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By Earl S. Greiner, Eric R. Jette

The constitution of the iron-rich iron-nickel-silicon alloys at 600ºC. has been determined by X-ray crystal structure methods. If the results are represented isothermally by a conventional equilateral triangular plot, the single-phase solid solution (body-centered cubic structure) of the iron-rich alloys at 600°C. is bounded by four two-phase areas. Three three-phase areas separate the respective two-phase areas. Three of the phases in the polyphase areas have the same structures as the corresponding phases in the binary iron-nickel or iron-silicon alloys. There are two phases, however, that have not been reported previously and whose structures have been tentatively identified in the present investigation. The solid solubility limit of the iron-rich solid solution at 600°C. was determined by comparing the lattice constants of this body-centered cubic phase in the single and polyphase areas. It has been shown that the lattice constants of ternary alloys may be represented by surfaces whose distances above the ternary plot (representing atomic composition) are proportional to the lattice constants of the respective alloys. Then the intersections of the surface for the single-phase alloys with those for the corresponding phase in the polyphase alloys, at the respective temperatures, are the solid solubility limits of the single-phase alloys at these temperatures. The data of the present investigation were evaluated by this method, although by a system of orthographic projection it was possible to represent the three-dimensional model in two-dimensional diagrams. This investigation has shown that the solid solubility of nickel in the body-centered cubic phase of the iron-rich iron-nickel-silicon alloys at 600°C. changes from 3.7 weight per cent (Owen and Sulley) in the binary iron-nickel alloys to 5.0 per cent in the ternary alloys containing 4.8 per cent silicon. Further additions of silicon increase the solid solubility limit of nickel to 16.0 per cent at 12.6 per cent silicon. From this point the limit decreases to 9.5 per cent nickel at 15.5 per cent silicon and is, of course, zero at 15.3 per cent silicon in the binary iron-silicon alloys.

Jan 1, 1943

By B. C. Allen

The solubility of chromium in liquid silver and that of molybdenum and tungsten in liquid tin have been determined by equilibrating Ike Liquid in a crucible of the solule metal. Generally the weight of solute in solulion was delermined both chemically and from crucible weight losses. The resulting weight percent solubilily of chromium in silver as a function of absolule temperature T is given by log w = -6660/T + 4.02 in the range 960° to 1445°C, and that of molybdenlttrl in tin by log w = —5050/T + 1. 79 in the range 1200° to 2200°C. The values appear unaffected by minor changes in solute composition. Calculations aye made of the parlial molar heal and excess entropy of mixing. The estimated solubility of tungsten in tin is 0.001 pct at 2000°C. Evidence is presented that molybdenwm dissolves in tin without compound for mation and that tungsten and tin form W10Sn. RECENT developments in coatings, heat transfer, and brazing require a knowledge of equilibrium solubilities in a variety of systems at elevated temperatures. Of particular interest are solid-liquid interactions in metal systems involving dilute solutions. This investigation was undertaken since little is known about the phase relations or solubility of chromium in liquid silver or of molybdenum and tungsten in liquid tin.1,2 EXPERIMENTAL WORK As indicated in Table I, high-purity metals or those of known impurity levels were used. Crucibles were machined from chromium, molybdenum, and tungsten bar stock 1.3 cm diam and 2.2 cm long. Chromium rod was prepared by arc melting and extruding iodide process crystals.3 The five chromium impurity alloys were in the form of 0.6-cm-diam swaged rod. A Cr-0.06N* alloy was prepared by exposure to NH3 at 1150°C in a closed quartz capsule for 24 hr. Because significant nitrogen was lost during annealing in argon at 1600°C,4 the estimated level was 0.01 pct N. Chromium and Mo,W crucibles were outgassed at 1200" and 1600°C, respectively, at 3 X 10-5 mm. The solvent was placed In crucibles of the solute metal- 6 g Ag in chromium, 0.4 g Ag in the five chromium impurity alloys, and 4 g Sn in Mo,W crucibles. The silver and tin series were outgassed at 900" and 110O°C, respectively, and held at the desired solution temperature in a tantalum resistance furnace under 1.05 atm 99.995 pct Ar. An equilibration time of 0.5 to 1 hr was chosen since 0.3-, I-, and 4-hr anneals at 1235°C for Cr-Ag yielded similar results. Temperatures were measured optically to +10°C under essentially black body conditions which were checked against the melting point of silver. The Cr-Ag and Mo-Sn,W-Sn series were respectively equilibrated at 990° to 1400°C and 1200" to 2000°C. Vaporization losses were generally held to <1 pct by the argon. In the 1800" and 2000°C anneals for 0.5 hr, the crucibles were placed in tightly sealed tantalum cans which kept tin losses to <5 pct. After the anneals, the specimens were furnace-cooled. One anneal was made on small Mo-0.02 C, Mo-0.5Ti-0.lZr, and tungsten crucibles containing tin at 2200°C. The metals were sealed in an electron-beam-welded tantalum can of minimum volume. After the anneal, tin was still present in each crucible, none of which was severely attacked. After equilibration, the saturated solvent metal was selectively removed from the crucible by acid leaching. Weight losses sustained by the crucibles were determined from the initial weight minus the final blank-corrected weight after acid leaching. Chemical analyses were performed on some of the samples. Hot 1:l HNO3-H2O removed silver from chromium with the blank running <0.1 mg or <0.002 pct on measured solubility. Tin was removed from molybdenum and tungsten by hot concentrated HC1 with the respective blanks amounting to -0.004 and -0.0004 pct in the presence of tin. The insoluble residues were weighed and checked for composition. The leaching solutions contained traces of the solute metals which were analyzed by colorimetric methods, Diphenyl Carbazide for chromium, Thiocyanate for molybdenum, and Dithiol for tungsten. The total of the two percentages found from the residue and leaching solution minus the blank gave the desired solubility values. On the basis of annealing empty molybdenum and tungsten crucibles, weight losses due to MoO3 or WO3 volatilization were found to be generally 10.4 mg and affect measured solubilities by <0.01pct. RESULTS AND DISCUSSION Chromium-Silver . Since no intermediate compounds are reported to form in the Cr-An system,&apos; crucible weight losses and chemical analysis of the entire melt should provide reliable measures of the chromium in solution. Values obtained by the two methods are in agreement as shown by the results for unalloyed chromium in Table II and plotted in Fig. 1. Furthermore, the results are consistent with those previously determined by melting equilibrated silver out of the crucible and analyzing calorimetrically for chromium.&apos; On the basis of limited thermal analysis data, the solubility of chromium in silver has been quoted as 8 pct at the monotectic temperature of -1445°C.1 This investigation indicates an extrapolated solubility of 1.38 pct which is believed to be more reliable. At the eu-

