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By E. A. Ernst, N. F. Carpenter

The benefits derived from an acidizing treatment are a function of the penetration achieved by the acid before complete spending. Additional penetration may be achieved by (1) controlling acid leak-08 into formation pores in the channel faces, and (2) retarding the reaction rate of the acid. A recently developed chemical additive consists of a synthetic polymeric material which absorbs hydrochloric-acid solutions, when suspended therein, swelling up to 40 times its original volume. These swollen particles have the ability to deform and seal-08 formation pores, providing fluid-loss control. In addition, they provide a diffusion barrier between the fracture face and the acid solution, prolonging the spending time of the acid. Field applications of this new technique have shown promising results. A method of conducting acid fluid-loss tests, using carbonate cores, is believed to provide fluid-loss data that are more representative of formation conditions than the conventional filter-paper determinations. INTRODUCTION The concept of oilwell acidizing has changed since its first commercial application, 30 years ago. Originally, it was visualized that the acid penetrated thousands of tiny pores and flow channels in the matrix rock, enlarging them by dissolving the carbonate walls. The resultant permeability increase was assumed to be the responsible factor in increasing production from the well. Recent laboratory studies,' however, have shown that this does not provide the complete picture. Although this type of individual pore penetration by the acid does take place during acid "soaks", designed to overcome "skin effect" due to mud invasion in the immediate vicinity of the wellbore, many years of experience have shown that considerable pressure is required to attain any appreciable injection rate into the fine capillary pores of the rock. During most acidizing treatments, the bottom-hole pressure build-up due to the restriction of flow into the formation exceeds the "breakdown" pressure of the rock so that a fracture is induced. In most cases, such fractures open up along natural, incipient fissures and zones of weakness in the rock and, therefore, tend to follow the natural stress pattern of the rock—whether it be horizontal, vertical or inclined. Because of the comparatively greater permeability of the channel in relation to that of the matrix, the bulk of the acid volume is diverted into the newly opened fracture. Here it quickly penetrates the formation, opening and ex- tending the fracture in much the same manner as a conventional fracturing fluid. Unlike the fracturing fluid, however, most acidizing solutions contain no propping agent; thus, the open fracture will again close when the injection pressure is relieved. Laboratory studies2 have shown that in many cases the etching of the fracture faces, resulting from the reaction between the acidizing solution and the carbonate rock, is nonuniform due to the heterogeneity of the rock structure. As a result, the two fracture faces no longer match when pressure is released, and support pillars and intermediate voids remain, forming a high-conductivity channel for well fluids. Unfortunately, this is not true over the entire area of the fracture, but only over that portion of the fracture where the rock has been partially dissolved by the acid. The acid solution spends as its travels away from the wellbore; once it has completely spent, even though it may provide additional mechanical fracture extension, no additional benefit due to etching of fracture faces can be expected. Studies of acid reaction rates under formation conditions,3 observing the effect of different variables upon spending time, have shown that the reaction was often so rapid that very little penetration of the formation occurred before the acid was spent. Study was undertaken to devise methods of increasing the penetration of the acid before spending, so as to provide greater benefit from the acidizing treatment by etching a greater portion of the fracture faces. Several techniques were devised to accomplish this purpose. First, chemical additives were developed which were designed to retard the reaction rate of the acid, causing it to penetrate a greater distance from the wellbore before finally becoming spent. Another method was to increase the injection rate of the acid. However, it was found that the resultant increased shear tended to accelerate the reaction rate of the acid, partially offsetting the benefits of the higher injection rate insofar as achieving increased penetration before spending was concerned.' Another approach to the problem of achieving increased penetration was the development of fluid-loss additives for acid solutions, which would minimize the volume of acid lost into formation pores in the fracture faces and provide maximum fracture extension for the volume of acid injected during the treatment. The use of fluid-loss additives is now considered the most effective method of providing maximum fracturing-fluid efficiency.~ Unfortunately, this latter technique does not solve the problem of rapid reaction rate, with consequent limitation of the fracture area benefited by reaction with unspent acid. A newly developed acid additive overcomes many of these limitations by providing the dual benefits of fluid-loss control and mechanical retardation of acid reaction

By H. B. Probst

Mechanical twins were produced in zone-refined tungsten single crystals by explosive working at room temperature. These twins are parallel to (112) planes and have irregular boundaries rather than the classical plane twin boundaries. These boundaries aye grooved surfaces in which the grooves themselves are parallel to a <111> direction and the sides of the grooves appear to be par-allel to (110) planes. TWINS were produced in tungsten single crystals by explosive working at room temperature. These twins differ in character from any previously reported for tungsten; however, they are similar to those found in molybdenum after compression at -196°C.1 Deformation twins "resembling Neumann bands in ingot iron" have been observed in tungsten by Bech-told and Shewmon.2 This observation was made with sintered polycrystalline tungsten pulled in tension to fracture at 100°C and using a strain rate of 2.8 x 10-4 sec-1. More recently Schadler3 found deformation twins in zone-refined tungsten single crystals pulled in tension at -196"&apos; and -253°C. These tests were conducted using a strain rate of 3.3 x l0-4 sec-1, and the twin bands were found to be parallel to a (112) plane. Deformation twins in tungsten&apos;s sister metal, molybdenum, were observed by Cahn.4 These twins were produced by compressing small (0.7 mm) vapor-deproducedposited molybdenum single crystals at -183°C. The compression was performed &apos;by impact." By the use of precession X-ray techniques, Cahn was able to identify the twin plane as {112} and the shear direction as <1ll>. Mueller and Parker1 produced deformation twins in polycrystalline electron-beam-melted molybdenum by compression at -196°C. Their "loading rate" was 5000 psi per min which, judging from their stress-strain curve, corresponds to a strain rate of approximately 0.3 x 10-4 sec-1. These twin bands were found to be parallel to (1 12) planes; however, they differed in appearance from previously observed twins. In place of straight and parallel twin boundaries they were found to be irregular, jagged, and sawtoothed. The sides of the saw teeth were identified as (110) planes and irrational planes of a (111) zone. The twins observed in the present work in tungsten single crystals are similar in appearance to those of Mueller and Parker in polycrystalline molybdenum. The starting material used in this investigation was 3/16-in. diam commercial tungsten rod produced by powder-metallurgy techniques. This material was converted to a single crystal by the electron-bombardment floating-zone technique.= The process was carried out in a vacuum of 10-5 mm of Hg using a traversing speed of 4 mm per min. Segments (=2 in. long and 3/16 in. in diam) of two crystals (A and B) produced in this manner were studied. Crystal A received one zoning pass, while crystal B received two passes. The two crystals were explosively worked at Bat-telle Memorial Institute in the following manner. A 1/2-in.-thick layer of plastic was applied to the crystals to serve as a buffer in an attempt to prevent cracking. The composite, crystal and buffer, was then wrapped with 1/8-in.-thick DuPont sheet explosive EL506A2 and detonated in water at room temperature. Metallographic samples of the worked crystals were prepared, and back-reflection Laue X-ray patterns were obtained using unfiltered molybdenum radiation. RESULTS AND DISCUSSION Blasting the crystals as described above failed to prevent cracking. The crystals fractured into several fragments about 3/16 to 1/2 in. long; however, the fragments were of sufficient size to be useful for the subsequent study. The diamond pyramid hardness of the crystals after blasting was in the range 430 to 450 as compared with 340 for the as-melted material, which shows a definite hardening resulting from plastic deformation. These hardness values were obtained using a 1000-g load and taking readings only in sound portions of the crystals free of cracks. The crystals exhibited profuse twinning as shown in Fig. 1. No such structure is present in the as-melted condition. Most of these twins have jagged twin boundaries and are similar in appearance to those found in molybdenum by Mueller and Parker. The twins in both crystals were found to be parallel to {112} planes. This identification was made by using the conventional two-trace method. Subsequent efforts to describe these twins more fully were carried out on crystal A. If the longitudinal axis of crystal A is placed in the (001)-(011)-(Il l) basic triangle of the standard cubic stereographic projection, as in Fig. 2, then the two sets of twins shown in Fig. 1 are parallel to the (112) and (121) planes. Fig. 3 shows a schematic representation of a twin with jagged boundaries. This type of twin with a <111>

