Search Documents

By Walter Maxson, Fred C. Bond

Since the last publication of tabulated results of grindability tests by the authors1 the total number of ball-mill tests made has more than doubled, and rod-mill tests have become increasingly important. Nearly all of the standard closed-circuit ball-mill and rod-mill tests made to date are included in the present tabulation, together with additional information not previously published. Standard Ball-mill Grindability Tests The method of conducting these closed-circuit tests has not been altered. A sample of the ore or other material is stage-crushed in rolls set at 1/16-in. opening with a 6-mesh screen until all of it has passed the screen. The combined minus 6-mesh screen undersize is mixed, sampled and screen-analyzed, and its apparent specific gravity is determined by packing and shaking in a standard container, and weighing. The apparent specific gravity ordinarily is about 60 per cent of the true specific gravity. The unit volume present in the mill .in all tests is 700 C.C. of the packed minus 6-mesh roll product, and the number of grams occupying 700 C.C. is the unit test weight. This weight is placed in the mill dry, ground for the number of revolutions estimated to be necessary, discharged, and screened mechanically in three testing sieves of the mesh size at which the test is to be conducted. The oversize, or circulating load, is weighed, sufficient fresh minus 6-mesh feed is added to bring the total weight up to that of 700 c.c., and the charge is returned to the mill for a second grinding period. All standard ball-mill tests are conducted at 250 per cent circulating load, and the number of revolutions in the mill necessary to obtain this circulating load at any grinding period is estimated from the results of the previous period. The number of net grams of screen undersize produced per revolution of the mill approaches its final equilibrium value after several grinding periods, and this is recorded as the grindability, or relative ease of grinding, of the sample. A cylindrical ball mill, 12 by 12 in. inside, with a smooth lining and rounded corners, is used with a revolution counter Under standard grinding conditions it runs at 70 r.p.m., and contains a charge of 285 iron balls weighing 20,125 grams, ranging from I½ to ¾ in. in diameter. Tests are conducted at the mesh size to which the ore is to be ground in practice, and the capacity- to be expected from a given mill is calculated by comparing the grindability with that of a similar standard ore whose performance characteristics arc known. No satisfactory general formula for convertillg standard grindabilities into corn; rnercial power-consumption values has been

Jan 1, 1943

By H. R. MacMichael

THE rate, at which zinc ore can be sintered varies widely, depending on the characteristics- of the ore. A coarse table concentrate already roasted to a very, low sulfur content may sinter at the rate of-8.6 tons per hour per 100 sq.ft., of wind-box area. At the other extreme is a finely ground flotation, concentrate, with considerable sulfur remaining in the ore as delivered to the sinter plant, for which corresponding rate may be only 3.8 tons per hour, or 44 per cent of the first rate. A usual, size-of sinter machine in zinc plants is one with wind-box 3.5 by 33 ft., making 115.5 sq. ft. of area, and using a 100-hp. fan. Such a machine with the coarse concentrate described would produce 10 tons of sinter per hour; and with the fine, flotation concentrate would produce only 4.4 tons per hour. Ore of the latter class is -sintered at the plant of the Mexican Zinc Co. (subsidiary of the American Smelting & Refining Co.) at Rosita, Mexico. ; Consequently, an unusually large sintering capacity per ton of ore to be treated is required, and provision was made for treatment of 9000 tons concentrates per month, preroasted to. 5 to 9 per. cent sulfur. For this duty three of the usual 33-ft., 100-hp. sintering machines would have been needed.. Considering the auxiliary equipment. necessary for proportioning, handling, and mixing charge, as well as means for disposition of the sinter, a three-unit plant becomes-very complicated, especially if the capacity should be doubled later. After much study it was decided to build a single sintering unit of ample capacity for the work, the plant thus being much simpler and cheaper in construction as well as in operation (Figs. 1 and 2). The sintering machine was built with a wind-box 5.25 by 60 ft. (315 sq. ft.) and provided with a 300-hp. fan.

Jan 1, 1932

By L. H. Reiss, J. Groult, L. Montadert

Many fields, where the reservoir is composed of sandy layers, show great complexity because of the lack of continuity which results from a particular type of seditnentation. This complexity may be a factor of prime importance in studying reservoir drainage, especially in the displacement of oil by water or gas. This paper describes work done to study this heterogeneity and its influence in the exploitation of the fields of the lower Devonian sandy shale in the Fort Polignac Basin in Algeria. Two lines of research were followed. 1. On the border of the hasin, a detailed sedimentological study established rules for the distribution of the sandy zones, which furnished hypotheses necessary for devising a geologic model of the reservoir. A certain number of sandy zones were recognized and characterized by their sedimentation. 2. In the interior of the basin, production logs deterrnined the productivity profile. Study of these point data led to a description of the distribution of oil and gas phases in the reservoir so that the continuity of different zones between the wells could be deduced. These techniques were applied to the study of the lower Devonian sandy shale reservoir of Zarzaitine. The sedimentological study guiding the correlation work, with the aid of the conventional static data (logs and cores), led to the establishment of a useful geologic model of the reservoir. Production logs, thanks to the contribution of dynamic elements, permitted corrobora-tion of the model. In particular, the logs confirmed separation, as proposed by the sedimenfological study, of two comparable sandy groups by revealing the continuity of a thin impermeable layer, invisible on conventional logs. INTRODUCTION The heterogeneity of sandstone reservoirs considerably influences the distribution of fluids, as well as their displacement in the reservoir. Rational development of a field and maximum recovery of hydrocarbons are closely tied to an exact and precise knowledge of the intern architecture. We will try to show how it is possible I refine the knowledge of the reservoir by seeking qui different and varied sources of information. Productic geology, the group of techniques which leads to idei concerning reservoir architecture and its geometric descris tion, furnishes the framework of this undertaking. Of these techniques. those used so far are qui varied. Descriptive as well as quantitative, the technique partake as much of classical geology as of reservo science. They include (1) geologic studies of the enti basin, especially a stratigraphic and sedimentologic study of the reservoir member at its outcrop, detail1 observations of cores taken from wells and, finall oriented cores; (2) interpretation of geophysical dat (3) interpretation of electric logs, nuclear and sonic lo and dipmeters: (4) petrophysical study of the reserve on samples taken from wells; (5) reservoir dynamics (stu of the production history and, in particular, interpretati of production logs); and (6) analysis of the fluids a thermodynamic studies. This incomplete list could also include (1) study the migration of the fluids, (2) geochemistry (chemi characterization of the oils), (3) regional hydrodynamii and (4) injection and detection of tracers. If most of these techniques are classic, attention i nevertheless be called to applying two of the techniql in conjunction: the study of the reservoir member its outcrop and the study of dynamic distribution fluid phases in the reservoir through use of product logs, used in the study of the sandy shale reservoirs the lower Devonian of the Fort Polignac Basin, Alge in particular, their application to the study of the 2 zaitine field where they have led to a new and satisfact model of the reservoir will be described. THE ZARZAITlNE FIELD Situated in the interior of the Fort Polignac 8 about 150 kilometers (km) north of the lower Devol outcrop, the Zarzaitine Devonian oil field was discovc On Ian. 28, 19'' (Figs. 1 and l8). It is in the fern a monocline truncated by two perpendicular fault: great throw, with a downthrown section to the north\: Erosion at the top of the structure has caused one of it to disappear. It is a reservoir with a gas cap an m coefficient on the order of 1:5, open to the aql On the northeast, with the oil zone covering an Of 106 sq km and with initial oil reserves of 222 mi

