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By R. B. Full

DURING the summer of 1949, the United Nations Scientific Conference on the Conservation and Utilization of Resources met at Lake Success. As summed up by one writer, the purpose was: "That everyone will try to learn how to feed more people from an acre of land." That objective is a stimulus to the phosphate rock industry as well as other industries. Phosphate rock was mined first in the Canadian province of Quebec where it was found as an apatite in pockets. In 1863, mining began near the Rideau Canal, first in open trenches and quarries until shafts were sunk. These operations were difficult since hard rock had to be drilled and blasted. Steam was the source of power. The phosphate rock was hand-selected and recovery ranged from 6 to 10 pct, averaging 70 pct to 85 pct bpl (bone phosphate of lime). Costs, f.O.b. Montreal, were $14 a ton, and selling prices averaged $17 a ton. In 1885, about 29,-000 tons were sold on this basis, but by 1892 the output fell to 8000 tons at an average price of $15 a ton at Montreal. South Carolina Deposits Phosphate rock was first discovered in South Carolina in 1837, but a mining company was not formed until 1867. It was immediately successful, and the industry continued to develop and ship an increased tonnage of land and river rock, shipments increasing from approximately 20,000 tons in 1868 to 537,000 tons in 1890, when the shipments began to decrease. In 1892 only 350,000 tons were shipped. In this field there were two types of deposits: land rock and river rock. More land rock tonnage was mined until about 1887 when river rock production began to equal it. The land rock was mined adjacent to and in the river marshes by pick, shovel, and wheelbarrow. Each worker dug daily a space 6 ft wide, 15 ft long and 6 ft deep, which was overburden, and then loaded about 3 cu yd of phosphate rock into wheelbarrows. His helper would wheel the barrows to small steam railroads up to 250 ft distant. Water was kept from the pits by steam pumps connected to railway cars. The phosphate rock was loaded into railroad cars, hauled and dumped into a log washer to which water was added, and run over trommel screens to eliminate the clay and loose sand. Steam shovels for loading were introduced about 1891. River rock was mined by floating dipper dredges and at low tide was loaded into barges by laborers using oyster tongs, the large lumps being hand-loaded. Barges and bateaus were unloaded by wheelbarrows which were wheeled to the same type of land rock washers. The washer product was piled with stacked pine cordwood in alternate layers and fired for drying. Usually this was done under sheds along the river bank. The dried rock then was transferred to ships, usually three or four-masted schooners. Cost of production of both land and river rock, including $1 a ton royalty paid to the state for state-owned property, was approximately $4.25 a ton averaging 60 pct bpl. The selling price was approximately $7 a ton f.O.b. In 1886 the courts handed down a decision permitting anyone to mine, and many additional companies were organized immediately. As a result, the selling price was lowered to a profit margin of 50Ø or less per ton. The South Carolina mines were eventually shut down when a more profitable source of phosphate rock was discovered in Florida. Florida Deposits Phosphate rock was discovered in Florida in 1882, and in 1883 a small quarry was opened near Hawthorne. In 1886 phosphate rock of good grade was discovered on Peace River, and the Arcadia Phosphate Co. was formed. In 1888 the company began shipments amounting to 3000 tons that year. During the latter part of 1888, hard-rock phosphate was discovered near Dunnellon. Exploration extended rapidly, and deposits of hard-rock phosphate were mined over a wide area so that shortly after 1900