Jan 1, 1968

Discussion of the paper of S. E. Bretherton, presented at the Butte meeting, August, 1913, and printed in Bulletin No. 80, August, 1913, pp. 1481 to 1487. S. E., BRETHERTON, San Francisco, Cal.:-Since preparing my paper, our chemist, F. L. Wilson, has been experimenting with Leadville, Colo., oxidized zinc ore and Leadville zinc sulphides containing lead, operating upon 40-1b. lots at a time. His results on Leadville sulphides are as follows: The raw ore contained a trace of gold, 5.04 oz. of silver, 19.1 per cent of zinc, and 7.24 per cent of lead. After roasting the ore contained a trace of gold, 5.72 oz. of silver, 0.75 per cent of copper, 21.7 per cent of zinc, and 8.29 per cent of lead. The loss in weight by leaching was 38.8 per cent. The extraction was 42.8 per cent of the copper, 80.6 per cent of the zinc, 7.1 per cent of the lead, 0.64 per cent of the silver, and no gold. The little .silver that is extracted would be thrown down with the copper when precipitating it from the zinc and both thrown hack into the residue to be smelted. There would be no trouble in smelting the residue with nearly all the sulphur eliminated and containing less than 7 per cent. of zinc. On zinc carbonates containing silicates the results are not satisfactory. H. O. HOFMAN, Boston, Mass. (communication to the Secretary*) :¬The paper by Mr. Bretherton upon the application of ammonia and carbonic acid to the recovery of zinc as oxide from roasted sulphide ore is of interest to me, as the first installation of the Schnabel process for the extraction of zinc from a mixture of zinc oxide, lead oxide (with some cupric oxide), and pellets of enriched lead, produced in steaming liquated zinc-silver-lead crust of the Parkes process, was made at Lautenthal, Harz mountains, where I was chemist. In that capacity I carried out numerous experiments on the Schnabel process, was active in the erecting of the plant, and later had charge of it. After the Lautenthal plant had been in operation for some time a similar one was built at Hoboken-les-Anvers, but both have been abandoned, as the wet process could not compete with the distillation of the liquated zinc-silver-lead crust, now used universally in the Parkes process.

Jan 11, 1913

By A. Raman, P. P. Singh

The constitution of the La-Rh system was studied by powder X-ray diffraction, metallopaphic, and differential thermal analysis techniques and an equilibrium diagram is presented. Eleven intermediate phases occur in the system and the crystal structural data for nine of them were determined. La3Rh crystallizes in an orthorhombic structure of undetermined type, whose unit cell is obtained by doubling the &apos;a; and &apos;c,,&apos; edges of an FesC type unit cell. The other intermediate phases of the system are LarRh-3( undetermined structures also occur in the system. LaRh, undergoes a polymorphic phase transformation at 1240°C. LaRh3 and La2Rh7 also exhibit polymorphisnz. The phases Laah and LazRh7 melt congruently. The latter undergoes a eutectoid transformation into LaRh, and Rh at 1205°C. Laah3 is formed by a peritectoid reaction between Laah and La,Rh,,. The other Phases result from peritectic reactions between the liquid and the adjacent rhodium-rich phases. The intermediate Phases of the La-Rh system are compared with those of the La-Co and La-Ni systems. DURING the course of a detailed investigation to study the occurrence of CrB, FeB, A1B2, and related structures in the rare earth alloys it was found that much information is lacking for the rare earth noble metal systems. Although the structures of several rare earth alloys containing the noble metals at the AB and AB2 stoichiometries have been reported, the occurrence of related structures at other stoichiometries has not been studied. We have initiated a project to study the crystal structural features of selected rare earth-rhodium alloys and to map the equilibrium diagrams of representative systems with conventional methods. The results of our investigation in the La-Rh system are presented in this paper. Two phases were known in the La-Rh system. LaRh has the CrB-type structure.&apos; LaRhz is a MgCu2-type Laves phase.z EXPERIMENTAL PROCEDURE Alloys weighing less than 1 g were prepared from commercially pure lanthanum (99.9 pct +), supplied by Lunex Company, Pleasant Valley, Iowa, and rhodium (99.92 pct +), supplied by Engelhardt Industries, Newark, N.J., in a conventional arc melting furnace under argon atmosphere. The buttons were turned upside down and remelted three times to insure homogeneity in the samples. Since negligible loss of material was encountered during melting, a chemical analysis of the alloy buttons was not undertaken. Powder specimens for X-ray diffraction studies in the as cast state were then prepared. The buttons were wrapped in thin molybdenum foils and homogenized by heating in vacuum at suitable high temperatures for more than 1 week. They were then broken into three or four pieces for annealing experiments. The pieces were wrapped in molybdenum foils and annealed at various temperatures in evacuated quartz capsules. The annealing was carried out for 2 hr at or above 1200°C, 1 day at temperatures close to llOO°C, 2 days at 1000°C, and for 1 week at temperatures below 1000°C. After annealing the alloy pieces were again broken and powder specimens for X-ray diffraction were prepared. The powders of the lanthanum rich alloys with more than 80 at. pct La were prepared by filing. The filings were sealed in molybdenum tubings and stress-relieved at 600°C in vacuum. It was not deemed necessary to stress-relieve the powders of the other alloys, since the alloys were very brittle and were ground easily. POWDER X-RAY DIFFRACTION X-ray diffraction photographs of powders (-325 mesh size) of the alloys in the as cast and annealed states were prepared in a Guinier-de Wolff focussing camera with copper K, X radiations. These patterns were studied to identify the stoichiometries and the crystal structures of the intermediate phases. The lattice parameters of the phases were calculated after minimizing the differences between the observed sin2 6 values, calculated from the diffraction angles 8, and the sin2 8 values, calculated using approximate lattice constants obtained from a few lines. These differences were minimized manually to less than 0.0005. The latLice constants are judged to be accurate to *0.005A for values less thp about 10A and to k0.01~ for values greater than 10A. The relative intensities of the lines were calculated using a computer program written by Jeitschko and Parthk.~ No attempt was made to refine the atomic positional parameters in the phases. METALLOGRAPHY The phase equilibria in the investigated alloys in the as cast and annealed states were also studied by metallographic examination. The polished specimen surfaces were etched with 10 pct picric acid in alcohol (alloys up to 25 pct Rh), concentrated picric acid (from 25 to 37.5 pct Rh), 2 pct nital (40 to 50 pct Rh), 10 pct nital (from 50 to 66.7 pct Rh) and with concentrated 48 pct HF for the other rhodium-rich alloys. Selected microstruture~ were then photographed using a Po-laroid Land camera. THERMAL ANALYSIS Differential thermal analysis of the alloys was carried out in DTA-668 Stone differential thermal ana-