Jan 1, 1967

SULPHUR is a very well known mineral and apparently is produced in many places. It is engendered from an unctuous earthy and powerfully hot substance so that it is considered among experienced workers to bear a resemblance to the element of fire. This is called by these same men the masculine seed and the prime agent of Nature in the composition of metals. Because of its great dryness and heat it has, as experience shows, affinity with fire, which is easily introduced in it when brought near. And once introduced the fire is quenched only with difficulty before its unctuousness is wholly consumed. Although it shows itself to us of such a hot and dry nature, we must not think that sulphur is a substance so pure that it can stand by itself or that in order to attain its form it has no need of its portion and part of moisture like every other mixture. This is shown to us by its easy fusion, in which it resembles the metals. This thing is found in many places and in many varieties of color; some are white, others yellow and orange, and some others between gray and black. They also say that some red is found. It does not occur in that arrangement of veins like other minerals, but most liberal Nature makes whole mountains of it as is seen in the Aeolian islands, [26] near Sicily, in Etna or Moncibello which is the island of Sicily, at Pozzuoli, in the territory of Rome, in the Sienese domain at San Filippo, and in many other similar places. It is a substance that has a very strong odor and is most unpleasant to smell. It is very fixed in its substance so that it is never corrupted by time, by boiling, or by the moisture of water. For although it may stand in water a long time it never becomes soft nor does it ever decrease or increase in weight. It is brittle and easily crushed very fine. It is said that a trace of garlic rubbed on the mortar helps anyone who wishes to reduce it to impalpability. Its ores have a form more like earth than like stone although they some¬times have the color of a certain soft rock. The places where it is found are easily recognized by the great odor that they give out and by the very hot baths that often spring forth near it. What I have already told you about quicksilver that is contrary to the

Jan 1, 1942

By C. E. Brukl, E. Rudy, St. Windisch

The ternary-alloy system Nb-Mo-C was investigated by means of X-ray, melting point, DTA, and metallo-graphic techniques; a complete phase diagram for temperatures above 1500°C was established. Above 1960°C, niobium monocarbide and the cubic (Bl) high-tenzperature phase in the Mo-C system form an uninterrupted series of solid solutions. The ternary range of the pseudocubic q MoCl-, is very restricted. Dimolyb-denum carbide dissolves up to 44 mol pct Nb2C (2240°C), whereas the maximum solid solubility of MO2C in Nb2C does not exceed 5 mol pct. The order-disorder transformation temperatures in Mo2C and Nb2C are lowered by the mutual metal exchanges. Six invariant (p = const) reactzons occur in the ternary system; three correspond to class 11-type four-phase reactions involving a liquid phase, one to a class I (eutectoid)-type, and two further isotherms are associated with limiting tie lines. The results of the Phase diagram investigation are discussed, and the thermodynamic interpretation identifies the low relative stability of the binary sub-carbides in conjunction with the large stability diflerences between niobium and molybdenum carbides as the cause for the formation of a stable equilibrium between the monocarbide and the metal phase in the ternary reson. Due to their refractoriness, the carbides of the high-melting transition metals have received increased interest in recent years as base materials in composite structures for aerospace applications at high temperatures and for the development of self-bonded cutting tool materials; other novel fields of application include power reactors, where operation at high temperatures becomes essential for attaining high power efficiencies. In these applications, the increased reaction rates at high temperatures require a close consideration of the chemical interactions between the alloy constituents. As a consequence, a detailed knowledge of the phase relationships in the alloy systems is required in order to provide a sound basis for developmental -type work. Partly as a result of the considerable experimental difficulties associated with the investigation of this high-melting alloy class, no complete studies of ternary metal-carbon systems have been performed until recently. Even the high-temperature phase relationships in the binary transition metal-carbon systems have been delineated only during the past few years to a degree of accuracy required for a more detailed study of ternary or higher-order alloys. In recent investigations of binary and ternary systems of refractory transition metals with carbon, boron, and silicon,&apos; alloys from the ternary systems Nb-Mo-C became of interest because of the demonstrated possibility2,3 of obtaining compatible composites based on metal + monocarbide combinations. In the meantime, however, studies in other, but related, ternary metal-carbon systems, such as Ta-W-C, have shown that the solid-state equilibria may change significantly toward higher temperatures (>2000°C), and that extrapolations based on low-temperature equilibrium data are, in general, not very reliable. Although the lower-temperature (<2000°C) phase relationships in the Nb-Mo-C system are similar to those found in Ta-W-C, a cursory thermodynamic analysis of the equilibria indicated4 that complete solid-solution formation between Mo2C and Nb2C should not occur at higher temperatures. The present work was conducted in order to experimentally verify these expectations and, in addition, to provide phase equilibrium data in the melting range of the alloys. In the boundary systems, niobium and molybdenum are known to form a continuous series of solid solutions.576 The continuous solubility was also confirmed by Kornilov and Polyakova,7,8 who also observed a minimum melting point at 22 at. pct Mo and 2345°C. The phase diagram investigations of the Nb-C system by Storms and Krikorian9 and Kimura and Sasaki10 were recently supplemented by Rudy et al.11,12 The system contains a high-melting monocarbide with the B1 structure, Table I, and a subcarbide, Nb2C, which exists in at least two different states of sublattice order at low temperatures47"-&apos;3 and a disordered state above approximately 2500°C.11,12 The melting-point measurements by Rudy et al .11,14 are in close confirmation of the data by Kimura and Sasaki.10 The rather complex phase relationships in the Mo-C system were only recently Clarified.15,18 The system is characterized by three congruently melting, intermediate phases, MozC, ? MoC1-x and a Mol-,, Table I, of which only Mo2C is stable at temperatures below 1650°C. Substoichiometric MozC exists in several states of sublattice order which interconvert in homogeneous phase transitions. Hyperstoichiometric compositions cannot exist in the ordered state. Upon cooling through a critical temperature range, the