Jan 1, 1962

By J. J. Gilman, T. A. Read

FOLLOWING the deformation 01 zinc monocrys-tals, sharply bent basal planes are observed near several types of inhomogeneities. Three of these in-homogeneities have characteristics which are quite regular so that they can be studied and analyzed. These are compressive kink bands, "deformation bands," and the inhomogeneities near end restraints. The present paper describes experiments in which "deformation bands" were artificially produced, and bend plane phenomena are discussed in terms of dislocation theory. Also, two new bend plane phenomena are described. The importance of bend plane phenomena in the deformation of crystals is not widely recognized. Many phenomena may be explained in a manner similar to the discussion in this paper. Jillson1 has pointed out that the "punching effect" in zinc is a bend plane phenomenon and is not caused by prismatic slip.&apos; Bowles3 has suggested that they may be involved in diffusionless phase changes. Cahn4 has discussed the role of bend plane formation in the polygonization of zinc. Experimental Work Tensile Kink Bands: Because of the geometrical similarity between "deformation bands" and "kink bands" (compare Fig. 1 of this paper with Fig. 1 of the paper by Hess and Barrett"), the band shown in Fig. 1 of this paper will be called a "tensile kink band," and that shown by Hess and Barrett will be called a "compressive kink band." It is felt that the term "deformation band" should be reserved for banded structures in polycrystalline materials such as iron." Tensile kink bands seem to form spontaneously in aluminum crystals deformed by tensile loading.7-10 In zinc and cadmium crystals they do not form in good, carefully loaded specimens.&apos;." However, tensile kink bands can be produced artificially in zinc crystals. The present authors did this by scratching one of the flat surfaces of triangular crystals transversely with a sharp needle. Natural tensile kink bands caused by inhomogeneities sometimes appeared in deformed crystals which were identical in appearance with the artificially produced ones. Zinc monocrystals were grown by the Bridgman method in graphite molds. Chemically pure zinc (99.999+ pct Zn) was used and the molds were sealed inside evacuated pyrex tubes during growth. The crystal cross sections were equilateral triangles with a typical base of 0.210 in. The artificial kink band shown in Fig. 1 is typical of tensile kink bands in zinc. The band lies between two bend planes which run from upper right to lower left and is inclined oppositely to the slip bands which are sharply bent at the two bend planes. The general form of the artificial tensile kink bands was independent of the scratch depth (1 to 5 mils deep) and also independent of which side was scratched. These variables did cause variations, however. Deep scratches produced more localized kink bands than light scratches. Also, if the angle between the slip plane traces and a transverse scratch varied appreciably among the three sides, then localization of the resulting kink bands also varied. Furthermore, if the slip direction lay nearly parallel to the scratched side, the band was more developed near the scratched side than at the opposite edge. Scratches produced tensile kink bands for crystal orientations from xo = 15" to x, = 75". Fig. 2 shows a scratched crystal after deformation. One triangular side lies in the plane of the photograph. The right hand tensile kink band was produced by a transverse scratch on the upper right side. The next two kink bands were the result of scratches on the front surface. The kink band at the left was caused by a scratch on the lower back side. All four bands have the same general form. A longitudinal scratch was also made on the crystal shown in Fig. 2 to determine the effect of a scratch on the critical shear stress. The critical shear stress of the scratched region was 33.9 g per sq mm compared to 24.4 g per sq mm for the un-scratched region above it. Fig. 3 shows Laue patterns of the crystal shown in Fig. 2. Fig. 3a shows the pattern of the undeformed crystal. The orientation was x, = 21°, A, = 31". After deformation, Fig. 3b was made of the homogeneously deformed portion of the crystal. The spots are compact but split into two halves. This region was elongated 45 pct and its orientation was x = 14", X = 20"; the sine law predicts x = 14", A = 20.5". Fig. 3c was taken near the center of the middle tensile kink band of Fig. 2. The pattern shows a range of orientations and polygonization in this region. The spread in orientation was due to the fact that the basal planes were curved (see Fig. 1) rather than flat as in the ideal case. Some may also have been the result of elastic distortions and "local curvatures." The orientation range was x = 23" to 32", A = 30" to 42". It is apparent from Fig. 3 that the material inside and outside the kink band rotated in opposite directions with respect to the tension axis during deformation. The orientation calculated from the ideal configuration of Fig. 9,

Jan 1, 1954

By M. J. Spendlove, H. W. St. Clair

The problem frequently arises, particularly in refining metals or smelting scrap metals, of separating metals in the metallie state. Many metals may be separated by taking advantage of their difference in vapor pressure. Such separations can be made at atmospheric pressure, but the separations are much more selective and can be carried out at considerably lower temperatures if the distillation is done at pressures of a few millimeters or less in an evacuated enclosure. Until recently, this has not been considered feasible as a metallurgical operation, but the recent improvemcnts that have been made in vacuum technology have broadened the applicability of vacuum processes and have prompted re-examination of low-pressurc distillation of metals as a practicable process. The distillation of zinc from lead is one separation that has already been reduced to practice.l This paper is the first of a series of studies being made on separation of nonferrous metals by distillation at low pressures. Although these experiments were confined to the separation of zinc from aluminum, the significance of the results is by no means confined to these two metals. The purpose has been to investigate a metallurgical technique rather than merely to devise a means of separating specific metals. The experimental work on distillation of zinc from zine-aluminum alloys at reduced pressure grew out of earlier work on distillation at atmospheric pressure.2 The earlier work indicated that it would not be practicable to decrease the zinc in the alloy much below 10 pct owing to the high temperature required. Therefore attention was turned to distillation ah low pressures, at which lower temperatures are required. After preliminary tests were made in a small, evacuated tube furnace, a larger furnace having a capacity of 100 to 150 Ib of metal per charge was constructed. Distillation tests were first made on pure zinc and then on aluminum-zinc alloys of various composition. Particular attention was given to the limit to which zinc could be reduced in the residual metal. Data were also taken on the rate of evaporation, and heat balances were made for both the crucible and the condenser. Distillation Furnace The vacuum-distillation unit is illustrated schematically in Fig 1. The major components are the induction furnace, the condenser, the vacuum system, and the power-conversion unit. Power is supplied to the induction furnace from a 50-kw 3000-cycle motor-driven alternator. The pressure in the furnace is reduced by a vacuum pump having a nominal pumping speed of 10 liters per sec. When in operation, the metal vapors travel upward from the furnace to the water-cooled condenser where they are collected in amounts of 50 to 100 lb. The condenser is removed with aid of an electric hoist. When the system is under vacuum, the condenser is made self-sealing by a rubber gasket between the smooth-faced, water-cooled flanges at the top of the furnace and the bottom of the condenser. The pressure of the atmosphere is more than sufficient to insure sealing. At the conclusion of an experiment, the residual metal can be removed from the furnace by removing the condenser and tilting the furnace with the electric hoist. The metal was cast into the molds carried on a mold truck. A photograph of the furnace and auxiliary equipment is shown in Fig 2. The details of the vacuum furnace are illustrated in Fig 3. The furnace proper is made vacuum-tight with rubber gaskets placed at each end of a fused quartz cylinder. A clamping plate at the bottom and a ring at the top are made to squeeze the rubber between the metal and the end of the quartz tube. A large graphite crucible placed inside the quartz cylinder is thermally insulated and physically supported by refractory insulating bricks. A thermocouple in a quartz protection tube is located at the bottom of the crucible: the leads pass through a rubber seal in the bottom plate. The supporting structure for the furnace is an angle iron frame with transite board sides. The condenser is made in the form of a water jacketed cylinder with an opening to the vacuum line at the top. The bottom has a projecting skirt inside the machined flange to provide additional cooling for the rubber gasket. Condenser sleeves are made in the form of two semicylindrical pieces of sheet metal that fit snugly inside the cooling jacket. The split sleeve facilitates removal of the condensate. Measurement of Temperatare and Pressure The metal temperature was measured by a platinum-platinilm rhodium thermocouple inserted in a well extending up into the bottom of the graphite crucible. During rapid evaporation there is a wide difference in temperature between the surface and the main body of metal in the crucible because of the large amount of heat that must be conducted to the surface to supply the heat of evaporation. The heat of

Jan 1, 1950

By J. F. Peters

The term solution loss is discussed and defined. Examples are given showing that solution loss may either have a favorable or unfavorable effect on blast furnace performance. A theory is advanced explaining the contradictions encountered during earlier studies of the problem. MANY papers have been written and numerous theories advanced concerning the chemical reactions in the iron blast furnace. Richards, discussing the utilization of fuel in the blast furnace, said: "All the carbon burnt in the furnace should be first oxidized at the tuyeres to CO and all the reduction of oxides above the tuyeres should be caused by CO, which thus becomes CO,. This dictum is not Gruner&apos;s own words, but expresses their sense, and from the point of view of the present discussion, it is the correct principle upon which to obtain the maximum generation of heat in the furnace from a given weight of fuel." Richards also said the cubic feet of wind per pound of coke does not express how efficiently a furnace is running, but it will be shown later where Howland paid much attention to this figure. Richards pointed out a shortcoming of Grun-er&apos;s discussion in that Richards claimed there is not enough CO generated in a good working blast furnace to possibly combine with the oxygen of the ore, so some of it must be reduced by solid carbon. Mathesius2 made two important contributions. First, he wrote the direct and indirect reduction equations showing the heats of formation in each case and stressed that the indirect reduction reaction produces heat while direct reduction absorbs heat. Secondly, he substantiated Richards&apos; conclusion that there may be a deficiency of CO gas and that some direct reduction is required, but he pointed out that "direct reduction in the hearth has often been confused with direct reduction in the stack, which more correctly should be termed &apos;premature combustion&apos;." Then he said that "as this carbon is consumed partly by economical (direct reduction in the hearth) and partly by wasteful reactions in ever varying proportions, it is evident that the relation of this total carbon gasified above the tuyeres to the fuel consumption of a furnace cannot possibly have a direct bearing on the economy of furnace operation. The premature combustion of carbon (CO2+C?2CO) must therefore in all cases be considered a detrimental reaction." Johnson3 discussed this whole problem. Through experience Johnson found that when departure from Gruner&apos;s ideal working is considerable, the fuel economy is poor, and when the quantity of blast goes up, the fuel economy goes up in the same proportion. Johnson said that the wind per pound of coke is a measure of fuel economy, which will be shown to be wrong. Howland4 established the fact from operating data that there was no relationship between the coke consumption of a furnace and the percentage of coke burned at the tuyeres. But some of his calculations have led to questionable conclusions, such as "it is practically impossible to obtain low coke consumption unless we keep our wind low." Howland made an important contribution when, in his concluding paragraphs, he reached certain negative decisions as to why one coke works better than another. He said the reason one coke works better than another in a blast furnace is not because: 1—there is any difference in the percentage gasified at the tuyeres, or 2—there is any difference of wind required per pound of coke. Korevaar5 propounded a new theory on combustion which disagrees with Gruner because "as far as we can see, this is sufficient proof for the invalidity of Gruner&apos;s ideal, for if it was valid, the Low coke consumption should always be accompanied by a higher percentage of carbon burned at the tuyeres. This not being the case, we must deliberately give up our belief in Gruner&apos;s ideal." The part in italics will be proved to be wrong. Martin" contended that there was not enough reducing capacity in the blast furnace unless solution loss occurs. He came to a series of conclusions, such as: 1—The greatest efficiency of the blast furnace may not be attained when the reduction is performed entirely by carbon monoxide as demanded by Gruner&apos;s definition of the ideally perfect working of the blast furnace. 2—The so-called solution loss reactions, which should more properly be termed direct reduction reactions, promote furnace efficiency. 3—In modern blast furnace practice, the carbon consumption of the process is determined primarily by the carbon needed for reduction purposes, any thermal deficiency created by the reduction process being balanced in practice by the use of blast heat. From a practical standpoint further discussion of the theory of chemical reactions is of little moment. There is however one phase of the general theory of chemical reactions which is very important and that is the combustion of coke in the blast furnace. The purpose of this paper is to show that the reaction

