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BOSTON SECTION ALFRED C. LANE, Chairman, GEORGE A. PACKARD, Vice-chairman, E. E. BUGBEE, Sec'y-Treas., Mass. Inst. of Technology, Cambridge, Mass. R. L. AGASSIZ, FRED W. DENTON. Forty-sixth Meeting The forty-sixth meeting of the Boston section was held at the University Club on Monday evening, Jan. 28, 1918. Thirty-six members and guests were present, which is a record for this Section. After a good dinner, President Philip N. Moore, of the Institute, was introduced, and spoke of the many problems of the Institute. The topics covered were the progress and growth of the Institute during the administrative year, the efforts to stimulate the local sections and to bring them more in touch with the policies and problems of the parent body and, last but not least, of the constructive results and patriotic work of the Institute in the national service of the war. The other guest of the Section, Mr. Thomas Nelson Perkins, Counsel for the Priority Board, spoke most entertainingly of recent experiences in Russia, England and France. After considerable discussion during which Mr. Perkins most obligingly answered innumerable questions, the meeting adjourned at about 10 p. m.

Jan 7, 1918

By R. J. Hafsten, E. S. Davenport, R. A. Grange

THE influence of austenite grain size upon the hardenability of steel is now fairly well understood; for a given austenite, increasing the grain size increases the depth of hardening, and, since the latter is related to the fundamental transformation behavior of austenite, it follows that austenite decomposition must, in general, be retarded by increasing grain size. To date, very little quantitative information bearing directly on the effect of austenite grain size upon the progress and result of isothermal transformation has been published. Davenport and Bain,1 and Zimmerman, Aborn and Bain2 have shown that the transformation is definitely retarded in hypereutectoid plain carbon steels at temperatures near and above the upper region of rapid transformation (Ar') when the austenite grains are increased in size. More recently, Griffiths, Pfeil and Allen3 investigated the effect of a series of different heating (austenitizing) temperatures upon the progress of isothermal transformation at 1245°F. (675°C.) and 745°F. (395°C.) in a medium carbon, low-alloy steel; while they reported no measurements of the austenite grain size, their results suggest that increase in grain size retards austenite transformation at 1245°F. (675°C.), but has no appreciable effect at 745°F (395°C.). The present paper deals with the results of a study of the isothermal transformation of austenite in a single commercial alloy steel in which a fine-grained or a coarse-grained state was induced by appropriate selection of austenitizing temperature. The data are presented in the form of isothermal transformation diagrams, which clearly show that coarsening the grains lessens the rate of transformation of austenite to proeutectoid ferrite and pearlite-that is, at temperature levels above about 1000°F. (540°C.)-but is without significant effect at lower temperatures where the transformation product is bainite or martensite. It is to be presumed that the foregoing statement would be valid, at least in principle, for many other grades of steel; it confirms again that the mechanism of transformation to a pearlitic structure is fundamentally different from that to an acicular structure. A number of related observations were made, which seem to throw some light on the transformation process as it takes place in alloy steels, particularly at temperature levels in the vicinity of 1000°F. (540°C.), where the transition from lamellar to acicular structures occurs. These data are discussed briefly in the belief that they may lead to a clearer understanding of this aspect of the subject. MATERIAL AND EXPERIMENTAL PROCEDURE The material investigated was a commercial grade of S.A.E. 4140 steel, the composition of which was as follows: 0.37 per cent C, 0.77 per cent Mn, 0.98 per cent Cr and 0.21 per cent M0. It was hot-rolled to a 1 1/8-in. round bar and normalized by air-cooling from 1800°F. (980°C.).

Jan 1, 1941

By James K. Gluck, William E. Galloway, Gary E. Smith, John P. Morton, Christopher D. Henry

INTRODUCTION Surface mining and in situ leaching of uranium have the potential to alter ground-water quality around mines and leach sites. Of particular concern is the fate of uranium and its associated trace elements: molybdenum, arsenic, and selenium. We wish to under- stand the natural processes that control trace element concentrations in ground water and how these processes will influence dispersion of the elements from a mineralized zone, both naturally and during and after mining or restoration. For example, it is commonly recognized that the trace elements are soluble in oxidizing ground water but are insoluble, and can be precipitated, in reducing ground water. Thus oxidizing, metal-bearing water leaving a deposit could re- enter reduced ground, causing the water to be re- reduced and the trace elements to be, reprecipitated. In a sense, this is recreating the original mineralization process. To accomplish the above goals, we have (1) examined the theoretical controls of concentrations based on the available geochemical and thermodynamic data, (2) determined the major ion composition and oxidation-reduction status of Oakville waters because of the influence of these factors on trace element solubility, and (3) determined trace element concentrations and distribution in Oakville ground water. The last approach is used to evaluate how well actual behavior follows predicted behavior. This report focuses on two elements, uranium and molybdenum, because they exemplify the results obtained. The report also is restricted to a regional study of Oakville ground water. Results of more de- tailed study in and around major uranium districts in the Oakville and much of the raw data that support the conclusions in this report are presented in Galloway, Henry and Smith (1980). This report is part of that larger study, which concerned the depositional systems, hydrology, and geochemistry of the Oakville. The U.S. Environmental Protection Agency funded the study, under grant numbers R-805357-01 and R-805357-02. Theoretical controls were determined by reviewing the available literature on aqueous chemistry and behavior of uranium and molybdenum. To aid in under- standing water chemistry, Oakville water analyses were run through a modified version of the computer model WATEQF (Plumer, Jones, and Truesdell, 1976). WATEQF calculates speciation of dissolved ions and determines saturation with respect to a variety of minerals. In the discussion below, ion activity products (IAP) are compared with the equilibrium constant (KT) for various reactions and mineral products. Values of log IAP/KT near zero indicate that the water is in equilibrium with a mineral. Values less than -1 indicate considerable undersaturation and values greater than +1 indicate oversaturation. Galloway, Henry, and Smith (1980) give a more complete discussion of the application of this approach to Oakville water chemistry. Eh-pH diagrams have been constructed or adapted from the literature to predict what form -- dissolved ion or stable mineral species -- uranium and molybdenum assume under various conditions. Construction of the diagrams has followed procedures described by Garrels and Christ (1965). This approach is particularly appropriate because the solubility of the elements is Eh-dependent, and Eh varies greatly within the Oakville aquifer. A number of assumptions or approximations are inherent in the use of Eh-pH diagrams and chemical models such as WATEQF and in the interpretation of water chemistry in general. Both Eh-pH diagrams and chemical modeling rely entirely upon available thermo- dynamic data, including free energies of formation and dissociation constants for various reactions. These values are known to varying degrees of accuracy. Most major ions and minerals are relatively well control- led; however, data for trace metals are much poorer. Thermodynamic data are not available for some minerals, and for other minerals, two or more divergent values exist. By necessity, we have relied on the judgment of others to evaluate thermodynamic data. Calculations by WATEQF and constructions of Eh-pH diagrams are based on an assumption of equi1ibrium. Equilibrium may not be comnon in low-temperature aqueous environments; at best, ground-water composi tion may be in a state of dynamic equilibrium, continuously changing due to changes in environmental conditions. Eh-pH diagrams show what phases are stable at equilibrium under given conditions; they do not prove that the phases actually exist. Many minerals persist or form metastably under conditions outside their equilibrium stability field. The kinetics of reactions, which cannot be evaluated here, are important in determining what phases occur. Kinetics may be less of a problem for ground water that travels and evolves slowly through a semihomogeneous matrix than for many other natural systems. Eh-pH diagrams show equilibrium fields only of phases included. They do not indicate anything about stability relative to phases not included in the diagram. WATEQF, obviously, cannot calculate the degree of saturation of a mineral not included in the program or for which the appropriate ions were not analyzed. Thus, a mineral that was not considered may be the most stable phase under a given set of conditions and may control the solubility of a trace element. Also, this study is limited exclusively to in- organic compounds. Organic material is known to be an