Jan 1, 1968

By J. L. SILSBEE

THE existence of large saline deposits in that flat arid basin, known as the Great Salt Lake Desert, has long been recognized, but the extent and great commercial value of these deposits has not been generally appreciated. That a very considerable interest has recently developed is indicated by the fact that about 263 prospecting permits-each covering about 2560 acres of land in this Desert-have been issued under the Potash Leasing Act (Act of 0ct. 2, 1917-40 Stat. 297), and patents granted on approximately 40,000 acres. It is reported that patents and leases have, or will be applied for on an additional 80,000 acres in the basin. , Tests have been made on probably less than 10 per cent. of the total area of the basin, but these have proved a practically unlimited supply of brine available -sufficient, probably, to. furnish the potassium and magnesium chloride requirements of this country for at least a hundred years.

Jan 1, 1925

By John E. Hockett

The need for basic information about the relationship between resistance to dejormatim (flow stress), temperature, strain, and strain rate, for the solution of metal-fovming problems, is pointed out. Some early attempts to satisfy the need are mentioned. A brief description of a machine called a Cam Plastome-ter is given, including the processing of results of testing in the machine. Next, a program of testing aluminum in compression over a wide range of constant true-strain rates (0.1 to > 200 sec-1) and a modevately broad range of temperatures (223° to 673°K) is described. The results of this program are presented in the form of true-stress us true-strain cuvves. The data in these curves are presented in the form of relationships between true stress and true-strain rate in sermilogarithmic and log-log form, as functions of temperatures—for several true strains. Finally, true stress, for a given true strain, is displayed as a function of log true-strain rate and absolute temperature, i.e., a surface. SINCE the publication of Hill&apos;s book1 on plasticity theory, increasingly rapid advances have been made in analytical solutions to forming problems. Methods in common use are the slab method, the uniform deformation energy method, the limit analysis method, the slipline field method, and the semiexperimental method called "visioplasticity". These approaches, and examples of their use, are covered in detail by Johnson and Melloor2 and by Thornsen, Yang, and Kobayashi.3 unfortunately, practically all of the above work has had to rely upon a number of basic assumptions, e.g., homogeneous deformation and Coulomb (sliding) friction; and virtually no dependence of the mechanical behavior of the deforming metal, upon temperature, strain, and strain rate, has been included. It is with this dependency that this paper is concerned. Loizou and sims4 looked at the above relationship for lead using both a constant compression-speed Cam Plastometer and a constant true-strain-rate Cam Plastometer designed by Orowan.5 Alder and Phillips8 investigated the relationships for aluminum, copper, and steel in the same constant true-strain-rate Plastometer. cook7 tested twelve steels at temperatures ranging from 1170° to 1473°K in the same Plastometer over a range of true-strain rates of 1.5 to 100 sec-&apos;. The writer tested commercially pure aluminum at room temperature and depleted uranium at temperatures from 573" to 873°K at strain rates from 10-3 to 1.0 sec-1 in a Cam Plastometer designed and built at Los Alamos Scientific Laboratory.8 During the past few years the Cam Plastometer at Los Alamos has been used for a number of relatively minor experiments on brass, mild steel, duralumin, Armco iron, and depleted uranium. In this time the machine has been continually modified and improved. It was decided to utilize fully the wide range of strain rates now available with the machine by testing a material about which some information was already in the literature. Prior work on commercially pure aluminum6 appeared to deserve confirmation and expansion. So it was decided to explore the resistance to compression of this material over the currently available range of constant true-strain rates and a convenient range of temperatures. This paper is a report of that exploration. I) EQUIPMENT AND PROCEDURE The principal item of equipment is, of course, the Cam Plastometer, the working part of which is shown in Fig. 1. In this figure, an aluminum specimen may be seen between two tungsten carbide platens, in position for a room-temperature test. Above the upper carbide platen is a load cell. Output from the load cell is amplified and applied to a galvanometer in a recording oscillograph. Auxiliary equipment includes the oscillograph, the control console, a counter, and a time-mark generator. The counter is first used in adjusting the time base. Then it registers the count, of sixty pips per revolution generated by the rotating cam, in each second. These pips and the time base are also applied to galvanometers in the oscillograph. Thus, load, time, and cam position are recorded simultaneously. Heating of a specimen prior to compression is done

