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By S. R. B. Cooke, H. S. Choi, I. Iwasaki

In magnetic taconite of the East Mesabi range, quartz, magnetite, and cummingtonite (OH}(Fe, Mg)(Si4O11)2) are the three major ore-forming minerals.' Fine grinding and magnetic separation of the taconite leads to a concentrate analyzing 64 to 65 pct Fe and 8 to 9 pct insoluble; the gangue content of the concentrate arising from the presence of locked magnetite-gangue particles, and of occluded gangue particles in magnetic floccules. To further upgrade such taconite concentrates, flotation appears to offer a particularly promising approach. To evaluate its effectiveness in separation, a knowledge of the flotation characteristics of each mineral comprising the concentrate becomes desirable. The floatability of magnetite and of quartz have already been reported in a previous paper,2 in which the flotation characteristics of the two minerals were correlated with the electrokinetic properties at the mineral-solution interfaces. Since cummingtonite has somewhat intermediate chemical composition between the above-named oxides, the electrokinetic and the flotation behavior of this mineral becomes of interest for both practical and fundamental reasons. In the present article the results of electrophoretic mobility measurements and of Hallimond tube flotation tests on cummingtonite are reported and are compared with those of quartz and magnetite. MATERIALS AND METHODS A sample of cummingtonite-rich taconite, received from the Mitchell pit, near Babbitt, Minn., was crushed and ground through 100 mesh, and the product was fractionated magnetically into three portions. First, a drum-type laboratory cobber was used to separate a highly magnetic portion, consisting mainly of free, but also containing locked, magnetite particles. The cobber tailing was then passed in two stages through a Franz isodynamic separator. The middle fraction was isolated, and, upon examination petrographically, was ascertained to be principally cummingtonite. The 100 to 150 mesh fraction of this sample was selected for the present study. It was cleaned by leaching briefly with hot, dilute HC1 and then by rinsing repeatedly with demineralized water. A ground and elutriated fraction of -5p cummingtonite was used for the electrophoretic measurements.

Jan 1, 1961

By Frederick N. Rhines, Robert F. Mehl

Relatively few data have been collected on the rates of diffusion in ternary solid solution systems. In general it does not seem worth while to gather extensive data for such systems until the factors that determine the rates of diffusion in binary systems are better understood. But the rates of diffusion in a ternary system of limited solid solubility, in which the terminal solid solution is in equilibrium with an intermediate phase and may thus be represented as a quasibinary system, represents a case of special interest. In this case, where by reason of the association of the two solute metals in an intermediate phase a special attraction of the two solute atoms may be safely assumed, some interaction of these atoms during simultaneous diffusion in the solvent lattice might be expected. It has frequently been proposed that molecular solid solutions exist in quasibinary and in other ternary (and binary) solid solutions. There is no direct evidence for such solid solutions; arguments based upon the supposed anomaly in certain physical properties1 are hardly important; unexpected density relationships have proved deceptive;2 suggestions of molecular solid solutions have been made on the basis of solution laws in their application to certain systems, but without result;3 attempts to find lattice parameters in the solid solution Al-MgZnz, which cannot be calculated from the separate effects of magnesium and zinc and might therefore be taken as evidence for a molecular aggregation, have given negative results.* Fundamentally, methods for the determination of the state of molecular aggregation in solution depend upon the kinetic behavior of the particles constituting the system. If it could be shown that the two solutes in a quasibinary ternary system diffuse in molecular proportions, despite a radical difference in the separate rates of diffusion of each solute when diffusing alone in the solvent lattice, some basis for the concept of molecular solid solution would be afforded. Freche5 examined the rates of

Jan 1, 1940

By D. M. Collingwood

The year 1928 has been marked by scattered but considerable wildcat drilling considering the overproduction prevalent in the oil industry. This wildcatting has resulted in the discovery of oil in one interior salt dome, several small fields producing from serpentine, Taylor marl, and Austin chalk, and two new fields producing from the Edwards limestone, one of which in the Bruner or Salt Flat area of Caldwell County promises to be of major importance if not quite as large as the Luling field. Fig. 1 shows production by fields for 1928, compared in the case of the older fields, with that for 1927. New Production From Salt Domes Boggy Creek Dome This salt dome was discovered in 1927 by core drilling and deep tests recommended on the evidence of surface geology. A large block of acreage controlled by the Humble Oil & Refining Co. covers the dome. Since the discovery of the deep-lying salt plug and oil on the flank of this dome in 1927, development has proceeded slowly. Since most of the acreage is controlled in a large block by one company there has been no intensive competitive drilling. Approximately 8 producers and 16 dry holes have been drilled. Production is from the eastern equivalent of the Woodbine and Eagleford horizons of the Mexia district, and has been found extending in a northeast-southnlest direction across the River Neches on the southeast flank of the dome. One of the best wells, the T. H. Jones No. 2, had an initial production of 90 bbl. per hr., and others have made from 100 bbl. per day to 50 bbl. per hr., but production of the better wells has either been shut in or choked down to less than 2000 bbl. per day. At the end of the year the daily average, production was about 2500 bbl. from 8 wells. Considerable gas pressure is present and one well, the Elliot Clark No. 2, drilled in 1927, was shut in with an estimated wet gas production of over 5,000,000 cubic feet.

Jan 1, 1929

By Robert W. Hunt

The American Institute of Mining and Metallurgical Engineers was the first American technical organization to consider steel-rail specifications and sections. If I am not mistaken, the first contribution on those subjects to the Proceedings of the Institute was presented by Prof. Thomas Egleston, in 1874, and was entitled "Investigations on Iron and Steel Rails made in Europe in the Year 1873.1" The next contribution was Dr. C. B. Dudley's report of his investigation for the Pennsylvania Railroad of the wear of steel rails in their tracks, in 1878, which contribution led to a very memorable discussion that extended through several meetings.2 In 1880, Mr. C. P. Sandberg3 presented an article on rail specifications and inspection and rail sections. This was followed, in 1881, by A. L. Holley,4 in his exhaustive report on the then used American rail sections. The next contribution, I believe, was made by myself5 at the Buffalo meeting, in 1888, on steel-rail specifications, which I followed,6 in 1889, by a paper on steel-rail sections. Since that time there have been numerous contributions to the Transactions of the Institute in relation to the manufacture of steel rails. The American Society of Civil Engineers officially took up the investigation of the relation of rail sections to the wear of car wheels in 1888. In the prosecution of this work, the Society appointed a special committee, whose labors extended through several years and resulted in an exhaustive report to the Society in 1893.7 Later, another committee was appointed to consider specifications for the manufacture of steel rails; the labors of this committee extended over a period of several years, a final reporta being made in 1910. It was my fortune to be the secretary of both of those committees. The American Society for Testing Materials, organized in 1899, took up the consideration of rail specifications and is still giving attention to that subject through its established procedure of investigations by committees. In 1900, the American Maintenance of. Way Association was organ-

Jan 1, 1920

Illinois is rightly known as containing the coal deposit of which the first record in the United States was made. Joliet and Marquette saw "charbon de terre" along the Illinois River in 1673;1 LaSalle saw "mines de charbon" along the same stream in 1679; and Hennepin reported having "found in Several Places Some Pit-Coal" in 1682.2 In 1687 Jontel, in writing of the Illinois River country said, "along the river are coal pits, the coal whereof has been try'd and found very good."3 J. Carver reported the northern parts of the Mississippi River valley, "containing lead, copper, iron, coals, etc." in 17664 evidently referring to the Illinois country. Croghan, in the same year, saw coal along the Wabash River, near the mouth of the Vermilion in Indiana, but not far from Danville. Capt. Philip Pittman in his journey up the Mississippi in 1770 reported that "great pieces of coal are constantly found on the sand-banks, from whence it may be concluded, that there are coal-mines on the upper parts of the Mississipppi."5 This coal could only have come from the Illinois or Big Muddy Rivers. Morse, in 1798, refers to great quantities of coal on the upper branches of "the Father of Waters," and "on the Illinois River is a coal mine, which extends for half a mile along the middle of its banks."6 While, undoubtedly, some coal was used locally in various places, no records have been left of it and the first definite one is that in 1809 William Boon shipped some coal from the Big Muddy River, in Jackson County, to New Orleans, which was the first coal exported from Illinois.1 There has been considerable discussion as to this first shipment, whether it was made in 1809, 1810 or 1812. The records of arrivals of: boats in New Orleans in those early years is specific as to date, name of craft and name of captain, but no mention is made of cargo carried. William Boon arrived