Jan 1, 1980

By G. M. Yocom

Blowing acid Bessemer converters with oxygen-steam produces steel of below 0.002 pct N2 content. This method of blowing, combined with a dephosphorizing treatment in the steel ladle, results in low-carbon steels of low nitrogen and low phosphorous (under 0.035 pet) contents, which has physical properties equivalent to open-hearth steels of similar analysis. Using a 50-50 mixture of oxygen and steam, the refinitzg rate is increased 25 pct over blowing with natural air, and scrap charge increased from 3 to 10 pet. Bottom life is normal with proper tuyere area and arrangements, fumes are decreased, yields increased, and hydrogen content is normal. THE acid Bessemer plant at the South Works of Wheeling Steel Corp., consists of two 15-ton bottom blown converters with a monthly capacity of 57,000 N.T. The product of the shop is skelp billets for continuous welded pipe and slabs for ordinary drawing and forming quality sheets. Approximately 50 pct of ingot production is regular Bessemer steel of natural Phos content and the remainder is a dephosphorized grade of steel made by a special treatment of the blown metal as it is poured into the steel ladle. The low Phos grade of steel has certain advantages over the higher Phos grade but since both grades were produced by blowing natural air, the N2 content was in the range of 0.015 pct which limited its application. In 1954 it was decided to explore the possibilities of blowing with a steam-oxygen mixture for the production of steel of both low N2 and low Phos contents. The necessary equipment was installed to operate one converter in this manner and early in 1955 an experimental run of 160 heats was made by blowing with a steam-oxygen blast and excluding natural air entirely. During this period the proper operating techniques were established, such as blast pressures, steam-oxygen mixtures, valves and instrumental control equipment, tuyere arrangement in the bottoms, blowing times and production rates, and a thorough study made of the final steel quality. Also during this experimental period the dephosphorizing practice was improved by the use of a tap hole below the lip of the vessel. This provided a clean separation of the acid converter slag and blown metal which made the dephosphorizing treatment more effective. The results of this experimental run dictated further development of this practice and a second run of 720 heats was made in 1957. The quality features and conversion cost results were in line with expectations and accordingly a 400-ton per day oxygen plant is now being installed. The plant is scheduled for completion in September of this year. This will provide sufficient oxygen to operate both vessels on steam-oxygen blast and delete natural air blowing entirely. The steel will then be below 0.002 pct N2 bar content and the dephosphorized grades will be between 0.015 and 0.040 pct Phos. STEAM-OXYGEN BLOWING The steam for the process is fed to the plant at 220 psig pressure through a 6-in. line. The high-purity oxygen is compressed to 200 psig and conducted through an 8-in. line. The oxygen from the main line is valved down to 100 psig and passed through a steam heated heat exchanger. The heat exchanger is regulated to supply oxygen at 300°F to the steam-oxygen mixing station. It is essential that the incoming oxygen be held at this temperature to avoid condensation of the steam with resulting excessive erosion of the clay tuyeres in the vessel bottom. Oxygen is admitted to the mixing chamber by a 6-in. hydraulically operated valve driven by the ratio control regulator on impulse from the flow of steam. Steam is admitted to the steam-oxygen mixture station through a 2 1/2-in. hydraulically driven valve. The ratio control regulator acts to increase or decrease oxygen input as the steam flow increases or decreases with changing positions of the Blower's control lever. The important point to note here is that steam flow always precedes the oxygen flow as a safety measure. The control valves have sufficient capacity to afford protection should blow pipe trouble develop. A 50-50 mixture for these 15-ton heats demands an oxygen flow of 3800 standard cu ft per min along with 317 lb of steam. The Blower's stations is provided with an indicating blast pressure gage, and indicating steam and oxygen flow meters. Signal and warning lights indicate the valve positions and line pressures. A control room at the real of the Blower's pulpit room houses the ratio control and pressure regulators, as well as the various meter bodies. The hand actuated wheels used to change the conditions are mounted on a panel on the front of the meter control house. The recording steam and oxygen meters used for totalizing and accounting purposes are also mounted on this panel.

Jan 1, 1962

By J. Howard Boyce

INTRODUCTION The Tsumeb Mine is situated in the northern part of South West Africa, approximately 335 air miles north-east of the port of Walvis Bay on the Atlantic Coast, and 240 air miles north of Windhoek, the capital of South west Africa. The mine came into operation in the early 1900's and has operated continuously since then except during the two world warm and the depression years of the thirties, Tsumeb Corporation purchased the property in 1947, installed a new plat, and ham built a modern tom. In addition to the concentrator which began operations in 1948, a copper smelter was started in 1962 and a lead smelter in 1963, Tonnage of ore treated in the concentrator has varied through the years, reaching a maximum of 70,000 tons per month with a current mined tonnage of approximately 50,000 tons per month. Copper concentrates, lead concentrates and zinc concentrates are produced, and for a number of years the production of a germanium enriched concentrate was an important feature of the operation, Both copper and lead concentrates are smelted in Tsumeb, while the zinc concentrates are shipped overseas. Ore change, and the installation of the smelters in Tsumeb have necessitated modifications in the concentrator operation. This paper will describe current plant practice in some detail and will touch briefly on the techniques used from 1954 to

Jan 1, 1970

By A. G. Preban, R. J. De Angelis, J. B. Cohen, P. Gordon, D. C. Brillhart

Changes in stored energy, resistivity, density, X-ray line broadening, and dislocation arrangements (from transmission electron microscopy) have been mensured on copper specimens shock-loaded at 75 to 435 kbars. The stored energy in shock-loaded copper and changes in density and resistivity are larger than those after conventional deformation. A high density of twins makes an appreciable contribution to all the measured quantities. Dislocation densities rise to values of 1011 lines per sq cm at the highest pressure. The resistivity, density, and X-ray In a study of the effects of shock loading on copper, changes in stored energy, resistivity, density, hardness, lattice parameter, line broadening, and dislocation structure have been examined after subsequent annealing of specimens shocked to various pressures. This paper presents certain aspects of the data, taken by two groups, one at the Illinois Institute of Technology, and the other at Northwestern University, and compares the changes in these properties, and what they indicate about the effects of shock deformation. measurements yield dislocation densities in substantial agreement at all pressures, after correcting for the effects due to the twins. At the higher pressures, electron microscopy yields dislocation densities that are lower by a factor of about 5 to 10. Dislocation densities from stored -energy measurements are high; this finding can be rationalized by the consideration that the deformation twin boundaries present are largely incoherent. There is no evidence for homogeneous dislocation nucleation during shock loading. Shock loading any metal or alloy, such as copper, causes a large increase in yield strength, with dimensional changes of the order of only a few percent."' Other properties have not been extensively studied, but this result clearly indicates the importance of examining the substructure. Previous studies of the substructure of conventionally deformed specimens have been made using electron microscopy,3-6 X-ray diffraction,7,8 and etch pitting.9,10 The dislocation density can reach values of about 10" per sq cm; and the dislocations are arranged in cells approximately 0.15 in size. The substructure varies widely, however, from point to point in the specimen. The substructure of explosively loaded copper consists of cells of about the same dimensions6,11 but their size is more uniform.6 Profuse twinning occurs12,13 and the twins are often too small to be detected optically.6,13 The boundaries outlining these twins appear to be incoherent.14 There is now ample evidence 13-15 that twinning is a strengthening mechanism in fcc metals but this is only part of the strengthening effect in shock-loaded copper.* From the twinning systems which operate in shock-loaded single crystals,13 and from changes in texture on shock loading,17 it has been shown that both the