Jan 1, 1968

By George Stern

THE purpose of this investigation has been to study one method for making ordinary carbon steel by the powder metallurgy technique. This method con- [ ] sists of pressing and subsequently sintering an intimate mixture of iron powder and graphite in a suitable atmosphere to produce a steel. Two types of iron powder have been studied, electrolytic and reduced. The sintering of mixtures of loose carbonyl iron and carbon black for several hours at 1000° to 1100°C. (1830° to 2010°F.), followed by forging of the resultant ingots to bars, has been studied by Offermann.1 Trail&apos; has reported on the properties of hydrogen-reduced and electrolytic iron powders, both containing 0.5 per cent carbon. Methods for making steel bodies by powder metallurgy have also been sug- [ ] gested by Goetzel,3 who recommends the addition of steel powder of well-defined carbon content to pure iron powder, and Volterra,4 who suggests compaction and subsequent sintering of a mixture of steel powder, iron oxide and iron powder. Additional methods for making steel bodies by powder metallurgy have been reported in the literature by Koehring,5 Allen,6 Balke,7 Wulff,8 Boegehold,9 and others. 10,11,12 It is felt that the method of pressing iron powder and graphite, as described herein, lends itself best to production because of the lubricating action of the

Jan 1, 1946

By R. W. Hertzberg, F. D. Lemkey, J. A. Ford

The effect oj unidirectional solidification upon the microstructural, crystallographic and mechanical characteristics of high -purity A1-A13Ni eutectic a1loy specimens has been investigated. varying the growth rate caused the morphology of the Al3Ni phase to gradually change from platelets to rods. In addition, the effect of increasing the rate of planar liquid-solid interlace movement is to decrease the Al3Ni rod diameter and inter-rod spacing. The platelets and rods grow approxinlately parallel to each other and are aligned in a faulted substructure. A single unique crystallographic relationship hetween the A1 and ASNi phases was found and may he descrihed as : interface {001} Al3Ni {331} Al and growth direction <010>Al3Ni || <110> Al The platelet interlace and the aligned rows of rods are both uniquely defined by the above statements. Alignment of the Al3Ni phase by unidirectional solidification has given rise to a threefold increase in strength over that exhibited by specimens with an as-cast microstructure. These results illustrate the possible use of this eutectic alloy as a zohisker -reinforced structure. It has been demonstrated by Winegard et al.,1 Kraft and Albrighht,2 Chilton and winegard,3 Chad-wick,4 Yue,5 Tiller,6 and others that a planar liquid-solid interface may be established in binary eutectic alloys by proper control of heat flow during the solidification process. The unidirectional movement of such an interface results in a eutectic mi-crostructure consisting of an essentially parallel array of the two phases over an entire ingot. Two dominant phase micromorphologies have been produced using this technique: that of substantially parallel alternating lamellae of each phase or long thin parallel rods of one phase imbedded in a continuous matrix of the other phase. While not all eutectics can be controlled in this manner, it has been shown that a "normal" eutectic may, under properly controlled conditions of purity, liquid and solid thermal gradients, and solidification rate, be forced to solidify with essentially parallel phase particles. Kraft and Albright2 have demonstrated that, when the growth rate is too rapid or when the thermal gradient in the liquid at the growing interface is too low, a layer of constitutionally supercooled liquid, formed by impurity build-up, will stabilize a cellular rather than a planar interface. The unidirectional movement of the cellular interface forms the macrostructure of eutectic colonies.7 Chilton,3 Chadwick,4 and Tiller,&apos; studying eutectics of zone-refined Pb-Sn, A1-CuAl2, Zn-Sn, and Cd-Zn systems, noted that in the absence of impurities a planar interface is stabilized even when rapidly solidified. Although it is generally agreed that impurities break down a planar interface to form the eutectic-colony macrostructure, at present the origin of the varying micromorphologies (i.c., plates or rods) in a given normal eutectic alloy is not completely understood. Tiller8 has predicted that the micromor-phology produced may be dependent on solidification rate (i.e., at fast rates a rodlike structure is preferred whereas at slow rates a lamellar structure should form), and also suggested that rod formation may be favored at large phase-volume ratios. yue5 has experimentally verified Tiller&apos;s prediction in the eutectic Mg-Mg17Al12 by observing a lamellae-to-rod transition with increasing growth rate where the phase-volume ratio was approximately 2.3:1. However, Hunt and chilton9 more recently unidirec-tionally solidified over a wide range of growth velocities six different eutectic systems with phase-volume ratios between 12:l and 2.7:1 and observed no lamellae-to-rod transition. chadwickl0 has proposed that the change in micromorphology in some eutectic alloys is due entirely to the presence of impurities and further states that the lamellar structure is the characteristic structure of pure eutectic alloys even when the phase-volume ratio is as great as 12:l. Kraft11,12 has shown that the parallel lamellae in certain eutectic alloys assume a unique crystallographic relationship during unidirectional growth. This preferred crystallography may be developed

Jan 1, 1965

By AIME AIME

ONE unique distinction can be claimed by the new Chairman of the Coal Division, C. E. Lawall: he is the only A.I.M.E. member also a university president, though several engineering colleges are headed by Institute members. Only last year he was promoted from director of the school of mines to president of West Virginia University, following a term as acting head of that institution of over 3000 students. He was born in Catasauqua, in eastern Pennsylvania, on Nov. 21, 1891, and went to Lehigh, receiving an E.M. in 1914, and an M.S. seven years later. His engineering career began in the Pittsburgh Testing Laboratories. After three years there he came to Palmerton for