Jan 1, 1942

By A. Lubinski

In drilling operations, attention generally is given to hole angles rather than to changes of angle, in spite of the fact that the latter are responsible for drilling and production troubles. The paper presents means for specifying maximum permissible changes of hole angle to insure a trouble-free hole, using a minimum amount of surveys. It is expected that the paper will result in a decrease of drilling costs, not only by avoiding troubles, but also by removing the fear of such troubles. SUMMARY, CONCLUSIONS AND RECOMMENDATIONS Excessive dog-legs result in such troubles as fatigue failures of drill pipe, fatigue failures of drill-collar connections, worn tool joints and drill pipe, key seats, grooved casing, etc. Most of these detrimental effects greatly increase with the amount of tension to which drill pipe is subjected in the dog-leg. Therefore, the closer a dog-leg is to the total anticipated depth, the greater becomes its acceptable severity. Very large collar-to-hole clearances will cause fatigue of drill-collar connections and shorten their life, even in very mild dog-legs. Another finding regarding fatiguing of collar connections in dog-legs is that rotating with the bit off bottom sometimes may be worse than drilling with the full weight of drill collars on the bit, mainly in highly inclined holes when the inclination decreases with depth in the dog-leg. Means are given for specifying maximum dog-legs compatible with trouble-free holes. An inexpensive technique proposed is to take inclinometer or directional surveys far apart; then, if an excessive dog-leg is detected in some interval, intermediate close-spaced surveys are run in this interval. The application of the findings should result in a decrease of drilling costs, not only by avoiding troubles, but mainly by removing the fear of such troubles. The result would be much more frequent drilling with heavy weights on bit, regardless of hole deviation. Because of errors inherent to their use, presently available surveys are not very suitable for detecting dog-legs. There is a need for instruments especially adapted to dog-leg surveys. Crooked hole drilling rules should fall into two distinct categories—(1) those whose purpose is to bottom the hole as desired, and (2) those whose purpose is to insure a trouble-free hole. Three kinds of first-category rules in usage today are as follows. 1. A means to bottom the hole as desired is to prevent the bottom of the hole from being horizontally too far from the surface location; this may be achieved by keeping the hole inclination below some maximum permissible value such as, for instance, 5. 2. Another means to achieve the same goal is to limit the rate at which the inclination is allowed to increase with depth. A frequently used rate is 1/1,000 ft. In other words, a maximum deviation of l° is allowed at 1,000 ft, 2 at 2,000 ft, 3 at 3,000 ft, etc. 3. Whenever application of the first two means precludes carrying the full weight on bit required for most economical drilling, then the best course is to take advantage of the natural tendency of the hole to drift updip, displace the surface location accordingly and impose a target area within which the hole should be bottomed. This method has already been successfully applied,'.' and its usage probably will become more frequent in the future. Means for calculating the amount of necessary surface location displacement are avail-able.3'5'6 If in high-dip formations the full weight on bit should result in unreasonably great deviations, the situation could be remedied by increasing the size of collars and (if needed) the size of both hole and collars,351 or in some cases by using several stabilizers. Rules which would fall into the second category (i.e., rules whose purpose is to insure a trouble-free hole) are seldom specified today. It is vaguely believed that following Rules 1 and 2 of the first category will automatically prevent troubles. Actually, this is not true. If at some depth the only specified rule is that the hole inclination must be less than 4", the hole may be lost if the deviation suddenly drops from 4 to 2, or if the direction of the drift changes, etc. Rule 3 of the first category is generally used in conjunction with a rule belonging to the second category, namely, that the hole curvature' (dog-leg severity) must not exceed the arbitrarily chosen value of 1½ /100 ft. Moreover, when using this rule, the industry is not clear over what depth intervals the hole curvature should be measured. All this results in a frequent fear

By G. R. Leverant, C. P. Sullivan

To evaluate the effect of a dispersion of nondeform-able, incoherent, second-phase particles on high-strain cyclic deformation and fracture, recrystallized TD-nickel (Ni-2ThO2) and a commercially pure nickel, Ni-200, were fatigued under strain control at total strain ranges varying from 0.009 to 0.036. Relative to the Ni-200, the slip at the surface of the TD-nickel was more wavy and discontinuous due to the presence of the thoria particles. This made crevice formation (incipient cracking) within slip bands more difficult in TD-nickel than in Ni-200. Both materials cyclically hardened to a constant (saturation) flow stress which increased with increasing plastic strain amplitude. Cellular substructures were developed in both materials during cycling. The cell size in TD-nickel was controlled by the thoria particle distribution and was independent of plastic strain amplitude over the range investigated. The cell size in Ni-ZOO was larger than that in TD-nickel at similar plastic strain amplitudes and was a function of plastic strain amplitude. These results, together with the cyclic stress-strain curves for both materials, are discussed in terms of a model for fatigue strain accommodation at saturation recently proposed by Feltner and Laird. NUMEROUS fatigue investigations have considered the interrelation of slip character, dislocation substructure, and cracking in pure metals and solid-solution alloys. However, except for the studies of the low-strain fatigue of internally oxidized copper alloys1 and cast, dispersion-strengthened lead,&apos; little is known about the effects which small, incoherent, nondeform-able, second-phase particles have on cyclic deformation and cracking processes. Effects due to the particles alone are often obscured by a dislocation substructure introduced during thermomechanical processing of dispersion-strengthened metals. In the present study, recrystallized TD-nickel and a commercially pure nickel, Ni-200, were employed to evaluate the effect of a thoria dispersion on high-strai fatigue deformation and cracking at room temperature. I) MATERIAL AND EXPERIMENTAL PROCEDURE The TD-nickel was supplied by DuPont as a 5/8-in.-thick stress-relieved plate which had been subjected to a proprietary schedule of thermomechanical treatments, and the Ni-200 as 3/4-in. bar which was subsequently annealed for 2 hr at 850°C in argon resulting in an average grain diameter of 0.05 mm. The compositions of these materials are given in Table I. The microstructure of the TD-nickel consisted of elongated grains parallel to the primary working direction with an average width of 0.16 mm, Fig. l(a). Many fine annealing twins were present indicating that the starting material was in a recrystallized condition; this supposition was confirmed by the absence of of any extensive dislocation substructure, Fig. l(b). Sheetlike stringers parallel to the rolling direction were occasionally seen both within grains and at grain boundaries. Some approximately spherical particles about 2 in diam, which may correspond to exceptionally large thoria particle aggregates, were also present. The average Young&apos;s modulus of the plate material in the rolling direction was 21.8 X 106 psi which is consistent with a {100}<001>recrystalliza-tion texture3&apos;* being prominent. In transmission microscopy, the 2.3 vol pct of thoria particles generally appeared to be uniformly distributed although some clusters, 0.1 to 0.3 in diam, of larger particles were observed as previously reported for TD-nickel sheet,5 and stringering of particles was present in some areas as welt. The average diameter of the thoria particles was 450A with a calculated mean planar center-to-center spacing of 2100A, as determined by quantitative metallographic analysis.= The 0.2 pct offset yield stress was 36,000 psi which agrees with the value predicted by the modified Orowan relation7 for edge dislocations bowing between thoria particles of the size and spacing observed in the present investigation. Fig. 2 illustrates the specimen design employed for the axial high-strain fatigue testing. Adapters were screwed onto the threaded portions of each specimen so that testing could be performed in the same manner as that reported for buttonhead specimens.8 Stressing was coincident with the working direction for both materials. The gage section of each specimen was electropolished and lightly etched prior to testing. The total strain was controlled, being varied between zero and a maximum tensile strain ranging from 0.009 to 0.036. In addition to these tests, a circum-ferentially notched TD-nickel specimen was cycled over a total strain range of 0.0075. The same strain