Jan 1, 1952

By Peter Beardmore, Michael B. Bever

The effects of the temperature of deformation and the degvee of deformation on the annealing spectrum of a Au-Ag alloy have been determined by vesistance measurements. Specimens were deformed in torsion at 4°, 78K and room temperature up to strains nd/l of about 1.0 and were isochronally annealed. Discrete annealing stages similar to those found for pure metals were observed. Stage A occurred below room temperature; the decrease in reszstance associated with this stage increased as the temperature of deformation was lowered or the strain was increased. Stage B manifested itself as an increase in resistance; this is attributed to the restoration of short-range order promoted by vacancy or vacancy aggregate migration. At temperatures above 48S0K, recrystallization or Stage C occurred. Stages A, B, and C in the annealzng curves of this alloy may be considered to correspond to Stages II, Ill, and V, respectively, in the terminology used to describe the annealing spectrum of pure metals. ThE annealing behavior after deformation below room temperature has been thoroughly investigated on pure metals, in particular gold and copper. The occurrence of several discrete stages in the annealing spectrum of these metals has been established, but the interpretation of some of these stages in terms of specific mechanisms is still controversial.' The annealing of alloys deformed below room temperature has been the subject of only a few investigations. In the investigation reported here, a Au-Ag alloy was annealed after deformation at and below room temperature. This continues a research program on the deformation and annealing behavior of Au-Ag alloys.2-8s EXPERIMENTAL PROCEDURE An alloy rod of composition 82.6 wt pct Au-17.4 pct Ag was reduced to wire, 0.0428 in. diam, by swaging and drawing. After a recrystallization treatment, the grain size was 0.031 mm, as obtained from the average linear intercept.7 The ends of wires about 6 in. long were flattened and sealed into short brass tubes with cold setting cement to form the grips of the torsion specimens. The equipment used for deformation has been described elsewhere.6 The specimens were deformed at room temperature, 78", and 4°K to various strains nd/l, where n is the number of turns, d the diameter of the specimen, and 1 its length between grips. The strain rate ?max was 0.1 min-' at the sur- face of the specimen except where noted otherwise. The specimens were annealed isochronally for 30 min at intervals of approximately 30°K from the temperature of deformation to approximately 800°K, except that after deformation at 4°K the lowest annealing temperature was 78°K. A potentiometric method was used to measure the resistance: all measurements were made at 78°K with the exception of some measurements at 4°K mentioned below. Data obtained with a single specimen yielded a complete annealing curve: in several cases duplicate specimens were used to confirm the reproducibility of the curves. Since the surfaces of the specimens were rough after deformation, their diameters could not be measured accurately. The resistance increments after deformation and annealing are therefore reported as percentages of the resistance of the same specimen measured after completion of the annealing, that is, after full recrystallization. The corresponding resistivity values can be calculated by obtaining the effective diameter of a deformed specimen from the ratio of its resistance after recrystallization to the true resistivity, which was determined with a smooth specimen of known dimensions. This method of calculation was used to obtain data shown in several of the figures. RESULTS AND DISCUSSION The isochronal annealing curves after deformation to various strains at 4", 78K, and room temperature are shown in Figs. 1, 2, and 3, respectively. The resistance increment of the as-deformed specimens as a function of strain is shown in Fig. 4. For the specimens deformed at 4°K. values measured at 78°K are plotted with the exception of those for the as-deformed