Jan 1, 1968

By John F. Elliott, John R. Rawling

The rate of reduction of silica to silicon by carbon at 1550° to 1700°C in iron blast-furnace type slag-metal systems has been investigated. In the tower portion of the temperature range oxygen transport in the metal boundary layer at the slag-tal interface is probably rate limiting. A second step, which apparently involves a chemical process in the slag at the slag-metal interface, appears to be rate limiting when the oxygen-transport limitation is avoided. Above 1650°C rate control of the reduction by carbon probably is mixed between the two steps. INVESTIGATIONS1.2 into the equilibrium distribution of silicon between graphite-saturated slags and metals typical of iron blast-furnace practice show that, under the experimental conditions specified, the reduction of silica to silicon is relatively slow compared with the rates of other slag-metal reactions under similar conditions. For this reason, the reduction of silica in blast-furnace slag-metal systems was considered to be an interesting subject for further study. Investigations of the kinetics of the reaction3-5 suggest that several possible steps may be rate controlling, and that the rate-controlling step depends not only on the physical state of the systems but also on the configuration of the apparatus. The main objects of the investigation reported here were to sug- gest plausible reaction paths for the reduction of silica under the experimental conditions chosen and to determine the step or steps which control the rate of reduction of silica. EXPERIMENTAL Two types of experiments involving the reduction of silica from the slag were carried out. In the first, the only reducing agent present was carbon (graphite), for which the over-all reaction is (SiO2) +2C(gr) —Si + 2Co(g) [l] The second type was a graphite-saturated ferro-silicon-iron concentration cell. In it silica in the slag could be reduced either by carbon according to Eq. [I], or by a reaction of electrochemical nature in which silicon was oxidized at the ferrosilicon-slag interface and silica was reduced at the slag-iron interface without the formation of CO. The two types of experiment are described below. Reduction of Silica with Carbon as the Only Available Reducing Agent. The reduction of silica was accomplished by allowing a liquid lime-alumina-silica slag and graphite-saturated iron to react in an induction-heated graphite crucible. The crucible was 2.25 in. diameter and 5 in. deep; it was fitted with a graphite stirrer which penetrated through the slag and into the metal, Fig. 1. The slag and metal charges weighed 350 g each. Pure carbon monoxide at a pressure of 1 atm was maintained above the melt by passing the purified gas through the crucible at a rate of 0.5 liter per min. The gas did not stir either the slag or metal. Temperature measurements were made by sighting an optical pyrometer down the hollow stirrer rod into a 1/4-in. hole drilled in the stirrer blade. The crucible temperature was maintained within *15"C of the required value by manual control of the power supply to the

Jan 1, 1965

By R. R. Jones, R. W. Kraft

Zn-Mg2Znn eutectic alloys nzay freeze willr either rodlike or lanzellar rnorphology. Alloys with slighlly more than /he eutectic arrzount of rnagnesillrn usually contain three-cnned dendrjles of MgzZnll in a eutec-lic ttlulris. All three morphologies haue the same cryslallographic orientution relationship: (0UOl) zn - 11 (111) Mg2Znll and (2310)Zn 11(101) Mg2Znll, but u3ith different prej-erred groulth direclions. The lurnellae lo rods transifion in con/rolled ingols qf euleclic cotnposition occurs because lhe large kinelic undercooling due to MgzZnll minirrzizes /he ejj-ecl of the solid-solid inlerface energy. The eutectic morphology is influenced by the presence of lhree-nned dendrites 0-f MgzZn11 which may conlrol /he rricroslrccture by acting as nuclealion sites. In recent years there has been much interest in eutectic solidification and several theories have been proposed. One of the confusing factors is the existence of various morphologies in which the solidified phases may form. The lamellar microstructure seems to be most common in metal eutectics, and it has been claimed' that all regular eutectics should be lamellar if sufficiently pure. However, there still remain eutectic alloys which are not lamellar or which change their morphology as a function of growth conditions. The eutectic between zinc and the intermetallic phase Mg2Znll was chosen for this investigation because it has been found to solidify in more than one morphology. The diagram in anssen' locates the eutectic point at 3.0 wt pct Mg and 367°C. lliott gives 364°C as the eutectic temperature, leaving the phase compositions unaltered. Since the growth conditions determine the micro-structure of the solidified alloy, the factors controlling the transition from one morphology to another could be studied. The lamellae to rods transition is of particular interest. PROCEDURE Alloys were prepared from carefully weighed portions of 99.999 pct Zn and 99.97 pct Mg by melting in Pyrex containers under argon and casting into graphite boats. The resulting ingots were remelted under argon and solidified unidirectionally in a horizontal tube furnace at growth rates ranging from 2.0 to more than 50 cm per hr under a temperature gradient, measured over a 5-cm length, of 9" to 14°C per cm. The solid-liquid interface appeared to be planar at all growth rates although no attempt was made to confirm this by decantation or quenching. A few ingots were allowed to freeze uncontrolled. Most alloys were of the nominal eutectic composition, 3.0 wt pct Mg according to Hansen2 and lliott, but some contained as much as 3.35 wt pct Mg. Chemical analyses were not run since metallographic examination confirmed that the desired composition was achieved. Specimens were cut from the middle portion of the ingot normal to the growth axis, polished mechanically, and etched with 2 pct Nital. Suitable areas were selected for the determination of crystallographic orientation relationships by a tiontechniqueof described previously by one of the authors.4 The (2310) planes of zinc and the (8701, {944}? (1032) planes of Mg2Znll were found suitable for orientation determination; experimental error was on the order of 2 or 3 deg. RESULTS Three different morphologies were found in the unidirectionally solidified alloys: lamellar eutectic, rod-like eutectic, and a structure whose most predominant characteristic was the presence of three-vaned (cellular) dendrites of Mg2Znll. These dendrites were only found in alloys with more than the eutectic amount of magnesium. In some ingots fine hexagonal needles of Mg2Znll surrounding a core of MgZn2 were observed. They were probably due to incomplete alloying and seemed to have no effect on the eutectic morphology. In addition hexagonal spirals like those discussed by Fullman and wood5 and Hunt and acksonh ere observed in some ingots frozen without directional control. Both MgZZn,, and MgZnz were detected by X-ray diffraction in these alloys. Since the morphology could not be grown unidirectionally and no characteristic orientation relationship between the phases was found, further study was limited to the lamellar: rodlike, and three-vaned dendrite morphologies. Alloys of Eutectic Composition, No Dendrites. The mcrostructures of allovs with no three-vaned dendrites were either lamellar or rodlike depending on the growth rate. At rates below 10 cm per hr the morphology was lamellar, consisting of two sets of parallel plates intersecting at about 54 deg like the Mg-MgzSn eutectic described by raft.7 At growth rates faster than 14 cm per hr the microstructure showed rods of zinc in a matrix of MgnZnll, while intermediate rates yielded mixtures of rods and lamellae in small groups. The lamellar "grains" were often several millimeters in cross section, but contained small irregular areas which divided each grain into perfect islands 100 or 200 p in diam. Lamellae were parallel to each other throughout the grain in spite of these defects in the structure, Fig. 1. Rods, on the other hand, could only be produced in small groups arranged like fish scales and separated by irregular areas of appreciable thickness, Fig. 2. Alloys Not of Eutectic Composition, With Dendrites. In alloys with 3.1 to 3.35 wt pct ME,-. three-vaned dendrites bf MgzZnll were usually found surrounded by eutectic. At growth rates slower than about 10 cm per hr the dendrites were separated from each other by small areas of both lamellar and rod eutectic, Fig. 3.