Jan 1, 1969

By N. A. D. Parlee

The diffusion rate of hydrogen in liquid iron has been measured by a gas-liquid metal diffusion cell technique. The diffusion cell was formed by immersing an alumina tube containing hydrogen gas at 1 atm in a bath of stagnant liquid iron. The change in the composition of the melt in the cell was determined by measuring the rate of absorption of the gas in the cell. The appropriate solution to Fick&apos;s second law was used to examine the data and calculate diffusivi-ties. The absorption of hydrogen in stagnant pure liquid iron has been found to be diffusion-controlled. The results show that the chemical diffusion coefficient, D, of hydrogen in pure iron in the range of 1547" to 1726°C can be represented by the following Arrhenius relation: D(sq cnz per sec) = 3.2 x X exp(- 3300 i 1800/RT) where the uncertainty in the activation energy corresponds to the YO pct confidence level. Oxygen in the melt (above 0.015 pct 2) increased the apparent rate of absorption of hydrogen. The importance of diffusion data on liquid metals for predicting the rates of certain metallurgical processes has been recognized for a long time. Moreover, these data are much needed to test and develop theory for diffusion in liquid metals. Despite this practical and theoretical interest, however, relatively little reliable information about diffusion in liquid metals is available in the literature. This is particularly true for gas components such as hydrogen, oxygen, and nitrogen in liquid metals, where almost no data on chemical diffusion coefficients are to be found. This is probably due to a multitude of experimental difficulties particularly associated with high-temperature melts. In an effort to fill this gap in information, a research program was undertaken to study the diffusivities and rates of solution of gases in liquid metals. This paper presents the results of a study of the diffusion of hydrogen in liquid iron. EXPERIMENTAL METHOD Two methods for the study of the kinetics of dissolution of gases in liquid metals are being employed in this laboratory. Both involve the measurement of the volume of gas absorbed by the melt as a function of time and as such both avoid the uncertainties involved in chemical analyses of quenched samples for relatively small amounts of gas. In the first method, the gas dissolves in an inductively stirred melt and, in the absence of a slow surface reaction, the results are often interpreted in terms of mass transport across a liquid "boundary layer" between the homogeneous gas phase and well-stirred part of the melt. Other interpretations of the results of such experiments have also been described in the literature.1&apos;5 In the second method a gas-liquid metal diffusion cell is used.&apos; The gas dissolves in a cylindrical column of stagnant liquid metal and, in the absence of a slow surface reaction, the results are interpreted in terms of a non-steady-state diffusion solution to Fick&apos;s second law. The weakness of the first method is that while it gives information on the mechanism of absorption by stirred melts it yields an overall rate constant which even in the simplest cases depends on the nature and the thickness of the "mass transport layer" or "boundary layer". It yields no values of diffusion coefficients. The second method was used in this research because in many cases it is possible to determine the diffusion coefficient of the gas component in the liquid metal. In this research it has been utilized to measure diffusion coefficients of hydrogen in liquid iron. The apparatus used was essentially the same as that described by Mizikar, Grace, and par lee but certain modifications have been introduced to meet the elevated temperatures and special conditions of this research. Fig. 1 is a schematic drawing of the apparatus and Table I gives the identification of various parts in this figure. The diffusion cell, shown in detail in Fig. 2, was formed by immersing an impervious alumina tube (hereafter called absorption tube) in a bath of pure liquid iron contained in an alumina crucible. Two types of tubes were used, Morganite triangle RR and McDanel AP35. The crucible was contained in a vertical impervious alumina combustion tube (32 mm ID by 914 mm long) which was closed at both ends by water-cooled brass heads employing O-ring compression seals, Fig. 1. A protection tube enclosing a Pt, 5 pct Rh-Pt, 20 pct Rh thermocouple was introduced through the lower end of the combustion tube

Jan 1, 1968

By J. B. Clark, A. J. McEvily, R. L. Snyder

The nonuniform distribution of precipitate particles has been recognized as a leading factor contributing to the relatively low fatigue resistance of aluminum alloys. The structure of many of these alloys is characterized by narrow precipitate-free zones adjacent to the grain boundaries. Alloys with such zones exhibit a tendency for brittle inter crystalline fracture. The interrelation between this type of structure and mechanical properties was investigated in an Al-10 wt pct Mg alloy. It was found that deformation during fatigue occurs preferentially along these zones and cracks initiate there. In Al-10wt pct Mg, the zones were found to be supersaturated even after extensive general precipitation and are due to the absence of proper precipitate nuclei in the region near the grain boundaries. Cold working the alloy prior to aging improves the mechanical properties by inducing precipitation within the zones and also by jogging of grain boundaries. The mode of fracture is changed from brittle inter crystalline to more ductile trans granular fracture. THE process of fatigue is highly structure sensitive, with the strength of the whole often dependent upon some localized discontinuity, either geometrical or metallurgical in nature. Much has been learned about the role of geometrical discontinuities, e.g., notches, in fatigue, but with the exception of the effects of inclusions or the shapes of carbides, relatively little is known about the specific effects of discontinuities in metallurgical structure such as nonuniform precipitation. In most age-hardening aluminum alloys, metallo-graphic studies have shown that the extent of precipitation adjacent to grain boundaries is much less than that which occurs in the interior of the grains. The width of these almost precipitate-free regions, which are sometimes called denuded zones, and the extent of solute depletion within them, are dependent upon the particular alloy and its aging treatment. It has been observed1 that these zones are relatively soft with the result that plastic deformation takes place preferentially within them. It has also been shown 2-4 that there exists a tendency for intercrys- talline cracking in fatigue when such zones are present. It is of interest to note that Broom et al.2,3 were able to reduce the incidence of this type of failure in an A1-4 wt pct Cu alloy by stretching the material 10 pct prior to aging. In the present study, the effects of precipitate-free regions on the fatigue properties of an A1-10 wt pct Mg alloy were studied in detail, and the effects of deformation prior to aging on the nature of the precipitation process as well as on fatigue properties were also investigated. MATERIAL AND PROCESSING An A1-10 wt pct Mg alloy was selected for this study, because it was known that well-defined precipitate-free regions along the grain boundaries are readily obtained in this alloy after aging at 200oC.5 The starting materials were 99.998 pct A1 and singly sublimed magnesium of about 99.9 pct purity. The aluminum was induction melted in a graphite crucible, and then the magnesium addition was immersed until dissolved. Chlorine gas was then bubbled through the molten alloy for 4 min to degas the melt, after which the melt was cast at a pouring temperature of 730" to 760°C into a cold, graphite-coated, tapered steel mold. Since A1-Mg alloys are difficult to homogenize,5 special care was taken to obtain a uniform composition. Two-in. cubes were cut from the ingot and heated at 446°C for 30 min. These cubes were then hot forged approximately 35 pct in each of the three cube directions and homogenized for 16 hr at 446°C. Sheet specimens were then obtained by pressing 40 pct and rolling 35 pct per pass with reheating between reduction steps to a final thickness of approximately 0.10 in. The sheet was then solution treated for 16 hr at 446°C and water quenched. The age hardening behavior of this material at 200°C was then determined, and the results are shown in Fig. 1. The age hardening of this alloy when subjected to cold work prior to aging is also shown in this figure. Preliminary work indicated that extensive deformation after quenching was required to affect drastically the precipitate-free regions in this alloy, and a rolling reduction of 50 pct was chosen. For purposes of comparison the following three conditions were studied: a) Solution treated, quenched, and aged 20 hr at 200°C