Jan 1, 1970

By J. M. GLEDHILL

IT would doubtless have been felt by many but a few years back that there was little left to be said on the subject of crucible tool-steel, and that something akin to finality had been arrived at in its manufacture and general treatment. Probably such feeling was justifiable when it is remembered that the making of steel in crucibles is by far the oldest method known, dating back from time immemorial, it being indeed impossible to accurately trace its origin and earliest development; but it seems certain that carbon-steel was made and used thousands of years ago for cutting-tools. Proof of this may be seen by the marvelous carvings and workings on the intensely hard stone-work of the ancients, for it would be difficult to conceive by what means, other than with steel tools, such work could have been executed; and it is wonderful to contemplate that steel cutting-tools should have been used so long ago, whilst the principle of manufacturing them-that is, by fusion of iron and charcoal in crucibles-was then in a measure on the sane lines as we work on at the present day. Archaeologists have discovered that the Chinese made steel in crucibles long before the Christian era. " Wootz " steel fabricated in India centuries ago was crucible steel, as was also the celebrated Damascus steel, produced at the forges of Toledo, and, curiously, this latter steel furnishes&apos; yet another proof that "There is nothing new under the Sun," for it is recorded that Damascus steel contained certain percentages of tungsten, nickel, manganese, etc., some of the very elements, in fact, contained in the present modern high-speed steel, so that a latent high-speed steel may be said to have existed

Mar 1, 1905

PRESIDENTS DAVID THOMAS 1871 R W RAYMOND 1872-1874 A L HOLLEY1875 ABRAM S HEWITT 1876 T STERRY HUNT 1877 ECKLEY B COXE 1878 - 1879 WILLIAM P SHINN 1880 WILLIAM METCALF 1881 RICHARD P ROTHWELL 1882 ROBERY W HUNT 1883 JAMES C, BAYLES 1884-1885 ROBERT H RICHARDS 1886 THOMAS EGLESTON 1887

Jan 1, 1923

By Walter Crafts, C. M. Offenhauer

IN a previous study&apos; of the carbide phase of chromium steels, it was shown that chromium carbide (Cr7C3) is a more stable carbide than cementite (Fe3C) at tempering temperatures above about 500°C. in quenched steels containing from about r1to 7.5 per cent chromium. Below this temperature the carbide phase was found to consist of cementite in all of the steels examined. In the present investigation it was found that a similar change in the carbide phase takes place in molybdenum and in chromium-molybdenum steels at approximately the same tempering temperatures. The approximate ranges of occurrence of the alloy carbides Cr7C3, Cr4C, and Mo2C have been determined in steels containing up to 5 per cent chromium, 1 per cent molybdenum and 0.60 per cent carbon. The experimental steels were made in small high-frequency furnaces from Armco ingot iron, low-carbon ferrochromium and calcium molybdate and were poured into 2-in. ingots, which were forged to 1-in. round bars. Sections of the bars were heated to the carbide-solution temperature and cooled rapidly to induce austenite transformation at relatively low temperatures. The steels were subject to some segregation, therefore comparatively high temperatures were necessary to dissolve all of the carbides. The microstructure of the quenched specimens was examined to verify the belief that the carbides had been dissolved, and also to determine the nature of the nonmartensitic structure of incompletely hardened steels. Ferrite and pseudomartensite were present in some of the specimens, but all specimens were quenched at such a rate that pearlite was not formed, in order to ensure the formation of the carbides at temperatures lower than about 500°C. It had been found previously&apos; that in chromium steels the chromium carbide could be produced by transformation of austenite in the pearlite range above 500°C., and, in order to start the tempering treatment with the carbide present as cementite, the absence of pearlite was considered to be essential. The specimens were tempered for 64 hr. to permit the formation of the stable carbide. After tempering, the specimens were submitted to microscopic and X-ray examination. The carbides were separated either electrolytically in a 5 per cent HCl electrolyte or with an acidified copper-potassium chloride solution. The crystal structure of the carbide phases was determined by the Debye-Scherrer type of X-ray patterns. The X-ray patterns obtained from specimens near the boundaries between two types of carbides were generally weak and diffuse as compared with the other patterns. For this reason, and because the more soluble components may be masked, the boundaries between the phases could not be determined precisely. However, the X-ray analysis of electrolytically separated carbides has given consistently reproducible results

