found to be the case; the textures are found to differ from those of magnesium only in the details of the scatter from the average orientation. This report covers not only samples rolled at room temperature, but some rolled at elevated temperatures. Since magnesium has been suspected by some investigators of altering its crystallo-graphic deformation mechanism at elevated temperatures, it was considered possible that beryllium might do so and alter its textures accordingly. No pronounced alterations were found, however. Unfortunately, the theory of deformation textures is not in a state of development that permits one to deduce the deformation mechanism from a knowledge of the textures, which means that the similarity of textures at different rolling temperatures, reported here, cannot be taken as definite evidence that the deformation mechanism is actually the same at all temperatures. The general similarity of the deformation textures of magnesium and beryllium also extend to the recrystallization textures of the two metals, judging by the pole figures for recrystallized sheet presented in this report. Samples were prepared in the form of composite sheets made up of small pieces stacked in a pile. Each piece was trimmed with scissors so that an edge was parallel to the rolling direction, dipped in paraffin, and assembled into the pack by aligning it under the cross hair of a microscope. As the desired orientation was obtained on each piece it was secured in place by touching with a hot wire to melt the paraffin. A stack of ten or fifteen pieces was built up in this way, then trimmed to the shape of a T; the portion to be X rayed was then etched to the shape of a wire about 0.045 in. diam with 6N HCl. This method of shaping the sample is a modification of that used by Bakarian on magnesium.&apos; The absorption of the rays in the sample was so slight that it caused no difficulty in interpreting the films. Exposures were made with a 0.030 in. diam pinhole, using molybdenum radiation (40 kv, 25 ma, Type A film at 5 cm, 2 to 3 hr exposures). With the recrystallized specimens it was found necessary to oscillate the specimen so as to reduce the spottiness of the lines. A range of oscillation of 5" was SUB- cient to produce reasonably satisfactory patterns, though the quality was somewhat inferior to that of the deformation texture patterns, and only two degrees of intensity were read from the arcs on the films. Typical photo-grams for each of the deformation textures and the recrystallization texture are assembled in Fig 1. The pole figures were plotted in the usual way with the intensity of the various portions of the diffraction rings estimated by eye. Seven to nine films were made of each sample and each was carefully read in plotting the pole figures. Typical series included exposures with the beam normal to the rolling direction and at 11, 26, 41, 56 and 71" to the cross direction, plus two exposures with the beam normal to the cross direction, and at 11 and 79" respectively to the rolling direction. The rolling was in each case considered sufficient to develop the final texture: the reduction by cold rolling was 84 pct (from 0.0045 to 0.0007 in. thickness), following prior hot rolling in longitudinal and transverse directions and recrystallization; the reduction by hot rolling at 800°C was 90 pct (0.010 to 0.001 in.), following similar prior treatment; the reduction by rolling at 350°C was 88 pct (from 0.005 to 0.0006 in.) after similar prior treatment. The recrystallization texture was determined on a sample rolled at 350" to a reduction of 88 pct (0.0165 to 0.002 in.) after similar prior treatment, then mounted between steel strips to keep it flat and annealed at 700" in an atmosphere of argon. Discussion of Results The results of the X ray determinations are assembled in the pole figures of Fig 2, 3, 4 and 5 for rolling at

Jan 1, 1950

Aamot, Olav Crone, (Member &apos;29) Met , Det Norske Aktieselskab for Elektrokemisk, Industry, Inc , Oslo, Norway Aase, Glenn D , (Junior Member &apos;40) Met Engr , Western Research Dyv , American Smelt & Refin Co , Selby, Calif Abadye, Henry G., (Member &apos;43) Petr Engr , Sr Prod Engr , Long Beach Oil Development Co, 255 So Santa Clara Ave, Long Beach, Calif Abbott, Clarence E , (Member &apos;04) Civil Engr , &apos;P , Tennessee Coal, Iron & R R Co , 1542 Brown Marx Bldg Birmingham, Ala Abbott, Fred E , Jr , (Member &apos;44) Petr Engr , c/o Arabian American Oil Co , A P 0 816, c/o Postmaster, New York, N Y Abbott, John H , (Member &apos;46) Resident Mgr , Hickman, Williams & Co, 230 No Michigan Ave, Chicago 1, Ill Abbott, R W R., (Member &apos;47) Asst Power Supt, Cananea Consolidated Copper Co , Cananea, Son , Mexico Abbott, Robert H? Jr., (Junior Member &apos;47) Petr Engr, Producing Dept, The Texas Cc, P 0 Box 951, Houma, La Abel, Raymond L., (Member &apos;28) Chem Engr, Prof of Chem Engrg, Univ of Pittsburgh, Pitts¬burgh Pa Abel, Walter D , (Member &apos;32) Min Engr, Div of Corporations, 403 State Bldg , Los Angeles 12, Cal if