Jan 1, 1963

By C. D. Hall, F. E. Dollarhide

Fluid Ioss agent.s for crude oil and for water have been studied in dynamic tests. A treatment using a spearhead with a fluid loss agent followed by plain fluid appears feas ible in crude oil, but not in water. An equation for spearhead depletion shows that spurt loss relative to fracture width must be low, if the portion of spearhead fluid in the treatment is to be small. The presence of colloidal matter in crude oils aids the fluid Ioss agent. Unlike in kerosene, where flow limited the agent deposition, in crude oils the filter cake continually formed and leak-off declined. The volume-time relation varied somewhat for different crudes, but was best described by a square root of time function. Spurt loss was inversely proportional to agent concentration. After the fluid loss agent initiated the filter cake, the crude oil colloids built on it effectively. A 2-minute or a 5-minute spearhead with double the normal agent concentration gave the same fluid Ioss curve as the same concentration did for a 30-minute test. The agents tested in water gave fluid Ioss plots on which, for the first few minutes, volume was proportional to the square root of time, but later became proportional to time. For fracture area calculation the customary square root of time function is a satisfactory approximation. Leak-off rates and spurt losses were higher in water systems than in oils. The spurt Ioss tended to be inversely proportional to concentration. In spearhead tests, the filter cakes were not eroded by water flow. However, the rather high spurt loss values make spearhead treatments impractical for water-based fluids. Introduction The effects of dynamic testing conditions on the performance of fluid loss agents in kerosene have been studied previously.&apos; We have extended the work to include crude-oil- and water-based fracturing fluids. An understanding has been gained of the mechanisms of formation and functioning of the filter cakes of fluid loss agents. The practical aspects of evaluating performance of agents in relation to fracture area calculations also are considered. The feasibility of using the fluid loss agent in a spearhead stage of the treatment is examined further for both types of fluids. Experimental Procedure The dynamic fluid loss tests were performed in an apparatus similar to the high-pressure apparatus described in a previous publication.&apos; A fracturing fluid was circulated over a rock surface located in a closed pressurized loop. The fluid flowed axially over the cylindrical surface of a core 2 in. in diameter X 3.5 in. long, mounted (with the flat ends sealed off) in a pipe, with 0.117 in. annular clearance. The filtrate was collected in a central hole in the core and led through valves to graduated cylinders. Provision was made for changing quickly the circulating fluid during the test (spearhead runs) without interrupting the filtration pressure. The only modifications were to add heating tapes and water jackets for the tests with crude oils, all conducted at ISOF, and to change all parts exposed to the test fluid to stainless steel for the tests with water-based fluids. The latter tests were made at room temperature, 80F. Three crude oils were tested. A mixed crude, obtained from a local refinery, contained a considerable amount of light ends. For safety reasons, it was stripped to 250F vapor temperature before use in the fluid loss tests. The other two oils were used as obtained from lease tanks. One was a greenish-brown, 37" API paraffinic crude, and the other was a black, 32" API asphaltic crude. The fluid loss agent for oil, here designated for brevity as Agent A, was Adomite@ Mark II*, a granular solid commercial agent, the same as previously tested in kerosene.&apos; Three different compositions of fluid loss agents were tested in Tulsa tap water. Agent B was adomit& Aqua*, a solid commercial fluid loss agent, comprising clays and hydrophilic gums principally derived from starch. Agent C was a mixture of three parts of Agent B with two parts of silica flour. Agent D was Dowel1 J137, a mixture of guar gum and silica flour. The test cores were cut from contiguous blocks of Berea or Bandera sandstones. For the oil tests, the cores were oven dried, evacuated, saturated with kerosene, and the kerosene permeability was measured. The cores used with the water-based fluids were pretreated by saturating with 3 percent calcium &loride solution to minimize pemeability damage by the fresh water due to clay migration. The

Jan 1, 1969

By E. W. Johnson, M. L. Hill

Cold working of iron-carbon alloys was found to increase greatly the hydrogen solubility and to decrease the diffusivity at temperatures up to 400° C. These effects are increasing functions of both the carbon content and the degree of deformation. The hydrogen behavior is consistent with the idea that cotd working creates "traps", which are concluded to be microcracks in which the hydrogen is chemisorbed. Hydrogen embrittlement is explained by the Petch theory of metal crack surface energy loss due to hydrogen adsorption. HYDROGEN embrittlement of steel has been studied for many years and has been the subject of an extensive literature, but the mechanism of the effect has not been completely understood. The embrittlement is unusual in that the ductility loss is not accompanied by an increase of the yield strength, being primarily a decrease of the fracture strength alone. The loss of fracture strength is usually most severe in the temperature range between 0°and 100°C. Here the solubility of hydrogen in the iron lattice at ordinary H2 pressures is extremely low while the diffusivity is still quite high. From the relationships between the ductility and the hydrogen content, test temperature and strain rate, it is apparent that the hydrogen atoms causing the ductility loss difbse to and concentrate in small regions of the metal which are especially susceptible to the initiation and propagation of fracture. This hydrogen segregation apparently occurs after plastic straining has begun. Below 0°C the ductility loss persists only at low strain rates in confirmation of the view that the embrittlement is diffusion .controlled. The tendency of the embrittlement to disappear above 100°C can be explained by the increasing lattice solubility of hydrogen with rising temperature. A common view of hydrogen embrittlement of steel is that the hydrogen initially dissolved in the metal lattice diffuses to structural discontinuities and there precipitates as H2 gas at very high pressures which assist the external stress in causing premature failure.1,2 The idea of a high H2 pressure in equilibrium with ordinary amounts of hydrogen in steel at room temperature is due to observations of hydrogen behavior in fully annealed material, for which the Sieverts&apos; law constant relating solute concentration to H2 pressure is extremely small. Hydrogen-embrittled steel, however, is always plastically deformed to some extent, and therefore it is important that hydrogen embrittlement be explained primarily in terms of hydrogen behavior in plastically deformed material. Such an explanation is attempted in this paper. Previous studies of hydrogen in cold-worked steel have shown that both the solubility and the diffusion rate are significantly chaned when the steel is cold worked. Darken and Smith discovered that the amount of hydrogen absorbed from acid by cold-rolled steel at 35°C is many times greater than that absorbed by hot-rolled steel. They found also that the hydrogen permeability of the steel is unaffected by cold working. Keeler and Davis4 confirmed the high apparent solubility of hydrogen in cold-worked iron-carbon alloys at temperatures up to and even beyond the recrystallization temperature. They also found that this solubility increase accompanying cold work is a sensitive function of the carbon content, being absent when no carbon is present. The present experimental study was undertaken primarily to obtain an improved understanding of the behavior of hydrogen in cold-worked steel. Data were obtained on the effects of temperature, H, pressure, carbon content, and degree of cold work on the hydrogen solubility and diffusivity in iron-carbon alloys. These data have been helpful in elucidating the nature of the cold-worked steel structure as well as in providing information on the mechanism of hydrogen embrittlement of steel. EXPERIMENTAL Cylindrical specimens for hydrogen absorption and diffusion rate measurements were prepared from three iron-carbon binary alloys and a commercial SAE 1010 steel. The iron-carbon alloys were prepared by vacuum melting electrolytic iron with graphite in a magnesia crucible. The alloys were cast in vacuum as 2 1/2-in. sq ingots weighing about 20 lb each. The ingots were hot rolled (above 1900°F) to 5/8-in.-diam round bars and then cooled in air to room temperature. The resulting metallographic structure consisted of islands of fine pearlite surrounded by free ferrite. Chemical analyses of the materials are given in Table I. The 5/8-in. diam bars were turned to diameters such that cold reduction to the desired final specimen diameters would result in either 30 or 60 pct reduction in area (RA). The machined bars were then cold worked by swaging at room temperature

