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By Ursula E. Wolff

The precipitation of carbides from an alloy containing 33 pct Ni, 21 pct Cr, balance iron, was investigated electron microscopically by means of extraction replicas and thinned metal foils. Annealing temperatures ranged from 565°to 870°C and up to several thousand hours. M23C6 precipitated in pain boundaries, incoherent and coherent twin boundaries in that sequence. The orientation relationship between carbides and austenite matrix was determined and correlated with the morphology of the carbides and with the type of boundary in which precipitation occurred. In large-angle grain boundaries, as well as in coherent twin boundaries, the carbides had the same orientation as one of the adjacent pains. These carbides formed sheets of individual flakes with shapes related to the orientation of the boundary. In incoherent twin boundaries carbides precipitated in ribbons composed of pavallel rods. An unidentified subcarbide was found to precede precipitation of M23C6 in these boundaries. The M 23 C6 rods had a kind of fiber texture with (110) parallel to the long dimension of the rods and ribbon, and with orientations of both of the adjacent twin-related austenite crystals Predominant in the texture of the carbide. A hard sphere crystal model has been used to discuss orientation and morphology of the carbides in terms of free volume and vacancies available in the boundaries. A number of papers have dealt with the morphology of chromium carbide (M23 C6) precipitated in austenitic stainless steels.1"7 In all these investigations, the carbides were examined in the electron microscope by means of extraction replicas. With this technique, the carbides retain the spatial distribution they had in the bulk sample. However, since the matrix is dissolved in the process, the particles can turn in an unpredictable way; and the orientation relationship between matrix and carbides cannot be established. In this paper the results of studies on extraction replicas and on thinned metal foils are reported. These studies were undertaken to determine the matrix-to-car bide orientation relationship, and to correlate the orientation of the carbides with their morphology. PROCEDURE The material used was an austenitic alloy with 33 pct Ni, 21 pct Cr, balance iron, containing approximately 0.05 pct C. Coupons of 1.25-mm sheet were first solution-annealed at 1050°C for 15 min and air-cooled. Then, to precipitate the carbides, samples were isothermally annealed in the range from 565" to 870°C for times up to several thousand hours. All further specimen-preparation procedures were carried out after the final anneal. Carbon extraction replicas from polished and etched surfaces were made with 10 pct bromine in methyl alcohol.' Thin foils were prepared from punched-out 3-mm-diam disksg which fit into the electron-microscope holder. The disks were prethinned by grinding to approximately 0.5 mm thickness, and then electro-polished in a polytetrafluoroethylene holder1' with a solution containing 5 pct perchloric acid in acetic acid to which 10 g per 1 Cro3 and 5 g per 1 nickel chloride were added (etchant modified from that of Briers et al."). This solution dissolves neither the carbides nor the austenite around the carbides preferentially. By using extraction replicas, electron micrographs and selected-area electron-diffraction patterns were taken from the same carbide arrays. By using thin foils, electron micrographs were made from a grain boundary area containing carbides. Electron-diffraction patterns were then taken from the same area and from each of the adjacent grains separately. In this manner, the orientation of each grain could be determined without interference by the carbide pattern. A peculiarity of extraction replicas should be pointed out. After the matrix is etched away, the carbide arrays float freely in the etching and washing solutions, and are held in place only at the anchoring points in the carbon replica. When the replica is picked up with a screen the carbide arrays tend to flip to one side. Thus, while the surface features are preserved, the original arrangement of the carbides may severely and unpredictably be disturbed whenever the specimen contains large amounts of interconnected carbides. Nevertheless, it is possible to correlate the different morphologies of the carbides with the type of boundary in which they have precipitated. RESULTS 1) Extraction Replicas. Fig. 1 shows that the grain boundaries usually are curved, multicornered surfaces of random orientation. The coherent twin boundaries (which are (111) planes) cut a grain into parallel slices. Incoherent twin boundaries occur at the ends and on the steps of twins and are often narrow, parallel-sided strips which are much longer than they are wide. Different morphologies can clearly be distinguished for the M23Ce carbides precipitated in each of these types of boundaries, and agree well with those observed by kinzel.2 The kinetics of this precipitation has been investigated." The first carbides precipitate in junctions of three grain boundaries and fan out from there into the adjoining boundary surfaces, Fig. 2(a). These carbides are oriented randomly, Fig. 2(b), and become coarser and thicker as annealing time increases. The large-angle grain boundaries are next to fill

Jan 1, 1967

By C. C. Templeton, J. C. Rodgers

A key step in feed water treatment for generating wet steam for thermal oil recovery is the removal of calcium and magnesiunt hardness by cation-exchange series softening. Knowing the solubility of any scale forming salts in brines at elevated temperatures is necessary for fixing the level to which the feed water must be softened. Such calcium sulfate solubility data, previously not available above 392F, were determined by the authors in a flow equilibrium apparatus mud will be reported elsewhere. These data were used to develop a method for predicting the solubility of anhydrite in hot water or steam droplets for saturated steam pressures as high as 2,000 psig (637F). (The calcium sulfate solubility product is represented by a combination of two factors, one reflecting the effects of ionic strength and the other accounting for the effects of complex ion formation in either calcium-magnesium-rich or sulfate-rich brines.) The method is applied to a calcium-magnesium-rich brine If moderately high salinity from a pilot hot-water flood, I nd to several sulfate-rich, low-salinity feed waters and l lowdown (cooled steam droplets) samples from steam s ak operations. The predicted calcium hardness levels corresponding to the calcium sulfate solubilities agreed reasonably well with the results of laboratory solubility determinations run on the field samples. Further testing of the method is needed for brines of other composition classes. Existing field cation exchange softeners in the cases tested are performing adequately since all the samples were found to be undersaturated with respect to calcium sulfate at their operating temperatures. Introduction Prevention of scaling caused by precipitation of calcium sulfate (anhydrite) is of considerable concern in connection with thermal recovery processes using wet steam or hot water. To avoid anhydrite precipitation in a heated system, an engineer must keep the product of the calcium and sulfate concentrations in the water or steam droplets below the value of the solubility product of anhydrite for the temperature and brine composition in question. Usually it is most practical to keep the concentration product lower than the solubility product by keeping calcium low in the presence of high sulfate, or by keeping sulfate low in the presence of high calcium. This can be done by a choice of combinations of natural waters and water treatment processes (such as series cation exchange softening to remove calcium). Until recently, few anhydrite solubility data, particularly for solutions containing other salts, were available for temperatures above 392F (211 psig steam); Marshall, Slusher and Jones' studied the CaS0,-NaCI-H,O system up to 392F and surveyed the work of previous investigators. To model natural brines, one needs to study the solubility of anhydrite in aqueous solutions of sodium chloride, sodium sulfate, calcium chloride, magnesium chloride and their mixtures. Since steam pressures as high as 2,000 psig (637F) may be involved in thermal oil recovery projects, a solubility study was conducted between 482 and 617F.' Discussed in this paper is the application of these data to the prediction of anhydrite precipitation in some practical steam soak and hot-water injection projects. Any simple method for predicting the solubility of an inorganic compound over a wide range of temperatures and solution compositions must be based on some assumptions, and therefore must yield approximate results. On the one hand, natural brines contain too many ionic species for all to be included in a simple scheme; on the other hand, there is no adequate theoretical basis for the exact prediction of solubility in even simple solutions of mixed electrolytes. However, it is possible at a given temperature to base a reasonable prediction scheme on two phenomena:'-' the increase in solubility with increasing total concentrations of all ions (as measured by the ionic strength; see the Appendix), and an increase due to formation of cornplexes between calcium ions and sulfate ions and between sulfate and magnesium ions. Stiff and Davis' developed such a scheme for predicting the solubility of gypsum (CaSO, ¦ 2H,O) in brines at temperatures u~ to 212F. However, for the higher temperaturei of present interest, the stable solid phase is anhydrite (CaSO,). This study involves a two-part method for predicting anhydrite solubility products. First, one predicts a value K, for a given ionic strength I and a given temperature T corresponding to mo./msr, = 1 from data of the CaS0,-NaCI-H,O system. Second, one determines a group of factors F .= F(Ca) • F(Mg) • F(SO,), where the individual factors account for increases in the solubility product due to complex formation by high concentrations, respectively, of calcium, magnesium and sulfate. Combining the two parts, one obtains the solubility product in molalities as