Jan 1, 1952

By H. E. Brauer

PART I INTRODUCTION Among the zinc base alloys used for casting in metal moulds, particularly die casting, those alloys containing aluminum usually together with copper, are probably the most widely used. The reason lies in their desirable mechanical properties and the fluidity and absence of hot shortness which facilitate casting. Objection to the use of these alloys may bee raised on account of the rather frequent instances of failure of a type variously described as disintegration, exfoliation and intercrystalline corrosion. In all cases which have come to our knowledge where the conditions of service were known, failure has occurred upon exposure to moisture and heat varying from a hot moist atmosphere to high pressure steam. Moreover, the alloys described in all such cases contain impurities or alloying metals which the present authors have found to have enormous accelerating effects on intercrystalline oxidation in zinc-aluminum alloys. Thus Williams1 experimented with an alloy containing a large percentage of tin exposed in Cuba. Bauer2 reports failure of a boiler plug containing over 1 per cent. of impurities, doubtless lead for the most part, and Rawdon&apos;s3 examinations were made on an alloy containing 0.4 per cent. lead also exposed to the Tropics. No thorough investigation of the nature and cause of this type of failure has been made heretofore, so far as the authors are aware, although Rawdon correctly classified it as internal oxidation. We have, therefore, carried out this investigation in order to determine the exact nature of this phenomenon and the effect of other metals present as impurities or intentional additions in accelerating or-retarding it. In the appendix we include a considerable amount of data which may be valuable in selecting an alloy combining good physical properties with high resistance to oxidation.

Jan 8, 1922

By Gilbert Soler

PRODUCERS of alloy steels recognize the importance of chemical composition in rela-tion to the hot-working properties and the typical surface defects found in their prod-uct. Each analysis of steel has its own peculiar characteristics. Under conditions of standard mill practice each analysis is susceptible to certain types of defects. Mill practice must be varied to obtain the best combination of surface and internal quality in the product. Chemical composition influences the cast structure and crystallization character-istics of the ingot. It also determines the rate of heating and cooling, the plastic hot-working range, and the phase structure of the steel at various temperatures as well as the tendency toward scale formation and decarburization, This paper endeavors to emphasize the manner in which chemical composition affects the various properties of steel, and to indicate the relative impor-tance of these factors in relation to the hot-working properties and surface charac-teristics of killed steel. The influence of chemical composition may be outlined as follows: I. Effect on the cast structure of steel, including: A. State of deoxidation, and type of inclusions. B. Gas content of steel. C. Freezing point and melting point of steel.

Jan 1, 1940

Abbey, Robert Graham, District Mgr.. The W. W. Sly Mfg. Co., 50 Church St.. New York, N. Y. &apos;21 Abbott, A. N., Mines Supt., Mazapil Copper Co., Ltd Concepcion del Oro, Zac., Mexico. &apos;28 Abbott, Argyle Campbell, Dev. Dept., E. I. du Pont de Nemours & Co., Wilmington, Del. &apos;29 Abbott, Clarence E., V.P., Charge of Raw Materials, Tenn. Coal, Iron & R. R. Co., 1242 Brown-Marx Bldg.. Birmingham, Ala. &apos;04 Abbott, Louis Benjamin, Chief Engr., Elkhorn Div., Consolidation Coal Co., Jenkins, Ky. &apos;16 Abeel, George H., Jr., Cons. Min. Engr Box 427, Santa Monica, Cal. &apos;18 Abel, Albert E., Asst. to Pres., Valley Mould & Iron Corp Hubbard, Ohio. &apos;19 Abel, Raymond L., Asst. Prof., Dept. of Pet. Refin., Univ. of Pittsburgh Pittsburgh. Pa. &apos;28 Abeel, Leo M., Research Engr., Sullivan Concentrator Chapman Camp, B. C., Canada. &apos;31