Jan 1, 1969

By Pierre J. Goossens

The coastal Cordillera around the Panama basin is made up of a Cu-rich tholeiitic volcanic belt representing either immature, late Mesozoic, island arc, or oceanic ridge. The same belt is exposed (uplifted) along the western slope of the Cordillera Occidental of Ecuador and Colombia. Ni-lateritic, Cr-podiform, and Pt-(Au) placer deposits are formed there after or within ultramafic bodies of oceanic association. Younger calc-alkaline igneous activity provided the mechanism for porphyry copper and associated ore formation. This Cu (+Mo) mineralization occurs along a narrow belt exposed by the erosion of the volcanic cap. Ag-Pb-Zn(-Au) and Hg volcanogenic mineralizations occur within the calc-alkaline belt, but slightly eastward and only above the copper belt. To the east and in the Venezuelan Andes, Paleozoic and older sediments, volcanics, and metamorphics contain volcanogenic massive sulfide, Pb-Zn, Au, and Cu sedimentary and vein deposits, and sedimentary iron ore and phosphate. The Cenozoic Cu, Ni-Pt-Cr and a portion of Ag-Pb-Zn, Hg, and Au mineralizations occur within or just above the oceanic basement whereas the eastern- most Cenozoic Ag-Pb-Zn mineralizations occur within and above a pre-Cenozoic sialic basement. This continental crust is the northern extension of the Peru-Chilean Andes where no oceanic crust has been preserved, however. Cenozoic metallogeny and subduction relationships are unclear in the Northern Andes. Yet strong structural volcanogenic ore controls exists apparently due to deep strike-slip movement between two plates.

Jan 1, 1977

By J. W. Rutter, K. T. Aust

Observations are presented of the formation of annealing twins during the growth of a recrystallized grain into a striated, melt-grown crystal of high-purity lead. The formation of an annealing twin resulted in replacement of the original migrating boundary by a large-angle grain boundary of the caincidence-site type or by a near-twin boundary. The results support the Fullman-Fisher theory of annealing twins and demonstrate the usefulness of the coincidence-site model. FULLMAN and Fisher1 have advanced the following theory to account for the formation of annealing twins during grain growth. Whenever a grain A grows into a matrix of grains S1 and S2, a twin A' of grain A may form, as shown in Fig. 1, if the following energetic condition is satisfied: where "E" and "a" refer to the specific interfacial free energy and the area, respectively, of the boundaries indicated by the subscripts. In effect, the twinned grain will persist when the energy of the boundary between the twin(A') and the matrix grain (S, or S,) is less than that of the boundary between the original growing grain (A) and the matrix grain by an amount which exceeds the energy of the coherent twin boundary between A and A'. Experimental observations on the changes in grain boundary free energy accompanying the formation of twins1 are in agreement with the theory outlined in Eq.[l]. Direct observations of the formation of annealing twins at moving grain corners1"3 and the depend-

Jan 1, 1961

By C. E. Mills

The Verde mining district is in Yavapai County, in north-central Arizona. Jerome, the principal town, has a population of 6000 and the two important mines of the district—the United Verde and the United Verde Extension—lie almost within its limits. Perched high on the southwest side of the Verde valley, at elevations of 4800 and 5500 ft. (1463 and 1676 m.) respectively, both mines have long tunnels from their lower levels through which the ore is hauled to the smelters. Clarkdale, the smelter town for the United Verde mines, lies in the Verde valley 4.1 miles (6.5 km.) distant in an air line from Jerome and at an elevation of 3560 ft. The two towns are connected by the standard-gage railroad of the Verde Tunnel and Smelter R. R., a subsidiary of the United Verde Copper Co. This road is 11 miles long and has a compensated grade of 4 per cent. At Clarkdale it connects with the Verde valley branch of the Santa Fe railroad. Clemenceau, the smelter town for the United Verde Extension mine, lies 3.5 miles (5.8 km.) in a southeasterly direction from Clarkdale, at an elevation of 3420 ft. It is connected to the Verde Valley branch of the Santa Fe railroad at Clarkdale by the Arizona Extension R. R., a subsidiary of the United Verde Extension Mining Co. The portal of the haulage tunnel, at an elevation of 3757.5 ft., is also connected with the smelter town by 4 miles of standard-gage track belonging to the same company. In the vicinity of Jerome, the great southern escarpment of the Arizona plateau forms the northeast side of the Verde valley. The general level of this escarpment, or "rim" as it is called, is about 4000 ft. above the Verde River. A corresponding elevation is reached only by the highest peaks of the Black Hills, which bound the valley to the southwest. Two lava-capped mesas, known as "Mingus" and "Woodchute mountains," form the culmination of the Black Hills. That part of the Old Black Hills mining district which lies between their crests and the Verde River is now known as the Jerome, or Verde, mining district. If the smelter towns are not included, a strip about 7 miles long by 3 miles (11 by 5 km.) wide, embraces practically all of the properties belonging to the Jerome district. T-bis strip lies parallel to and along the side of the Verde valley from First View, 2 miles northwest from Jerome, to the

Jan 1, 1925

By J. Stuart Warner

The reaction COSO4(c)?CoO(c) + So3(g)was investigated from 950° to 1170OK by two different methods. The sulfate was decomposed in a previously evacuated space and Pso3 calculated from the measured total pressure on the assumption that Pso2 = 2P02. The sulfate was also decomposed under conditions of controlled oxygen pressure. The difference in the results obtained by these experimental techniques is attributed to the effects of thermal diffusion. The standard free energy change for the reaction is given by ?F° = 66,292 + 17.273T log T — 102.57 T (±300 cal/mole). It was found by DTA and X-ray diffraction studies that CoS04 undergoes a transformation at 617°± 4°C. THE metal sulfates may decompose directly to the oxide or pass through one or more intermediate basic sulfates. Provided the oxidation state of the cation does not change in the course of the reaction, the decompositions may be described in general as follows: One of the major objectives of this work was to determine the validity of Eq. [3] as much of the existing sulfate data has been calculated from decomposition pressures by Eq. [5]. The type of apparatus generally used for these measurements consists of a reaction vessel at high temperature which communicates with a manometer at room temperature. A temperature gradient is thus imposed on the gas phase causing thermal diffusion which brings about a partial separation of the different gaseous species present in an initially uniform mixture. The basic theory and experimental facts associated with this phenomenon are summarized elsewhere1 and it will suffice to state here that the heavier gaseous species diffuses down the temperature gradient. Since thermal diffusion destroys the homogeneity of the gas phase, it is not necessarily true that Pso2 = 2Po2 in any particular volume element of the gas phase even though there are twice as many molecules of SO, as of 0, in the gas phase considered as a whole. Under these conditions, the decomposing sulfate gives off SO, which dissociates to SO, and 0, and all three species are distributed throughout the system by thermal diffusion. However, the sulfate will continue to dissociate until the equilibrium value of PSo3 is established on the hot zone. Thermal diffusion may distribute SO2 and O2 throughout the gas phase in any way so long as Eq. [2b] is satisfied in the hot zone; i.e. Pso2(Po2)1/2=K3pso3. When a steady state is reached, we can characterize Pso2 in that