Jan 1, 1943

By R. Taggart, D. H. Polonis, J. C. McMillan

Thin foil electron microscopy techniques have been used to study the micros tructure of a Ti- 7.5 wt pct vanadium alloy in the quenched and aged cond_itions. Selected area diffraction has identified a (1011) twinning mode accompanying the martens ite t9.an.sfo?rn~ation. Precipitation of the 0 phase during the aging of quenched alloys has been obsevced to occuv in the form of plates and needles with preferential growth along the (1120) direction in the a&apos; matrix. In the present work, the micro structures of quenched and aged alloys of the 0 isomorphous Ti-V system have been investigated by means of thin-foil electron-microscopy techniques. Most of the previous studies of martensitic phases in titanium-base alloys have utilized optical microscopy and X-ray diffraction techniques. The application of transmission electron microscopy is more recent, and only a few such studies have been reported.1-5 This paper is concerned particularly with the crystallographic nature of substructures formed in the martensitic a&apos; plates. The early stages of aging or tempering have also been examined to determine the modes of precipitation from the martensitic solid solutions. Such information is essential in studying strengthening mechanisms and in determining the feasibility of high-strength rnartensitic alloys based on titanium. Liu and Margolin6 reported several types of substructures within the martensitic plates of quenched Ti-Mn alloys. It was later suggested by Liu7 that the structure of martensitic plates might be very complex and not single crystals as normally assumed. The sub-markings could not be identified with the two possible habit-plane orientations {334}ß and {344}ß, and it was not positively determined whether they are associated with mechanical twinning. The recent studies by Armitage5 on titanium martensites in a variety of commercial alloys have indicated that twinning does not generally accompany the transformation. This observation is in accordance with the Mackenzie and Bowles model for the transformation in pure titanium8 which suggests that twinning is not to be expected. However, a number of other recent investigators have reported twinlike substructures in the martensitic plates of several titanium systems including commercia1 titanium,2,3 Ti-Cu,5,9 Ti-Ta,1 and the commercial alloy Ti-8A1-1Mo-1V. Blackburn* has reported that the structure of quenched Ti-8A1-1Mo-1V depends on the prior solution temperature. He identified two types of martensitic structures, one of which is described as a heavily faulted hcp or twinned fcc structure and the other an hcp structure with a high internal dislocation density. The latter type exhibited some twinning with a (1011) twin plane. EXPERIMENTAL PROCEDURES AND RESULTS The alloy used in the present work was a 7.5 wt pct V alloy that was prepared by conventional arc-melting techniques. This alloy composition was selected

Jan 1, 1968

By Francis H. Crockard

DURING the past five years we have probably witnessed greater technological advances than in any similar period. Industry and science have steadily marched ahead. The makers of iron and steel products have watched with great interest the development of stream-lined trains, faster and safer airplanes, the entirely new science of air-conditioning and the widespread use of stainless steels. Since 1930, we have witnessed the erection of the world&apos;s largest suspension bridge, the largest and fastest ocean liner and the spectacular Boulder Dam, while even greater things are planned in these particular fields. Other industries, too, have been moving ahead rapidly. This advance in a few years has been caused primarily by the fierce competition of the depression. Now, more than ever, is the emphasis placed upon quality and service. Makers of southern merchant pig iron have been no exception and have been confronted with increasingly strict chemical specifications and with much greater stress on physical characteristics than ever before. For example, sulfur in the iron must be held very low, competition demanding that it be in the 0.020&apos;s and low 0.030&apos;s; silicon content must be held within a very narrow range, and definite requirements of phosphorus and manganese are an essential part of the specifications. Physical characteristics are becoming of increasingly greater importance. The introduction of triple-skimmed, slow-cooled iron several years ago was the first step in this direction. Now there are a few customers who buy iron by chemical analysis and physical appearance, having found that the analysis alone is an insufficient guide to their requirements.

Jan 1, 1935

By Richard Birk

Rf sputtered quartz has been successfully used in multilayer interconnects for MOS devices. The target potential was 2000 v peak to peak with a frequency of 13.56 Mc per sec. The sputtering rate was approximately 60A per min with an operating pressure of argon gas at 10 to 15 p. The dielectric strength is 8.0 x 10&apos; v per cm and the pinhole density is less than 2 x 10-2 per sq mm. Infrared absorption peaks for the Si-O vibrational bands we at 9.4, 12.4, and at 22.5 p. The etch rate of the sputtered SiO2 is 1200A per min. MOS devices have been fabricated using multilayer interconnects with no degradation of electrical characteristics. MULTILAYER interconnects require an insdating film between the metal layers which can be deposited below any potential eutectic temperature. Sputtered Si4, which can be deposited with substrate temperatures below 200°C, satisfactorily meets the multilayer insulation requirements and has been used to prepare MOS circuits with multilayer interconnects. The physical and electrical properties of rf sput-