Jan 1, 1932

By A. B. Kinzel

IT is with sincere appreciation and a deep sense of responsibility that I accept the honor of delivering the Howe Memorial Lecture. In our time metallurgical research has delved into phenomena ever more complex in character, so that much of this work must be done by organized teams. We, of the Union Carbide and Carbon Research Laboratories and the Electro Metallurgical Co., are truly a team and the lecture which is to follow is due largely to a few of the outstanding members of that team here cited and well known to you in their own right: John Lamont, whose work deserves special recognition, Walter Crafts, William Forgeng, William Binder, Robert Fowler, and David Swan. Thus it is a fitting pleasure to accept this honor for the research institution which I represent. Henry Marion Howe was truly inspirational and his breadth of view was such that he could use either the method of rigid proof or implied mechanism to further the broad understanding. But perhaps his outstanding characteristic was his sound judgment and his ability to weigh observations. The following is replete with observations and judgment as to their implication. We would like to feel that it is the type of dissertation which Howe, himself, would have enjoyed. Chromium carbides in stainless steel have, in a sense, been a key to the progress of the stainless steel industry. Brearley&apos;s first stainless steel, containing approximately 0.30 pct C, was successful because it differed from previous Fe-Cr alloys in that the carbides were present in such amount and form as to provide the stainless and cutting properties. Again, the 12 pct Cr die steels with a carbon content of 1.5 pct or more achieved their properties by virtue of the form and distribution of the chromium carbides. In the straight chromium steels with from 12 to 25 pct Cr, chromium carbide amount and distribution are once more the determinants with respect to physical properties and corrosion resistance. In each of the cases just cited the role of the carbide is in essence similar to that of iron carbide in carbon steel and this story is so well known that we need not dwell on it. In the austenitic stainless steels, the chromium carbide plays only a minor positive role with respect to mechanical properties, and for corrosion resistance the objective is to achieve maximum homogeneity. Thus the tendency has been to ever lower the carbon and correlatively the chromium carbide content of the austenitic steels. Today stainless steels with carbon less than 0.03 pct are beginning to find their true place, and this is being aided in major degree by our Company&apos;s introduction of a new type of ferrochromium having negligible carbon. The role of carbon in these austenitic steels has long been recognized, and a truly enormous amount of research has been devoted to the subject. Much has been learned about corrosion resistance, but the attack on the problem from the point of view of carbide precipitation and the nature thereof has left much to be understood. Because of its fundamental importance to the future of stainless steels, further study of the mechanism of carbide formation in austenitic steels is of interest. In the 18-8 Cr-Ni steels, which are typical of the austenitic class of steel, carbon is soluble in the aus-tenite to the extent of approximately 0.15 pct at 1832°F (1000°C). The solubility becomes less with decreasing temperature, so that at 1652°F (900°C) the solubility limit is approximately 0.06 pct, at 1472°F (800°C) it is about 0.03 pct, and at 932°F (500°C) it is approximately 0.01 pct. Obviously this is the ideal situation for the classical precipitation phenomenon. It is true that, in a sense, 18-8 austenite resulting from annealing is in a metastable state, and the phenomenon of ferritic precipitation from this austenite should be considered. However, in a so-called fully austenitic 18-8 steel this is secondary with respect to the carbides and can be neglected. A wide variety of chromium carbides is known. Only one chromium carbide, Cr,,C,,, frequently referred to as Cr,C, has been found in unmodified low carbon 18-8 steel. This, fortunately, simplifies the study of the precipitation phenomena. For this study an A.I.S.I. type 304 stainless steel of commercial quality was selected. Analysis showed C 0.07 pct, Mn 0.49 pct, Si 0.33 pct, Ni 9.34 pct, Cr 18.91 pct, and N 0.033 pct. As expected with this composition the steel was homogeneous after annealing except for a few pools of ferrite, which were disregarded in that their location was not coincident with subsequent carbide precipitation areas. On corrosion testing the steel was normal and typical of good commercial quality. Specimens 1/2xlx4 in. were annealed and subjected to precipitation heat treatment for 100 hr at 100" temperature intervals from 1000° to 1500°F and for ten selected intervals ranging from 5 min to 64 hr at 1300°F. These blocks were then cut into small pieces for metallographic studies. The first problem