Jan 1, 1969

By John C. Martin

The discovery of a new gas reservoir demands that the planning of a sottnd well-spacing program be initiated early in the development stage. It is the purpose of this discussion to illustrate by actual field examples the application of basic well-spacing principles, previously developed for oil reservoirs, to the problem of well spacing in natural gas fields. These studies are presented for the field use of geologists and engineers who are concerned with the initial planning of the proper development of the newly discovered gas reservoir. INTRODUCTION The phenomenal growth of a vigorous natural gas industry emphasizes the increasing importance of natural gas as a source of energy, fuel, and raw materials to our nation's economy. Since 1945 marketed production of natural gas for the U. S. has increased 21/2 times to a record high of 10.6 trillion cu ft during 1957. As major participants in the gas industry, we share an added interest to develop and produce our natural gas reserves with constantly improved efficiency. The subject of well spacing is vitally important to the gas industry, for the well itself plays a significant role in the development of the natural gas reservoir and in control of the recovery process. Maximum utilization of wells is an integral part of sound conservation practices. The discovery of a new gas reservoir demands that careful choice of well location and well spacing be made and that the planning of a sound well spacing program be initiated early an the development stage. With the drilling of the first development well, efforts of the geologist and engineer must be directed toward acquisition of adequate technical evidence upon which a firm recommendation for a spacing program may be based. With this technical appraisal as a foundation, operators and state regulatory agencies jointly can go far in providing a framework for sound development of gas fields to achieve a program of conservation that avoids the unnecessary well. WELL-SPACING CONCEPTS Through laboratory and field investigations of the mechanism of the recovery of oil and gas, of fluid behavior, and of effective control of reservoir and well, a crystallization of ideas regarding reservoir behavior has emerged as a well-developed technology. Associated with a better understanding of the fundamental principles underlying reservoir and well behavior has been the growth of concepts concerning the role of wells and their spacing in the development and operation of an oil or gas reservoir. In addition to serving as outlets for the withdrawal of fluids from the reservoir, wells are recognized as having two other important functions: (1) providing access to the reservoir to obtain information concerning the characteristics of the reservoir and its fluids, and (2) serving as a means by which the natural or induced recovery mechanism may be effectively controlled. Beyond a minimum number of wells required to fulfill these two functions, additional wells will not increase recovery. With particular emphasis upon well spacing in oil reservoirs, many studies of the well spacing-recovery relationship have evolved the concept that the ultimate oil recovery is essentially independent of the well spacing.' These fundamental concepts are no different when regarding the role of wells and their spacing in the natural gas reservoir. They are equally applicable to the consideration of well spacing in gas reservoirs. For the gas reservoir, the problem of well spacing then revolves around the question of drainage and the degree or extent to which a well may drain gas from its surrounding reservoir environment. Theoretical and Experimental Work During the past 30 years, theoretical and experimental work carried on to study the physical principles involved in the flow of fluids through porous media has shed light upon the matter of drainage. Fundamental mathematical equations have been derived to describe the mechanism of flow of oil and gas through porous rocks. With the recent advent of high-speed digital computers, attempts have been made with mounting success to develop solutions, employing numerical techniques, to mathematical expressions that describe more rigorously the physical behavior and mechanism involved in the unsteady-state flow of compressible fluids, such as a gas, through porous rock. In 1953, Bruce, Peace-man, Rachford and Ricez published a stable numerical procedure for solving the equation for production of gas at constant rate. The results of these calculations are significant with respect to this matter of drainage, for they indicated (1) that depletion of the gas reservoir resulted in a drop in pressure at the extremity of the

By Richard R. Zupp, David A. Stevenson

An experimental method to determine the carbon activity in various solid solutions was developed. Samples of the solid solutions and an Fe-C reference sample were equilibrated with respect to carbon content by annealing them in a Hz-CHI gas mixture in a closed cell. The carbon content of each sample was determined, and the carbon activity was obtained from the carbon content of' the reference sample and well-established data for the dependence of the carbon activity on the carbon content in Fe-C alloys. The influence of vanadium on the activity of carbon in austenitic Fe-C-V alloys at 1000°C was determined by this technique. A general semiempirical correlation was developed to describe the dependence of the activity coefficient of- carbon on the composition of ternary austenitic iron-base alloys containing interstitial carbon and a substitutional solute. The correlation is particularly useful in predicting the partitioning of- carbon between various alloys. THE influence of substitutional solutes on the activity of interstitial solutes in metallic solid solutions is important in the theory of alloy phases. Solid solutions of particular interest are the fcc austenite phases of ternary Fe-C-Z systems, in which carbon is the interstitial solute and Z represents any substitutional solute. Carbon-activity data are also needed in these solid solutions to solve practical problems; the quantitative treatment of diffusion-controlled transformations' is an examvle. Experimental information is available for ternary austenitic solid solutions containing Mn,2 Si,2 and A16 as the substitutional solute. Because qualitative preliminary studies indicated that vanadium appreciably decreases the carbon activity,14 the present investigation was undertaken to determine the influence of vanadium on the activity of carbon in austenite. Although Flender and ever" measured carbon activity over a range of compositions in the Fe-C-V system, they were primarily interested in determining phase boundaries in the ternary system, so most of their measurements were outside the composition range for stability of single-phase austenite. Petrova and Shvartsman' also made limited measurements of the carbon activity in Fe-C-V alloys. The dependence of the carbon activity on the carbon content of an alloy may be determined by passing COCOz or H2-CH4 gas mixtures of carefully measured composition over the alloys, determining their equilibrium carbon content, and utilizing the relevant equilibrium relationships. smith16 used this technique with both CO-COZ and HZ-CH4 gas mixtures to determine the dependence of carbon activity on carbon content in pure Fe-C alloys. His data, which are in good agreement with those of some previous and subsequent investigators,6'17 but not with those of others,lfl'ao are usually considered to be the most reliable. Although the same method can be applied to Fe-C-Z alloys, it is difficult experimentally, because the appropriate gas mixtures are almost pure CO or Hz. A simpler technique is to simultaneously equilibrate the Fe-C-Z alloys and pure Fe-C alloys with the gas mixtures, and then use Smith's data to determine the carbon activity in the pure Fe-C alloy. A number of investigators have employed this technique. others Have used the partitioning of carbon at the interface of a diffusion couple prepared from an Fe-C alloy and a Fe-C-Z alloy. In the present study, a closed-cell method was developed by which carbon was selectively transferred via a gas medium between pure Fe-C alloys and Fe-C-V alloys which were not in physical contact. A similar method was developed independently by Heckler and winchel14 and used to study the Fe-C-Ni system. EXPERIMENTAL METHOD A schematic diagram of the experimental arrangement is shown in Fig. 1. In a typical experiment, a 5-g Fe-C sample was placed in one compartment of a quartz cell, and a 5-g Fe-V sample was placed in each of the other compartments. The cell was evacuated, outgassed, backfilled with approximately 125 Torr of hydrogen, and sealed. After being placed in a fur-