Jan 1, 1960

By T. F. Kassner

The dissolution kinetics of type 304 stainless steel in the Bi-Sn eutectic alloy have been investigated under the well-defined hydrodynamic conditions produced by the rotating-disc sample geometry. In addition, the mutual solubilities of iron, chromium, nickel, and manganese from 304 stainless steel in the eutectic alloy were determined over the temperature range 450" to 985°C. The convective -diffusion model for mass transport from a rotating disc was used to interpret the experinlental dissolution data. The dissolution process was found to be liquid-diffusion-controlled under specific conditions of temperature and Reynolds number. Liquid penetration into the 304 stainless steel resulted in a reduction of the di,ffusion-controlled mass flux and thus precluded the calculation of the diffusion coeficients of the four components from 304 stainless steel in the Bi-Sn eutectic alloy. The convective-diffusion model for diffusional limitations of electrode reactions and mass transport at the tationssurface of a rotating disc set forth by Levich 1,2 has found wide applicability in the investigation of electrochemical and dissolution phenomena in aqueous systems. Riddiford 3 and Rosner have reviewed the model and also include numerous references on work of this nature. More recently the rotating-disc system has been applied to the investigation of hetereogeneous reactions in liquid-metal systems. Shurygin and Kryuk 5 have measured the dissolution rates of carbon discs in molten Fe-C, Fe-Si, Fe-P, and Fe-Ni alloys. Shurygin and shantarin6 also studied the dissolution kinetics of iron, molybdenum, chromium, and tungsten, and the carbides of chromium and tungsten in Fe-C solutions with a rotating-disc sample geometry. In these systems it was possible to distinguish between diffusion and reaction control mainly through experimental confirmation of the velocity dependence of the dissolution rate predicted by the model. However in the absence of dependable solubility data and the virtual lack of diffusion data in these systems, a quantitative check of the magnitude and the temperature dependence of the rate was not possible. In many instances, estimates of the activation energy for solute diffusion and the diffusion coefficient based upon the experimental dissolution data are not credible. A recent study by this author7 has resulted in a critical test of the model in a liquid-metal system. The solution rates of tantalum discs in liquid tin were measured over a wide range of temperature and velocity conditions. In addition, the solubility and diffusion coefficient of tantalum in liquid tin were determined as a function of temperature. The latter data were used with the model to predict both the magnitude and the temperature dependence of the dissolution flux. In that work it was also deemed necessary to reevaluate the solution to the convective diffusion equation to incorporate the effect of the lower range of Schmidt numbers encountered in liquid-metal systems. Good agreement between the model and the experimental dissolution data in the region of diffusion control was obtained in the Ta-Sn system. The Bi-Sn eutectic alloy is used as a seal between the reactor head and the reactor vessel in the Experimental Breeder Reactor-11. The alloy is fused periodically prior to fuel-handling operations. In that connection, it was necessary to investigate the compatibility of the liquid alloy with the type 304 stainless-steel containment material. The results of a rotating-disc study in this multicomponent system are presented. EXPERIMENTAL METHOD The 5.08-cm-diam discs were machined from 0.317-cm-thick plate. Chemical analysis information for the type 304 SS material is given in Table I. The discs were ground flat on metallographic paper and given a final polish on Linde B abrasive. A thin support rod was threaded into the disc and the region around the threads was fused under an inert gas. The support rod was fitted with a quartz protection tube and then was attached to a supporting shaft which passed through a rotary push-pull vacuum seal. The disc and supporting shafts were dynamically balanced prior to insertion into the furnace tube. The apparatus is shown schematically in Fig. 1. The 58 pct Bi-42 pct Sn eutectic alloy melts were prepared from 99.995 pct pure Bi and Sn by fusing the components in a 7-cm-ID Pyrex crucible. The system in which the melts were made was evacuated to a pressure of 1 x 10-6 Torr and back-filled with purified argon several times before melting the charge. The ingot was reweighed and placed in a slightly larger-diameter Vycor crucible used in the dissolution runs. A run was started by lowering the disc into the liquid

Jan 1, 1968

By C. H. Samans, G. F. Tisinai, J. K. Stanley

The addition of nitrogen (0.10 to 0.20 pct) to Fe-Cr-Ni alloys of simulated commercial purity results in a real displacement of the u phase boundaries to higher chromium contents. The effect is small for the (Y + s)? boundary, but is pronounced for the (y + a +s)/(y + a) boundary. Although there is an indication of an exceptionally large shift of the n boundaries to higher chromium contents, especially in steels with nitrogen over 0.2 pct, the major portion of this apparent shift results from the fact that carbide and nitride precipitation cause "chromium impoverishment" of the matrices. The effect of combined additions of nitrogen and silicon to the Fe-Cr-Ni phase diagram is demonstrated also. Nitrogen can nullify the effect of about 1 pct Si in shifting the (y + o)/? phase boundary to lower values of chromium at all nickel levels from 8 to 20 pct. NItrogen can nullify this U-forming effect of about 2 pct Si at the 8 pct Ni level, but not at the 20 pct Ni level. The alloys studied were in both the cast and the wrought conditions. There are indications that the u phase forms more slowly in the cast alloys than in the wrought alloys if both are in the completely austenitic state. The presence of 6 ferrite in the cast alloys accelerates the formation of U. Cold working increases the rate of o formation in both cast and wrought alloys. THE major improvement in Fe-Cr-Ni austenitic alloys in recent years has been in the addition or removal of minor alloying elements to facilitate better control of corrosion resistance, sensitization, and heat resistance. One shortcoming of the austenitic Fe-Cr-Ni alloys, which never has been completely circumvented, is their propensity toward u formation. In the AISI-type 310 (25 pct Cr-20 pct Ni) and type 309 (25 pct Cr-12 pct Ni) steels, sufficient amounts of u phase can form, if service or treatment is in a suitable temperature range, to cause severe embrittlement. Also, there is a growing conviction that this phase may be contributory to some unexpected decreases in the corrosion resistance of certain of the 18 pct Cr-8 pct Ni-type steels. The present paper discusses the effect of nitrogen additions on the location of the (r+u)/d and the (y+a+u)/(y+a) phase boundaries in the ternary Fe-Cr-Ni system, for cast and wrought alloys of simulated commercial purity, and in similar alloys containing up to about 2.5 pct Si. The objective is to define compositional limits for alloys which will not be susceptible to u formation when used near 1200°F (650°C). An excellent review of the early studies of the u phase in the Fe-Cr-Ni system has been compiled by Foley.1 Rees, Burns, and Cook2 have determined a high purity phase diagram for the ternary system, whereas Nicholson, Samans, and Shortsleeve3 are- stricted themselves to a portion of the simulated commercial-purity phase diagram. Both groups of investigators show almost an identical position for the commercially significant (y+u)/y phase boundary. Further comparison of the work of the two groups indicates that, below the 8 pct Ni level, the commercial alloys have a decidedly greater propensity toward u formation than the high purity alloys. The two groups of workers agreed that both the AISI-type 310 (25 pct Cr-20 pct Ni) and the type 309 (25 pct Cr-12 pct Ni) steels are well within the (y+~) region and that the 18 pct Cr-8 pct Ni-type alloys straddle the U-forming phase boundaries. Nicholson et al.3 showed, in addition, that these boundaries shift toward lower chromium contents if greater than nominal amounts of silicon or molybdenum are added. The effect of nitrogen on the location of the s phase boundaries in the Fe-Cr-Ni system has not been known with any certainty. In 1942, an approach to this problem was made by Krainer and Leoville-Nowak,&apos; but at that time they apparently were unaware of the slow rate of s formation in strain-free samples and aged their samples for insufficient times, e.g., 100 hr at 650°C (1200°F) and 800°C (1470°F). For this reason, it would be expected that their (y+ u) /y boundary would be shifted toward lower chromium contents (restricted ?-field) when equilibrium conditions were approximated more closely. Procedure for Studying the Alloys The alloys used were prepared in the following way: Heats of 200 lb each were melted in an induction furnace. A 5 lb portion of each heat was poured into a ladle containing an aluminum slug for de-

Jan 1, 1955

By R. H. Jacoby, V. J. Berry

When dry gas is injected into a reservoir containing a volatile crude oil, a significant amount of the reservoir liquid phase may become vaporized. The resultant rich gas phase, when subsequently produced, con-tributes to tank oil production. This contribution assumes greater importance, the more volatile the oil in-olved. Oil recovery may be sub-stantially greater than that predicted by conventional frontal-drive methods, which do not consider the vaporization equilibrium between the reservoir oil phase and the injected gas. A calculation method has been (developed to account for vaporiza-. tion of the reservoir liquid phase during gas-injectiorz operations, and for the tank oil production which results from this factor. Recovery performance calculations are presented for a reservoir containing a highly volatile oil. Tank oil recovery is calculated to be about twice that predicted by the use of the conven-rional frontal-drive equations. In contrastto usual pressure maintenance performance results, in which the gas-oil ratio rises at an increasing rate after gas breakthrough, the pre-dicated gas-oil ratio rises rapidly to about 12,000 scf/bhl and then rises much less rapidly. During gas inje-tion, most of the reservoir liquid phase contacted is evaporated by the dry injection gar. The gas-oil ratio during this period is dependent upon reservoir pressure. The higher the operating pressure, the lower the gas-oil ratio. The predicted behavior i., in accordance with laboratory PVT tests made to .simulate the vaporization behavior. In addition to recovery performance predictions, results of the calculation procedure include complete wellstrearn composition data of value in the design of gasoline plant facilities often used in con-rrction with gas-injection operations. INTRODUCTION In the cycling of gas-condensate reservoirs, dry gas is injected to maintain reservoir pressure during wet gas production and to thereby eliminate or reduce ultimate loss of liquids due to retrograde condensation within the reservoir. Gas injected into crude oil reservoirs has a dual function. It displaces oil to the producing wells and at the same time serves to partially or fully maintain reservoir pressure. Oil shrinkage which would occur upon pressure reduction is thereby minimized or eliminated. Accepted calculation methods are available&apos;= for predicting recovery performance of either gas-condensate reservoirs or crude oil reservoirs which are being subjected to gas injection. In gas-condensate reservoirs, any retrograde liquid formed does not flow, and it is necessary to account only for the vaporization equilibrium between this liquid and the injected gas. Conversely, in normal crude oil reservoirs, both the oil and gas phases flow, but it has not been considered necessary to account for any vaporization of the reservoir liquid which might occur upon contact with the dry injection gas. Recently, high shrinkage reservoir fluids known as "volatile oils" have been found in increasing amounts. These oils are characterized by tank oil gravities above 4.5" API, solution gas-oil ratios above 1,000 scf/bbl, and reservoir volume factors above two. Special techniques have been devised for predicting depletion performance of reservoirs containing such oils.74.5.1; One of the characteristics of reservoirs producing volatile oils is that the reservoir gas phase carries a significant amount of oil which is recoverable as stock-tank liquid. This unusual vaporization be-havior implies that an appreciable amount of reservoir liquid would be vaporized upon contact of the oil with dry injection gas. In a gas-injection operation, tank oil recovery would he obtained not only through frontal displacement of the reservoir liquid by the injected gas, but also through production of the rich gas phase. This means that not only are improved methods needed to predict recovery of such oils from reservoirs undergoing gas injection, but it would also be expected that high oil recoveries might be obtained by such operations. When relatively dry gas is injected in to a volatile oil reservoir, phase equilibrium between the injected gas and the reservoir oil will tend to bc established. Initially the most volatile components, such as propane, the butanes and the pentanes, will account for most of the material transferred from the oil phase to the gas phase. As the partially stripped oil phase is contacted with additional dry injection gas, the heavier intermediate components, such as the hexanes, the heptanes and the octanes, will gradually transfer to the gas phase in increasing amounts. This is because the supply of lighter components in the oil phase dwindles due to the stripping action of the injected gas. This stripping action will usually continue to be effective down