Jan 1, 1955

By L. W. Emery

Undergrorlnd combustion operations were initiated in a 60-acre Bartlesville sand "shoe-string" reservoir in Allen Connty, Kans., in 1956. Tests in separate patterns were conducted using various co~nbinations of air and recycle gas to propagate combustion fronts from the injection toward the producing wells. These patterns were made up of 6 injection and 20 prodrrcing wells Gas and liquid prorluctiorz from each pattern was measured on an individual-well basis, and comparisons were made between the three patterns to ascertain the relative effects of injected gas composition on production behavior. Breakthrough of the combustion front at a well was characterized by an increase in water production from the well followed by an increase in bottomhole temperatrrre to approximately 250" F. After burning fronts had broken through at five producing wells, operations were terminated in 1960. From the total project approximately 79,000 bbl of oil were produced during thermal operations at a cumulative produced GOR of 23 Mcf/bbl. No appreciable change in the character of the produced crude was observed. Combustion in the reservoir was maintained with injected gas compositions ranging fronz 6 per cent oxygen in recycle gas to 100 per cent air. lnjectiotz of large quantities of recycle gas resulted in higher producing GOR&apos;s from offset wells than were measured from a pattern into ~vhich straight air ~vas injected. The air required to move the combustion front through I acre-ft of reservoir was computed to be 20 MMscf. This valrre was found to be relatively independent of the quantities of recycle gas injected. The recovery efficiency from the swept area was esti~izated to be about 59 per cent. Areas swept were similar in shape to tlzose obtained with a laboratory potentiometric model. Samples of sund taken from behind the burning front by coring indicated almost total oil removal from the sand. Petrographic analysis of the core samples indicated that the sand had been heated to peuk temperature of rlbout 1,200" F. No rignificant difference in peak temperature was forrnd in two areas where compositions of injected gas were quite different. Compression costs for thermal recovery were estimated to be $1,20/bhl of produced oil. INTRODUCTION The use of the "forward combustion" process as an oil recovery method has received a great deal of attention. This method involves ignition of the formation in an injection well, followed by propagation of a combustion front through the reservoir. Combustion is maintained by the injection of an oxygen-containing gas to react with reservoir hydrocarbons. As the flame front progresses through the reservoir, oil and formation water are vaporized, driven forward in the gaseous phase and recondensed in the cooler part of the formation. In turn, the condensed fluids push oil into the producing wellbores. Completed field tests of the process were first reported by Kuhn and Koch,&apos; and by Grant and Szasz.&apos; Results from other tests have since been reported by Walter,3 by Moss, White and McNeil,&apos; and by Gates and Ramey." ach of these tests essentially utilized a single injection well surrounded by four or more producing wells. Sinclair Research, Inc., elected to do field experimental work using a number of test patterns in a single field in order that comparisons between various operating schemes could be made. The site selected and purchased in 1955 for this experimental work was a 60-acre Bartlesville sand reservoir located in Allen County, Kans. Combustion operations were initiated in mid-1956. Between that time and termination of the project in mid-1960, combustion fronts were propagated from injection wells to producers in three separate well patterns, using different mixtures of air and recycle gas. The test was terminated before sweep of the three patterns was complete so that information about the effect of combustion on the swept areas could be obtained by coring. Results from the test in the form of injection and producing well performance have been carefully recorded, and these form the general basis for this paper. DESCRIPTION OF RESERVOIR The reservoir in which the combustion tests were conducted is a Bartlesville sand "shoe-string", typical of a number of small reservoirs in Southeastern Kansas. Average reservoir characteristics are shown in Table 1. Fig. 1 is an isopachous map of the producing sand showing the reservoir to be approximately 400-ft wide and 2,500-ft long. Maximum net productive sand thickness is 21 ft. Fig. 2 shows a typical core analysis obtained by coring with water-base mud. The reservoir has no appreciable dip and is closed on the sides by degradation of sand into shale. The main body of sand is heavily laminated with shale stringers, which are not continuous between wells. The main reservoir is overlain by 30 to 40 ft of laminated low-permeability sand and shale streaks. No information is available on the original properties of

By G. Bhat, J. F. Libsch

lsoembrittlement curves depicting the influence of time and temperature in the range 800&apos; to 1260°F (425&apos; to 680°C) on the development of embrittlement in a commercial chromium alloy steel and a commercial molybdenum alloy steel are presented. Two distinct regions of embrittlement occur in the chromium alloy steel: I—at 800&apos; to 1000°F (425&apos; to 540°C) and 2—in the region just below the lower critical temperature. Embrittlement is most pronounced at 800&apos; to 1000°F, decreasing very rapidly with increasing temperature above this region, only to increase again as the lower critical temperature is approached. The data suggest two distinct modes of embrittlement with possible superposition of the two modes at extended embrittling times in the temperature range 1100° to 1150°F (590&apos; to 620°C). While the molybdenum alloy steel shows little susceptibility to embrittlement at 800&apos; to 1000°F (425&apos; to 540°C), considerable embrittlement may occur just below the lower critical temperature. THE subject of temper embrittlement in alloy steels has received considerable attention in the last few years. Points of view on the mechanism of embrittlement differ, however, resulting in part from the incompleteness of the data developed and in part from the speculation regarding the susceptibility of plain carbon steel to temper embrittlement. Libsch, Powers, and Bhat1 carried out short-time embrittling treatments on an AISI 1050 steel and demonstrated that hardened plain carbon steels are quite susceptible to embrittlement when tempered in the range from 850°F (455°C) to the lower critical temperature. The isoembrittlement diagram,&apos; representing the embrittling characteristics of this steel, is reproduced in Fig. 1. It is evident from the shape of the curves shown that embrittlement in plain carbon steel increases progressively with both temperature and time in the embrittling range. A comparison of the isoembrittlement diagram for AISI 1050 steel with that presented by Jaffe and Buffum&apos; for an SAE 3140 steel shows that up to 930°F (500°C) the isoembrittlement characteristics of the plain carbon steel are similar to those of SAE 3140 steel, although the embrittlement is much more severe in the latter steel. Above 930°F (500°C), the rate of embrittlement in the plain carbon steel increases continuously with increasing temperature; whereas, in the SAE 3140 steel, the embrittlement rapidly decreases. The influence of alloying elements upon embrittlement during tempering thus appears to cause a decrease in embrittlement above the region of maximum embrittlement, i.e., 850" to 1000°F. The question naturally arises as to what effect individual alloying elements have upon the embrittling characteristics of the plain carbon steel. Current knowledge on the influence of alloying elements on temper brittleness may be found in the review papers of Hollomon" and Woodfine. Hollo-mon," from the results of other investigators, has shown that, in general, the amount of embrittlement increases with increasing alloy content (except for molybdenum and possibly tungsten and columbium). Jaffe and Buffum," by a comparison of the embrittlement in a plain carbon steel with that of a SAE 3140 steel postulated that the presence of alloying elements in moderate amounts tends to retard the development of temper brittleness. It is difficult to determine what effect chromium has upon temper brittleness, since most of the information available has been based on the combined effect of other elements with chromium, particularly nickel and manganese. However, Wilten, and recently Jolivet and Vidal,&apos; Vida1, and Woodfine have reported that chromium steels are temper brittle, that the embrittlement is reversible with a maximum rate of embrittlement at approximately 975°F (525"C)," and that the susceptibility increases with increasing amounts of chromium. Taber, Thorlin, and Wallacel" have found a large embrittling effect with increasing chromium content in a medium C-Mn-Ni steel. But Hultgren and Chang," from their experiments conducted on synthetically prepared ternary Fe-C-Cr alloys, could not conclude that these alloys are susceptible to temper embrittlement. However, on addition of manganese or phosphorus, these Fe-C-Cr alloys became susceptible, from which fact they concluded that the embrittlement developed in chromium-bearing Fe-C alloys is due chiefly to the presence of these elements. Considerable data are available to show that molybdenum decreases the susceptibility of steel to temper embrittlement. However, its effectiveness in preventing or decreasing embrittlement appears limited to its presence in small amounts. Vidal" has shown that a plain 2 pct Mo steel was susceptible. Hultgren and Chang" also have shown that molybdenum additions in excess of 2 pct to synthetically prepared Ni-Cr steels did not prevent embrittlement. Jolivet and Vidal&apos; and Lea and Arnold found that molybdenum reduced temper brittleness. Lea and Arnold further stated that molybdenum decreased the rate of embrittlement rather than the total amount of embrittlement, whereas Preece and Carter" have shown that the presence of molybdenum greatly reduces the equilibrium extent of the change at a given temperature but does not appear to influence the rate of embrittlement. There appears to be very little information as to how molybdenum by itself affects the temper brittleness susceptibility of a plain carbon steel.