Jan 1, 1962

By W. R. Kanne, P. R. Strutt, R. A. Dodd

A study has been made of the creep behavior of ß-NiAl of stoichiometric composition in the temperature range 475" to 775°C. Single crystal tensile specimens were deformed under a constant applied load. From slip line studies and tensile axis rotations it was concluded that the crystals deformed by pencil-glide on (100) and (110) planes with a (100) slip direction. Electron microscope thin-film studies showed evidence for profuse cross-slip, and the dislocation Burgers vector was determined to be (100). There was no tendency for the dislocations to lie in definite slip planes, although a strong alignment of segments of edge dislocations was frequently seen. The activation energy for creep in the temperature range 475" and 775°C was judged to be much lower than that for creep at temperatures above 775°C. It is possible that the rate-controlling process is the thermally activated cross-slip of screw dislocations. In recent years considerable interest has developed in the possible use of high-melting-point intermediate phases for high-temperature applications necessitating good creep resistance. Ordered alloys such as Ni3A1, modified by alloying additions, have received particular attention, culminating in the recent development by Pratt and Whitney of single crystal turbine blades.' The intermediate phase NiAl has considerable theoretical interest from a creep standpoint because not only is it ordered up to the melting point (-1650°C at stoichiometry) but it contains a high concentration of structural vacancies on the aluminum-rich side of stoichiometry. These vacancies are known to exert a strengthening effect under some conditions,' but in addition they might be expected to affect high-temperature creep rates through dislocation climb. The present paper deals with the slip line and dislocation structures developed during creep of stoichiometric NiAl between 475" and 775C, and with possible attendant creep mechanisms. MATERIALS AND TEST TECHNIQUES NiAl ingots of stoichiometric composition were prepared by induction melting 99.996 pct A1 and 99.98 pct Ni (kindly provided by the Kaiser Aluminum Co. and the International Nickel Co., respectively) in 100 cc pure alumina crucibles in a helium atmosphere. No chemical melting flux was used in order to avoid con- tamination of the alloy. Each ingot was slowly direc-tionally solidified to avoid porosity and to produce very large grains up to 2 in. in length. In this way, ingots were obtained which were suitable for sectioning into single crystal creep specimens. Chemical analysis indicated that all alloys were within the range 49.6 0.3 at. pct Ni. For the present investigation, identical rectangular single crystal blanks were cut vertically from one particularly large crystal in an ingot using a series of three diamond cut-off wheels with appropriate spacers. Thus, all sliced crystals had the same orientation with the tensile axis parallel to the long side of the rectangle; the initial orientation of the tensile axis is shown as S, in Fig. 1. The blanks were homogenized at 1250 to 1300°C for 48 hr and subsequently cooled slowly to room temperature. A gage length of rectangular cross section and about 0.75 in. in length was prepared on each specimen using an electrolytic lathe and a dilute nitric/phosphoric acid electrolyte. The gage length cross section was 0.250 0.0005 in. by 0.040 ± 0.0005 in. Holes to accommodate tensile grips were electrolytically machined using a rotating acid jet. After creep testing, this same technique was used to cut discs for transmission electron microscopy. Thus, all shaping was done electrolytically, the only exception being that (122) sections were spark-machined from some speci-

Jan 1, 1970

By J. C. Shyne, T. D. Gulden

The steady-state creep behavior, in compression, of indium containing a dispersion of atomized glass particles was studzed over a range of temperature, stress, and composition. The observed behavior was not consistent with a simple model of dislocation climb over dispersed particles. From room temperature to about 100°C the temperature dependence could be described by an activation energy about one and a half times that of self-dijfusion in indium while at higher temperatures a much higher ternperature dependence was observed. THE value of dispersion-hardened and precipitation-hardened alloys has long been recognized. The mechanism of hardening, however, is not completely understood, especially at elevated temperatures where dispersed particles increase creep resistance. Following the suggestion of Schoeck1 that the rate of creep in dispersion-hardened materials is controlled by dislocation climb over second-phase particles, Ansel and weertman2 derived steady-state creep equations for dispersion-hardened materials. Their derivation was based on dislocation climb over non-deformable spherical particles in a metallic matrix. For low stresses they predicted a linear relation between applied stress and steady-state creep rate. For stresses large enough to cause dislocations to pinch off around the precipitate particles, the theoretical creep rate was shown to be proportional to the fourth power of stress. For very high stresses an exponential stress dependence was predicted. In all cases the theoretical creep rate is linearly proportional to the diffusion coefficient. This provided the primary temperature-dependent factor. Subsequently, Ansel and Lenel3 reported some experimental data on SAP, a composite of aluminum and aluminum oxide, that was in good qualitative agreement with the above theories. Their experimental creep rates, however, were four orders of magnitude less than the theory predicts. This was attributed to a low density of active dislocation sources in the material; somehow the dispersed particles inactivated the dislocation sources instead of merely interferring with dislocation motion. Meyers and sherby4 have recently reported on the creep behavior of SAP. Their results were inconsistent with a simple dispersion-strengthening model and led them to conclude that a continuous network of oxide controls the high-temperature mechanical characteristics of SAP. Because of the rather complex mechanical behavior of SAP and the lack of agreement about the morphology of its oxide phase, it appeared desirable to study the creep of a dispersion-strengthened material of controlled morphology. For the present work a dispersion of spherical glass particles imbedded in a matrix of indium was chosen. The microstructure of this material was similar to the idealized uniform dispersion of hard spheres in a deforming matrix assumed by Ansel and Weertman in their theoretical treatment. SPECIMEN PREPARATION For experimental convenience, it was decided to use a low melting-point metal in this investigation. It was necessary that the metal not be oxidized readily up to its melting temperature and that it wet the second-phase particles. It was desired that the material chosen for the dispersed phase be available as a fine powder with closely controlled size and shape. The material chosen to satisfy these requirements consisted of a matrix of indium metal containing a dispersion of atomized particles of a soft soda-lime-silica glass ranging from 5 to 30 µ in diam. The glass particles had a highly regular spherical shape. The composite was made by stirring the glass powder into liquid indium maintained at a temperature just above the melting point of indium, 156°C. The creep specimens were prepared by a hot pressing operation. Five to ten g of the mixture were placed in a steel die and pressed at 125°C for 5 min under a pressure of 30,000 psi. The resulting specimens were cylinders about 3/8 in. in diam and ranged from 0.4 to 0.7 in. high. By following this procedure a uniform dispersion was obtained in the creep specimens. Typical microstructures of specimens containing twenty and forty volume percent glass powder are shown in Figs. 1 and 2. Unfortunately, grain size could not be determined by the metallo-graphic techniques used. CREEP TESTING PROCEDURE Creep tests were performed in compression under constant stress. The specimens were placed in