Jan 1, 1940

By J. R. Hauber, N. R. Borch

A study uras made on the creep behavior of cast and extruded SR grade beryllium. It is shown that, for stresses below about 1000 psi in the temperature range 760" to 85o° c, the creep behatior is nearly exactly described by the Nabarro-Herring mechanism. The activation energy is apparently that for self-diffusion, 42 2 kcal per mole, below 1000 psi, but it rises to 88.6&apos; kcal per mole at higher stresses. The slress exponent is 4.5 at the higher stresses. Tertiary creep in this material is clearly related to void formation apparently caused by grain boundary sliding. BERYLLIUM is becoming an increasingly important technological material. A search of the literature has revealed a lack of fundamental studies on the creep of beryllium. This investigation was therefore undertaken to gain insight into the creep behavior of this material. EXPERIMENTAL PROCEDURE Tensile creep samples were prepared from cast and extruded SR grade beryllium. The vendor&apos;s chemical analysis of the beryllium tested is shown in Table I. The test specimens were machined as shown in Fig. 1 from i -in.-diam rods. The machined creep specimens were annealed to remove surface damage and to establish a variation in grain size. The heat treatments and their resulting grain diameters are shown in Table 11. The average grain diameters were determined by the intercept method. True average grain diameters were estimated by multiplying the measured average grain diameters by 1.5.&apos; The creep tests were performed at temperatures from 700" to 850°C in a vacuum of 105 Torr or better in a tantalum-element resistance furnace. Loads were applied by an Instron testing machine such that the stress was held constant to within 20 psi. The test temperature was measured by thermocouples on the sample. With one exception the deviation from the desired test temperature was 2"C, and in that case the deviation was 4"C. Elongations were measured by an extensometer within the vacuum chamber which was fastened to the creep specimens at the grooves in the grip ends, as shown in Fig. 1. The precision of measurement of elongation was 1 X loe4 in. The creep tests were begun only after thermal equilibrium had been established in the sample and the extensometer to insure that there was no strain rate component due to transient thermal effects. Two kinds of creep experiments were performed. First, at constant grain size, the effects of varying stress and temperature were studied. In the second set of experiments, the temperature was held constant and the grain size was varied. The activation energy for creep was determined by varying the temperature at constant stress, as described by orn, during the first set of experiments. The variable grain size experiments revealed that creep at low stress apparently occurred by the abarro-errin&apos; mechanism. In all of the creep tests, the stress was held constant only until sufficient strain had occurred to accurately determine the creep rate. The stress was then raised to a new value and the new creep rate determined. This process was reported until tertiary creep occurred. In this manner several data points were obtained from each specimen. The assumption was made that the structure was not changing appreciably between the incremental stress levels. This was verified by reducing stresses to previously tested levels and comparing the creep rates with the creep rates obtained at the lower strains. The re-producibility was within experimental error over strains on the order of 1 pct. EXPERIMENTAL RESULTS A typical series of creep tests is shown in Fig. 2. Primary creep was not observed in any test. Transient creep, which we define as creep rates that are nonlinear with time, was not observed for any change in stress except where tertiary creep occurred. The absence of both primary and transient creep has been previously reported for hot-pressed powder beryllium.= The data for the creep behavior at constant grain size are summarized in Fig. 3. The data show two distinctly different stress dependencies. At stresses below about 1000 psi, in the equation:

Jan 1, 1969

By A. Kelly, S. Sato

Tensile tests have been performed on aluminum single crystals of 99.99 pct purity at 196°K. The resolved shear stress when the stress-strain curve becomes concave to the strain axis depends on orientation, being always higher for crystals with an initial orientation close to the sides of the unit triangle. Observations have been made, with the optical microscope, of the appearance of slip lines on polished surfaces of the crystals. Particular attention has been paid to the first appearance of cross-slip. It is confirmed that marked cross-slip is first seen when the stress-strain curve becomes concave to the strain axis. Measurements are reported of the lengths of the slip lines observed on top and side faces of the crystals at various stages of the deformation. The lengths of the slip lines vary inversely as the flow stress for all crystals, independent of crystal orientation. From measurements of the lengths and separations of the slip traces it is possible to estimate the number of dislocations which are annihilated at each cross-slip site. It is concluded that the annihilation of screw dislocations of opposite sign is the most important cross-slip process occurring. FRIEDEL1 recently pointed out that the stress-strain curve of many pure metals of face-centered-cubic structure can be considered as composed of three stages. Plastic flow commences with a region showing a small rate of hardening which is approximately linear. This has been called "easy glide*. It is followed by a region of more rapid hardening in which the slope of the curve is again approximately constant and this, in turn, is superseded by a final part of the curve which is concave to the strain axis. Following Diehl2 we may designate these regions as stages I, II, and III respectively. The onset of stage 111 occurs at lower stresses at higher temperatures. The appearance of the slip lines formed during these various stages has been described by a number of workers?-5* The object of the present work was to make detailed observations of the lengths and separations of slip lines for a number of crystals of different orientations deformed at the same temperature and strain rate. Although much previous work has been published on the deformation of aluminum single crystals the importance of measurements of these quantities has only recently been realized—they are essential for experimental tests of recent theories of work hardening.&apos; 9 In aluminum, deformed at room temperature, stage 11 of the stress-strain curve is very weakly marked, as the stress for the onset of stage 111 is attained soon after the end of easy glide. In the present experiments it was decided to use a temperature lower than room temperature so that stage II of the curve was better developed. However, at temperatures as low as 90°K slip lines are not measured easily with the optical microscope.5 For this reason 1960k was chosen as a suitable temperature. Even at 1960k however, stage III of the curve commences at glide strains of 6 to 10 pct and hence