By B. E. Eakin, R. T. Ellington

An apparatus and a procedure for determining the viscosity behavior of hydrocarbons at pressures up to 10,000 psia and temperatures between 77 and 400° F are described. The equipment is suitable for measuring viscosity of either the liquid or vapor phases or the fluid above the two-phase envelope for systems exhihiting retrograde phenomena, according no the phase state of the system within these ranges of temperature and pressure. Equations are developed for calculation of viscosity from the experimental measurements, and new data for the viscosities of ethane and propane at 77° F are reported. INTRODUCTION With the advent of higher pressures and temperatures in industrial processes and deep petroleum and natural gas reservoirs, demand has increased for accurate values of physical properties of hydrocarbons under these conditions. Proportionately, more frequent occurrence of natural gas and condensate-type fluids is encountered as fluid hydrocarbons are discovered at greater depths. This increases the importance, to the reservoir engineer, of being able to predict accurately the physical properties of light hydrocarbon systems in the dense-gas and light-liquid phase states. Reliable gas viscosity data are limited primarily to measurements made on pure components near ambient temperature and at low pressures. Few investigations have been reported for high pressures, and except for methane, data on light hydrocarbons are subject to question. This is demonstrated by the large discrepancy between sets of data on the same component reported by different investigators. For mixtures in the dense gas and light liquid regions and for fluids exhibiting retrograde behavior there are very few published experimental data. Viscosity data for methane have been reported by Bicher and Katz,1 Sage and Lacey,12 Comings, et al,3 Golubev,3 and Carr,3 with good agreement among the last three sets of data. Comings, Golubev and Carr utilized capillary tube instruments for which the theory of fluid flow is well established. The theory permits calculation of the viscosity directly from the experi- mental data and dimensions of the instrument alone. Sage and Lacey, and Bicher and Katz used rolling-ball viscometers. The theory of the rolling-ball viscometer has not been completely established, and these instruments presently require calibration by use of fluids of known viscosity behavior before viscosities of test fluids can be measured. To obtain accurate data it is necessary that the rolling-ball viscometers be calibrated by use of fluids of density and viscosity similar to the test fluids, a difficult selection for the gas phase. From the methane data and experimental tests on various natural gases, Carr developed a correlation for predicting the PVT behavior of light natural gases.2,3,4 This correlation was based on data for a very limited composition range; its application to rich gases and condensate fluids is questionable. The object of this investigation is to develop an instrument which can be used to obtain viscosity data at reservoir temperatures and pressures, for rich gases, condensate-type systems above the two-phase envelope and light liquid mixtures. These data will be used in an effort to develop correlations to represent the viscosity behavior of these fluids. APPARATUS In a previous viscosity study Carr2 utilized a modified Rankine capillary viscometer configuration," Fig. 1. In this instrument the gas to be tested is forced through the capillary tube in laminar flow by motion of a mercury pellet in the fall tube, the measured displacement time being that required for the mercury slug to move between the brass timer rings. The viscometer is constructed of glass and mounted in a steel pressure vessel. The test gas pressure in the viscometer is balanced by an inert gas (usually nitrogen) in the vessel. Excellent results have been obtained with instruments of this type, with Carr2 and Comings5 reporting repro-ducibilities of 99.5 to 99.3 per cent and an estimated absolute accuracy of 99 per cent. However, these instruments have limitations which have precluded their use for liquids. The need for maintaining a balance between pressures of the test fluid and inert gas in the viscometer vessel presents operating problems, and requires charging the test fluid to the viscometer very slowly. The principle drawback to the Rankine unit is behavior of the mercury slug which provides the pressure differential across the capillary. When even trace quantities of propane or heavier hydrocarbons are present in the test gas, the mercury tends to subdivide

By A. Ghosh, T. B. King

The kinetics of the discharge reaction: 20&apos;- (in silicate melt) = O,(g) + 4e- at a platinum anode in lithium silicate melts have been studied al 1350°C by galvanostatic methods. Plots of the sleady-state overpotential, q, as a function of the logarithm of the current density, i, showed injlections and were linear only at high current densities. The value of the overpotential was influenced by bubbling gas through the electrolyte. The ocer potential was also studied as a function of time. The rise and decay of overpotential were very slow processes. At low current densities transport is the likely rate-controlling process but at high current densities passivation of the electrode, Presumably by an oxide film on the surface, seems to be a contributory functor. IT is well-established that molten silicates behave as electrolytes&apos;5 and, except in a few cases,6 conduction is entirely ionic. Moreover, it is supposed that a possible, and perhaps predominant, mechanism for phase boundary reactions between metals and slags is similar to that in corrosion whereby anodic and cathodic processes occur at unrelated sites, the metal serving to conduct electrons.1&apos;8 Thus electrochemical studies of some slag-metal reactions would seem to be a useful way to diagnose the rate-controlling steps in the overall reaction. The electrochemical method is, in principle, a better diagnostic tool than the direct chemical method for the following reasons: 1) The partial electrochemical reactions, which are simpler than the overall reaction, may be studied individually. 2) The rate of reaction can be controlled at will and independently of the concentrations of reactants. 3) Fast reactions can be studied by relaxation methods.&apos; Esin and his coworkers5&apos;10"12 have pioneered such studies in silicates and have deviloped some ingenious techniques. Not all of their findings, however. can be accepted without a good deal of further work. In this investigation, the kinetics of the oxygen discharge reaction: 202- (in silicate melt) = Oz(g) + 4e- [I] at a platinum electrode were studied by both steady-state and transient galvanostatic techniques. Interest in this reaction was first developed as a result of the findings of Fulton and chipman13 that the reduction of silica, in a silicate slag, by carbon, dissolved in liquid iron, is a very slow reaction. Subsequent work, for example, by Rawling and ~lliott,&apos;~ has demonstrated that the reaction under these conditions must be slow, because the rate is limited by diffusion of oxygen in the iron to the metal-crucible phase boundary at which a CO bubble may be nucleated. Further work by Tarassof,&apos;~ in which the reduction of silica by aluminum dissolved in copper was studied, has shown that under these conditions, where carbon monoxide evolution is not involved, control of the reaction rate resides in diffusion of silica in the slag phase. However, there is no practical way of inducing sufficient convection in the system to make it clear that the phase boundary reaction is indeed fast. The overall reaction of silica reduction involves the discharge of silicon ions at cathodic sites and oxygen ions at anodic sites. In the examples cited, the discharged ions are dissolved in a liquid metal. In the present study of oxygen ion discharge, gaseous oxygen may be evolved at high current densities or oxygen may simply dissolve, possibly as oxygen molecules, in the silicate at very low current densities. The discharge of an oxygen ion at an anode must, in silicates less basic than the orthosilicate composition, be preceded by a reaction in the vicinity of the electrode, such as: which makes oxygen ions available. Platinum was chosen as the working electrode since it is comparatively inert to oxygen and is, therefore, expected to come rapidly into equilibrium with the electrolyte and with gaseous oxygen. Minenko, Petrov, and Ivanova16 have measured the electromotive force at a platinum electrode in molten silicates as a function of the partial pressure of oxygen in the atmosphere, the concentration of oxide ions in the melt, and the temperature. They found platinum to behave as a reversible oxygen electrode. At two different oxygen pressures, Po2 (I) and Pq (11). the electromotive force is given by: where F is the Faraday constant, equal to 23,060 cal per v equivalent, indicating that the electrode reaction is as written in Eq. [I.]. Platinum has been similarly used in molten silicates by other inve~ti~ators. "&apos;~~ In this investigation platinum was used only as an anode, since a current deposits other elements on its surface and changes its characteristics.