Jan 1, 1956

By F. W. Schremp, J. W. Chittum, T. S. Arczynski

This paper describes a laboratory study of causes of external casing corrosion and the test work that led to the use of oxygen scavengers to prevent this attack. External casing failures are classified as water-line, casing-casing, collar and body failures. A corrosion mechanism based on principles of differential oxygen availability is developed that is consistent with facts known about each kind of failure. The field use of oxygen scavengers is depicted as a direct result of the laboratory study. A part of the paper is devoted to reporting on the field use of hydra-zine to control external casing corrosion. Results of field measurements made over a period of several years are presented as evidence of the efectiveness of the hydrazine treatment. The first conclusion reached is that the use of hydrazine materially reduces the cathodic protection requirements for treated wells. This result is interpreted to mean that a reduction is taking place in the amount of corrosion on the casing. Results indicate also that hydrazine shows its greatest usefulness within the first 12 to 18 months after a well is completed when pitting corrosion is likely to be most active. INTRODUCTION According to surveys sponsored by the National Association of Corrosion Engineers,&apos; the cost of repairing casing leaks caused by external corrosion may exceed $4 million per year. In addition, well damage and lost production resulting from casing leaks probably costs the petroleum industry an additional $5 to $6 million per year. Concern about the cost of external casing corrosion led to an extensive laboratory study of factors causing this external corrosion and to the development of a new approach to its prevention. This paper presents a discussion of various causes of external casing corrosion, details of laboratory studies and the results of the field use of an oxygen scavenger in well cementing fluids to prevent the external corrosion of oil-string casing. Measurements on test wells over a period of several years show that cathodic-protection current requirements are greatly reduced when hydrazine is used in cementing mud. Reduction of current requirements can be interpreted to mean that removal of oxygen by hydrazine has greatly suppressed corrosion cells on the external surface of the casing and thereby, has reduced corrosion. To date, hydrazine has been used by the Standard Oil Co. of California in more than 200 well completions. KINDS OF CASING FAILURES A survey of a large number of casing leaks disclosed four types of external casing failures — water-line, casing-casing, collar and body failures. These types are identified largely by their location on the casing. Water-line failures are found just below the surface of water or mud in the casing annulus. Casing-casing failures occur on the oil string just below the shoe of the surface string. Collar failures are found in the threaded ends of casing joints where they are screwed into casing collars. Body failures may occur at any point on the body of a casing joint. Ex- amples of each kind of failure have some of the general characteristics that are shown in Fig. 1. Water-line failures usually result in the circumferential severance of an oil-string casing. The corrosive action causing a water-line failure usually is sharply defined and is limited to a short length of the casing. Casing-casing failures usually are accompanied by pitting corrosion distributed around the oil-string casing for distances up to 100-ft below the shoe of the surface string. Casing-casing failures may also sever the casing. Collar failures seem to start on the first thread at the bottom of recesses between collar and casing joint. Corrosion proceeds across the threads by what appears to be a normal pitting mechanism. Both casing and collar are severely attacked. Body failures are the result of highly localized pitting at any point on a casing wall. Besides the pit that perforates a casing, a large number of other pits usually are found along one side of the casing joint. The pits occasionally are filled with corrosion products consisting largely of oxides and sulfides.&apos; Frequently, the mill scale is largely intact on the rest of the casing. Examination of a casing failure does not always reveal the cause of the failure. Frequently, the necessary details are destroyed when the failure occurs. For example, formation water flowing through a perforation at high velocity may enlarge the hole and destroy any remaining evidence of the cause of the failure. One way to obtain undistorted information about a failure is to study the nature of other pits on the casing in the vicinity of the failure. A study of such pits frequently suggests that they are characteristic of an attack resulting from the differential availability of molecular oxygen.

By R. D. Daniels, T. G. Oakwood

A study was made of the effects of small quantities of hydrogen on the mechanical properties of colum-bium. Tensile specimens, hydrogenated to concentrations of 20 to 200 ppm, were tested at temperatures of 300°, 191°, and 77°K. Although hydrogen was found to have little effect on the strength of columbium, the ductility of Cb-H alloys was found to be quite sensitive to both hydrogen concentration and temperature. At 300°K, an abrupt loss in ductility occurred at a critical hydrogen concentration, although some ductility was observed beyond the tolerance limit. A similar result was found at a lower hydrogen concentration at 191°K. At 77°K, however, a more gradual loss in ductility with increasing hydrogen concentration was observed. Hydrogenated columbium was thus observed to undergo a ductile-brittle-ductile transition. Metallographic examination of fractured specimens revealed extensive porosity at both 77° and300°K which was a distinct function of hydrogen content. At 191°K, although some secondary cracking was noted, the amount of observed porosity was minimal. These observations are interpreted in terms of hydrogen solubility and mobility as a function of temperature and in the role of hydrogen in promoting growth of microcracks. lHE effect of hydrogen on the mechanical properties of the refractory metals is not, at present, completely understood. A number of studies have shown these materials to be susceptible to hydrogen embrittlement. Roberts and Rogers1 have found that vanadium can be embrittled by hydrogen. It was further demonstrated that fracture undergoes a ductile-brittle-ductile transition as the temperature is lowered from 150° to -196°C; i.e., there is a ductility minimum observed at a certain temperature. The ductility is increased by either raising or lowering the temperature from this point. A more complete study by Eustice and Carlson2 on vanadium containing 10 to 800 ppm placed the ductility minimum at about -100°C with variations reportedly due to hydrogen content and strain rate. Ductility minima have also been found at certain temperatures for tantalum containing 7 ppm H3 and 140 ppm H.4 At hydrogen concentrations above 270 ppm, however, the ductility return at low temperatures was considerably reduced.4 In the case of columbium, some disagreement exists in the literature. Eustice and Carlson,5 Wilcox et al.,6 and Imgram et al.4 failed to find a ductility minimum although a composition-dependent ductile-brittle transition was observed. Hydrogen concentrations in these investigations were 20 ppm,5 1 to 30 ppm,6 and 200 to 390 ppm.4 However, Wood and Daniels7 observed a rather pronounced ductility minimum at hydrogen contents ranging from 19 to 252 ppm. Those theories of hydrogen embrittlement involving the precipitation of diatomic hydrogen which have been applies to ferrous metals8-12 do not seem to be applicable to the case of columbium and other exothermic occluders. Such theories propose that extensive crack formation and propagation occurs by the precipitation and expansion of diatomic hydrogen at internal voids and microcracks. However, photomicrographs of hydrogenated columbium do not show any evidence of damage introduced by the sorption and precipitation of diatomic hydrogen; rather, at high hydrogen concentrations, a hydrogen-rich second phase is precipitated.13&apos;14 In addition, a number of these theories require the development of high hydrogen pressures at voids in the structure.8&apos;10&apos;12 This does not appear to be feasible in the concentration ranges discussed in the aforementioned paragraphs. The possible interaction of atomic hydrogen with microcracks resulting from dislocation pile-ups15,16 remains in doubt since pile-ups have not been observed in bcc metals17 including columbium.18 Wood and Daniels7 have put forth the possibility that a hydride precipitation could be responsible for crack nucleation in columbium. Work by Longson19 has shown that hydrogen embrittlement of columbium parallels the bulk solubility limit; i.e., as the solubility increases, for instance with temperature, the amount of hydrogen necessary to cause embrittlement also increases. Although a hydride precipitation appears attractive as a means of nucleating microcracks in columbium, what require more intensive study are the low-temperature anomalies which have been observed, i.e., the ductile-brittle-d&apos;ictile transition characteristics. Also, the hydrogen concentrations where embrittlement occurs are often below the bulk solubility limits determined by Albrecht et al.13,14 and Walter and Chandler.20 This work is an attempt to determine more definitively the effects of concentration and temperature on the mechanical properties of dilute Cb-H alloys. EXPERIMENTAL PROCEDURE Ultrahigh-purity columbium rods, obtained from the Wah Chang Corp., were cold-reduced by rotary swaging. A chemical analysis is given in Table I. The material was cut into cylindrical blanks 1.50 ±0.005 in. long. Individual specimens were either given a stress relief anneal at 750°C or recrystal-lized at 1200°C. Resulting microstructures were either a "bamboo" structure characteristic of a wrought material or a recrystallized structure with a grain diameter of approximately 100 n. All heat treatments were carried out in a vacuum of 10-5 Torr or less.

Jan 1, 1969

By Rudolph G. Wuerker

MUCH descriptive matter has appeared on the subject of suspension roof supports, or roof bolting, as it is more commonly called. The widespread introduction of roof bolting into coal mines and metal mines is truly phenomenal. Mine operators were quick to recognize the advantages of supporting wide openings without hindrance to machine maneuverability and ventilation. Although suspension roof support has long been installed at St. Joseph Lead Co. mines in southeast Missouri,&apos;" its application to coal mining presented new problems, such as proper anchorage and bearing for the bolts, bolt diameter, and spacing of bolts. After continuous testing and experimenting at the mines, standard roof-bolting materials were determined.&apos;!&apos; The study reported in this paper is not concerned with such details as bolt diameter, which may be considered already solved in practice. In the tests discussed here, small models patterned on actual bolts were found to function in the same way and as satisfactorily as their prototypes. The aim of these tests was rather to investigate the influence of roof-bolting systems on the stress distribution around mine openings and to study the fracture patterns obtained in actual testing. Little was found about this in the literature, as testing of suspension roof methods and quantitative measurements are only now coming to the fore. Several suggestions and actual measurements have been made to evaluate critically the functioning of roof bolting systems, single roof bolts, and parts thereof. Outstanding among them is Bucky&apos;s outline of structural model tests.&apos;" Since none of the suggested testing equipment was available, however, for the experiments discussed below, a different approach was chosen. The response of a mine roof under stress has often been compared to that of a beam. The slow coming down and bending through of beam or plate-like banks of shale, sandstone, or top coal is a familiar occurrence, extensively cited in the literature." It was felt that testing of roof-bolt systems installed in a concrete beam which was loaded in bending would be a fair approximation of the behavior of a mine roof underground. Another school of thought considers the roof behavior over an underground opening in connection with the stress distribution all around a circular or rectangular opening. This is more accurate, and leads to the concept of a dome-shaped zone of material destroyed under tensile stress. This is likewise a common sight in unsupported roadways where the continuous fall of roof results in what has been called the natural outline of roof fracture. This theory could not be tested and is treated separately in Appendix B. It is important to note that according to both assumptions the immediate roof fails in tension; the use of a beam in these tests, therefore, should give information valid for either of the two theories. With the testing equipment at hand it was possible to load concrete beams 6xlx0.5 ft under two-point loading, giving an equal bending moment over the center part in which the model bolts were installed. A comparison was made of the ultimate loads needed to break plain beams and beams in which roof bolts were installed. Arrangements were made with: 1—plain beams; 2—bolts with plate washers, some with holes drilled at 90" angles and others with holes drilled at 45" angles; 3—bolts with channel irons underneath; 4—bolts in holes filled afterward with cement; and 5—bolts anchored in a stronger stratum. The foregoing arrangement is made in order of increasing strength, as assumed from the theory of reinforced concrete. Likewise, laminated beams with wooden model bolts and with combinations of the foregoing set-ups were tested. All in all, 21 experiments were made out of the much greater number of combinations possible. There were, too, some trial tests. Enough observations from this limited number were made to interpret the behavior of mine roof, supported by various types of suspension bolts, at fracture. In present-day concepts, which have been proved by mathematical derivations and stress analyses, any opening driven underground will change the distribution and magnitude of the stresses existing around it. It does not matter whether the stresses become visible, as in rocks whose strength is less than the forces acting upon them, or whether they are invisible, as in the gangways lacking evidence of rock pressure. In this latter case the rocks can withstand changes in stress-distribution. To consider the mine roof as a beam, there are, with transversal loading, tensile stresses in the lower fiber and compressive stresses in the upper layers above the neutral axis of the beam. Beams of brittle material such as rock and concrete fail exactly as shown in Fig. 1. Nearly all model beams showed the same fracture pattern as that of a tension crack. The influence of support, by roof bolting or conventional