Jan 1, 1969

By Francis Crockard

DURING the past five years we have probably witnessed greater technological advances than in any similar period. Industry and science have steadily marched ahead. The makers of iron and steel products have watched with great interest the development of stream-lined trains, faster and safer airplanes, the entirely new science of air-conditioning and the widespread use of stainless steels. Since 1930, we have witnessed the erection of the world&apos;s largest suspension bridge, the largest and fastest ocean liner and the spectacular Boulder Dam, while even greater things are planned in these particular fields. Other indus-tries, too, have been moving ahead rapidly. This advance in a few years has been caused primarily by the fierce competition of the depression. Now, more than ever, is the emphasis placed upon quality and service. Makers of southern merchant pig iron have been no exception and have been confronted with increasingly strict chemical specifications and with much greater stress on physical characteristics than ever before. For example, sulfur in the iron must be held very low, competition demanding that it be in the 0.020&apos;s and low 0.030&apos;s; silicon content must be held within a very narrow range, and definite requirements of phos-phorus and manganese are an essential part of the specifications. Physical characteristics are becoming of increasingly greater impor-tance. The introduction of triple-skimmed, slow-cooled iron several years ago was the first step in this direction. Now there are a few customers who buy iron by chemical analysis and physical appear-ance, having found that the analysis alone is an insufficient guide to their requirements.

Jan 1, 1936

By H. Hooykaas

A few years ago Marsal published a method of predicting oil recovery as a function of time for an edge-water-drive reservoir with several rows of wells.&apos; The method is based solely on oil- and water-production data obtained during a short period for the first row of wells and the gross productions (past and future) of all wells. This paper presents a streamlined version of the original publication and experimental evidence showing the reliability of this version. For this purpose scaled experirnents were performed in a large model (750 X 500 X 23 cm, 24.4 ft X 16.2 ft X 8.9 in.) in which edge-water-drive reservoirs with three rows of production wells were simulated. From the results of the experiments it is concluded that the method gives a fairly accurate prediction of the future production behavior of the reservoir and enables us to select the most profitable production plan. Results obtained with the method were always slightly conservative. INTRODUCTION The calculation method developed by Marsal aims at forecasting future reservoir behavior.&apos; It can be applied to floods in which a number of down-dip wells are rapidly going to water. Essentially it is a method for predicting the future behavior of the up-structure wells on the basis of the observed behavior of the down-structure wells. As such, the method is especially useful for predicting behavior of edge-water drives toward several rows of wells. The requirement to be met for justifying the approach is identity of production mechanism in all regions of the reservoir. The attraction of the method then lies in the fact that it applies even when we have no complete description of the prevailing production mechanism. The original publication was in the German language and for this reason, perhaps, did not attract widespread attention. In our opinion it deserves to do so because it does offer in a simple way results which are in fair agreement with observed data. The aim of the present paper is to bring the method to the attention of a wider audience. The method in its most simplified form is reviewed and experimental data in support of it are presented. PRINCIPLES To show Marsal&apos;s principle. we will first write the continuity equation for very simple conditions, linear flow (in the positive x-direction) in an incompressible porous body [cross-section A(x), porosity +(x)], saturated with incompressible fluids (oil saturation S). When the oil in this formation is displaced by water, injected at x = 0 (inlet face), we get where q(t) is the total gross displacement rate and f the fractional oil flow. The equation can be solved if we make the assumptions: (1 ) that the fractional oil flow (f) is a univalent function of saturation (S); and (2) that the function is the same throughout the porous medium. Then: f = f(S) Solving this equation with the method of characteristics, we obtain

Jan 1, 1966

By S. Wiederhorn

The energy and force of interaction between a crack and a slip band have been calculated. When the distance between the crack and the slip band is greater than the dislocation spacing of- the slip band, the force on the cgrack is perpendicular to the slip band. The energy needed to propagate a crack against this force can be lawe enough to pevmit crack propagation parallel to noncleavage planes. SEVERAL recent experiments have demonstrated the occurrence of crack-slip band interaction in ionic crystals having a NaC1-type structure. Stokes, Johnston, and Lil have studied the nucleation and growth of cracks in MgO. From their observations, they conclude that slip bands have the ability to halt and stabilize cleavage crack nuclei. Gilman2 has shown that a crack, propagating along a (100) cleavage plane in a previously deformed LiF crystal, tends, at certain points within the crystal, to switch to (110) planes. He reasons that plastic deformation reduces the cohesive energy of the (110) planes below that of the (1130) planes, making possible (110) crack propagation. Keh3 has demonstrated that when indentations are made in MgO, crkcks form along (110) planes at the indentations. Gilman2 attributes the formation of these cracks to plastic cieformation. There is some doubt as to the formation of cracks on (110) planes. Kuznetsov4 in his discussion of crack propagation in rock salt, shows that splitting along (110) planes requires 1.76 times as much energy as splitting along a stepped (100) surface. Therefore, the observed (110) cracking in LiF and

Jan 1, 1962

By Horace Pops, L. L. Isaacs

SOLID-STATE transformations at low temperatures are frequently studied by optical metallography. Greninger and ooradian&apos; and Hull and Gar-wood2 have examined the martensitic-type transformations in 0-phase Cu-Zn alloys during continuous cooling below room temperature. The lowest range of temperatures in the apparatus described by these investigators is limited to approximately 110°K. Present studies of reversible first-order phase transitions near liquid-helium temperatures have necessitated the design of a new cold-stage cryo-stat for optical metallography. The features of interest of this instrument are as follows: a) temperatures down to 4.Z°K are obtainable; b) the specimen is always under vacuum, thereby avoiding the condensation of moisture; c) specimens can be rapidly cooled or heated at varying rates, thus permitting studies of transformation hysteresis. Fig. 1 shows the details of the cryostat. All major parts are made of type 304 stainless steel. Both helium and nitrogen cans are Heliarc-welded to the helium filling tube, permitting the use of a common shielding can. Hence, the over-all weight and the area of the evacuated surface are mini- mized. The copper sample holder is screwed into the bottom of the helium can. To assure good thermal contact between the sample holder and the refrigerant, an indium ribbon is placed on the threads