Jan 1, 1970

By J. Lewis, J. Harper, M. F. Ashby

A grain boundary in a metal interacts with second-phase particles, which exert a pinning force (first estimated by Zener) on the boundary opposing its motion. We have computed the shape of boundaries which interact with more or less spherical second-phase particles and have constructed a soap-film model to reproduce the shape of the boundary surface. An important result is that measurement of this shape allows the pressure, or driving force, on the boundary to be measured. We hare applied this technique to grain boundaries in two alloys and hate measured the pinning force exerted by single second-phase jwrticles on the boundaries. It is in good agreement with Zener&apos;s estimate. J\. boundary between two grains, or two bulk phases, interacts with small inclusions or particles of a second phase, whether they are gas or solid. This interaction means that the boundary, forced to migrate by a difference in free energy between the material of the two grains or phases which it separates, exerts a force on a particle which it touches, tending to drag it forward. (The movement of inclusions through metals under the influence of this force, has, in fact, been observed. 1-3) Equally, the particle can be thought of as exerting a pinning force on the boundary, tending to hold it back. Zener (in a celebrated private communication4) first realized that this interaction, and the resulting pinning force, existed. His calculation of its magnitude was crude but adequate: a spherical inclusion of radius r blanks off an area nr2 of the boundary on which it sits; since the boundary has an energy of rMM x per unit area, the blanlung-olf decreases the energy of the system by MM: this energy is returned to the system if the boundary is pulled free from the inclusion— a forward movement of the boundary by a distance r will do this—so that the maximum pinning force is Trrym.M- A similar argument can be made for inter-phase boundaries. The nature of the particle itself did not enter this, or two subsequent treatments.5,6 When it is considered, tic leifthe energyoftheb a) The boundary may enter and pass through the particle if the energy of the boundary is lower within the particle than in the matrix. Fig. l(r/). Certain coherent precipitate particles behave like this. h) More usually, the boundary will bend round the particle, enveloping and bypassing it. Fig. l(b). In doing so, it changes the structure and energy of the interface between the particle and its matrix. This means that the boundary does not touch the particle surface at right angles, as early treatments assumed,5&apos;9 but at some angle a which depends on this change in surface energy of the particle, and can be calculated from the equilibrium of surface tensions. Most precipitate particles and inclusions behave like this. Gas bubbles or liquid drops can be regarded as belonging to either group. The progress of bypassing is conveniently measured by the angle shown in Fig. 1. When the nature of the particle is ignored, its maximum pinning force is exerted when - 45 deg. When it is considered, this critical value of is found to depend on a and thus on the nature of the particle. The maximum pinning force lies between nryMM and 2jtjMM (Appendix 1). not very different from Zener&apos;s result. In reality, a boundary between crystals has a specific energy and tension which varies with the orientation of the boundary. Furthermore, recent experiments7 indicate that such a boundary is not atom-ically smooth, but has steps on it: migration of the boundary corresponds to the sweeping of these steps across the boundary surface, like the Frank model of crystal growth from the vapor. This means that the interaction of a boundary with particles should really be considered in terms of the way in which particles hinder the movement of these steps. To suppose a grain boundary or interphase boundary to be smooth, and to ignore the variation of its energy with orientation, is to liken it to a soap film. The advantage of this soap film approximation, which we have used throughout this paper, is that interaction energies and boundary shapes can be calculated easily. We have done this by numerical computation and by using a soap film model, and have compared the results with grain boundaries in an aluminum-based and a copper-based alloy. It turns out that the shape of the boundary which bulges between particles allows the pressure an it to be calculated; that is, the local driving force an the boundary can be measured. This has allowed us to check the Zener relationship experimentally.

Jan 1, 1970

By J. F. Enrietto

Internal friction measurements were used to determine the effect of manganese on the solubility and precipitation kinetics of nitrogen. Manganese, in concentrations up to 0.75 pct, has little effect on the solubility at temperatures above 250°C. On the other hand, at Concentrations as low as 0.15 pct, manganese inhibits the formation of iron nitrides, especially Fe4N, even though it may not form a precipitnte itself. The precipitation and solubility of carbides and nitrides have been extensively investigated in the pure Fe-C and Fe-N systems.1-3 In recent years, some effort has been ispent in studying the influence of substitutional alloying elements on the behavior of carbon and nitrogen in ferrite.4 -7 In particular Fast, Dijkstra, and Sladek have investigated the effect of 0.5 pct Mn on the internal friction and hardness during the quench aging of Fe-Mn-N alloys.&apos;, &apos; They found that at low temperatures (below 200°C) the presence of 0.5 pct Mn greatly retarded quench aging. For example, after 66 hr at 200°C very little precipitation had taken place in the iron alloyed with manganese, whereas precipitation was complete after a few minutes in a pure Fe-N alloy. The effect of varying the manganese content and the details of the precipitation process were not mentioned in these papers. Fast&apos; postulated that manganese causes a local lowering of the free energy of the lattice with a resulting segregation of nitrogen atoms to these low energy sites. The segregated nitrogen atoms are bound so tightly to the manganese atoms that they cannot form a precipitate. The internal friction measurements of Dijkstra and Sladek tended to confirm the concept of segregation of nitrogen around manganese atoms, and the increase in free energy on transferring a mole of nitrogen atoms from a segregated to a "normal" lattice site was computed to be - 2800 cal. Dijkstra and Sladek9 distinguished between two types of precipitates: ortho, a nitride of appreciably different manganese content than that of the matrix, and para, a nitride with a manganese content essentially that of the matrix. With each type of precipitate a solubility, again designated ortho or para, can be associated. Since the internal friction maximum in alloys which were aged several hours at 600" C dropped almost to zero, Dijkstra and Sladek9 concluded that the ortho solubility must be very low. The effect of temperature on the ortho and para solubilities has no1: been investigated. There are obviously several gaps in our knowledge concerning the influence of manganese on the behavior of nitrogen in a-iron. It was the purpose of the experiments described in this paper to determine the following: 1) The ortho and para solubilities of nitrogen as a function of temperature. 2) The details of the precipitation process at elevated temperatures. 3) The effect of varying the manganese concentration on the above phenomena. EXPERIMENTAL PROCEDURE Internal friction is conveniently employed in studying the precipitation of nitrides and/or carbides from a -iron because it is one of the few parameters, perhaps the only one, which is not affected by the presence of the precipitate itself. For this reason, internal friction techniques were heavily relied upon in the present experiment. A) Preparat of -. All specimens were prepared from electrolytic iron and electrolytic manganese. Alloys containing 0.15, 0.33, 0.65, and 0.75 wt pct Mn were vacuum melted and cast into 25 lb ingots. After being hot rolled to 3/4 in. bars, the ingots were swaged and drawn to 0.030 in. wires. The wires wen? decarburized and denitrided by annealing at 750° C for 17 hr in flowing hydrogen saturated with warer vapor. To obtain a medium grain size, - 0.1 mm, the wires were then heated to 945oC, allowed to soak for 1 hr, furnace cooled to 750°C, and water quenched. Subsequent internal friction measurements showed that this procedure reduced the nitrogen and carbon concentrations of the alloys to less than 0.001 wt pct. The wires were nitrided by sealing them in pyrex capsules containing anhydrous ammonia and annealing them for 24 hr at 580°C, the nitrogen being retained in solid solution by quenching the capsule into water. Immediately after quenching, the wires were stored in liquid nitrogen to prevent any precipitation of nitrides. By varying the pressure of ammonia in the capsule, it was possible to produce any desired nitrogen concentration. B) Internal Friction. The internal Friction measurements were made on a torsional pendulum of the Ke type,&apos;&apos; a frequency OF 1. or 2 cps being used. For