Jan 1, 1954

By S. W. Nabbs, W. D. Bachman, A. W. Last

A study has been made of wet, autogenous grinding of disseminated copper ores, including testing of a large number of samples from. Kennecott Copper Corp.&apos;s Chino mine. The grindability of the various samples was found to vary over an extremely wide range, even among samples taken from diflerent locations within the same ore body. These variations in grindability can be qualitatively correlated with the composition of the rock and the fracturing and alteration that had occurred. Because of the large-scale nonhomogeneity of disseminated copper ore bodies, extensive and careful testing of autogenous grinding is needed. An accurate correlation of the grinding data from various samples with the geology of the ore body and with mining plans will indicate if variations in grinding mill throughput can be avoided in practice by mixing or blending of mill feed and if, therefore, autogenous grinding can be applied successfully. In the past 15 years, comminution by primary autogenous grinding has been investigated and applied economically in the minerals industry, particularly in iron-ore processing. Successful application of primary grinding is described by A. A. Dor,l.&apos; P. B. Dettmer,334 R. R. Turner, and W. F. McDermott." other authors describing the advantages of primary grinding include H. Hardinge,&apos; R. T. Hukki,&apos; D. S. Coyle" and F. C. Bond,"," to mention only a few. For the past several years, Kennecott Copper Corp.&apos;s Metal Mining Div. Research Center has been investigating primary grinding of disseminated copper ores. Both fully autogenous and semi-autogenous grinding has been explored. Ores from three of Kennecott&apos;s western operating properties and from several development properties have been tested. Host rocks include diabase, schist, silicated and argillaceous limestones, shale, granite porphyry, quartzite, and granodiorite. The object of this paper is to supplement the general knowledge in the field of primary grinding by presenting information gained from the studies with the various disseminated copper ores. Data are presented which show that carefully considered test programs are required to determine the practicality of applying autogenous grinding to a specific mining-beneficiation operation. Particularly this is so if more than one rock type occurs within an ore body or if there are wide variations in the degree of rock alteration within a deposit. Extensive sampling and subsequent testing of the ores within a mine are required if serious pitfalls are to be avoided and good engineering data acquired. We hope that this presentation will stimulate studies to increase the understanding of the fracturing process applicable to large pieces of rock, a factor which is important in autogenous grinding. Test Facilities and Procedures Both pilot-scale and full-scale test facilities were used to investigate primary grinding characteristics of the disseminated copper ores. A 6-ft Hardinge Cascade mill, a 24-ft Hardinge Cascade mill, and a 5-ft Aerofall mill were used in the studies. The Cascade mills were operated wet, while the Aerofall mill was operated both wet and dry. The Cascade mills were operated both for fully autogenous and for semi-autogenous grinding; the Aerofall mill was operated only semi-autogenously. Fully autogenous grinding, as the term is used in this presentation, is defined as the direct grinding of coarsely crushed rock by the action of the rock mass. Semi-autogenous grinding is defined as the grinding of coarsely crushed rock with the rock mass supplemented by a steel ball charge. Most of the test work was done using the 6-ft Cascade mill, operated under fully autogenous grinding conditions. Only limited testing was conducted with the 5-ft Aerofall mill and only one ore was tested using the 24-ft Cascade mill. This presentation is concerned, therefore, primarily with fully autogenous, pilot-scale grinding in the Cascade mill. Fig. 1 presents a diagram of the test facilities used with the 6-ft Cascade mill. Principal test equipment consists of a 6-ft Cascade mill, an 18-in. spiral classifier, an elevator, a 2 x 9-ft belt feeder, a conventional integrating kilowatt-hour meter, and a recording wattmeter. In testing, feed charges, adjusted as to weight, were deposited as evenly as possible on the belt feeder every 15 min. The weight of each charge was adjusted to maintain a mill charge of 27 to 30% of mill volume. The classifier was adjusted to obtain the desired size-consist of the final product, generally at 20% + 100 mesh. After equilibrium was established in the circuit, the feed, mill discharge, classifier sand, and classifier overflow were sampled and their size distributions determined. During the sampling periods, the power drawn by the mill was measured to determine the energy consumed in the reduction of the size consist of the ore solids. Fully Autogenous Pilot Grinding Tests on Various Kennecott Ores Typical test results obtained in the 6-ft Cascade mill with samples from Kennecott&apos;s Ray Mines Div., Chino Mines Div., Utah Copper Div., and a development property designated as Mine A are presented in Table 1. As shown, the ores that were ground most easily in the pilot-scale mill were composite mine samples obtained

Jan 1, 1971

By W. C. Leslie

The textures of cold-rolled and of annealed iron are compared with those of an iron-0.8 pct copper alloy in which the amount of precipitation after cold rolling was controlled. Previously published pole figures -for cold-rolled and for annealed iron are confirmed. The effects of precipztatiotz after cold rolling are to retain the cold-rolled texture after annealing, to inhibit the formation of the usual allnealing texture, and to produce elongated recrys-tallized ferrite grains. It is suggested that the inhibition of new textures by precipitation after cold rolling is a general phenomenon. A great deal of attention has been paid to the development of texture during the secondary or tertiary recrystallization of ferritic alloys, but very little work seems to have been done on the control of texture during primary recrystallization. If such control were attained, it might be possible to simplify the processing of oriented materials or to change the characteristics of current cold-rolled and an-nealed products. From previous experience, it seemed likely that texture could be controlled by recrystallizing a supersaturated solid solution. Green, Liebmann, and Yoshidal found that the formation of preferred orientation in aluminum (40 deg rotation about <111> relative to the deformed matrix) was inhibited when iron was retained in supersaturated solid solution in the strained aluminum. The authors attributed this inhibition to iron atoms in solid solution. There is, however, an alternative explanation. Green et al, took a highly supersaturated solution of iron in strained aluminum and heated it to an unspecified temperature for recrystallization. It is probable that precipitation occurred prior to and during recrystallization, and it is proposed that the inhibiting agent is this precipitate, rather than the iron atoms in solid solution. It is important to note that precipitation before cold work is ineffective; the effective precipitate is that formed after cold working and either before or during recrystallization. The location and distribution of the precipitate are critical. Precipitation in such a manner has been found to have profound effects upon kinetics of recrystallization and the microstruc-ture of the recrystallized alloys.2-4 It would be surprising, indeed, if this were accomplished with no change in texture. Because of the relative simplicity of the system, and because of previous experience,4-7 it was decided to determine the effect of precipitation on texture in an alloy of iron and copper. Bush and Lindsay5 found an unspecified change in texture in cold-rolled and annealed low-carbon rimmed steel sheets when the copper content exceeded 0.1 pct. MATERIALS In earlier work, the rate of recrystallization of a low-carbon steel was greatly decreased by 0.80 pct copper, and, after the proper treatment, the recrystallized ferrite grains were greatly elongated.4 Accordingly, it was decided to investigate the effect of precipitation on texture at this level of copper content. The iron and the iron-copper alloy were made from high-quality electrolytic iron and OFHC copper, vacuum-melted in MgO crucibles, cast, hot-rolled to 0.2 in., then machined to 0.150 in. The compositions are given in Table I. The plates were heated to 925°C and brine quenched, twice. This produced a ferrite grain size of ASTM 0 in the iron and ASTM 1 in the Fe-Cu alloy. Disk specimens were cut from the heat-treated plates, repeatedly polished and etched, then used to determine (110) and (200) pole figures by reflection. Despite the complication of large grain size, these pole figures strongly indicated a random texture. PROCEDURES The copper content in solid solution in ferrite before cold rolling and recrystallization, and hence, the amount that could precipitate during the recrys-tallization anneal, was controlled at three levels by heat treatment. The specimens as quenched from 925° C were presumed to have all the copper, 0.80 pct, in solid solution. Other samples of the quenched alloy were aged 5 hr at 700°C to retain about 0.5 pct Cu in solid solution.6 A third set of quenched specimens was reheated to 700°C, then slowly cooled in steps, to reduce the amount of copper in solid solution to a very low level. All specimens were cold-rolled 90 pct, from 0.150 to 0.015 in. thick. The rolling was done in one direction only, i.e., the strip was not reversed between passes, with a jig on the table of the mill to keep the short specimens at 90 deg to the rolls. The rolls were 5 in. in diameter and speed was 35 ft. per min. Machine oil was used as a lubricant. In a supersaturated alloy, the maximum effect of the copper precipitate on microstructure and on recrystallization can be developed by a treatment at 500°C, after cold rolling and before recrystallization.&apos;