Jan 1, 1967

By C. S. Giggins, F. S. Pettit

The oxidation of Ni-Cr alloys in 0.1 atm of oxygen has been studied at temperatures between 800" and 1200°C. For alloys with 30 wt pct or more Cr, continuous layers of Cr2O3 are formed during oxidation. In the case of alloys with chromium concentrations between approximately 5 to 30 wt pct, external scales of Cr203 are formed over grain boundaries whereas internal precipitates of Cr2O3 and external layers of NiO are formed at other areas on the alloy surface. When such conditions are present on the alloy surface, chromium diffuses laterally from those areas covered with a continuous layer of Cr2O3 to areas where a Cr2O3 sub scale exists and it is possible for the sub-scale zone to become separated from the alloy by a continuous layer of Cr2O3. Whether such a state will be attained depends upon the initial grain size of the alloy and the oxidation time. When the concentration of chromium in the alloy is less than 5 pct, Cr2O3 is formed internally both at grain boundaries and within the interior of grains and the alloy is covered with an external layer of NiO. MECHANISMS which describe the growth of oxide scales on nickel-base superalloys are complex and the effects produced by the various elements in these alloys on the oxidation behavior of superalloys are not clearly understood. In order to determine the influence of the different elements on the oxidation behavior of superalloys, it is first necessary to examine the oxidation properties of binary nickel-base systems which contain the principal elements present in the superalloys and then progressively more complex systems until compositions typical of the superalloys are attained. Chromium is present in virtually all nickel-base superalloys and the purpose of the present studies was to examine the selective oxidation of chromium in Ni-Cr alloys. The oxidation characteristics of Ni-Cr alloys have been extensively studied1-" to date principally as a result of the high oxidation resistance exhibited by some of these alloys. Ni-20Cr* has long been known *All compositions are given as wcight percent unless specified otherwise. to be oxidation resistant and is commonly used as resistance heating elements for service temperatures up to 1100°C. This alloy cannot be used for extended periods of time at higher temperatures because of the apparent reaction of the external scale with oxygen to form gaseous CrO3. In spite of the considerable work cited above some important aspects of Ni-Cr oxidation still remain unresolved. Virtually all of the previous studies agree that small additions of chromium to nickel, e.g., <10 wt pct Cr, result in increased oxidation rates as compared to that of pure nickel, whereas larger additions, e.g., 20 to 30 wt pct Cr, form alloys with substantially lower oxidation rates. The controversial aspects of the oxidation mechanisms for these alloys that still remain unresolved are as follows: 1) A description of the oxidation mechanism for the low chromium alloys. 2) A description of the oxidation mechanism for the high chromium alloys, particularly with respect to the composition of the external scale which results in the lower oxidation rates. 3) The specific alloy compositions at which the oxidation mechanism changes from that obtained for low chromium contents to that of the high chromium alloys and the reason for this transition. EXPERIMENTAL The Ni-Cr alloys listed in Table I were prepared from high purity metals by nonconsumably arc melting and casting as buttons. These alloys were then given a preliminary annealing treatment in argon at 815°C for 100 hr to promote homogeneity. Each button was cut into 0.250 in. thick sections that were subsequently cold-rolled to 0.050 in. thicknesses and annealed in argon at 815°C for 48 hr to provide a twinned, equi-axed grain structure. The grain size for these alloys was not uniform and the limits, within which the average grain size lies, are given in Table I for the single-phase alloys. All the alloys were single phase with the exception of the Ni4OCr alloy in agreement with the Ni-Cr phase diagram.&apos;&apos; Rectangular specimens were cut from the sheet to provide surface areas of approximately 2.5 sq cm. Exact areas were determined with a micrometer after surface preparation was completed. All of the specimens except the Ni-40Cr alloy and pure chromium were polished through 600-grit Sic abrasive paper, ultrasonically agitated in ethylene trichloride, rinsed with ethyl alcohol, and electro-polished. The specimens were electropolished in a 10 vol pct H2SO4 (conc), 6 vol pct lactic acid, methyl alcohol solution at 70" to 80°C for 2 min at a current density of 0.8 to 1.2 amp per sq cm. This electro-polishing procedure did not produce acceptable surfaces on the Ni-40Cr alloy nor on pure chromium and the oxidation properties of these materials were obtained for specimens polished through 600-grit Sic

Jan 1, 1970

By Harold M. Feder, Robert V. Schablaske, Irving Johnson, Ewald Veleckis

The stoichiometry, structure, and stability of the internzediate phases formed between zinc and some of the rare earth (RE) metals were systematically exarnined by means of a recording effusion balance and X-ray diffraction analyses. In the La-, Ce-, PY-, Nd-, and Y-Zu systems, at or below about 600 C, the following sequences of phases (REZnx) were found: La, x = 1, 2, 4.0, 5.25, 7.3, 17/2, 11, and 13.0;&apos; Ce, x = 1, 2, 3, 11/3, 4.3, 5.25, 7.0, 17/2,* and 11; Pr,x = 1,2, 3, 11/3 ,* 4.3, 5.3(?), 7.0, 17/2,* and 11; Nd,x = 1,2, 3,* 11/3,* 4.3, 6.5, 8.5,* and 11; Y,x = 1,2,3, 11/3, 4.5, 5.0, 17/2,* and 12.* The structure types of all these phases were classified. In addition, lattice parameters were obtained for the first time for the pluses denoted by asterisks. In the absence of de tectable valency or electronegativity effects the systesnatic trends in the results have been ascribed to the effects of&apos; the lanthanide contraction. For example, the maximum number of zinc atoms in the coordination polyhedron surrounding the RE atom decreases from twenty-four to twenty-two to twenty as the size of the RE atom decreases. THE structures and compositions of a great many intermetallic phases (e.g., the Laves phases) are known to be based primarily, but not exclusively, on the space-filling efficiency of various modes of packing together atoms of different sizes. The valencies and electronegativities of the constituent atoms are, however, also influential. In extreme cases hypothetical intermetallic phases which fulfill the efficient spacefilling requirements may not be present in the constitutional diagram because of thermodynamic instability brought about by the operation of valency or electronegativity factors. Hence, for a detailed study of the influence of atomic size on alloy structure and composition, it would be desirable to minimize variations of valency and electronegativity. The intermetallic phases formed by the rave earths (RE) with some common partner offer an excellent opportunity for isolating the effects of size from those of valency and electronegativity. The rare earths exhibit a large, but smooth, decrease in size (the lanthanide contraction) in the series from lanthanum to lutetium when inter comparison is made for a common valence state, e.g., isolated atoms or trivalent ions. The elements yttrium and scandium are frequently included as pseudo rare earths; their sizes place them in the vicinity of dysprosium and lutetium, respectively. The electronegativities of RE elements vary by less than 10 pet. The trivalent state is the most common; however, the well-known tendency of cerium, praseodymium, and terbium to achieve higher valencies, and of samarium, europium, and ytterbium to seek lower valencies, requires that caution be exercised in the assumption of equal valencies. In the present study the existence, constitution, and structure of each of the numerous intermediate phases formed by zinc with lanthanum, cerium, praseodymium, neodymium, or yttrium were examined systematically and in detail. The investigation was conducted by a recording effusion balance technique and by X-ray diffraction analysis. The results enrich our knowledge of the phase diagrams of these systems. In addition, they present some clear-cut evidences of the operation of the size factor alone. EXPERIMENTAL PROCEDURE Apparatus. The mode of operation of the recording effusion balance and its application to phase studies have been discussed in detail elsewhere.&apos; In this work, an effusion cell containing a finely divided alloy was suspended within an evacuated tube from the beam of an analytical balance. The tube was immersed in a massive molten salt bath whose temperature was controlled to within 0.5o C during each experiment. The loss in weight of the alloy owing to effusion of zinc* was continuously recorded. Two effusion cells, 1/2 in. diam by 1 in. high, were machined from tantalum rods. Two orifices were drilled laterally into the walls of each cell. The orifice areas were determined by calibration with pure zinc: cell A had a total orifice area of 6.5 x 10-41 sq cm, and cell B an orifice area of 9.8 x 10-3 sq cm. By appropriate choices of orifice area and temperature the wide range of volatilities from pure zinc to pure rare earth metal could be investigated. X-ray diffraction powder photographs were made at room temperature with a 114.6-mm Debye-Scherrer camera with both filtered CuKa radiation and filtered CrKa radiation. Lattice parameters were refined by a computer-programmed least-squares analytical treatment which incorporated appropriate extrapolation techniques.2 Frequent use was also made of a special computer program3 designed to generate a powder pattern from an assumed structure in order to verify structural assignments. Materials. Lanthanum, neodymium, and yttrium were purchased from the Lunex Co., cerium from the Cerium Metals Corp., and praseodymium from the St.