Jan 1, 1960

By Louis A. Rosales, Kanji Ono

The anomalous behavior of hydrogen in iron is examined and a new model is proposed It is postulated that hydrogen single and di-interstitials are in equilibrium with each other. When the binding enthalpy of the di-interstitial, B, is assumed to be about 10 kcal per mole, the predicted temperature dependence of the solubility agrees reasonably well with the reported data. Effective diffusiuity is calculated assuming the di-interstitials to be immobile. Since the number of mobile single interstitials is greatly reduced at lower temperatures, the effective diffusivity is decreased. A good agreement in the temperature dependence of diffusivity between theory and experiment is provided with B - 10 kcal per mole. The predicted diffusivity is also a function of total concentration of hydrogen and this appears to be Partly responsible for the large scatter in the reported data. HYDROGEN is known to form an interstitial solid solution with iron.&apos;&apos;&apos; At elevated temperatures, it has been well-established that hydrogen enters iron in the dissociated form and occupies the interstices, and that diffusion of hydrogen in iron is via the interstitial mechanism. The concentration, c, permeability, P, and diffusivity, D, of hydrogen in iron have been separately determined and are consistent with each other under the steady-state condition for temperatures above approximately 400°C; that is, the relationship P = c-D holds. At lower temperatures, however, anomalous behavior of hydrogen in iron has been observed by many investigators. The results of experimental studies4-&apos;&apos; of the diffusion of hydrogen in iron are summarized in Fig. 1, in which the values of D are plotted against the reciprocal temperature. Although there is an increasing divergence in observed values of D with decreasing temperature, the values of D at temperatures below approximately 200°C are definitely small as compared with the extrapolation from high-temperature data. The lines shown bracket the region containing the most probable data. The observed values of c6," are plotted against the reciprocal temperature in Fig. 2, where the dashed lines represent Geller and sun&apos;s&apos;&apos; most probable isobars at 1 and 100 atm. At high external hydrogen pressure and low temperatures, the values of c observed are higher than those expected from high temperature data. It is generally accepted that certain trapping mechanism is responsible for the anomalies, but the nature of the trapping mechanism has not been clarified. Various types of hydrogen traps were suggested in the literature, such as "rifts" in the lattice dislocations, voids, microcracks, and so forth. However, no quantitative account of the observed behavior has been given in terms of these trapping mechanisms. Formation of methane in hydrogen-charged steel was found by ~odgurski." This appears to explain the existence of the so-called "residual" hydrogen,4 but not the observed anomalies of hydrogen in iron. McNabb and Foster&apos; proposed a mathematical analysis in terms of a trapping mechanism. However, the physical significance of the traps was not clarified. The necessity of evaluating five parameters in their theory made comparison with experiment difficult. In the present paper, a model is proposed in which the anomalous behavior of hydrogen is attributed to the formation of hydrogen di-interstitials at lower temperatures. In this model, it is postulated that single and di-interstitials are in equilibrium with each other. (The hydrogen di-interstitial is to be distinguished from a hydrogen molecule which may exist in voids.) It can then be shown that most of the anomalies can be explained by assuming that the di-interstitials are virtually immobile and that the binding enthalpy of the di-interstitial, B, is approximately 10 kcal per mole. Predictions of the theory agree with the available experimental data reasonably well. THEORY The basic postulate of the present analysis is that single and di-interstitials are in equilibrium with each

Jan 1, 1969

By C. H. Li, R. D. Carnahan, R. J. Stokes, T. L. Johnston

When silver chloride deforms by pencil glide at temperatures of 26ºand 72°C, grain size has no effect upon the proportional limit and the material necks down to a knife edge under tension. At -196ºC, deformation takes place on fewer slip systems to produce straight slip traces, the flow stress becomes sensitive to grain size and fracture occurs by cleavage from an intergranular source without any reduction in area. THERE are significant differences between the appearance of the slip bands formed in silver chloride and in the alkali halides at room temperature. The bands in sodium chloride, e.g., are fine and geometrically straight, whereas in silver chloride they are coarse and wavy.&apos; Such differences in the slip behavior prompted us to inquire what the effects of grain boundaries might be on the mechanical properties of these solids. As an introduction to a general study of ionic poly-crystal behavior, we have investigated the effect of grain size on the stress-strain relationships in silver chloride at different temperatures. We have chosen silver chloride chiefly because polycrystals can be easily prepared free from cracks and macroscopic voids by conventional techniques. Moreover, as we shall see, the slip behavior can be readily modified by changing the temperature. It will be shown that the effect of grain size on the yield stress of silver chloride is strikingly dependent upon the temperature of deformation and that this temperature effect comes about through a change in the number of slip systems which operate at the onset of plastic flow. EXPERIMENTAL PROCEDURES Polycrystalline tensile specimens were made by recrystallizing cold-rolled AgCl sheet. Thirty-gram quantities of AR grade silver chloride powder were melted in a low heat capacity furnace, cast into cylindrical forms 3/4 in. in dim and approximately 1 in. in length and then extruded with a 16:l reduction ratio into rods 3/16 in. in dim. Extrusion temperatures were varied from room temperature to 350°C. In order to impart various degrees of cold&apos;work to the extruded rod. All extrusions were then cross rolled to sheet 0.040 in. in thickness at room temperature and machined to form tensile specimens having a gage length of 0.750 in. with a width of either 0.250 or 0.145 in.* Annealing was carried out in an air oven at 200°C for periods up to 20 hr. specimens having average grain diameters ranging from 0.024 to 0.40 mm were prepared by varying the extrusion temperature and annealing time in the conventional manner. To obtain specimens with grain sizes larger than 0.40mm, lightly deformed thin prisims of silver chloride single crystals (supplied by Harshaw Chemical Co.) were recrystallized. Grain size determinations were made by an intercept method and revealed that the grains were equi-axed. Laue back-reflection X-ray patterns showed little tendency to form a preferred orientation texture. Specimens which were to be used for the slip band studies were polished by immersion in fresh boiling concentrated NH,OH for 60 sec. They were then etch polished for 10 sec on each side by hand lapping on a stationary velveteen lap wetted with a dilute Hypo solution (5 g of Kodak Acid Fixer in 100 cc H20). Finally the specimens were rinsed in warm water and alcohol, blown dry and carefully mounted in the tensile machine for testing. Tensile tests were performed at three temperatures, namely, 26º, -72º, and -196°C. A screw-driven hard tensile machine was used at a constant strain rate of 0.006 in. per in. per min. Loads were determined with a light- proof ring fitted with A-18 strain gages and extensions were measured with a 0.0001 in. dial gage. Constant level baths of alcohol with dry ice and liquid nitrogen were used to maintain the two lower test temperatures. EXPERIMENTAL RESULTS Reproducibility—In view of the light sensitivity and reactive nature of silver chloride, it was necessary to determine whether reproducibility depended upon either exposure to light or to the condition of the surface. However, during the course of the work it became clear that variation of laboratory lighting