Jan 1, 1953

By T. E. Leontis

The specific effect of various rare-earth metals on the room- and elevated-temperature properties of magnesium has been evaluated. Alloys containing didymium exhibit the highest tensile and compressive strengths at room and elevated temperatures. All the rare-earth metals increase the creep resistance of extruded magnesium at temperatures in the range of 400° to 600°F, but the degree of enhancement depends on the temperature and on the concentration of the added metal. THE effects of rare-earth metals on the properties of sand-cast magnesium were discussed in some detail in earlier paper by the author.&apos; The present paper deals with the effect of the same alloying elements on the properties of extruded magnesium. This investigation also had as its aim the development of a wrought alloy having a better combination of room-temperature strength and ductility and elevated-temperature strength and creep resistance than is found in magnesium-Mischmetal-manganese alloys, which have been reported earlier.2-5 The only known attempt to study wrought magnesium alloys containing pure cerium instead of Mischmetal was made by Mellor and Ridley.6 They found that in the form of rolled bars there is a definite, optimum cerium content for creep resistance at 200 °C and that the creep resistance of these alloys at 200 °C is significantly Improved by heat treatment at 550" to 580"C. In the present investigation the compositional variation in mechanical properties of the following alloy systems is presented: I—magnesium-Misch-metal. 2—magnesium-cerium-free Mischmetal. 3— magnesium-didymium. 4—magnesium-cerium. 5— magnesium-lanthanum. Alloys containing predominately praseodymium are not included in this series because of the lack of this material. Experimental Procedures The alloying ingredients used in preparing the alloys described herein are the same as those reported in the earlier paper.&apos; Cerium-free Mischmetal is the rare-earth mixture remaining when the cerium is removed from Mischmetal, which contains all the rare-earth metals as they occur naturally in mon-azite sand, the ore from which Mischmetal is produced. Removal of both cerium and lanthanum from Mischmetal leaves what is commonly called "didym-ium," consisting predominantly of neodymium and praseodymium. Although the composition of the particular batch of each metal used may differ somewhat from the analysis presented previously, these differences are not great enough to warrant repeat- ing the specific composition of each material. The electrolytic magnesium used as the starting material in these alloys has the same typical analysis as that given in the earlier paper.&apos; The alloys were prepared in small laboratory melts applying all the necessary precautions for alloying rare-earth metals with magnesium described by Marande.&apos; Most melts were large enough to cast one 3 in. diam billet 10 in. long. In a few cases, particularly the didymium-containing alloys, the lack of sufficient amounts of the rare-earth metal limited the size of the billet to 6 to 8 in. All billets were scalped to a diameter of 2 15/16 in. and faced to a length of 9 ¼ in. as limited by the size of the extrusion container. The alloys were extruded into ½ in. diam rod on a 500-ton direct-extrusion press using a 3 in. container. The details of the extrusion step are: billet preheat, 925°F (2 hr); container temperature, 900°F; die temperature, 900°F; extrusion speed, 10 ft per min; reduction ratio, 36:1; and percent reduction, 97.3. The lower melting point of alloys containing didymium&apos; necessitated reduction of the extrusion speed to 5 ft per min in order to prevent hot shorting during extrusion. Adequate lengths were cropped from both ends of each extruded rod to assure that all the material used for tests was extruded under uniform conditions. Tensile and compressive properties at room temperature are reported in the several conditions of heat treatment. The ASTM designations are used to distinguish these conditions as follows: T5—Direct age at 400°F (16 hr) T4—Heat treat at 950°F (4 hr) for alloys containing didymium T4—Heat treat at 1050°F (4 hr) for all other alloys T6—T4 + age at 400 °F (16 hr) The lower heat-treating temperature for alloys containing didymium is necessitated by the lower melting point of these alloys. All heat treatments were conducted in electrically heated, circulating-air furnaces. A protective atmosphere containing 0.5 to 1.0 pct sulphur dioxide was used for the high temperature heat treatments. Tension and creep specimens 6 Yz in. long and compression specimens 1½ in. long were cut from the extruded rod. A reduced section of ? in. diam was machined on the tension specimens, whereas on the creep specimens a reduced section of 0.450 in. diam

Jan 1, 1952

By Richard A. Kesler, W. Michael Reed

Grindability or resistance to fine comminution is the fundamental issue involved in the application and sizing of grinding mills. Over the years, many theories for determining the amount of energy required for size reduction to a specified fineness have been proposed. The first published theory, introduced by P. R. von Rittinger in 1867, stated that the energy required for crushing or grinding is proportional to the increase in surface area produced. The surface area of a ton of particles of diameter (d) is proportional to (1/d), therefore the work input per ton for particles of diameter (d) would be proportional to (1/d). The second significant theory is that of F. Kick published in 1885. Kick said that the work required is proportional to the reduction in volume of the particles concerned. If (f) is the diameter of the feed particles and (p) is the diameter of the product particles, the work input relationship would be: [W =Ioq(f/p) 9] The "Third Theory" of F.C. Bond, introduced in 1951, states that the work required is proportional to the length of the new crack tips produced in particle breakage during crushing or grinding and equals work represented by the product minus that represented by the feed. This takes the form: [W=]

Jan 1, 1982