Jan 1, 1968

By C. R. Whitney, G. N. Maniar, D. R. Muzyka

Previous results have shown that Pyromet 860, an Fe-Ni-base heat-resistant alloy, is stable at temperatures as high as 1500°F for aging times as long as 100 hr. This Paper describes the results of long-time creep-rupture testing at 1050" to 1400°F at various stress levels. Times as long as 37,660 hr were employed. The effects of time, temperature, and stress on the precipitates and their morphologies were studied by optical and electron microscopy, X-ray and electron diffraction, and microprobe techniques. phase, containing cobalt, nickel, and molybdenum, was detected after extended exposures from 1200" to 1400°F and careful study was performed to describe the kinetics of its formation in this alloy. µ phase formation apparently has little effect on the elevated-tem-perature properties of Pyromet 860. For times as long as 500 hr at 1300°F and below, with µ phase present, m significant effects on ambient temperature properties were noted. For longer times at 1300°F and after 1400°F exposure, the effects of u phase on ambient temperature tensile strength properties are not clear due to y&apos; effects and grain boundary reactions. Electron-vacancy, N,, numbers were calculated using different methods described in literature and correlated with the present findings. In the selection of alloys for use in gas turbine applications, structural stability ranks as a primary criterion. High-temperature strength and cost are also of major concern. With these factors in mind, Pyromet 860 alloy, an Fe-Ni-base superalloy was designed. This alloy combines the cost advantages of Fe-Ni-base alloys such as A-286, 901, and V-57 with improved strength and structural stability&apos;1,2 and no tendency to form the embrittling cellular 77 phase. A previous study3 reported on the stability of Pyro-met 860 at temperatures from 1375" to 157 5°F and times up to 100 hr. That study showed that the y&apos; precipitates increased in size and separation and decreased in number with an increase in time or aging temperature. No deleterious phases were found to occur. In the present work, samples from four production heats were subjected to long-time creep-rupture testing at 1050" to 1400°F at various stress levels. Various heat treatments were used on the starting samples and tests were run up to 37,660 hr. The effects of time, temperature, and stress on the precipitates and their morphologies were studied by optical and electron microscopy, X-ray and electron diffrac- tion, and microprobe techniques. Electron vacancy numbers, Nv , calculations were made by TRW.4 Experimental results are correlated with the Nv data used to predict occurrence of intermetallic phases such as a phase. EXPERIMENTAL PROCEDURE Mechanical Tests. Material for the present study came from four production size heats of Pyromet 860 alloy, weighing from about 3000 to about 10,000 lb. All of these heats were made by vacuum induction melting plus consumable electrode vacuum remelting. The nominal analysis for this alloy is compared with the actual analysis of the four heats in Table I. Sections of these heats were forged to 9/16-in. round bar,3/4-in. square bar, 3-in. round bar, 4-in. square bar, and a gas turbine blade forging about 16 in, long, about 6 in. wide, and weighing about 20 lb. In general, all forging of this alloy is done from a 2050°F furnace temperature. Longitudinal test blanks were cut from the centers of the smaller bars, from mid-radius positions for the 3- and 4-in. bars, and from the air foil of the gas turbine blade and heat-treated according to the procedures outlined in Table 11. Heat treatment A is the "standard treatment" recommended for this alloy for best all-around strength and ductility. Heat treatment B is a modification of treatment A for improved tensile strength at moderate temperatures. The treatment coded C was designed for treating large sections according to a procedure previously described.&apos; Heat treatment D was developed to yield optimum stress relaxation characteristics at 1050°F for a steam turbine bolting application. After heat treatment, the test blanks were machined either to plain bar creep specimens with a gage diameter of 0.252 in., to combination smooth-notched stress-rupture bars with a plain bar diameter of 0.178 in. and a concentration factor of Kt 3.8&apos; at the notched section, or to notch-only specimens. All specimens conformed to ASTM requirements. Metallography. Most of the creep-rupture tests were continued to failure. A few bars were fractured as smooth or notch tensiles after creep-rupture exposures. After fracturing, ordinary metallographic sections were made primarily in gage areas adjacent to fractures to represent a "high-stress" region and through specimen threads to represent a "low-stress" region. All metallographic sections were made in a longitudinal direction with respect to the test specimen axes. For optical microscopy, the samples were etched in glyceregia (15 ml HC1, 5 ml HNO,, 10 ml glycerol). For XRD analysis, the phases were extracted electrolytically in two media: 20 pct &Po4 in H20 for selective extraction of y&apos; and 10 pct HC1 in methanol for carbides and other phases.

Jan 1, 1970

By J. H. Henderson, L. A. Wilson, R. L. Gergins, R. J. Wygal, D. W. Reed

This paper presents a method of predicting the production history of an underground combustion recovery process. A rigorous solution of the thermodynamics and hydrodynamics involved is beyond the scope of this report. However, a practical scheme to appraise combustion recovery performance has been Worked out. By prudent assumptions, regarding the attainment of steady-state conditions, a trial-and-error solution, which satisfies both material balance and three-phase relative permeability requirements has been evolved and tested. As a combustion zone moves through a formation the interstitial water, the water of combustion and a portion of the oil wil1 be vaporized. These vapors will move downstream to a cooler region of the formation where they will condense. Part of the remaining oil will be displaced by the imposed gas drive and what remains behind will be consumed as fuel. The fluids will distribute themselves downstrean? in a manner which will satisfy the three-phase relative permeability characteristics of the formation. Three distinct zones will exist ahead of the combustion zone: (I) a zone termed a water bank which contains three mobile phases, oil, water and gas; (2) a region containing two mobile phases, oil and gas, called the oil bank; and (3) a zone having the fluid saturution distribution which existed prior to combustion. The mobile water will impose a water flood on the oil in the region containing three flowing phases. The process can be visualized as a simultaneous water .flood and combustion-supported gas drive. It is possible to estimate the saturation distributions in the water bank and oil bank as well as the production history for any combustion temperuture if the porosity, three-phase relative permeability characteristics, oil viscosity and initial saturations are known. This has been done for injection pressures ranging from I to 100 atm, oil viscosities of 10 to 1,000 cp and porosities from 20 to 40 per cent. Several of these calculations have been checked in the laboratory, with the result indicating rather good agreement between the predicted and observed production histories. The analytical and experimental techniques are described and the comparison of predicted and observed performance presented. INTRODUCTION Considerable research by the oil industry during the past 10 years has been directed toward the thermal recovery of crude oil. Of all schemes considered, the one receiving the most attention is in situ combustion. This is a process in which part of the oil is actually burned within the reservoir rock. The uniqueness of the process lies in its potential. Theoretically, this potential generates from the following: (1) it may open the door to vast reserves previously not economically producible, (2) it may induce increased production rates where low rates are now realized and (3) it may give high over-all recovery efficiency. The general idea of the process has received considerable attention in the literature &apos;-&apos; and the following brief description is intended for orientation only. In situ combustion can be visualized by considering a linear system, one end of which represents the injection well and the other the producing well. The formation at the inlet end is raised to ignition temperature by placing a gas or electric heater at the face of the formation or by the ignition of thermite, charcoal or similar material. While heating, the water and part of the oil in the heated region is vaporized and moved out into the reservoir by the injection of gas containing oxygen. In addition, part of the oil is displaced as liquid by the imposed gas drive. The oil not displaced is modified to an immobile, coke-like material (hereafter referred to as coke) by the high temperatures existing immediately ahead of the combustion zone. This coke is used as fuel, and the combustion front advances through the formation at a temperature which is probably in excess of 800°F. As the combustion front advancek it will drive ahead all interstitial oil not rendered immobile and all interstitial water plus the water formed by the combustion of hydrogen. As a result the sand behind the combustion zone is freed of all oil and water. Although the general description of the process has been outlined in the literature, there has been no consideration of the behavior of the fluids moving downstream from the combustion front. It is the purpose of this paper to present a technique

By J. V. W. REYNDERS

IN THE remarks which I am about to make concern¬ing the safety work of the Bureau of Mines, I want first of all to disengage myself from a disposition, which is frequently in evidence, to give spectacular emphasis to the hazards of mining operations. It is of the&apos; utmost importance to the industrial welfare of our country that the enormous annual requirement of fuel be produced and be made available in our economic activities, with a minimum of distress and hardship. The safety of coal mining is a condition with which. we are all concerned in a fundamental way. And yet we should not rush to accept hastily conceived notions that laxity among employers in respect to safety measures exists to any outstanding degree. In fact in a large number of instances, the precautionary efforts of min¬ing companies, directed to the elimination of accidents are beyond all criticism, albeit the industry suffers in reputation from the carelessness of, its less effi¬cient constituents. The Bureau of Mines has been a leading influence in disseminating the doctrines of safety, in securing the cooperation of the miners themselves, and in stirring up &apos; the laggards, both among the employers and the employed. It is much to the interest of well-conducted mine operations that the industry, as a whole, observe safety regulations to the utmost.