Jan 1, 1963

By F. H. Vitovec, D. W. Hoeppner

Pusk-pull fatigue tests were conducted on copper tricrystals of 99.988 pct purity to ascertain the role of grain boundaries in the initiation and propagation of fatigue cracks. Significant differences in behavior were found for specimens which possessed different transverse-boundary misorientation. In speciwens with low boundary angles cracks initiated within the transverse boundary, while higher angles led to transcrystalline fatigue failure. It is suggested that at low angular misorientation moving dislocations may interact with dislocations of the boundary or dislocations present in adjacent pains on favorably oriented glide planes, thus initiating a fatigue crack. MANY fatigue studies have been concerned with fatigue-crack initiation within grains and the mechanism causing initiation and propagation1-3 Although the initiation of fatigue cracks in or near grain boundaries of pure metals has been observed and reported in the literature, the mechanism of this phenomenon has received little attention.4"11 EXPERIMENTAL PROGRAM Testing Procedure. To investigate the role of grain boundaries regarding initiation and propagation of fatigue cracks, copper tricrystals were tested in push-pull. Axial-stress tests were used to avoid the stress gradients introduced by stressing of some other nature. Copper was selected as the most suitable test material since extensive work has been done on the formation of fatigue-induced slip in copper. Tricrystal specimens were used to provide a grain boundary geometry with one boundary transverse to the principal stress axis and one or more boundaries nearly coincident with the direction of maximum resolved shear stress. The boundary energy and the slip characteristics in the vicinity of the transverse boundary depend on the relative orientations of grains across the boundary. Testing was done at room temperature, at a frequency of 700 cpm, in an atmosphere of either air or argon. It is known that the incidence of grain boundary cracking increases with increasing temperature and decreasing frequency.4 The fatigue machine applied a uniaxial tension-compression load to the specimen by a flat spring which was actuated by an adjustable eccentric. Use was made of an adjustable head and a load cell which had strain gages mounted at 120-deg intervals around its periphery, thus providing a means for eliminating any detectable superimposed bending moment. A clip gage was mounted between the gripping heads to record the cyclic strain amplitude applied to the specimen. Each specimen's hysteresis characteristics were recorded by supplying the load and strain signals to an oscilloscope. Microstructural changes were observed and recorded with an optical microscope which was mounted on the fatigue machine. Immediately prior to insertion in the machine, each specimen was chemically polished. Extreme care was exercised while inserting the specimen in the test machine to avoid either bending the specimen or introducing a mean load. Each specimen was stressed in the tensile direction first and subsequently the load was reversed. Specimen Preparation. Tricrystal fatigue specimens of 99.988 pct Cu were grown from the melt using a modified Bridgman technique. Ingots about 1 in. high were grown to the configuration shown in Fig. 1. Upon sectioning the resultant ingot, several specimens were provided with an identical shape and grain boundary orientation. Except for plane sectioning and final polishing, this method eliminated machining the specimens. The apparatus used for growing the tricrystals is shown in Fig. 2. A spectrographically pure graphite mold, Fig. 3, was inserted in a vycor tube which was mounted on a vertical zone-refining table. Prior to insertion in the tube the mold was assembled as follows. The insert and graphite rods were fixed in position in the mold. The copper was then placed in the mold and the entire mold assembly was positioned in the vycor tube, Fig. 2. At this point the tube was sealed, evacuated, purged with helium gas, and finally placed under a slightly positive helium pressure. The heating coil was then adjusted to melt the copper charge, after which it was raised at a speed of 4 in. per hr. It is also possible to use

Jan 1, 1964

By Carl F. Lutz, Robert F. Mehl

A layer of brass was formed on 0 brass using a vapor-solid reaction technique. The variation in composition with distance within the phase layer and across the a -ß interface was determined by an electron probe. The diffusivities were calculated as a function of concentration; the diffusion coefficient in brass was found to change by a factor of twenty-five over a compositional range of 8 at. pct. An explanation is proposed to account for the large change in the diffusivity, based on possible stresses in the diffusion zone; it is suggested that the concentration dependence of the thermodynamic factor might decrease the activation energy with increasing zinc content. ALTHOUGH the body of data on diffusion coefficients (D- values) in terminal solid solutions and in isomorphous systems is quite large, there is but little quantitative information for intermediate alloy phases. This paper recounts measurements of D-values for the ? phase in the Cu-Zn system. The experimental method employed consists in the forming of alloy layers and determining and analyzing the concentration-distance (c-x) curve. The experimental method employed could have been applied to several systems; the Cu-Zn system was chosen because the D-values in the a and ß phases are well known, and can thus be used for purposes of comparison. The rates of growth, of intermediate alloy layers are known for a number of systems. In nearly all cases the thickness varies parabolically with time, as to be expected if the rate of growth is controlled by the D-values. In a few cases, nonparabolic behavior has been noted.1 Nonparabolic growth may be the result of a) a variation of the composition at the phase interface with time, b) interface-controlled kinetics, c) short diffusion times coupled with long vacancy lifetimes (in vacancy diffusion),2 and d) crystal reorientation during diffusion in anisotropic systems.3 In the Al-Ni system4 and in the Al-u5 system parabolic growth kinetics are slowly approached after an initial transient period. In general those phases stable at a given temperature in a system A-B appear as separate phase layers when A and B are put in contact and given a diffusion he at-treatment. In most cases the compositions at the interface of adjacent phase layers are those read from the phase diagrams for the termini of the respective phases. This discontinuity in concentration is taken as independent of time, and the growth of one phase at the expense of another is assumed to be independent of interface kinetics, i.e., the rate of interface movement is controlled by volume diffusion in the phases concerned. wagner6 has given many solutions for multi-phase diffusion processes, assuming that the chemical diffusion coefficient, D, is independent of concentration, as have others. Jost7 has pointed out that the familiar Boltzmann-Matano solution is as valid for a (c-x) curve exhibiting concentration discontinuities at phase boundaries as for the usual single-phase case. Heumann and associates8,9 have applied this solution to the multi-phase case, but lacking concentration data within the several layers were forced to assume linear concentration gradients, thus calculating only average D-values. The purposes of the present study were: 1) to determine the dependence of D on concentration, 2) to calculate the intrinsic diffusion coefficients D?Cu and D?Zn, 3) to establish the time-law for the movement of the interface, 4) to determine the concentration limits at the ß-? interface, and, 5) to attempt to clarify the mechanism of diffusion. EXPERIMENTAL PROCEDURE The ? phase is exceedingly brittle; conventional solid couples and conventional sectioning methods were thus inapplicable. Vapor-solid couples were accordingly employed. Unfortunately the equilibrium vapor pressures of Zn for the ? phase are unknown; to escape this awkwardness, samples were enclosed in a capsule containing powder of an alloy of known composition, sufficiently large in quantity as to be effectively an infinite source, maintaining the concentration of Zn at the sample surface constant with time. The sources of Zn-vapor which can be employed in reaction with Cu, or a brass, or ß brass, to form a layer of ? brass satisfactorily wide in concentration range, are alloys of the phases ?, or ? + E or ? + d (depending on the temperature).