Jan 1, 1962

By A. R. Orban, R. B. Engdahl, J. D. Hummell

The initial phase of a research program on smoke abatement from Bessemer converters is described. In work sponsored by the American Iron and Steel Institute, a 300-lb experimental Bessemer converter was assembled to simulate blowing conditions in a commercial vessel. Measurements of smoke and dust were also made in the field on a 30-ton commercial vessel. During normal blows the dust loading from the laboratory converter averaged 0.51 lb per 1000 lb of exhaust gas. This was similar to the exhaust-gas loading of a commercial vessel. The addition of hydrogen to the blast gas of the laboratory converter caused a decided decrease in smoke density. Smoke was also reduced markedly when methane or ammonia was added instead of hydrogen. The research is continuing on a bench-scale investigation of the mechanism of smoke formation in the converter process. DURING the past 2 years, on behalf of the American Iron and Steel Institute, Battelle has been conducting a research program on the control of emissions from pneumatic steelmaking processes. The objective of the research program is to discover a practical method for reducing to an unobjectionable level the emission of smoke and dust from Bessemer converters. PRELIMINARY INVESTIGATION Although conceivably some new collecting technique may be devised which would be economically practicable for cleaning Bessemer gases, no such system based on presently known principles seems feasible because of the extremely large volume of high-temperature gases involved. Hence, the research is being directed toward prevention of smoke formation at the source. A thorough review was first made of former work to determine the present status of the cleaning of converter gases. No published work was found on work done in the United States on collecting smoke or on preventing its formation in the bottom-blown, acid-Bessemer converter. In Europe, however, a number of investigations have been made on the basic-Bessemer converter. Kosmider, Neuhaus, and Kratzenstein1 conducted tests on a 20-ton converter to obtain characteristic data for dust removal and the utilization of waste heat. They concluded that because of the submicron size of the dust, special equipment would be necessary to clean the exhaust gases. Dehne2 conducted a large number of smoke-abatement experiments at Duisburg-Huckingen in a 36-ton Thomas converter discharging into a stack. A number of wet-scrubbing and dry collectors were tried unsuccessfully. A waste-heat boiler and electrostatic collector with necessary gas precleaners was felt to be the best solution for this particular plant. Meldau and Laufhutte3 determined that the particle size was all below 1 µ in the waste gas of a bottom-blown converter. Sel&apos;kin and zadalya4 describe the use of oxygen-water mixtures injected into a molten bath in refining open-hearth steel. They claim that with use of oxygen-water mixtures the amount of dust formed was reduced between 33.3 and 20 pct of its previous level, and emission of brown smoke almost ceased. Pepperhoff and passov5 attempted unsuccessfully to find some correlation between the optical absorption of the smoke, the flame emission, and the composition of the metal in a Thomas converter in order to determine automatically the metallurgical state in the melt. In a recent U. S. Patent (NO. 2,831,762)&apos; issued to two Austrian inventors, Kemmetmuller and Rinesch, the inventors claim a process for treating the exhaust gases from a converter. By their method the inventors claim that the exhaust gases from the converter are cooled immediately after leaving the converter to a degree that oxidation of the metal vapors and metal particles to form Fe2O3 is inhibited in the presence of surplus oxygen. Gledhill, Carnall, and sargent7 report on cleaning the gases from oxygen lancing of pig iron in the ladle. They claim the Pease-Anthony Venturi scrubber removed 99.5 + pct of the smoke, thereby reducing the concentration to 0.1 to 0.2 grain per cu ft, which resulted in a colorless stack gas after the evaporation of water. Fischer and wahlster8 developed a small basic converter and compared the metallurgical behavior of the blow with that of a large converter. Later work by Kosmider, Neuhaus, and Hardt9 on the use of steam for reduction of smoke from an oxygen-enriched converter confirmed that the cooling effect of steam is detrimental to production. From review of all of the published information on the subject, it was concluded that a practical solution to the smoke-elimination problem had not been found. Accordingly, it was deemed desirable to investigate the feasibility of preventing the initial formation of smoke in the converter.

Jan 1, 1961

By H. C. Theuerer

GeC14 may be purified by extraction with HCI and chlorine. The process is as effective for the removal of AsCI:, as the more cumbersome distillation methods usually used for this purpose. GERMANIUM for semiconductor use contains impurities at levels no higher than a few parts in ten million. Material of this quality is obtained from highly purified GeC1, by hydrolysis to the oxide and reduction of the oxide in hydrogen. When purifying GeCl,, AsC1, is the most difficult impurity to remove. This is usually accomplished by multiple distillation procedures.1-3 AsC1, may also be removed from GeC1, by extraction with HC1.1-4 Reducing the arsenic to low concentrations is not practicable, however, because of the large number of extractions needed. This paper discusses a new method for the removal of arsenic from GeC1, by extraction with HC1 and chlorine. The method is rapid, leads to little loss of germanium and is at least as efficient as the distillation procedures currently being used. Theory of Extraction Procedures In the simple extraction of GeC1, with HC1, the following reaction occurs ASCl8G8C1 D AsCl3rc1 at equilibrium CA/Cn= K, where K is the distribution coefficient, and C, and C,, are the molar concentrations of AsC13 in HC1 and GeCl,, respectively. The materials balance equation for this reaction is VACA + vncn = VnC,, where V, and Vn are the volumes of HC1 and GeCl4, respectively, and C, is the initial concentration of AsC13 in GeC1,. From this it can be shown that for multiple extractions where C,, is the concentration of AsC13 in GeC14 after n extractions, and r is the ratio of V, to V,,. It is assumed that r is maintained constant, that equilibrium is established during each extraction, and that K is independent of the AsCl3 concentration. By saturating the system with chlorine, the following reaction occurs in the aqueous phase AsCl3 + 4H2O + Cl2 D H5AsO4 + 5HC1 at equilibrium K&apos; = ------------ ai - a4 h2u - aet2 where a is the activity of the various components. The effect of this reaction is to reduce the concentration of the AsC1, in the aqueous layer and, therefore, to promote further extraction of this component from the GeC1, layer. If the arsenic acid remains entirely in the aqueous phase, the net effect of this reaction is to promote the removal of arsenic from the GeC11. The materials balance equation for extraction with HC1 and chlorine with the foregoing reaction is, then, VaCC + VACA + VACn = VnCo where C,. is the molar concentration of H3AsO, in the HC1. With the added assumptions that the activities of AsC13 and H8ASO4 in the aqueous phase are equal to their molar concentrations, it can be shown that for n extractions Cn/Cu = (1/rkK + rK + 1) n where k - K1 a4h2o - acl2/aoncl. It can be seen by comparing Eqs. 1 and 2 that if k is large, the removal of AsC1, by HC1 extraction will be greatly improved by the addition of chlorine. Dilution of the HCI used in the extraction with chlorine would also favor the separation. This, however, would increase the loss of GeCl,, which is undesirable. Experimental Procedure Germanium prepared from oxide of semiconductor purity is n-type with resistivities greater than one ohm-cm. The resistivity is controlled by the donor concentration, which is —lo-: mol pct. Germanium prepared from material with added arsenic will have lower resistivity commensurate with the arsenic concentration. With such material, at arsenic concentrations above 10-1 mol pct the resistivity is controlled by the added arsenic, and effects due to other impurities initially in the oxide are negligible. In this investigation GeO, of semiconductor purity was converted to GeCl,, and -0.01 mol pct As was added. This material was used for the extraction experiments and the purification attained determined by a comparison of the resistivity data for samples of germanium prepared from the initial and purified GeC1,. A method for calculating the arsenic concentration from the resistivity data is discussed later. The details of the experimental procedures used are as follows: Two hundred and thirty cu cm GeC1, were prepared by the solution of GeO, in HC1, followed by

Jan 1, 1957