Jan 1, 1925

By J. C. Williams, M. J. Blackburn

Several of the decomposition processes that can occur in supersaturated phases in a Ti:11.6 wt pct Mo and a Ti:20 wt pct V alloy have been studied by transmission electron microscopy. The deformation induced "marternsitic phase" in the Ti:Mo alloy has been found to have a bcc or bct structure rather than the previously reported hexagonal structure. The morphology of&apos; the transformed region is a rather complex asserrlblage of twins, twinning occurring in one or more systems; this internal twinning has been found to occur on (112). The w phase is formed in both alloys on aging and is present in the Ti:Mo alloy after quenching. The structure of this phase has been confirmed as hexagonal in both systems, however, differences in morphology and stability are found between the two alloys. Thus in the Ti-Mo alloy the w phase has an ellipsoidal morphology with the major axis lying parallel to <111>ß or [0001]w while in the Ti-V alloy the phase forms as cubes, the cube faces lying parallel to {100}ß or {2021}w Some observations on the particle sizes, volume fraction, and composition of the w phase in the Ti-Mo alloy are listed. The mode of formation of The a phase from the (ß + w) structures is also different in the two alloys. In the Ti-Mo alloy the a phase is formed by either a cellular reaction or by the growth of isolated needles, whereas in the Ti-V alloy the a phase is nucleated at an w:ß interface and grow to consume the w phase. Some of the difjerences in behavior of the w phase are attributed to the mismatch between it and the solute enriched ß matrix in which it forms. MaNY transition elements tend to stabilize the bcc or ß-phase when added to titanium. In general two types of phase diagrams are produced, either a ß-stabilized (ß-isomorphous) system, e.g., Ti:Mo, -Ti:V, Ti:Nb, or a ß-eutectoid system, e.g., Ti:Cr, Ti:Fe, Ti:Mn. In previous papers&apos;-4 the phase transformations in the a-phase and (a + ß)-phase alloys have been described and this work has been extended to ß-stabilized systems. Specifically, transformations in the alloys Ti:20 wt pct V and Ti:11.6 wt pct Mo have been studied; in both of these alloys the ß phase is retained at room temperature when quenched from the ß-phase field. A number of phase transformations can occur in such metastable ß phases and the two alloys were chosen to include most of the transformations reported for ß-stabilized systems. We list these possible phase transformations below. Ti:11.6 Mo quenched from >780°C to retain the ß phase: a) The w phase can form on quenching.5 b) Martensite can be produced by subzero cooling or deformation. Two martensite habit planes have been reported in Ti:Mo alloys; (334)ß and (344)ß=6 c) On aging at temperatures <-550° C the w phase is formed before the a-phase.5,7 d) On aging at temperatures >550°C the a phase is formed.7 e) The martensite can be tempered. It has been reported that the a phase rather than the ß phase is precipitated during tempering.&apos; Ti:20V quenched from >660°C to retain the ß phase:9 a) At aging temperatures <260°C separation into two bcc phases occurs. b) The w-phase is produced prior to the a phase on aging at temperatures <-400°C. c) At temperatures 2400°C the a phase is formed directly. T-T-T diagrams describing the temperature and time regimes for the formation of these phases have been published7,9 for a Ti:12 pct Mo and a Ti:20 pct V alloy. We have attempted to investigate these transformations using transmission electron microscopy, however thin foils undergo a spontaneous transformation in all conditions except the equilibrium (a + ß) structure. This transformation has been reported previ0usly10,11 and we will comment on its morphology and nature in the various sections of experimental results. EXPERIMENTAL The compositions in wt pct of the two alloys investigated were: Ti:11.6 Mo, 0.100 02, 0.006 N2, 0.0015 H2 Ti:20V, 0.0574 O2, 0.0111 N2, 0.005 H2 These alloys were cold-rolled to 0.020 in. thick sheet. Specimens were heat treated in vacuum or in inert gas at temperatures >500°C and in a circulating air furnace at temperatures <500°C. Thin foils were prepared using standard techniques, described in detail previously." Dark field micrographs were obtained using high resolution technique. RESULTS Martensitic Transformation in Ti:11.6 pct Mo. Detailed study of the deformation induced martensite is not possible due to a spontaneous transformation which occurs near the edge of thin foils as shown in Fig. 1. Similar transformations have been observed in iron-" and copper-base13 alloys as well as other titanium alloys, but some observations specific to the Ti:1l.6 Mo alloy are listed below. a) The boundaries of these transformed regions are glissile and move under the influence of the electron beam during examination. b) Selected area diffraction indicates the transformed regions have the same structure as the matrix, being separated by tilt boundaries. The misori-

Jan 1, 1969

By J. A. Herbst, D. W. Fuerstenau

This paper examines the zero order production phenomenon in the context of the size discretized batch grinding model. A restrictive interrelationship between the selection and breakage parameters of the model, which is mathematically sufficient to ensure zero order behavior, is delineated. Based on this interrelationship, a scheme is developed for predicting the values of the selection and breakage parameters for all size fractions from a minimum of experimental data. To test this scheme, dolomite was ground in a laboratory batch ball mill. Successful simulation of the comminution behavior of this system was achieved using the batch grinding model with the parameter values obtained from the proposed scheme. Historically, a preponderance of comminution research has centered on attempts to relate breakage energy to the performance of comminution machines. Concomitantly, grinding data were interpreted almost exclusively in terms of inherently empirical energy-size reduction relationships1-3 or "laws of comminution"4- 6 which were based on highly oversimplified descriptions of the fracture process. In some instances, these relationships provide a crude basis for the correlation of experimental data, but, invariably, this approach is inadequate for meaningful process simulation. The control and optimum design of comminution circuits require a mathematical model capable of depicting the size reduction behavior of every size fraction for grinding conditions of technological importance. Energy-size relationships do not provide this detailed information. For example, an energy-size analysis for batch ball-milled dolomite has recently been completed in the authors&apos; laboratories.7 This analysis provided a satisfactory correlation between the hypothetical size modulus, which characterizes only the finer portion of the size distribution, and the input energy. However, this relation did not constitute an adequate description of the complete system. In general, one of two distinct viewpoints toward simulation has been adopted in recent comminution research. The first approach focuses on the fracture of single mineral specimens, with the essential aim of representing the over-all process in terms of the breakage characteristics of individual particles and the characteristics of the stress field which the particles experience within the particular size reduction device. As illustrated by Harris&apos;s review,&apos; the actual incorporation of single-specimen fracture information into a description of the behavior of a multiparticle comminution system has seldom been attempted. The only model to accomplish this incorporation was the one developed by Schönert.8 The Schonert model treats single-passage grinding machines in terms of a distribution of effective loads acting on a particle, a distribution of required particle breakage energies, and the breakage product distribution. However, the use of single-particle fracture behavior to derive models of the size reduction process is at present limited to machines in which the particle residence time is approximately equal to the time required to apply stress. As a result, Schönert&apos;s analysis cannot be used to advantage to simulate the complex environment which prevails within a tumbling mill. For multiple-passage systems a second approach to the description of comminution processes has been more fruitful. This approach entails the formulation of a mathematical model which is phenomenological in nature in that it lumps together the entire spectrum of stress-application events which prevail in a system under a given set of operating conditions. The appropriately defined average of these individual events is then considered to characterize the over-all breakage properties of the device. Thus, to analyze the performance of a tumbling mill, the manner in which the particles of a particular size (or size fraction) are stressed need not be distinguished. Instead, a single parameter is assumed to represent the resistance of that size to fracture, given the average grinding environment which exists in the mill. The isolation of such a parameter and a related set of quantities, which constitute the breakage product size distribution for the average event in this size fraction, allows the formulation of physically meaningful descriptive equations capable of yielding precise and detailed information for simulation.

Jan 1, 1969

By R. A. Flinn, R. D. Pehlke, D. R. Glass, P. O. Hays

Knowledge of the depth of penetralion of an oxygen jet into the bath of the oxygen converter and of the correlation of penetration with driuing pressuve, lance heighl, and nozzle throat area is vital to the understanding of converter operation. If the penetration is too shallow, then severe and hazardous slopping takes place. On the other hand, if the jel penetrates entirely Lhvough lhe bath for an apprcciublc period of time, bottom damage occurs. In addilion to measurement of the penetralion of the jel, knou~ledge of the circulatory movement in the bath is also of interest in order to evaluale various theories oj-concerter operating behavior which have been published. In this investigation, experimental converters were buill of IOU- , 300-, and 4000-lb capacity. Four independent methods were used to determine penelralion: the onset of bottom marking, a nitrogen bubbler probe, observation througlt an optical syslew built into the oxygen lance, and direcl viewing of the jet issuing from the bottom of the vessel. Good correlation zuas obtained, and empirical relalions for pvedicling perletration were found. These relations were conjzrmed by bottom marking tests in 55- and 110-ton vessels. Within the operaling conditions employed in these tests, the depth to which a single oxygen jet penetrated zuas found lo vary according to the relatiorl ThE technical literature is replete with data concerning the successful use of the basic oxygen furnace or converter in steelmaking. Experimental data are lacking, however, on the vital factors of the depth of penetration of the jet into the bath and the induced circulation. Commercial operating conditions usually have been the result of cut and try experiments in lance manipulation until satisfactory results were obtained. There have been, however, two hotly argued opposing theories concerning desirable depth of penetration and these are exemplified by the Schwarz and Miles patents1,2 on one hand and the Suess patent3 on the other. The Schwarz patent teaches that the jet, issuing from the nozzle at supersonic speed, penetrates deeply "so that the reactions between the iron and the oxygen and between the oxygen and the rest of the smelting components take place in the center of the bath". Specific operating suggestions are given by Miles.2 By contrast, the Suess patent calls for surface Circulalion was investigated by lour methods: by direct observation in 200-lb open baths, by the use of graphite rudders in the 300- and 4000-lb converlers, by direct observalion through an oplical system in the lance, and by various models al room temperature. All were in excellent agreement and indicated that the motion of the bath ulas up at the center, radially outluavd at the surface, and down at the sides. Experi-ments in small and in commercial vessels indicate that it is essential to operate with a jet penetration of approximately 50 pct of the bath depth. Surface blowing results in low oxygen eficienty and in hazardous conditions which may render the process inopeuable. RejYactory dartzage al the bottom of the vessel is only encountered when the jet penetrates to the bottom, and this can be avoided by properly applying the penetration formula. The application of this en/pirical formula in commercial peraations is best when limited to combinations of lance size, pressure, and height which are typically encounteved in the use of a single-hole lance. blowing so that "...the oxygen jet does not penetrate deeply into the molten metal bath and is confined to an impingement area at the central portion of the bath surface". These references are given merely to illustrate the basic differences between the two schools of thought and to point out the need for measurement of penetration for the sake of the operator. For example, it is shown later that inefficient and even dangerous conditions can arise if improper blowing conditions are used. Differences are also evident between the two schools of thought as to the mixing, circulation, and agitation which is to be accomplished by the jet. The Schwarz patent states that "surface contact is not sufficient in most cases to bring about quick reaction, the same as the blowing of the gas over the bath surface or the mere blowing of the gas onto the bath surface". The patent goes on to call for active mixing. In contrast the claims of the Suess patent call for "discharging a stream of oxygen ... to an extent to avoid material agitation of the bath by the oxygen stream". In this patent the circulation is said to be downward in the center and up at the sides of the vessel. A number of investigators4-12 have explored penetration and circulation in transparent models. In general, it is agreed in these tests that the circulation is upward at the center (along the sides of the jet cavity),