Jan 1, 1962

By Ralph H. Sweetser

Carbon in pig iron is not only essential but, ordinarily, it is the most abundant metalloid present; iron without carbon could not be pig iron. Carbon in pig iron has been accepted, but seldom specified. How it gets into the pig iron as it is being smelted in the blast furnace, and why just about so much of it enters into the pig iron in certain blast furnaces, and more of it,; or less of it, enters the pig iron in other blast furnaces, or in the same blast furnace under varying conditions, are problems that have not yet been satisfactorily solved. During my own practical blast-furnace experience, I collected analyses of pig iron, including the graphitic and combined carbon, and recorded them in reference to the percentage of silicon present, having in mind the generally accepted explanation that "silicon drove out the carbon." To this actual experience, I have added some analyses given to me by J. B. Rogers, who was superintendent of the blast furnaces of the Ashland Iron & Mining Co., Ashland, Ky., and was making "silvery iron" at the time. These analyses are shown in Table 1. The analyses do not tell the whole story, but they show that with Pocahontas beehive coke, the total carbon was generally high. The data relating to the composition of the slag are lacking but my own customary practice was to have a rather basic slag. During my experience in making pig iron for the Bessemer and open-hearth departments, there were times when the steel men would complain about some mysterious quality in the pig iron that caused bad steel, in spite of the fact that the percentages of silicon, sulfur, phosphorus and manganese were within the specifications. Many reasons were advanced for this undetermined "poison" in the pig iron, but the favorite one offered by the steel men was that the blast-furnace man had "oxidized" the pig iron at the tuyeres in his eagerness to get big production and low cost. To these complaints of the users of steel-making pig irons were added those of the gray iron-foundry men and the malleable iron men. These groups are unaninlous in their conclusions that the trouble lies with the blast-furnace men; but there is not yet any harmonious conclusion as to the exact cause and remedy.

Jan 1, 1927

By AIME AIME

THE two illustrations below, furnished through the courtesy of J. P. Messervey, Deputy Inspector of Mines, Department of Public Works and Mines, Province of Nova Scotia, show a fourth-level stope in the salt mine at Malagash, N. S. Fig. I was taken when the stope was begun, and Fig. 2 shows how it looked two months later, the photographs being taken from the came point. This stope is at a depth of about 300 ft. and is on the MacKay seam, which is the central band of white salt running through the Malagarh deposit. As run be seen. the vein is vertical. and varies from a foot to 50 ft. in thickness. The

Jan 1, 1932

By R. Sprycha, J. Szczypa, W. Janusz, J. Neczaj-Hruzewicz

INTRODUCTION Because large aqueous suspensions of various fine solids originate during the beneficiation of ores, the Mining Industry constantly requires efficient and feasible dewatering methods. Among the techniques currently in use (1,2), filtration (vacuum, press) and centrifugation are fairly expensive and thus are limited to use with more valuable materials. Furthermore, installations of reasonable size produce only moderate outputs. Therefore, tailings from ore dressing plants are usually treated by sedimentation or settling techniques with a tailings dam as the ultimate sedimentation basin (3). Apart from growing environmental and economic questions, the solid product obtained from the tailings dam basin has in most cases, a rather high moisture content and a persistent semi-fluid consistency. Thus, it is difficult to transport and use further. On the other hand, the solids may be utilized as a filling material, fertilizer, etc. providing they are dry enough and properly granulated. Demands for an improved technique are met to a great extent, by the technique presented in this paper. AGGLOMERATION OF FLOCCIJLATED SLIMES Origin and Qualitative Description of the Process Phenomena of consolidation and pellet formation in flocculated fines were first observed by Yusa and Gaudin (4) while studying rheological properties of flocculated kaolinite. On this basis, methods of separating liquid from solids have been worked out (5-7) and many possibilities for application have been suggested (5,6). Yusa termed the technique "wet pelletization&apos;&apos;.

Jan 1, 1980

By W. C. PAGE

HARVEY S. MUDD has the unusual distinction of having virtually inherited a Directorship in the -A.I.M.E., for he was appointed to the Board to fill the unexpired term of his father in 1926. Since then he has served as a Director continuously until 1936, and now he is a candidate for Vice-President. One of the first words that young Harvey tried to say was "ore," for he was born in Leadville, Colorado, in the heyday of that camp, where Seeley W. Mudd for years ran the Small Hopes mine. Stanford and Columbia gave him his mining education, and the Shattuck-Arizona, at Bisbee, his first job. Afterwards he examined many