Jan 1, 1964

By M. N. Pathasarathi, P. A. Beck, T. J. Koppenaal

EARLIER work1 indicated that in rolled and annealed copper the volume fraction of the cube-texture component may increase on continued isothermal annealing. Merlini found2 that in rolled copper the well-known increase in the amount of cube-oriented material with increasing annealing temperature3 is concurrent with grain growth, which takes place after recrystallization is complete. It was found that both in copper2 and in aluminum4 the texture resulting from primary recrystallization, even if this is followed by some grain growth, may in general comprise, in addition to the cube-texture component, four symmetrical texture components of the general type (123)[412], i.e., components not very far from the ideal orientations of the rolling texture. The volume fraction of the cube-texture component increases on further annealing at the expense of these four components.&apos; The present work was undertaken to study quantitatively the increase in the volume fraction of the cube-texture component in the course of normal grain growth, under varying conditions. Since the grain-gr3wth behavior of A1-Mn alloys has been already investigated in some detaiL5 an A1 + 0.8 pct Mn alloy was chosen for this work, in which the extent of normal grain growth can be effectively controlled by the inhibiting effect of a dispersed second phase. EXPERIMENTAL PROCEDURE High-purity aluminum of the following composition was used: A1 99.997 pct, Cu 0.0004 pct, Fe 0.0005 pct, and Si 0.001 pct. A high-purity A1-Mn master alloy was prepared by dehydrating manganese chloride at 500°C and reacting it with molten high-purity aluminum at 750°C in a graphite crucible.5 The Mn metal reduced from the salt by the molten aluminum goes into solution in the latter. The master alloy? containing 13 pct Mn, was solidified under conditions suitable for minimizing segregation. An ingot of the final alloy, to be used in the experiments, was per pared by melting high-purity aluminum and the necessary amount of master alloy in a graphite crucible, and by pouring the molten alloy into a graphite mold preheated to 700°C through a box containing graphite baffles designed to intercept most of the aluminum oxide films floating in the melt. The melt was then solidified by placing the mold on a water-cooled copper plate and by keeping the top of the mold hot, in order to minimize piping. The 6 1/2 in. long, 11/4-in.diam ingot consisted largely of a single crystal. Chemical analysis showed that there was little segregation from the bottom (0.83 pct Mn) to the top (0.81 pct Mn). Spectrographic analysis of the ingot showed that the impurity content remained essentially the same as that of the aluminum used. The ingot was processed so as to obtain a uniformly small grain size. without banding. The ini-tial steps of rolling in the form of a flat bar and of annealing, down through a thickness of 0.4 in., are shown in Table I. At this point the bar was cut into four sections, each approximately 3 in. long, which were processed separately as shown in Table If. The table also gives the penultimate grain sizes obtained in each section by annealing prior to the final rolling of 90 pct RA. During the final rolling the strips were reversed end to end between passes. Annealing treatments were given in a salt pot furnace, with the temperature controlled to 1°C. All annealing of section "6s" was carried out above the solvus point to avoid the formation of second-phase precipitate. The penultimate annealing was sufficiently short to give a relatively fine grain size in spite of the high temperature. Section "6L" was processed in a similar manner, except that the penultimate annealing was much longer, so as to allow considerable grain growth. Section "4" was given the two final

Jan 1, 1961

By A. J. Ardell, O. D. Sherby

The elevated-temperature steady-state creep behavior of polycrystalline a Zr was studied in vacuo under constant tensile stress. The experiments were conducted from 660° to 845°C over the stress range 870 to 5400 psi. The steady-state creep rate, is, of a Zr was found to be related to the applied stress, a, and temperature, T, by the equation where C is a constant. Two interpretations of the data are considered, each requiring different values of the parameters ß, Q, and . It is shown that this formula can be derived using Gilman&apos;s phenomeno logical equation relating dislocation velocity to the shear stress for the case where dislocation mobility is limited by the nucleation of kink motion. From the fundamental dependence of strain rate on dislocation velocity and the density of mobile dislocations, the factor exp(ßo) can be related to the stress dependence of the dislocation density, and Q to the energy of kink formation. The two interpretations considered are as follows: Case I—the parameters are adjusted so that the equation above describes the data over the entire range of stress and temperature, implying that creep is controlled exclusively by dislocation glide; Case IIthe parameters are adjusted to allow for the possibility that creep is diffusion-controlled at high stresses but glide-controlled at low stresses. For both cases the values of the parameters are within the spread permitted by the data. Factors favorable to the observation of the unusual stress -dependence of i, for a Zr are discussed and it is shown that the anomalously rapid rates of self-diffusion and the low modulus on the principal glide system are primarily responsible for the anomalous creep behavior. Metal-lograflhic observations are presented and discussed with reference to the overall creep behavior. PREVIOUS investigations of the creep behavior of a Zr have been conducted from room temperature to about 500C The results of these investigations are primarily of engineering significance because of the variety of conditions under which they were carried out, and the low testing temperatures employed. Little information is available concerning the fundamental nature of the elevated-temperature creep behavior of a Zr. This investigation of the creep behavior of a Zr above half the melting temperature (the melting temperature, T,, of a Zr is estimated to be about 1860"K was initiated because there are two significant anomalies of the physical properties of a Zr: a) the rate of self-diffusion in polycrystalline a Zr is anomalously rapid: b) both Young&apos;s modulus for polycrystalline a Zr and the shear modylus as-sociated with the prismatic slip system { 1010)( 1120) (the principal slip system in a Zr 7) have anomalously large temperature variations The importance of these quantities to the understanding of creep behavior has been demonstrated, and it was hoped that this investigation would provide a critical test of the predictions based upon the analysis of creep data for other pure metals. However, as the investigation progressed it became increasingly clear that the unusual creep behavior of a Zr could not be analyzed in a conventional manner. 1) EXPERIMENTAL EQUIPMENT AND PROCEDURES Two grades of zirconium, designated A and B, were used in this investigation. The analyses, as supplied by the vendors, are presented in Table I. Material A, with a nominal purity of 99.84 pct, was supplied by the Wah Chang Corp. in the form of cold-rolled sheet from which specimens were machined without further fabrication. Material B, with a nominal purity of 99.95 pct, was purchased from the Nuclear Materials and Equipment Corp. in the form of iodide crystal bar.