Jan 1, 1968

By A. Gemperle

Electron diffraction and transmission electron microscopy on foils at room temperature were used to investigate the ordering of Fe-Si alloys containing 10 to 23 at. pct Si. A certain degree of DO3 order was found in all alloys. With the exception of the lowest silicon concentration for which the antiphase domains could not be clearly resolved, the alloys have a domain structure of two-domain type with boundaries having 1/4a01<111> displacement vectors for less than 12.3 at. pct Si and with boundaries having 1/2 a0<100> displacement vectors for more than 12.3 at. pct Si. The alloys with 12.3 at. pct Si have a domain structure consisting of fine domains with 1/2a&apos;o<100> boundaries within much larger domains with 1/4a&apos;o<111> boundaries. The development of these structures can be explained by transition of the alloy from the disordered state into the B2-type order and then into the D03-type order by the mechanism proposed previously for the FeSAl alloys. The existence of the B2 structure in the lower part of the investigated concentration range reported in some articles can be explained by fine domains with 1/2a&apos;o<100> boundaries formed by several disordered planes within large domains with 1/4a&apos;o<111> boundaries. The ordered structure predicted by the theory —with practically no domain boundaries —is found in the alloys having 12.3 at. pct Si where it develops in the B2 structure region. ORDERING in Fe-Si alloys was first studied by phragmenl who found that beginning with 13 at. pct Si the DO3 (Fe3Si) superlattice reflections appear in the diffraction patterns. The equilibrium diagrams constructed later by corson2 and Haughton3 from various measurements proposed the existence of a homogeneous solid solution (a phase) in the range from 0 to 25 at. pct Si. Jette and Greiner4 and Farquhar et al.5 measured the relation between lattice parameter and composition and they considered the break in the curve at 9 to 10 at. pct Si to be caused by the ordered solution a"(Fe3Si). Glaser and Ivanick6 Determined critical ordering temperatures of the alloys containing from 10.9 to 27.9 at. pct Si from the measurement of the electric resistivity of quenched samples. In all cases the critical temperature was lower than the melting point and it was highest for 25 at. pct Si. Lihl and Ebel7 measured the lattice parameter curves at various temperatures up to 1000°C. The region between two breaks on these curves, corresponding to 10 to 12.5 at. pct Si at room temperature, was considered by them to be two-phase (a + a"). They concluded by extrapolation of the measured values that a" in the alloy having 25 at. pct Si is stable up to the melting point. Davies8 studied superlattice reflections in the X-ray diffraction patterns of an alloy containing 8.7 at. pct Si. He found the B2 structure and short-range order in the slowly cooled samples and the DO3 A. GEMPERLE is Research Scientist, Institute of PhysicS, Czechoslovak Academy of Sciences, Prague, Czechoslovakia. __Manuscript submitted January 2, 1968. IMD structure in the quenched and annealed samples. This investigation first reports the presence of the B2 structure, phase a&apos;: in Fe-Si alloys. Meinhardt and krisement9,10 also found its existence in Fe-Si alloys by neutron diffraction. No order was detected by them in the alloy containing 8 at. pct Si. The onset of B2 order was observed at a composition of 9.2 at. pct Si. They found almost perfect B2-type order with partial DO3-type order at room temperature in the 10 to 12.5 at. pct Si range and almost perfect DO3-type order in the 12.5 to 25 at. pct Si range. They established the critical temperatures Tc of both the structures through measurement at higher temperatures. The critical temperature for the B2 structure was found to be always higher than the critical temperature for the DO3 structure of the same alloy. They extrapolated the curves of the critical temperatures and concluded that the alloys with more than 17 at. pct Si have the B2 structure up to the melting point and the alloys with more than 23 at. pct Si have the DO3 structure up to the melting point. The results of Meinhardt and krisement9,10 were confirmed by Dokken&apos;s measurement" of the temperature dependence of the electrical resistivity in the 10.8 to 15 at. pct Si range. On the other hand chessin12 detected by X-ray diffraction a considerably lower degree of order in the 12.7 at. pct Si alloy. ANTIPHASE DOMAIN STRUCTURE IN ALLOYS WITH B2- AND DO$-TYPE ORDER The ordered structures B2 and DO3 can be described in terms of a subdivision of the bcc lattice into four fcc sublattices with a parameter double that of the bcc lattice. Following Marcinkowski13 we will label them I. 11, 111, IV. The B2 structure in the AB alloy is formed by placing A atoms on sublattices I and II and B atoms on sublattices III and IV. In a non-stoichiometric perfectly ordered alloy having concentration (A) > (B), A atoms occupy sublattices I and 11. and A and B atoms distributed at random occupy sublattices III and IV. The DO3 structure in an A3B alloy is formed by placing A atoms on sublattices I, 11, and 111, and B atoms on sublattice IV. In a nonstoichiometric, perfectly ordered alloy having concentration (A)/3 > (B), A atoms occupy sublattices I, 11, and 111, and A and B atoms distributed at random occupy sublattice IV. As further shown in Ref. 13, two types of domains are possible in the B2 superlattice and the antiphase domain structure has associated with it boundaries with displacement vectors 1/4 a&apos;o<ll1> only. Four types of domains are possible in the DO3 superlattice and the antiphase domain structure has associated with it boundaries with displacement vectors 1/4a&apos;o<111> and 1/2a&apos;o<l00>. Bethe14 suggested on the basis of theoretical considerations that in a structure with two sublattices at low temperatures only one domain should be present in the whole crystal at equilibrium. Similarly Bragg15 concluded that at low temperatures the domain struc-