Jan 1, 1937

By R. A. Bergier

The Transactions of the Institute contain nothing, as yet, on the subject of pocket-compasses; and in the belief that American miners, explorers, geologists, and engineers will gladly welcome any infor mation about new forms of apparatus, simple in construction and adapted to rapid use, I offer a short description of an improved French combination pocket-compass, which, if once introduced into use, would seem likely to supersede similar instrumens of older pattern. The instrument in my possession I owe to the kindness of M. G. de.La Bouglise, of Paris. The frame of the compass is a double-bottomed box, 2 1/2 inches square and 9/16 of an inch high, made of white metal, and with its sur faces very carefully planed. The sides of the box make true right angles with each other. Two sides are parallel with, and two are normal to, the north and south line of the principal graduation. The bottom is accurately planed, so that the instrument will lie true on a plane surface. A glass, 2 3/8 inches in diameter, is set in the top plate. The graduation for the needle is engraved on the upper face of a ring, connected with the flat disk, in which the center-pin is in serted. Between this disk and the bottom of the box there is a hollow space a quarter of an inch deep. The needle is 1 7/8 inches in length, and can be balanced by means of a small sliding piece of metal. The graduation extends from 0° to 360°, reading from right to left. The stopping of the needle is effected by a screw, A, on the side of the box near the handle, H (Fig. 1). There is also a clinometer attachment, the graduation for which is engraved on the upper face of the silver-plated disk. The compass has two sights. The north sight, B, is a simple pin, a little more than an inch high, which, when not in use, is kept folded down, and is partially protected by the exterior sliding-plate, C. When raised for use, the sight is stopped in a vertical position by a projecting pin attached to the side of the box.. The south sight, D, is connected with the handle, H, and the sight and handle

Jan 1, 1890

By N. J. Grant, B. C. Giessen, Paul Predecki

ALLOYS of gold, silver, and aluminum with elements of the groups BII, BIII, BIV, and BV were prepared by a rapid quenching technique (splat) and were examined by X-ray diffraction. Five new intermediate phases were found and will be described briefly herein. For the gold and silver systems, the concentration ranges having an electron/atom ratio e/a of 1.4 to 1.5 ("3/2 Hume-Rothery phases") were studied primarily. Master alloys were prepared from high-purity metals (99.9+ pct or better) by melting either in evacuated fused silica capsules or by nonconsum-able-electrode arc melting in an argon atmosphere. Small pieces, 20 to 50 mg, of each alloy were blast-atomized to form a splat, by a technique similar to that described by Duwez and Willens.1 The technique used for this study is described in detail in Ref. 2; it utilizes a resistance-heated graphite crucible with a small hole at the bottom, directed toward a metal substrate or quenching plate. The prepared alloy rests over the fine hole, through which it is expelled by an explosion shock wave in the form of fine droplets (1 to 50 µ) of molten metal onto a copper or silver substrate, which is maintained at about -190°C. The resulting very high cooling rates (see Ref. 2 for quantitative measurements) can prevent the process of nuclea-tion and growth in many instances, resulting in the formation of metastable phases. The splat particles were transferred to a GE-XRD5 diffractometer and maintained at -190°C, where they were examined with CuKa radiation. The samples were then allowed to warm to room temperature or were heated to higher temperatures until the equilibrium structures formed. Of fifteen alloy systems considered, nonequi-librium structures were encountered in six; these are described below and summarized in Table I. In the system Au-Sb a metastable £ phase (A3 type, hcp, a = 2.898 + 0.002A; c = 4.731 * 0.004A; c/a = 1.633) was found in the concentration range Au + 13 to 15 at. pct Sb. This phase is isomorphous with the stable phases in the systems Au-Cd, Au-In, and Au-Sn, all at an average e/a ratio of 1.4 to 1.5. The concentration range of one-phase metastable was deduced from the small amounts of supersaturated gold solid-solution phase present in the splat product. It was found that ? could also be retained by splatting onto a substrate held at room temperature: however, decomposed into the equilibrium phases Au + AuSb2 after heating to 200°C for 1/2 hr, or on holding the powdered splatted alloy at 20°C for several months. Calorimetric measurements will be made in an attempt to decide the question whether ? is metastable at all temperatures or whether it is a stable phase at low temperatures. There is evidence that another phase, possibly also close-packed but with a different stacking sequence, can be obtained by rapid quenching of alloys with a different antimony content. Klement, Willens, and Duwez3 reported the existence of an amorphous phase on quenching Au-Si alloys (25 at. pct Si) to - 196°C. They found that on heating to room temperature another phase of unknown crystal structure was formed. This was confirmed (see Table I); however, the new crystalline phase, designated as ?, could also be formed simply by rapid quenching to room temperature, and even was found to exist already in the as-cast Au + 20 at. pct Si alloy. It was found that ? decomposed into Au + Si on the specimen surface at room temperature. This behavior, and the question whether or not there is an equilibrium-temperature region for ?, have not yet been resolved. It is probable that ? (Au + 20 to 21 at. pct Si) is cubic of the -brass type (D81-3) with a = 9.60, + 0.01A and N = 52 atoms per cell [compare 6 (CU-Sn)4]. Except for two very weak lines, the powder pattern of about thirty lines could be indexed on this basis; however, a determination of the atom positions has not yet been attempted. For Au-Ge the C phase was observed at about 21 at. pct Ge as reported by Luo et at.5 Lattice parameters a = 2.876A, c = 4.73,A, c/a = 1.64 were found. In the Au-Pb system, formation of a ? phase was not observed, but in the lead-rich region at 75 at. pct Pb, broad peaks belonging to an amorphous phase were found. The maximum diffracted intensity occurred at 28 = 32.4 deg which is about 1 deg larger than the position of the (111) line of lead (Cuka). For Ag-Pb, an amorphous phase analogous to the one found in the Au-Pb system was observed; this metastable phase exists probably at about 75 at. pct Pb. Since no lead-rich alloys were tested, all alloys consisted of silver + amorphous phase at -190°C. In A1-Ge alloys, line-rich and complex powder patterns were obtained at about 30 at. pct Ge; they bear similarities to those of aluminum and germanium, but are of lower symmetry; the existence of more than one intermediate phase is possible. The authors are grateful to the Kennecott Copper Corp. for Fellowship support, and ARPA (Contract

Jan 1, 1965