Jan 1, 1968

R. L. Fleischer (Department of Metallurgy, Massachusetts Institute of Technology)— The authors have presented valuable observations of oehavior of aluminum at intermediate and high temperatures. I should like to consider their interpretation of the appearance of cross slip plane traces in sample B-3 as partly due to "bending of the crystal due to grip constraint." Because bending moments might be set

Jan 1, 1960

By P. G. Shewmon, F. N. Rhines

THE determination of the self-diffusion coefficient of magnesium has been made possible recently by discovery1-1 of a radioactive isotope, Mg28 having a half-life of 21.3 hr,1 and subject to manufacture in useful quantity. In the present research this material was condensed from the vapor phase upon a surface of high purity magnesium. The progress of diffusion of the tracer atoms into polycrystalline magnesium was followed by machining layers and measuring the change in the intensity of radiation as a function of the distance of each layer from the surface. The self-diffusion coefficient was found to be 2.1 X 10-8 sq cm per sec at 627°C, 3.6 X 10-9 sq cm per sec at 551°C, and 4.4 X 10-10sq cm per sec at 468°C; the activation energy is about 32,000 cal per mol. Experimental Procedure Since there was no other published measurement of a diffusion velocity in any magnesium-base material, is was necessary to employ a number of new experimental techniques. The short half-life of Mg28 made it necessary to complete the entire experimental procedure within three or four days. This meant that the work had to be done where a cyclotron was readily accessible and that all operations, prior to the diffusion heat treatment, had to be so designed as to minimize their time requirements. Unusual problems were imposed also by the chemical reactivity of magnesium, its high vapor pressure, and the fact that no satisfactory method for elec-trodepositing magnesium on magnesium is presently available. Finally, the machining and handling of the easily air-borne radioactive-magnesium chips involved certain health hazards, resulting in the need for further experimental restrictions. Preparation of Mg28 The Mg28 was produced in the Carnegie Institute of Technology syncrocyclotron by the neutron spallation of chlorine.5 his involved bombarding a 2 gram crystal of high purity NaCl with a beam of 350 mev protons for a period of 2 hr, after which the crystal was dissolved in warm water and the Mg28 was concentrated and purified by chemical means (see Appendix). About 50 microcuries of Mg28 thus were obtained in the form of magnesium oxinate (8 hydroxyquin-olatc?), which was ignited in air to produce MgO. This in turn was reduced to magnesium metal vapor, by the method of Russell, Taylor, and Cooper," in the vacuum apparatus shown schematically in Fig. 1. Here the essential part is a tantalum ribbon, slightly dished to receive the MgO. The ribbon, pre- viously outgassed at high temperature, is heated to about 1700°C by passing an electric current through it, whereupon tantalum oxide is formed, magnesium vapor is released almost instantaneously, and condensed partly upon the diffusion sample. Diffusion-Sample Preparation: Hot-extruded magnesium rod, 21/32 in. round was used in making the diffusion specimens. The magnesium analyzed as follows: 0.004 pct Al, 0.027 pct Fe, 0.040 pct Mn, 0.0004 pct Cu, 0.0002 pct Ni, and less than 0.01 pct Ca, 0.0004 pct Pb, 0.0011 pct Si, 0.001 pct Sn, and 0.001 pct Zn. A brief study of the crystal texture of this material revealed a sharp fiber texture with the (001) plane roughly parallel to the extrusion axis. Cylindrical samples 1/2 in. long by 5/8 in. were machined from this rod, the end faces dressed on 3/0 emery, and lightly etched with 20 pct HC1 in water. These samples then were annealed for at least twice the intended time of diffusion, at the intended diffusion temperature, in order to stabilize the grain structure at about 1 mm average diameter. The annealing treatments were conducted in argon in the same apparatus and in the same manner as the subsequent diffusion treatments, which will be described presently. Thus, a strain-free plane surface was produced, but there remained a layer of MgO which had largely to be removed before the layer of Mg28 was deposited. Most of this layer was taken off by two light passes over 3/0 emery paper. The balance of the oxide and a thin layer of metal were then removed by etching 5 to 10 min in 4 pct nital (4 pct HNO3 and 96 pct ethyl alcohol) made with absolute alcohol. There followed immediately three quick rinses in: 1-49 1/2 pct methanol, 49 1/2 pct acetone, and 1 pct formic acid, 2-50 pct methanol and 50 pct acetone, and 3-pure benzene. This procedure is essentially that of Sturkey.7 The resulting surface, which was of almost elec-tropolished brightness, remained plane and was free of cold work. It could be kept clean by storing under benzene, or in a desiccator; short exposure