Jan 1, 1962

By Paul A. Farrar, Sanford Adler

The Tz-Zr system was reinvestigated using both metallographic and X-ray diffraction techniques. It mas found that titanium and zirconium are soluble in all proportions in both the a and 0 phases. The minimum in the transformation was found at 50 at. pct and 535°C. THE Ti-Zr system has previously been the subject of several investigations.1-6 A study by Hayes et al.5 using magnesium-reduced titanium and zirconium melted in graphite crucibles indicated that the system was a continuous series of solid solutions in both a and 0 phases with a minimum at approximately 545°C and 65 pct* Zr. These conclusions were substantiated by the work of Duwez4 as well as by the limited data of Craighead et al.,3 Fast,2 and deBoer.1 However, a more recent investigation of the titanium-rich region by Ence and Margolin6 indicated that the solubility of zirconium in a, titanium at 500°C is approximately 22 pct with the a + 0 field extending to approximately 47 pct Zr. Therefore in order to resolve this discrepancy the following investigation was initiated. EXPERIMENTAL PROCEDURES The alloys used in this investigation were prepared as 20- to 30-g buttons by nonconsumable electrode melting in a helium atmosphere. The starting materials were Bureau of Mines titanium BHN 71 and Foot Mineral iodide crystal bar zirconium. The titanium was premelted before the alloys were prepared to avoid excessive weight losses in the final melting. The compositions of the alloys prepared were as follows: Ti-12.0, 22.2, 29.0, 39.6, 50.0, 58.0, 72.4, 83.5. and 90.8 pct Zr. Prior to heat treatment the alloys were cold-worked 10 to 20 pct, stopping at the first signs of cracking. Specimens for heat treatment were wrapped in molybdenum sheet and annealed in argon-filled quartz capsules for the following times and temperatures: 750°C-40 days, 675°C-61 days, 600°C—109 days, 575°C-112 days, 550°C-112 days, 525°C—127 days, and 450°C —153 days. Quenching was accomplished by breaking the capsules in an iced-brine solution. The standard techniques used for polishing the specimens involved belt grinding, grinding on emery paper, polishing electrolytically, and etching with Remington "A" etch7 or "R" etch 8 Debye-Scherrer X-ray powder photograms were obtained for a number of samples using a 114.6-mm camera and CuK, radiation. Exposure times were from 19 to 23 hr. The powder samples were obtained by filing to 270 mesh size. After filing the samples were wrapped in molybdenum foil capsuled in quartz and reannealed at the temperature of original heat treatments; following the heat treatment the samples were quenched into iced brine without breaking the capsules. RESULTS AND CONCLUSIONS On the basis of the microstructural examinations of the heat-treated samples and of the X-ray diffraction data, the Ti-Zr system shown in Fig. 1 was constructed. The data indicate that a, and 0 titanium form a complete series of solid solutions with a, and ß zirconium. The minimum was found to occur at 65 pct (50 at. pct) Zr and 535°C, in excellent agreement with the earlier investigations,1-5 see Fig. 2. Typical microstructures are shown in Figs. 3 to 7. Figs. 3, 4, and 5 show the 22.2, 39.6, and 50.0 pct Zr alloys after heat treatment at 675°C for 61 days with the 22.2 pct Zr showing equiaxed a, and the 39.6 pct Zr alloy a + transformed ß, while the 50 pct Zr alloy shows only transformed B. Figs. 6 and 7 show the 29.0 and 39.6 pct Zr alloys after heat treatment for 109 days at 600°C. The 29 pct Zr alloy has an equiaxed a structure while the 39.6 pct alloy has a partially recrystallized a+ß structure. X-ray diffraction data obtained from the 50 pct Zr alloys after heat treatment at 525° and 450°C as well as from the as-cold-rolled material showed only those lineos which couid be indexed a: a titanium, good agreement with the previously published values. 2,4,5,9 Chemical analysis* of the 50 pct Zr alloy

Jan 1, 1967

By V. B. Kurfman

The temperature, strain rate, and orientation deDendence of defbrnzation of single-crystal MgAg has been examined. The crystals exhibit a tendency to single glide and little or no hardening at 25°C for many orientations. A much higher hardening rate is observed when multiple glide occurs, such as can be initiated by surface defects. The tendency for easy glide becomes less dependent on surface preparation and orientation as T — 100°C and bars so tested often fail after one-dimensional necking-. At T > 200°C (transition temperature for single-crystal notch sensitivity and poly crystalline ductility) single glide diminishes and two-dirnensionul necking begins. The crystals do not strictly obey a critical resolved shear stress law, but show the influence of {loo) cracks in determining the slip mode. The results are correlated with the difficulty of sciperdzslocation intersection and semibrittle behavior of this compound in single-crystal and poly crystalline form. Comparisons are made with the slip selection mode observed in tungsten, with the reported observations of easy glide in bee metals. and with the mechanical behavior of poly crystalline MgAg. PREVIOUS work on tensile deformation of polycrys-talline MgAgl and bending deformation of single-crystal MgAg2 has shown that the compound is semi-brittle (i.e., notch and grain boundary brittle). If this semibrittleness is supposed to result from the difficulty of multiple glide (associated with the problems of superdislocation intersection) one might expect single crystals deformed in tension to show pronounced single glide and strong orientation dependence of hardening rate. These experiments were done to examine this supposition and to study the tensile deformation of a highly ordered system which may be considered bcc if the difference between the two kinds of atoms is ignored (actual structure: CsC1). EXPERIMENTAL Single-crystal ingots were grown by directional freezing as previously described.' These ingots were sliced into a by a by 2 in, rectangular bars by electric discharge machining, then round tensile bars were conventionally machined to 1/8-in.-diam by 1-in.-long reduced section. The bars were typically tested without an anneal because of the problem of magnesium vapor loss and they were typically tested as mechanically polished. The analyses are within the same limits as those reported earlier; i.e., the average composition for each specimen is within 0.5 at. pct of stoichiometry, while the total range from end to end in a given specimen varies from 0.7 to 1.4 at, pct. There has been no indication in the results of any variation in slip or fracture mode attributable to the composition fluctuations. The slip systems were determined by two-surface analysis of the bars after testing to failure at room temperature. Single glide was so dominant that there was little difficulty in identification of the dominant slip system even though the tensile elongation to failure often approached 7 to 8 pct in room-tempera- ture tests. Elevated-temperature testing was done in a silicone oil bath and low-temperature testing was done in liquid Np or a dry-ice bath. All stress measurements are reported as engineering stress unless otherwise specified, and crosshead travel is used as the strain measurement. RESULTS The tendency toward single glide is best seen in the pictures, Figs. 1, 2, and 3, which depict deformation at fracture as a function of test temperature. While it is possible to find regions of secondary slip by careful microscopy, such regions are very small. The development of a ribbon-shaped configuration from an initially round section bar pulled at 100°C is typical, occurred by single glide, and illustrates the degree to which such glide continues. At temperatures =100°C the bars typically show elongation of 20 to 50 pct by predominently single glide. Despite the large elongation, fracture even at 150°C occurs in a brittle mode, Fig. 2, in the sense that it is an abrupt failure which shows no discernible necking in the second dimension of the bar's cross section (i.e., there is no appreciable action of any slip modes which would decrease the broad dimension of the cross section). Near 200°C the fracture mode changes slightly. Although most of the sample extension is by single glide, after the bar develops the characteristic ribbon shape it begins to neck in the second (i.e., broad) cross-sectional dimension. The bar becomes very thin in the "necked down" region, Fig. 3, and the reduction in area approaches 100 pct. Often there oc-