Jan 1, 1969

By Ravindra M. Nadkarni, Milton E. Wadsworth

The kinetics of cetnentation of copper with iron were observed to follow first-order kinetics and increase with speed of agitation to a limiting value. Maximum rates agree closely with theoretical values based upon a model of aqueous solution diffusion through a litniting boundary film. Back reaction kinetics are shown both theoretically and experimentally to be independent of ferrous iron concentration in solution. The inlportance of attnospheres of air, oxygen, nitrogen, and hydrogen was studied and the results have been correlated with several impovtant oxidation processes involving metallic iron and copper. The kinetics of the reaction of ferric ion with metallic iron were found to be slow in the absence of metallic copper and essentially proportional to the surface area of metallic copper present in the system. THE precipitation of copper on iron is classic as an example of a relatively ancient art applied successfully for centuries with little fundamental understanding of the important parameters involved. There is some indication that the process has been a commercial means to produce copper since the sixteenth century.&apos; The amount of fundamental work on the cementation of copper with iron is not great. Wartman and Roberson2 carried out a series of detailed copper cementation experiments using natural and synthetic mine water. The following were presented as the three principal reactions: Reaction [I] is the desired cementation reaction and accordingly 0.88 lb of iron would produce 1 lb of copper. In actual practice iron consumption would more normally fall in the range of 1.5 to 2.5 lb per lb of copper. Wartman and Roberson attributed the excess consumption of iron to Reactions [2] and [3]. They found that Reactions [I] and [2] proceeded at approximately the same velocity while Reaction [3] was much slower and would be diminished by controlling the contact time. It was also pointed out that increased agitation is beneficial in removing hydrogen bubbles and barren layers of solution at the iron surface as well as removing contaminants resulting from the hydrolysis of iron. Episkoposyan3 and Episkoposyan and Kakovskii4 studied copper and silver cementation on rotating iron disks in chloride solutions. The kinetics based upon a diffusion model were first order and varied linearly with surface area and with angular velocity raised to the one-half power according to the Levich equation. The experimental activation energy for both copper and silver was approximately 3 kcal per per mole. Excess iron consumption was found to increase with temperature. The rate of cementation first increased with increasing acidity and then diminished at high acid concentrations. sutolov5 has presented an excellent review of the Leach-Precipitation-Flotation (LPF) process including a discussion of copper cementation from an electrochemical point of view although few experimental results were presented. From voltage considerations he predicted that cementation should not be influenced by the concentration of ferrous iron in solution. He considered several secondary reactions including Reactions [2] and [3] and pointed out the importance of oxidation of ferrous iron to ferric with oxygen. In addition it was suggested that Reaction [2] was enhanced by the dissolution of metallic copper by ferric iron which in turn consumed excess iron by the cementation reaction, Eq.[1]. Cementation of copper on metals other than iron has been studied by several investigators but, as in the case of iron, the amount of fundamental work is not extensive. Bashkova and kovalenko6 and Bashkova7 studied the cementation of copper on indium from copper and indium sulfate solutions. The rate was found to be first order and to increase with acidity. This was associated with a decrease in potential (EIn — ECu) and the simultaneous reduction of hydrogen ions at low pH. The rate of cementation also decreased with increasing indium concentrations which was postulated to be due to the decrease in the rate of diffusion of the ions in solution. Below 97°C the experimental activation energy was found to have the unusually low value of 2 kcal per mole and was attributed to diffusional control. Above 97°C the rate increased suddenly and was explained as a change in the rate-controlling step to a chemical reaction. In Part I of this study Nadkarni et a1 .1 have reported on preliminary results obtained in a laboratory study of the kinetics of the cementation process. The rate was found to be first order, proportional to the surface area of the iron, and to increase with speed of stirring until a maximum rate was observed. At low stirring speeds the deposit was spongy and adherent. At medium speeds the copper peeled off in bright strips and at high speeds finely divided copper was produced and continually removed from the surface. The amount of excess iron consumed increased with speed of stirring and with temperature. The average experimental activation energy combining results from several types of iron was 5.8 + 1.6 kcal per mole suggesting diffusional control through a limiting boundary film. Traditionally copper cementation has been carried out over the centuries in gravity-fed launders of various design containing scrap iron. More recently rotating drum precipitators and activated launders8&apos;10 have been used. In the latter, copper-bearing solutions are

Jan 1, 1968

By F. C. Bond

Since the Third Theory of Comminution was presented eight years ago (I) it has found increasing use in crushing and grinding problems. The practical utility of its wok index equation is quite generally acknowledged (2). However, its theoretical basis has been questioned in at least three technical articles (3) (4) (&apos;). The purpose of this paper is to present experimental proof that it is scientifically correct. Particles under compressive stress are strained and deformed. They absorb strain energy, and when this locally exceeds the breaking strength, a crack tip forms. The surrounding strain energy flows to the crack tip, which rapidly extends and splits the rock, releasing the strain energy as heat. The initial energy flow causes additional crack tips in highly strained areas. If the compression is rapidly applied by impact, crack tips may form before the strain energy has reached equilibrium in the particle, thus decreasing the total work input required for breakage. The energy necessary to break is essentially the energy necessary to produce crack tips, since the energy necessary to extend the cracks to breakage is already present as strain energy in the deformed particles. After breakage nearly all of this energy appears as heat. The crack length cannot be measured directly. However, in particles of regular and similar shape the crack tip length is considered as equal to the crack depth, or crack extension necessary to break, so that the crack length equals the square root of one-half of the surface area. The Third Theory states that the useful work done in crushing and grinding is directly proportional to the total length of the new cracks formed. It can be confirmed by showing that a constant work input produces a constant length of new cracks when reducing the same material to different product sizes. This is done in the present paper on a wide variety of material. The constant work input was supplied by one revolution of the 12" x 12" laboratory ball mill used in making grindability tests by the Allis-Chalmers method (12) (13) The new crack lengths produced per mill revolution were measured from all available grindability test results at 28, 35, 48, 65, and 100 mesh on fifteen different ores, and were found to remain substantially constant for each ore at all mesh sizes. A new technique is used for the measurement of crack lengths. Size distribution plots of the mill feed and product are made by the Third Theory method (9) and the crack lengths are obtained from these plots by the procedure described in the present paper. The energy input required to produce a unit length is of fundamental importance in the size reduction of brittle solids. The crack length Cr is expressed in centimeters per cubic centimeter of solid material. It bears a definite relationship to the external surface area of the crushed or ground solid. A uniform particle shape must be assumed before the surface area and crack length can be evaluated. In this paper it is assumed that the relationship between the surface area and the particle volume of a particle d microns in diameter is the same as that of cube d microns on a side. The external surface areas of particles with a cubical breakage probably agree approximately with this rule, and correction factors can be applied when physical measurements of the surface areas are available for comparison. However, the assumption of equivalent cubes has been found satisfactory for most calculation purposes. Assuming equivalent cubes, one cubic centimeter of particles d microns in diameter will have a crack length Cr of v30.000/d centimeters, and a surface area of 60,000/d square centimeters. The specific crack length is thus equal to the square root of one-half the specific surface area. Where Sa is the surface area in square centimeters per gram and Sg is the specific gravity of the ground solid, then Cr = vSg . Sa/2 = 173.2/ vd (1)

Jan 1, 1961

By P. D. Merica

ALUMINUM and its alloys have been the subject of much investigation, during recent years, in the course of which the principal features of the constitution of most of the binary alloy systems with aluminum have been determined. Except in the case of a few metals-silicon, bismuth, cadmium, lead, zinc, and tin-an aluminum-rich compound is formed in each binary system, which forms a eutectic with the aluminum or its solid solution with this compound. Thus such compounds as FeAl3, CuAl2, Mg4Al3, and NiAl3 are formed, which are found in aluminum-rich alloys of their respective series as eutectics with the aluminum solid solution. These compounds are, in all cases, hard and brittle and their presence affects profoundly the physical properties of the alloys in which they occur. Within the zinc-aluminum system a compound, Al2Zn3, is formed that decomposes at lower temperatures. Silicon and tin each form a simple eutectiferous series with aluminum; bismuth, lead, and cadmium are only partly miscible in the liquid state with it. The extent of the solubility in aluminum, in solid solution, of these compounds or of the elements themselves, in the case of those series

Jan 7, 1919

By S. V. Krusiewski, J. W. Gabelman

Regional metallogenic zones and areas of equal paragenetic range have been interpreted from a compiled metallogenic map. of Mexico. Mineralization gradients of the pattern subsequently identified extend inland from coastal zones of high-temperature mineralization, and join in a central longitudinal zone of relatively low-temperature mineralization, apparently leaving no zonal room for lowest temperature deposits. The striking regional zoning pattern can be correlated with regional deformation gradients of similar position and orientation. The gradients represent at least three major orogenic cycles in two structural belts. Paleozoic structural belts existed along each coast with deformation gradients diminishing inland toward a narrow, central peninsular craton. The Mesozoic-Tertiary Mexican (Cordilleran) belt covered the entire country with a gradient oriented toward the Gulf of Mexico. Most ore deposits in this belt are of Tertiary age, although Tertiary mineralization gradients seem to have "inherited" the pattern left by Paleozoic tectonic gradients. Mineralization gradients can be used to identify the regions most favorable for the occurrence of selected metals. A metallogenic province is characterized by particular proportions of various metals, and the relative proportions can vary from province to province; but within any province the same intensity of mineralization would produce similar deposits in different places. Declining intensity of minetalization produces different kinds and sizes of deposits in zones that are concentrically arranged around sites of strongest rock alteration. On a regional basis, districts of similar average intensity of mineralization occur in zones that generally parallel some tectonic grain with the result that tectonic deformation has been designated a principal cause of mineralization. Variations in the type and intensity of deformation would determine the type as well as size of metal deposits. Mexico is especially favorable for such a study be- cause of its great abundance and variety of epi-genetic mineral deposits in various lithologic, structural and tectonic environments. The concept is also useful in recognizing portions of Mexico favorable for the occurrence of specific metals. METALLOTECTONIC ZONING IN MEXICO Theory: In studies of individual districts and groups of districts throughout North and South America, areal zoning of metal deposits generally seemed to respond to the paragenetic sequence of iron, gold, copper, zinc, lead, silver, antimony and arsenic, gold and fluorite. However, the definition of district zones is greatly obscured by contemporaneous or rhythmic deposition of metals, especially in the middle of the paragenetic range, and by the common superimposi-tions of different generations of metals. Zones could usually be distinguished only for groups of metals, such as early precious metals, base-metal sulfides, sulfosalts and late precious metals. The best defined zoning in a district is normally the spatial separation of the greatest volumes of altered and metallized rocks.&apos; Exceptions are districts representing the high- and low-temperature extremes of the mineralization range where the former creates a matrix for the latter. In most districts the total quantity of alteration products is greatly in excess of, and appears to be roughly proportional to, the total quantity of metallization products. Subsequent to the bicomponent nature of each district, the most important mineralogical feature is the great excess of one to several metals or minerals which results in a district identifiable as one of predominantly gold, copper, lead-zinc, fluorite, etc., although a great variety of other elements may be present. The para-genetic range for the abundant elements in any district is only a portion of the entire possible range for all the elements representing the complete process throughout the region. Early during investigations of this type, it became apparent that the consistent bicomponent character of districts allows two methods for studying epi-genetic mineralization patterns. One is to identify and plot altered rock "pipes,"* which are easily recognizable in the landscape by their anomalous