Jan 1, 1962

By L. C. Michels, O. G. Ricketts

The disorientations between s?nall grains, whose growth has been arrested at an early stage of recrys-tallization, and the deformed matrix in cold-rolled aluminum single crystals were determined using transmission Kikuchi line and electron diffraction patterns. The orientations of the recrystallized grains were found to be random, and the disorientations of these grains with the matrix weve found to be intermediate to large. This leads to the conclusion that the observed vecrystallization began in small areas of large disorientation present in the cold-worked structure. heavily cold-worked thin sections of aluminunz single crystals and of polycrystalline aluminum and nickel were produced directly by a mechanical technique. The specinlens thus prepared were heated with the electron beam to bring about vecrystallization during observation in the electron microscope. Motion pictures taken du.ring heating and the electvon, microg.raphs taken both before and aftev heating allowed the recrystallization process to be traced to its ovigin. Re cvystallized grains originated in very s,mall regions of the cold-worked structure and developed through rapid migration of high-angle boundaries. The boundaries either were present as such in the matrix or were formed out of dense dislocation networks. SIGNIFICANT advances have been made in recent years in the study of nucleation of recrystallization using the technique of transmission electron microscopy of thin metal foils. Bollman1 in a study of heavily rolled polycrystalline nickel found support for the Cahn-Cottrell2,3 theory of nucleation. According to this theory nuclei form by the initially slow growth of subgrains formed through polygonization. During this initial period of slow growth (the incubation period) the migrating boundary of the subgrain increases its disorientation with the cold-worked matrix and thereby increases its mobility to become a rapidly migrating high-angle boundary. Bailey4,5 investigated the annealing behavior of several metals deformed both in tension and by rolling and concluded that recrystallization took place through the migration of high-angle boundaries. With low deformations these boundaries were present in the metal before deformation. With high deformation it was not possible to tell whether the boundaries were pieces of the original grain boundaries or were produced either during deformation or by polygonization during ameal- ing. Direct observation during heating of metal foils indicated that subgrains form by polygonization and grow at an uneven rate. The grain size obtained decreased with decreasing foil thickness indicating that the foil surface resists boundary motion. Votava,6 in heating stage experiments on rolled copper, observed nuclei to appear suddenly and grow in jumps of differing magnitude. However, he found no special dislocation configurations where the nuclei appeared. Fujita,7 as a result of a study of subgrain growth in heavily worked aluminum, concluded that the boundary of a recrystallized grain initially forms from the boundary of a group of subgrains. This occurred by a process of deposition of vacancies and dislocations in the group boundary as the boundaries within the group disappear. HU8,9 directly observed a similar process in heating stage experiments on 70 pct rolled Si-Fe single crystals. The growth of subgrains appeared to proceed by a coalescence mechanism. The observed fading away of the boundary between two subgrains was explained by the moving out of dislocations from the disappearing boundary into the connecting or intersecting boundaries around the subgrains. The subgrain size and degree of disorientation with the surrounding structure were thus increased. With the increase in disorientation occurred a corresponding increase in boundary mobility, which eventually allowed the boundary to migrate rapidly. This process was observed to occur within "microbands" consisting of parallel narrow segments disoriented by a few degrees present in the as-rolled structure. The conclusion of Rzepski and Montuelle10 that growth is preceded by the coalescence of blocks through disappearance of their common boundaries supports this view. In contrast to Hu&apos;s coalescence model for nucleation were the conclusions of Walter and ~och.""~ Working with the same material as Hu, of the same orientation and rolled to the same reduction, they concluded that nucleation occurred by the Cahn-Cottrell mechanism. They observed, in agreement with Hu, that recrystallization began in the "microband" regions which they referred to as "transition" bands. Bartuska13 studied subgrain growth in heavily rolled nickel using a beam heating method in the electron microscope. He concluded that nuclei for recrystallization form from the largest most perfect subgrains present in the cold-worked structure by rapid intermittent migration of parts of subboundaries. In rare instances he observed subgrain growth by coalescence. EXPERIMENTAL PROCEDURE The materials used in this study were 99.999 pct A1 supplied by A.I.A.G. Metals, Inc., and 99.999 pct Ni supplied by Johnson and Matthey and Co., Ltd. The Hitachi HU-11 electron microscope, with uniaxial

Jan 1, 1968

By E. R. Dean

A wire-feed mechanism has been employed to fabricute metal alloy film resistors to various sheet resistivities on oxidized silicon substrates. The effect of several thousand hours storage in air at elevated temperatures on the resistance and temperature coefficient of resistance is presented. Unprotected films of- sheet resistivity between 100 and 500 ohms per sq fabricated at 300°C substrate tenzperatures were unstable when stored at 150°Cfor extended periods. The higher sheet resistivity films exhibited the greatest instability; however, even the 100 ohms per sq films drijted excessively for device application. When stored at still higher temperatures, the normalized resistance increases to maximum, then decreases until a minimum value is obtained, then finally increases in resistance until open. The use of a protective overcoating of SiO has had considerable benejicial effects on the film stability, so that 250 ohms per sq films deposited on 300°C substrates are now stable after 1000 hr storage at 250°C and possess a temperature coefficient of resistance less than 200 ppm per C. The use of a low substrate temperature during depositon (100°C) enables the preparation oj resistors with very low tenperature coefficients of resistance (10 to 20 ppnl per &apos;C). However, these films are less stable than their higher substrate temperatuve counterparts. During extended storage, the resistance of the protected films always decreases with lower substrate films exhibiting larger normalized resistance decreases. This resistance decrease is accompanied by a linearly related increase in the temperature coefficient of resistance. The electrical behavior of these films may be explained by postulating That the structure of the films is in the transition region between thick continuous films and ultrathin island structure films so that the conductivity is the restlt of both electron scattering and tunneling-activated charge carrier creation between neighboring grains. The annealing behavior when thermally aged is the result of defect anneal, grain growth, and selective oxidation. TheRE is a considerable interest in thin metallic films for use as resistor elements in microelectronic circuits.1,2 These resistor films must be stable when exposed to elevated-temperature storage or operating ambient, possess low temperature coefficients of resistance (hereafter referred to as T.C.R.), and be of a high enough sheet resistivity to be useful. The resistivity and T.C.R. of a thin metallic film are determined by the structure and thickness of the film. Hence the stability of the film when exposed to stress would depend on the stability of the structure. The resistivity of a thin film is the result of the electrons&apos; interaction with the lattice vibrations (electron-phonon scattering), scattering of electrons due to impurities, defects, and grain boundaries and specula reflection from the film surfaces.3&apos; All of these effects serve to reduce the electron mean free path and result in a resistivity higher than the ideal bulk. In ultrathin films possessing essentially an island structure consisting of a planar array of aggregates separated by a few to a few tens of angstroms, conduction is dominated by a combination of tunneling and activated charge carrier reation. Thin films would be characterized by relatively low resistivity, positive T.C.R., and reasonable structure stability whereas ultrathin films possess unstable structures, negative T.C.R., and exhibit high resistivity. Feldman6 has investigated films intermediate between the two regions and postulates that the resistivity of a film in the transition region is composed of two linearly additive parts: that due to the resistance of the grains and that due to the gaps. The grain resistivity would represent the scattering of the conduction electrons by defects, surfaces, and phonon interaction while the gap term represents the tunneling contribution. For gold and platinum films, he found as the film structure becomes progressively less continuous, corresponding to thinner films, the absolute value of the T.C.R. becomes progressively smaller until negative values are observed with respective zero T.C.R.&apos;s at about 35 and 150 ohm per sq, respectively. Recently the author7 investigated the behavior of a multicomponent alloy of nickel, chromium, and iron and found the T.C.R.&apos;s of these films decrease with decreasing substrate temperature during deposition. The lower T.C.A. associated with lower substrate temperatures was attributed to an increased contribution to the total resistivity from the negatively temperature-dependent gap tunneling since it is well-known that thin films deposited on low-temperature substrates consist of a higher density of smaller grains than the same sheet resistivity films deposited on higher-temperature substrates. Under conditions of thermal anneal, the author found that, when various sheet resistivity films fabricated at 300°C substrate temperature are exposed to extended storage in air at 300°C, the thinner films increase in resistance until open, while the thicker films

Jan 1, 1967

By K. Brack, G. H. Schwuttke

Annealing properties of subszerface amorphous lavers produced through high-energy ion implantation in silicon are studied. The buried layers are produced through the implantation of ions (nitrogen), ranging in energy from 1.5 to 2 mev. X-ray interference patterns, transmission electron microscopy, and resistivity profiling are used to study the annealing characteristics of the ion damage. The annealing experiments indicate a low temperature (below 700°C) and a high temperature (above 700°C) region. Significant changes occur in the amorphous layer during the high-temperature anneal. Such changes are corre-lated with the re crystallization of the amorphous silicon and the formation of subsurface (buried) silicon-nitride films. TODAY&apos;S main problems in the field of ion implantation are related to the accurate determination and prediction of 1) the distribution profiles of implanted ions, 2) the lattice sites occupied by the implanted ions, 3) the lattice damage produced through ion implantation, and 4) the annealing characteristics of damage centers in the lattice. This paper reports investigations concerned with the problems listed under 3) and 4). EXPERIMENTAL Our investigations cover the energy range of incident ions from 100 to 300 mev and from 1 to 2.5 mev. The emphasis of this study is on the energy range from 1.5 to 2 mev. The experiments are conducted with single charged nitrogen ions. To implant the ions a van de Graaff generator is used as described by Roosild et al.1 Accordingly, a gas containing the desired ion specie is passed through a thermome-chanical leak into a radio frequency activated source. The positive ions are driven into the van de Graaff with the help of a variable voltage probe. Emerging from the accelerator the ions drift into a magnetic analyzing system and here the desired ion specie is bent 90 deg into the exit port. The ion beam leaving the analyzer is defocused and drifts down a 4-ft long tube to hit the silicon target. At this position the 20 pamp ion beam has a circular cross-section of 2.1 cm. N2 is used as a source gas for nitrogen ions. The implantation target is silicon with zero dislocation density, 2 ohm-cm resistivity, (111) orientation, mechanically-chemically polished, and 1 mm thick. The target is mounted on a water-cooled heat sink and kept at room temperature. A fluence of 1015 to 1016 ions per sq cm is used. RESULTS 1) Silicon Perfection after Bombardment. High-energy ion bombardment of silicon has some striking effects on lattice perfection. Some results were reported in detail previously at the Santa Fe conference2 and are here briefly summarized for the benefit of the experiments described in the following. 1.1) Identification of Surface Films on Silicon. After bombardment all samples are found to be coated with surface films. The films on the silicon surface vary in thickness and color; they can be transparent, slightly brown, or opaque. The films are thicker and darker in the high-intensity area of the beam and they delineate the bombarded surface area of the crystal. The films produce electron diffraction patterns characteristic of carbon and of SiO2. Carbon is predominant. The presence of carbon in these films was confirmed by use of the electron microprobe. Formation of the films occurs independently of the ions used and is attributed to a contaminated vacuum of the high-voltage machine. The carbon is most likely the product of the pump oil which is cracked and polymerized under ion impact. The films stick tenaciously to the silicon surface and burn off in a low-temperature Bunsen flame. 1.2) Mechanical Perfection of the Silicon Surface. The mechanical perfection of the bombarded silicon surface was investigated through optical microscopy, electron microscopy in which the replica technique is used, and optical interferometry. No mechanical damage of the surface was visible after bombardment. However, if a bombarded sample is soaked for several minutes in hydrofluoric acid (HF), gas bubbles may develop in certain spots of the silicon surface. It is also noted that in these areas the surface film starts to peel off. Relatively large patches of film come off if the sample is soaked in HF during ultrasonic agitation. After HF treatment, pits may be present on the silicon surface. The pit dimensions are estimated to be as large as 50 µ. The pits appear in the region of most intense irradiation. 1.3) Lattice Perfection After Bombardment. No lattice damage is found on the silicon surface. Electron transmission micrographs and selected area diffraction patterns of the surface show no difference before and after bombardment. Measured approximately 2 µm down from the surface, the silicon lattice throughout this depth is of good perfection. Well-defined Laue spots and Kikuchi lines are obtained from the surface as well as from the indicated area below the surface. However, some radiation damage is dispersed in this top layer. A sharp boundary line separates this surface layer from a highly damaged layer which extends further downward into the silicon. Typical of this