Jan 1, 1955

By B. F. Decker, D. Harker

The recrystallization reaction in OFHC and spectroscopically pure copper has been followed by X ray diffraction determinations of the amount of material with the cold-worked and recrystallized textures in specimens which had been given various heat treatments. The heats of activation for recrystallization are found to be: 29.9 kcal per rnol for OFHC copper and 22.4 kcal per rnol for spectroscopic copper. A reliable and speedy method for measuring the amount of recrystallized material in a piece of rolled metal, together with a new scheme for low temperature heat treatment, have made possible a determination of the activation energy for recrystallization in rolled copper. This method for studying recrystallization rates is different from others reported in the literature.&apos; Oxygen free, high conductivity copper—OFHC copper—of purity 99.98 pct gave an activation energy for recrystallization of 29.9 kcal per mol with recrystallization data taken in the temperature range 208-245°C. Copper of purity 99.999 pct from the American Smelting & Refining Co. gave an activation energy for recrystallization of 22.4 kcal per mol with recrystallization data taken in the temperature range 43-135°C. The great change in activation energy for recrystallization as the small amount of impurity in the metal is decreased suggests that the motion of grain boundaries is conditioned by the impurities which must be concentrated in them. Thus the OFHC copper recrystallizes with an activation energy of the order of magnitude to be expected for the diffusion of the kind of impurities likely to be present. The high purity copper, on the other hand, recrystallizes with a much lower activation energy, in accord with the notion that the grain boundaries need not wait for the impurities to diffuse with them. Copper when strongly cold-rolled has a well-defined crystallite texture which can be described by saying that the [111] direction of the face- centered cubic crystals is aligned approximately in the cross-rolling direction and that faces of the form {110} lie approximately in the rolling plane. When such rolled copper is completely annealed the crystallite texture is quite different: the [loo] direction is in the cross-rolling direction and faces of the form (100) are in the rolling plane. At intermediate stages of annealing, the copper contains both textures. If a Geiger counter X ray diffraction spectrometer is set up so as to measure the intensity of the (200) Bragg reflection from the rolled surface of a piece of copper strip, this intensity increases from a very small value in the cold-worked state to a maximum as the specimen is annealed. The same would be true for (200) reflections from planes normal to the rolling or cross-rolling directions. In view of this fact, the amount of material in the copper which has recrystallized can be measured by comparing the intensities of one of these reflections just mentioned with the intensity from a similar specimen after complete annealing. In much the same way the intensity of (111) reflections from planes normal to the cross-rolling direction could be used to measure the amount of unrecrystallized material in the specimen. In the work to be described here the amounts of recrystallized and unrecrystallized material in partly annealed rolled copper specimens were determined in this way. Experimental Procedure Samples of oxygen free high conductivity (OFHC) copper, cold-rolled 99.7 pct to 0.002 in. thickness, were subjected to heat treatments for various chosen times at four different temperatures, and high purity copper samples, cold-rolled 98.0 pct to 0.0015 in. thickness, at six different temperatures. Each sample, after treatment, was placed in an orientation sample holder for use with a Geiger counter X ray spectrometer and a reading was taken of the intensity of the (200) reflection from planes perpendicular to

Jan 1, 1951

By H. Papazian, Robert A. Lad

SINCE it can be shown both theoretically&apos;&apos; : and experimentally that disordered Cu3Au has a greater internal energy than the ordered alloy, it might be expected that its chemical reactivity would be greater. Experiments were initiated&apos; to test this hypothesis. This note will describe some recent results obtained in the study of the rates of corrosion of ordered and disordered polycrystalline Cu3Au by aqueous FeCl., The work of Robertson&apos; has indicated that the attack by such solutions is intergran-ular and highly orientation dependent. It was of interest to see whether the order-disorder effect was great enough to overcome the variations from specimen to specimen which would arise from this type of attack. The polycrystalline alloy specimens (0.020 in. wire, 1¾ in. long) were preannealed in ½ atm H, at 700°C for 2 hr. The subsequent heat treatments to produce order and disorder were performed in M atm He. Disordered specimens were obtained by heating at 450°C for 2 hr and quenching in water. Ordered specimens were produced by heating for 2 hr at 450°, furnace cooling to 380°, and soaking at 380°C for 80 to 90 hr. Room temperature was reached by cooling to 150" at a rate never exceeding 7" per hr, and water quenching from 150°C. The specimens so produced exhibited electrical resistivities characteristic of the two states; of the order of 11.27x10 ohm-cm for the disordered, and 4.28x10 ohm-cm for the ordered specimens. The corrosion rates were studied at 30°C in ferric chloride solutions of several concentrations (?,¼, 1.5, and 1 molar). The specimens were attached to a stirrer during the tests. Tests were interrupted at intervals, and the specimen weight change was determined after washing and drying. Analysis of the solutions showed that the attack was limited to the copper. The higher corrosion rate for the disordered alloy, as illustrated in Fig. l, is typical of the results obtained. The effect of variation in ferric chloride concentrations is indicated in Table I for typical specimens. The variations in the ratio of reactivity for disordered and ordered specimens at the different concentrations are within experimental error, and the best value is estimated to be near 1.4. The experimental variation was due in part to the mode of attack, as mentioned, and in part to the fact that films were formed on the specimens under some conditions. Electron diffraction patterns of the films were very complex and indicated the presence of metallic iron, as did chemical analysis. In an attempt to eliminate the difficulties arising from the mechanism of attack, single crystal specimens were used. A greater reaction rate was again measured for the disordered specimens, but the reproduci-bility was no better than that for the polycrystalline metal. This would indicate that film formation was primarily responsible. It was concluded that the reaction with aqueous ferric chloride served to demonstrate that the reactivity of disordered Cu3Au is greater than that of the ordered alloy. However, the difficulties with film formation did not permit determinations of a precision great enough to warrant extension of the study to partially ordered alloy systems.

Jan 1, 1957