Jan 1, 1969

By H. D. Outmans

The paper presents a theoretical approach to the drilling problem based on rock mechanics and drilling fluid hydraulics at the bottom of the hole. The volume of the fractured rock around the vevtically penetrating bit tooth is expressed as a function of the tooth pressure, included tooth angle, drilling fluid pressure and formation characteristics on the assumption that fracturing is preceded by plastic flow. This assumption is verified by comparing the stress-strain state in the rock around the tooth to that in equivalent tri-axial compression tests described in the literature. An expression similar to the above is derived for the tooth penetration in the cuttings which have been retrained at the bottom. Combination of these two expressions leads to a relation between depth of tooth penetration in the virgin rock, the amount of cuttings retained and the other variables. By relating the cutting volume to the drilling rate, the hydraulic bit horsepower and the drilling fluid pressure, it is possible to arrive at formulas which express the drilling rate for viscous or turbulent flow at the cutting surface, for instant clearance, for retained cuttings, or for false tooth foundering as functions of such variables as the drilling fluid pressure, hydraulic horsepower at the bit, rotary speed and bit weight. Plotting the results as drilling curves and comparing them to published field and laboratory data justify the conclusion that the drilling model presented in this paper approximates the actual mechanism. The drilling curves serve to explain some characteristic features of the rotary drilling operation, and the equations may be used for numerical evaluation of the effect of changes in the magnitude of some of the variables on the drilling rate. INTRODUCTION Published laboratory data have established that consistent and relatively simple relations exist between the drilling rate and any one of the important variables. For instance, the drilling rate usually increases at an increasing rate with the bit weight. It increases at a decreasing rate with the rotary speed, decreases at a decreasing rate with increased drilling fluid pressure and plots as an S-shaped curve against the hydraulic horsepower at the bit. This relative simplicity seems difficult to understand considering that the drilling mechanism is a continuous process of generation and removal of innumerable individual cuttings, under different conditions for each. On the other hand, the drilling rate (although determined by this complicated process) only expresses the average rate at which cuttings are produced and cleared away. Ignorance of the drilling mechanism as far as individual cuttings are concerned, therefore, does not exclude the possibility of arriving at a concept of the drilling rate provided the mechanism is properly evaluated for the average cutting. It then may be concluded from the first sentence that this concept also should be rather simple in its final form, although not necessarily in its derivation. Two aspects of the drilling mechanism are difficult to evaluate — (1) the effect a penetrating drilling fluid has on the stresses around a bit tooth entering the formation and (2) the volume of rock fractured by a penetrating tooth which moves both vertically and horizontally. Thus, the scope of this paper has been limited to an evaluation of the drilling mechanism in formations which have a low permeability and are drilled with hard-formation bits. In the following sections, the drilling mechanism is analyzed in three steps. First, the effect of tooth pressure on the volume of fractured rock around the tooth is evaluated, assuming initially that the bottom of the hole is free of cuttings and, then, determining the effect of a layer of these cuttings. In the second section, the thickness of this layer is related to the drilling rate and the chip clearance time, and the latter to the hydraulic horsepower at the bit under conditions of laminar and viscous flow at the cutting surface. In the third section. the previous results are combined with the rotary speed into drilling equations which are then plotted and compared to experimental drilling curves. This section also contains some numerical examples of drilling rate computations. Some mathematical details are given in the appendixes. Finally, it should be mentioned that, although some drilling equations are already available, these equations (with the exception of one&apos; which attempts to relate drilling rate to drilling fluid pressure) are all mathematical expressions for observed laboratory data; none considers the effect of delayed drill-cutting removal, an essential feature of the drilling mechanism.

By A. H. Geisler, J. H. Keeler

Preferred orientations in rolled and annealed titanium sheets were determined by the Geiger counter spectrometer X-ray diffraction technique. Five annealing textures dependent upon the temperature range of annealing were found, and in order of increasing annealing temperature pendent upon the temperature range are: 1—a deformation like texture, 2—a rotated inorder a-recrystallization temperature texture, 3-a retained u-recrysraIlization texture, on annealing at lower temperatures of the ß-region, 4—a transformation texture based on recrystallized a and predicted by the Burgers&apos; relationship, and 5—a ,ß-cube texture. These results are examined in terms of current theories of recrystallization textures. UMEROUS investigators have described the tex- ture obtained by cold rolling the hexagonal metals, titanium, zirconium, and beryllium, which have c/a ratios less than that of ideal packing, 1.633. The basal planes are rotated out of the rolling plane, about the rolling direction, so that the basal poles are tilted toward the transverse direction as shown schematically in Fig. la. In all instances but one,&apos; it was also reported that the [1010] direction was parallel to the rolling direction (see Fig. lb). Hot rolling has been reported as causing a similar tilt of the basal poles in the transverse direction (see Fig. la) and causing the [1010] direction also to be parallel to the rolling direction as shown schematically in Fig. lb. Annealing after deformation does not appreciably change the tilt of the basal poles in the transverse direction." Beryllium2-7 continues to have the [1010] direction in the rolling direction after annealing, and similar observations for titanium and zirconium&apos; . have been reported for annealing at fairly low temperatures, again as in Fig. lb. At higher annealing temperatures, however, the recrystallized grains of titanium" and zirconium have an orientation such that the [1120] direction is approximately in the rolling direction, although the basal poles are still inclined in the transverse direction. Figs. la and lc show the resulting orientations schematically. This change in orientation has been described as a nominally ±30° rotation of the hexagonal crystallites about the basal poles of the cold rolled texture and is apparent from the results which are summarized in Table I for investigations with the X-ray diffraction technique employing film. The angles y, , and ß are indicated in Fig. 2 which represents the stereographic projection of (1070) poles for the mean orientation of a pole figure. Texture determinations for titanium using the Geiger counter spectrometer have provided similar results except that in some instances additional components of the texture were proposed, as shown by the summary of data in the upper half of Table 11. On the other hand, the spectrometer technique, when applied to zirconium,* has revealed a splitting Recently completed studies of the textures of annealed zirconium", show zirconium to possess textures very similar to those reported here for titanium. Therefore, much of this discussion will include zirconium by virtue of its close similarity to titanium in pref erred orientations. of the intense areas of the pole figure for samples annealed at 600°C. This splitting could be described by a 7" rotation of the tilt axis about the normal to the rolling plane. Such a splitting for the annealed texture relative to the cold rolled texture was not observed in other determinations for either zirconium or titanium using the less sensitive film X-ray methoe and makes the relationship between the two types of texture more complex than the simple rotation about the (0001) pole based on film work. The more precise investigations on zirconium permit the descriptions in the lower part of Table 11, which show that the texture depends quantitatively on the temperature of annealing. When zirconium is annealed at temperatures up to 400°C, the texture is similar to the cold rolled texture, while annealing in the range 500" to 900°C produces a texture which is only approximately described as [11%] in the rolling direction. More precisely described results for zirconium show that the two types of splitting ( 1—about an axis in the rolling plane through an angle given in the second column in Table II and 2—about the normal to the rolling plane through an angle given in the third column of Table 11) depend on annealing temperature. The [1120 is the rolling direction only when the annealing temperature is in the vicinity of 900°C

Jan 1, 1957

By Carle R. Hayward

WHEN lead production began to recede from the peak productions of 1929 many plants took advantage of the curtailed operations to make necessary improvements and repairs about the plant. There followed a period of adjustment to synchronize production and demand, which merged in the present period of marking time. The lead smelting and refining industry has produced little that is new during the past year except in what might be called makeshifts, which have been introduced to make operation at fractional capacity profitable. This is more particularly true in refineries than in smelters. The Parkes process had become standardized during the period of normal production, with softening, desilverizing, and dezink-

Jan 1, 1935