Jan 1, 1970

By James H. Courtright, Kenyon Richard

TOQUEPALA is a porphyry copper deposit in which mineralization is localized by a large breccia pipe formed in close genetic relation to intrusive rocks. The deposit is in southern Peru, 55 airline miles north of the small city of Tacna and the same distance inland from the port of 110. Quellaveco and Cuajone, geologically similar deposits, lie 12 and 19 miles north of Toquepala. Chuquicamata is 400 miles to the south. The deposit is high on the southwestern slope about 20 miles from the crest of the Cordillera Occidental of the Andes Chain. It lies in a mountainous desert where the steep southwesterly slope of the Andes is dissected by a succession of rapidly downcutting, deep canyons. Local topography is moderately rugged with a dendritic drainage pattern and an elevation of 8000 to 14,000 ft. Volcanic peaks along the crest of the Cordillera rise over 19,000 ft. Local precipitation, including a little snow, amounts to about 10 in. during January and February, but general runoff in the region is slight. Throughout southern Peru the springs and streams are widely separated. Crude canals irrigate small farms on terraced slopes along the streams and provide sparse subsistence to the semi-nomadic inhabitants. During the past decade, engineering and geological explorations of the region, as well as the mineral deposits themselves, have required construction of a network of several hundred miles of roads. Before this, roads extended only a few miles inland. Many areas still can be reached only by trail. Toquepala was briefly described in 19th century geographical literature as a copper deposit, and it received desultory attention from Chilean prospectors early in the present century. It was first recognized as a mineralized zone of possible real importance by geologist O.C. Schmedeman during an exploration trip for Cerro de Paso Copper Corp. in 1937. The discovery was late as compared to earlier recognition of Chuquicamata, Potrerillos, and Braden of Chile and Cerro Verde of southern Peru. This was due partly to the region&apos;s difficult accessibility but principally to the obscure character of the outcrop evidence of copper. From 1938 until 1942 Cerro de Pasco Copper Corp. partially explored the deposit by adits and diamond drillholes. This campaign was supplied by a 60-mule pack train continuously shuttling over a 30-mile trail. Northern Peru Mining & Smelting Co., a wholly owned subsidiary of American Smelting & Refining Co., undertook regional engineering stud- ies in 1945 and drill exploration in 1949. According to published data1 the deposit contains 400 million tons of open pit ore averaging a little over 1 pct Cu. It is currently undergoing large-scale development by Southern Peru Copper Corp., which is owned by American Smelting & Refining, Phelps Dodge, Cerro de Pasco, and Newmont Mining. Summary of Geology: The deposit is situated in a terrane composed of Mesozoic(?) and Tertiary volcanic rocks intruded by dioritic apophyses of the Andean Batholith. These formations are exposed in a northwesterly trending belt about 15 miles wide. Along the northeast they are unconformably overlain by Plio-Pleistocene pyroclastic rocks, which occupy much of the crest of the Andes, and along the southwest they are covered by the Moquegua formation of Pliocene(?) age. The mineralized area, oblong in shape and about 2 miles long, has been a locus of intense igneous activity. Several small intrusive bodies having irregular forms occur within and adjacent to a centrally located, large breccia pipe. The mushroom-shaped orebody consists of a flat-lying enriched zone of predominant chalcocite with a stem-like extension of hypogene chalcopyrite ore in depth within and around the pipe. This breccia pipe is relatively large and has been formed by repeated episodes of brecciation. Small satellitic pipes occur at random within a 2-mile radius of this central pipe. These too were individual sourceways of mineralization, although not always of ore grade. Within and around the zone of breccia pipes and mineralization there are a few faults and veins, but these are discontinuous random structures of minor significance. There are no regional or local systems of faults or other planar structures recognized which could account either for the mechanical development of the breccia pipes or for their localization as a group or as individuals. Hydrothermal alteration is pervasive in the zone of mineralization. Clay minerals appear to be abundant in places, but their percentages are undetermined. Quartz and sericite are the principal alteration products, and in many instances original rock textures are obliterated. The principal sulfides, hypogene pyrite and chalcopyrite and supergene chalcocite, occur mainly as vug fillings in the breccia and as small discrete grains scattered through all the altered rocks. Sulfide veinlets are relatively scarce. Sulfides are more abundant and alteration is more intense in certain rock units, such as the diorite and most of the breccias. Although the Toquepala mineral deposit is similar in most respects to the porphyry copper deposits of southwestern U. S., it most closely resembles the Braden deposit of Chile, as described by Lindgren

Jan 1, 1959

By R. Ebeling, M. F. Ashby

The paper is in two parts. The first develops the formulae and method needed to calculate the size, nu)nber, spacing, and volume fraction of hard or inert particles in the interior of a specimen from measurernents made on an extraction replica; throughout, the spread of- particle sizes in a single specimen is taken into account, and the extraction efficiency of the replica is considered. A method for testing whether the particles are randotnly distributed in space is described, and the spacing of randomly distributed particles is discussed. The second Part describes the application of the method to copper containing spherical particles of silica; the description includes an estimate of the errors involved, and a comparison of this with alternative methods. All parameters except uolume fraction can be determined satisfactorily, and the method offers advantages over other electron-,rzicroscopic techniques when dealing with hard particles in a softer matrix, or inert particles in a chem-ically less inert matrix. 1 HE experimental measurement of sizes and numbers of discrete particles of a second phase embedded in a matrix material from opaque, plane sections has been considered by many authors, notably by Fullman.&apos; It is assumed that, when the specimen is sectioned, the particles are also sectioned, so that the diameters of the circles of intersection of particles with the plane of the cross section ("surface diameters") are measured. It is difficult to derive the distribution of true particle diameters ("volume diameters") from measurements of the distribution of "surface diameters" on a plane section. The measurement of size and numbers of small second-phase particles by transmission electron microscopy through thin foils of matrix containing the particles has been described by Cahn and ~uttin~&apos; and by ~illiard.~ Because of overlap and the fact that particles which intersect the foil surface may be sectioned whereas those in the bulk of the foil are not, it is again difficult to measure the true distribution of particle sizes. The use of replicas for this sort of measurement has not been analyzed in detail. A replica made from an opaque cross section, thus replicating the circles of intersection of particles with the plane of section, is analyzed by Fullman&apos;s method and suffers from the same difficulties. Under certain circumstances, however, extraction replicas can be used, and offer a special advantage in that they make the determination of the true distribution of particle diameters particularly simple. It is the purpose of this paper to describe this extraction replica method. Consider the example of an alloy consisting of SiOz particles, whose mean diameter is about 1000A, dispersed in a copper matrix. When the alloy is sectioned and mechanically polished, the particles are not sectioned, because they are much harder than the matrix, and so will be either left in the surface or pulled out whole, and because the particle size is smaller than the abrasive size of even the finest abrasive. The particles are chemically inert in solutions which chemically etch or electropolish the matrix, so that chemical treatments do not section the particles either. (Experimental measurements of particle shape in this alloy confirm this—truncated particles are never seen.) Small, hard or inert, particles are therefore not sectioned when the specimen is sectioned and polished. An extraction replica taken from this polished surface shows the true "volume" diameter of each extracted particle; the distribution of diameters seen on the replica is the surface distribution of true "volume" diameters, and the volume distribution of true "volume" diameters is easily derived from it. The extraction replica method is not subject to the difficulties associated with sectioning of particles, and avoids the problem of overlap from which the transmission method suffers. The method is well-suited for studying hard particles in a soft matrix. Examples are provided by internally oxidized alloys, oxide-dispersion strengthened metals like T.D. Nickel, and carbides, nitrides, and strong intermetallic compounds in metals. It should be possible to adapt polishing techniques in alloys containing softer but inert particles to avoid sectioning the particles (say by electropolishing or etching). The method is tested and illustrated in Sections 8 and 9 by applying it to a dispersion of spherical SiO2 particles in a copper matrix. I) SYMBOLS Symbols are defined where they first appear in the text. Volume fractions of particles are denoted by the symbol f, particle diameters by x, center-to-center particle spacings by D, and numbers of particles by N. Mean particle diameters and standard deviations about these means are denoted by x and a. Both arithmetic (subscript A) and geometric (subscript G) means are employed. In this paper we measure the distribution of true diameters of particles which touch or pass through a random-plane section of the specimen. We may want to know the distribution of particle diameters in the