Jan 1, 1967

By E. J. Rapperport, E. D. Levine

The boundaries of the (a +ß) field in the Al-Li system were determined between 150°and 550°C utilizing quantitative metallography and lattice-parameter measurements. The solubility of lithium in aluminum decreases from 12.0at. pct Li at 550°C to 5.5 at. pct Li at 150°C. P Al-Li is saturated with aluminum at 45.8 at. pct Li and has this boundary value constant over the temperature range 150°to 550°C. THE solid solubility of lithium in aluminum has been determined by several investigators, 1-6 but, as shown in Fig. 1, there is little agreement among the various determinations. The earliest investiga-tions'-' are suspect because of the use of impure materials. Although high-purity materials were employed in more recent work,4'5 the experimental techniques may have led to contamination of the specimens. Probably the best work has been that of Costas and Marshall,6 who obtained close agreement between results obtained by two independent phase-boundary techniques: electrical resistivity and mi-crohardness. No detailed studies of the solubility of aluminum in the bcc ß phase, Al-Li, have been reported. Cursory investigations1,2,6 have indicated only that the (a+ß) -p boundary lies between 40 and 50 at. pct Li and is relatively independent of temperature. The present work was undertaken in order to provide an independent check on Costas and Marshall's determination of the solubility of lithium in aluminum, to extend knowledge of this solubility limit to temperatures below 225°C, and to make an accurate determination of the solubility of aluminum in Al-Li. EXPEFUMENTAL Alloy Preparation. In view of the difficulties encountered in previous investigations of the A1-Li system, close attention was paid to the use of methods of alloy preparation and treatment that would minimize contamination. Aluminum sheet (99.99 + pct Al) was vacuum-induction melted in a beryllia crucible to remove hydrogen. Lithium (99.9 pct Li) was charged with pre-melted aluminum into a beryllia crucible, in a helium-filled drybox. The crucible was sealed in a Vycor tube and transferred from the drybox to an induction furnace. Melting of alloys was performed by induction heating in a helium atmosphere. Solidification was accomplished by means of a suction apparatus, shown in Fig. 2, in which the alloy was forced by changes of pressure into a 3/16-in. inside diam closed-end beryllia tube. This technique produced rapid solidification of a small portion of the melt, resulting in alloys with a high degree of homogeneity. Typical lithium distributions are presented in Table I. Transverse sections 1/8 in. long were cut from the alloy rods, and each section was split in half longitudinally. One half of each section was analyzed for lithium, and the opposing halves were employed for phase-boundary determinations. Lithium contents were determined by flame photometry with an accuracy of 1 pct of the amount of lithium present. Thermal Treatments. Homogenization and equilibration heat treatments were performed in electrical-resistance furnaces with temperatures controlled to ± 2OC. Calibrated chromel-alumel thermocouples were employed to measure temperature. Homogenization was performed in helium-filled l?yrex tubes for 1 hr at 565°C. The encapsulated specimens were then transferred directly to furnaces maintained at lower temperatures for equilibration. Equilibration times were 2 hr at 550°C, 8 hr at 450°C, 27 hr at 350°c, 90 hr at 250°c, and 285 hr at 150"~. These times were chosen on the basis of conditions employed by previous investigators. Alloys were quenched from the equilibration temperatures by breaking the capsules into a silicone oil bath. By performing all possible operations either in sealed capsules or in a helium-filled drybox, the specimens were given minimum exposure to the atmosphere. Quantitative Metallography. Metallography of Al-Li alloys is difficult because of the atmospheric reactivity of the ß phase. It was found possible, however, to prepare surfaces of good metallographic quality by preventing contact with moisture during preparation. Grinding through 4/0 paper was performed in the drybox. The specimens were then transferred under kerosene to the polishing wheel. Three polishing stages were employed: 25-p alundum with kerosene lubricant on billiard cloth, 1-µ diamond paste on Microcloth, and 1/4-p diamond paste on Microcloth. Between stages the samples were cleaned by rinsing in trichloroethylene and buffing

Jan 1, 1963

By J. G. Richardson, J. W. Graham

The purpose of this paper is to present the results of theoretical and experimental studies of water imbibition. The imbibition processes are involved in recovery of oil from stratified and fractured-matrix formations in natural water drives and water flooding. An understanding of the role of inhibition in implementing the recovery of oil from such formations is deemed essential to proper control of these reservoirs to achieve maximum recovery. The theoretical studies involved development of the differential equations which describe the spontaneous imbibition of water by an oil-saturated rock. The dependence of the rate of water intake by the rock on the permeability, interfacial tension, contact angles, fluid viscosities and fluid saturatiorls is discussed. A few experiments were performed using core samples to determine the effects of core length and presence of a free gas suturation. The role of water imbibition in recovery of oil from a fractured-matrix reservoir by water flooding was investigated by use of a laboratory model. This model was scaled to represent one element of a frac-tured-matrix formation. Water floods were made at various rates with several fracture widths. Interpretations were made of the behavior expected in a system containing many matrix blocks. The presence of a free gas sntu.ration was found to reduce the rate of water imbibition. In the reservoir prototype of the fractured-matrix model, water imbibition rather than direct displacement by water was the dominant mechanism in the recovery of oil at low rates. INTRODUCTION Imbibition may be defined as the spontaneous taking up of a liquid by a porous solid. The spontaneous process of imbibition occurs when the fuid-filled solid is immersed or brought in contact with another fluid which preferentially wets the solid. In the process of wetting and flowing into the solid, the imbibing fluid displaces the non-wetting resident fluid. Common examples of this phenomenon are dry bricks soaking up water and expelling air, a blotter soaking up ink and expelling air and reservoir rock soaking up water and expelling oil. As increasingly better lithological descriptions have been made of the characteristics of petroleum-bearing formations, it has become obvious that imbibition phenomena which were once considered laboratory curiosities are of practical importance. For instance, in reservoirs composed of water-wet sand strata of different permeability in intimate contact, the tendency of water to channel through the more permeable stratum is offset by the tendency for water to imbibe into the tight sand and expel oil into the coarse sand. Also, in fractured-matrix formations the tendency of water to channel through the fractures is offset by water-wet matrix blocks. As some imbibition of the water into the of the largest fields in the world are fractured-matrix reservoirs, it has become increasingly important to understand all the factors involved in the imbibition process. Examples of fractured-matrix reservoirs are the Spraberry field in West Texas which produces from a fractured sandstone', the giant Kirkuk field in Iran', the Dukhan field in Qatar, Persian Gulf2, and the Masjid-I-Sula-main and the Haft-Kel fields in Southwestern Iran, which produce from fissured limestone3. Research into recovery of oil from fractured-matrix formations was stimulated by the rapid decline of oil productivity of wells in the Spraberry formation. One result of this research was the water imbibition process developed by the Atlantic Refining Co.4 Another idea was that much of the Spraberry oil could be recovered by conventional water-flooding procedures5. Subsequently, pilot floods were conducted in this field to test the feasibility of these ideas. It was felt that an understanding of the role played by imbibition processes in displacement of oil from a fractured-matrix reservoir could not be obtained from field data alone because of the many complicating factors and uncertainties involved. Therefore, theoretical and laboratory studies were undertaken to provide this understanding. Study of the equations which describe the linear, countercurrent imbibition process provided an insight into the role of various factors in the process, such as the permeability of rock and inter-facial tension. In addition to the theoretical studies, imbibition experiments were conducted with core samples to determine the effect on the rate of imbibition of such variables as core length and free gas saturation. The principal experimental studies were conducted by water flooding a scaled model of an clement of a frac-tu red-matrix reservoir to evaluate