Jan 1, 1967

By W. E. Stiles

A method is presented for predicting the performance of water flooding operations in depleted, or nearly depleted, petroleum reservoirs. The method makes use of permeability variations and the vertical distribution of productive capacity. From these two parameters can be calculated the produced water cut versus the oil recovery. Derivations of the mathematical analogy is shown and sample calculations and curves of prediction are presented. Comparison is made of the predicted and actual performance of a typical 5-spot in an Illinois water flood. INTRODUCTION The use of water as a flooding medium in both depleted and "flush" oil reservoirs is gaining greater recognition and acceptance. Many of the shallower fields, depleted by primary production, have been and are being subjected to water injection in order to obtain some part of the large volume of oil remaining after primary production. Some of the earlier water flood installation proved highly discouraging and the value of water flooding was often questioned. Many of these earlier floods were haphazardly selected and developed as little was known of the physical characteristics and contents of the producing formations. The prior evaluation of the flood performance was impossible. During the past decade the development of the required reservoir engineering tools-—core analysis, reservoir fluid analysis, electric logs, fluid flow formulae, etc.—has allowed the engineer to construct and apply the methods which are presently being used to evaluate the economic and mechanical susceptibility of a reservoir to flooding. This discussion will present a method for taking into account the effect of permeability variations in predicting the performance of water floods in depleted reservoirs. PERMEABILITY AND CAPACITY DISTRIBUTION It is generally agreed by most investigators that in a single phase system fluid will flow in a porous and permeable medium in proportion to the permeability of the medium. Producing formations are usually highly irregular in permeability, both vertically and horizontally. However, zones of higher or lower permeability are often found to exhibit lateral continuity. Thus, while structurally comparable stringers in adjacent wells may differ several fold in permeability values, they usually bear resemblance as being part of a general continuous higher or lower permeability section. It is generally agreed that where such stratification of permeability exists, injected water sweeps first the zones of higher permeability, and it is in these zones that "break-through" first occurs in the producing well. It is a basic assurnption of the presently described method that penetration of a water front follows the individual permeability variations as if such variations were continuous from input to producing well. This is admittedly not rigorously true, but can be justified as making possible a simplifying mathematical approach to an otherwise extremely complicated three dimensional flow problem. As a basis for study of the lateral flow of fluids in formations of irregular permeability, the irregularities may be conveniently represented by a permeability distribution curve and a capacity distribution curve. In obtaining these curves, the permeability values, regardless of their structural position in the formation, are rearranged in order of decreasing permeability. If these permeability values so arranged are plotted against the cumulative thickness, a permeability distribu- tion curve is obtained. This curve may then be likened to a "smoothed" permeability profile of the formation. In making comparison between different distribution curves it is convenient to state the permeabilities in terms of the ratio of the actual permeability values to the average permeability of the formation. These ratios termed "dimensionless permeabilities", are used in this paper rather than the permeabilities in terms of millidarcys. The capacity distribution curve is a lot of the cumulative capacity (starting with the highest permeabilities) versus the cumulative thickness. The capacity and thickness are given as fractions of the total capacity and thickness. Mathematically, the capacity distribution is the intergration of the permeability distribution curve. In practice it is convenient to first obtain the capacity distribution curve and derive from it a smoothed dimen-sionaless permeability curve. The method of obtaining the capacity distribution curve is illustrated in the successive column of Table 1, in which capacity and thickness are derived as fractions of their respective totals. If only a small number of permeability values are available, it is generally desirable to smooth the resultant curve. This has been done to give the capacity distribution curve shown in Figure 1. The differentiation of the capacity distribution curve to obtain the permeability distribution curve is shown in Table 2. Here, the capacity values are read from the smoothed curve at intervals of cumulative thickness, and the increments of capacity are divided by the increments of thickness to obtain the dimensionless permeability, K. Due to this stepwise procedure of calculation these premeability values must be plotted at the midpoints of the successive increments of thickness. The curve so obtained from these data is shown in Figure 1. The total area under the K&apos; curve is equal to unity.

Jan 1, 1949

By W. E. Stiles

A method is presented for predicting the performance of water flooding operations in depleted, or nearly depleted, petroleum reservoirs. The method makes use of permeability variations and the vertical distribution of productive capacity. From these two parameters can be calculated the produced water cut versus the oil recovery. Derivations of the mathematical analogy is shown and sample calculations and curves of prediction are presented. Comparison is made of the predicted and actual performance of a typical 5-spot in an Illinois water flood. INTRODUCTION The use of water as a flooding medium in both depleted and "flush" oil reservoirs is gaining greater recognition and acceptance. Many of the shallower fields, depleted by primary production, have been and are being subjected to water injection in order to obtain some part of the large volume of oil remaining after primary production. Some of the earlier water flood installation proved highly discouraging and the value of water flooding was often questioned. Many of these earlier floods were haphazardly selected and developed as little was known of the physical characteristics and contents of the producing formations. The prior evaluation of the flood performance was impossible. During the past decade the development of the required reservoir engineering tools-—core analysis, reservoir fluid analysis, electric logs, fluid flow formulae, etc.—has allowed the engineer to construct and apply the methods which are presently being used to evaluate the economic and mechanical susceptibility of a reservoir to flooding. This discussion will present a method for taking into account the effect of permeability variations in predicting the performance of water floods in depleted reservoirs. PERMEABILITY AND CAPACITY DISTRIBUTION It is generally agreed by most investigators that in a single phase system fluid will flow in a porous and permeable medium in proportion to the permeability of the medium. Producing formations are usually highly irregular in permeability, both vertically and horizontally. However, zones of higher or lower permeability are often found to exhibit lateral continuity. Thus, while structurally comparable stringers in adjacent wells may differ several fold in permeability values, they usually bear resemblance as being part of a general continuous higher or lower permeability section. It is generally agreed that where such stratification of permeability exists, injected water sweeps first the zones of higher permeability, and it is in these zones that "break-through" first occurs in the producing well. It is a basic assurnption of the presently described method that penetration of a water front follows the individual permeability variations as if such variations were continuous from input to producing well. This is admittedly not rigorously true, but can be justified as making possible a simplifying mathematical approach to an otherwise extremely complicated three dimensional flow problem. As a basis for study of the lateral flow of fluids in formations of irregular permeability, the irregularities may be conveniently represented by a permeability distribution curve and a capacity distribution curve. In obtaining these curves, the permeability values, regardless of their structural position in the formation, are rearranged in order of decreasing permeability. If these permeability values so arranged are plotted against the cumulative thickness, a permeability distribu- tion curve is obtained. This curve may then be likened to a "smoothed" permeability profile of the formation. In making comparison between different distribution curves it is convenient to state the permeabilities in terms of the ratio of the actual permeability values to the average permeability of the formation. These ratios termed "dimensionless permeabilities", are used in this paper rather than the permeabilities in terms of millidarcys. The capacity distribution curve is a lot of the cumulative capacity (starting with the highest permeabilities) versus the cumulative thickness. The capacity and thickness are given as fractions of the total capacity and thickness. Mathematically, the capacity distribution is the intergration of the permeability distribution curve. In practice it is convenient to first obtain the capacity distribution curve and derive from it a smoothed dimen-sionaless permeability curve. The method of obtaining the capacity distribution curve is illustrated in the successive column of Table 1, in which capacity and thickness are derived as fractions of their respective totals. If only a small number of permeability values are available, it is generally desirable to smooth the resultant curve. This has been done to give the capacity distribution curve shown in Figure 1. The differentiation of the capacity distribution curve to obtain the permeability distribution curve is shown in Table 2. Here, the capacity values are read from the smoothed curve at intervals of cumulative thickness, and the increments of capacity are divided by the increments of thickness to obtain the dimensionless permeability, K. Due to this stepwise procedure of calculation these premeability values must be plotted at the midpoints of the successive increments of thickness. The curve so obtained from these data is shown in Figure 1. The total area under the K&apos; curve is equal to unity.

Jan 1, 1949

By J. W. Byrkit

Design features and operating methods at the new Ajo smelter are described in detail. Successful operation of a novel method of handling and charging wet concentrates to a deep bath type reverberatory furnace contribute to the daily production of 200 tons of anodes with good results from the standpoint of both metallurgy and economy. THE New Cornelia Branch of the Phelps Dodge Corp. is located at Ajo, Ariz. Large scale mining operations were started at Ajo in 1917, when a 5000 ton leaching plant was put in service to treat the copper carbonate ore that overlaid the sulphide ore-body. In 1924 a 5000 ton flotation plant was built to treat the sulphide ore and the leaching operation was abandoned a few years later. Changes in practice and additions to plant facilities have resulted in a gradual increase in the milling rate which now approaches 30,000 tons daily. Prior to the erection of the Ajo smelter, flotation concentrate was shipped by rail to Douglas, Ariz, where it was treated in the Phelps Dodge smelter. Construction of a one reverberatory furnace smelter to produce an average of about 200 tons per day of copper anodes at Ajo was started early in 1949. Initial heating of the reverberatory furnace was begun on June 21, 1950. Charging the furnace was started on july 8, and the first anodes were cast on July 14. Neither the anode furnace nor the reverberatory furnace has cooled since the initial firing. Fig. 1 is a photograph of the new smelter. Built primarily to eliminate the &apos;300 mile haul of concentrate to the Douglas plant, the new smelter was designed to treat the Ajo concentrate, with a minimum of capital investment. A novel, and relatively simple, method of concentrate handling and furnace charging made unnecessary the installation and operation of expensive storage facilities and reclaiming equipment, and obviated the necessity of holding in storage large quantities of copper-bearing materials. (Patents are pending in the United States and abroad on the Ajo process.) In the selection and arrangement of equipment, consideration was given to the full utilization of all smelting facilities, reducing the amount of idle equipment to a minimum. The layout of the smelter is shown in Fig. 2. The important problem of internal transportation was minimized by arranging the reverberatory furnace, converters, and anode furnace in a compact group and providing storage bins inside the smelter building, within reach of overhead cranes, for fluxing materials and other supplies needed in the daily operation. This can be seen in the sectional view of the smelter given in Fig. 3. Incorporated in the design were many innovations in smelting equipment intended to facilitate operations and effect economies in maintenance expense and manpower requirements. Table I gives operating data of the entire plant. Metallurgy The metallurgical practice at Ajo is based essentially upon the use of a single, deep bath reverberatory furnace for smelting wet concentrate, without concentrate storage facilities. The large reservoir of slag and matte maintained in the furnace serves to equalize, to some extent, the day to day variations in the nature and grade of concentrate and permits

Jan 1, 1954