Jan 1, 1970

By H. E. Swanson

A description of the field apparatus has been published by D. G. Brubaker. Data from laboratory model studies of the in-line and broadside methods of operation are detailed. The conductor models include a single sheet conductor at several strike and dip angles and a schistose-type conductor. In addition, data on the effects of the strike length and of the depth of the conductor are presented for the broadside method of operation. All of the data show that the inclintion reverses direction over the top edge of dipping single sheet conductors. Differences between anomalies over conductors dipping between 90" and 30" are subtle,but flat-lying sheets can be readily distinguished from steeply dipping sheets. Schistose conductors are easily distinguished from single sheet conductors, and a procedure is given for determing the direction of schistosity. The depth of penetration under ideal conditions is 0.7 of the coil separation. In 1950, The New Jersey Zinc Co. (of Pa.) developed a lightweight apparatus for electromagnetic prospecting. A description of this apparatus has been published by D. G. Brubaker.1 It consists of a small, battery-powered vertical source coil and a direction-finding receiving coil. Source and receiver are kept a fixed distance apart and move as a unit along the survey lines. At each receiver station the inclination of the magnetic field created by the source is measured. Two methods of operation are used, in each of which measurements are taken with a constant source-to-receiver distance and with the plane of the source oriented to pass through the receiver station. In one procedure, called the in-line method, source and receiver travel in tandem along the same line. In the other procedure, called the broadside method, source and receiver move along adjacent lines, with the source-to-receiver direction oriented parallel to the assumed strike direction. The data recorded in either case are the inclinations, measured in degrees from the horizontal, of the magnetic field created by the source. These data are plotted at the receiver stations on a map. Scale-model studies of both methods of operation have been conducted in the laboratory and are reported in this paper. APPARATUS Small source and receiving coils wound with many turns of fine copper wire were used to produce and measure an alternating magnetic field in the same manner as the field apparatus. The model source coil was mounted in a vertical position and oriented so that the receiving coil was in the plane of the source. The receiving coil was mounted in a frame so it could be rotated to obtain the inclination measurements. The source coil was excited by a 100-kc signal generator. A tuned amplifier and a voltmeter were used to detect the null position of the receiving coil. The coils used in the field apparatus are usually spaced about 400 ft. The coils used in the model work were spaced from 4 in. to 4 ft apart. Thus the scale of the model work ranges between 1/1200 to 1/100 natural size. The model conductors were aluminum-coated building paper supported on wooden frames. The conductivity-thickness product was measured using a four-electrode system and was found to be about 300 ohms-'. CONDUCTIVITY CONSIDERATIONS The electromagnetic response from a thin sheet conductor is a function of the product of the conductivity of the sheet times the thickness of the sheet times the frequency of the alternating electromagnetic field times some linear dimension of the system such as the coil separation which indicates scale. The secondary electromagnetic field produced by the conductor increases in amplitude from zero at zero conductivity to a maximum value for infinite conductivity, while the phase of the secondary field with respect to the primary field shifts from 90" at zero conductivity to 0" at infinite conductivity. Several authors2-' have shown that most of both the increase in amplitude and the change in phase of the secondary field takes place over a three-decade range of the conductivity-thickness-frequency product. Only small changes in the amplitude and

Jan 1, 1961

By Margaret V. Doyle, R. W. Powers

INTERNAL friction measurements, carried out as functions of temperature, have been used extensively to obtain data on the mobility of interstitial impurities in the Group V metals, vanadium, colum-bium, and tantalum. For a given frequency of vibration, the internal friction peaks at a temperature which depends principally on the particular metal and interstitial impurity. Relaxation times and corresponding activation energies are determined from a knowledge of such peak temperatures for various vibration frequencies. These experimentally determined relaxation times are related to the average jump frequencies of the interstitial impurity atoms using the model of Snoek1 and can be related to macroscopic diffusion coefficients as shown by Polder.2 In these Group V elements, the impurity whose diffusion rate has been the hardest to pin down by the internal friction technique has been carbon. The experimental difficulties associated with carbon in these metals were not evident in the earliest investigation carried out by Ke on tantalum.3 The carbon peak was thought to occur near 140°C at 1 cps on the basis of the enhancement of the height of this peak measured on specimens which had been coated with colloidal carbon and heated in vacuo at 1200°C. Presumably on the basis of similar evidence, Wert identified a peak also near 140°C at 1 cps in columbium with the diffusion of carbon in that element." Later, the difficulties of admitting carbon into these metals without simultaneously introducing oxygen were pointed out by Ang and Wert- nd by Marx, Baker, and Sivertsen.6 The latter suggested that the peaks attributed to carbon in tantalum and columbium might arise from the presence of traces of oxygen, adventitiously absorbed during the carburization. Such suspicion was confirmed, at least in tantalum, when it was found possible to prepare C-Ta specimens without contamination by oxygen. Specimens sci prepared, aitnough high in carbon content, gave no evidence of a damping peak near 140°C.7 Recently, the authors were able to Show that carbon diffuses in tantalum at a rate only slightly slower than that for nitrogen.8 At 1 cps, the carbon peak lies at 34g °C, only a few degrees above the nitrogen peak at 341 °C. However, the two could be distinguished by a different stability with respect to aging at 400 °C and by a different dependence of peak height on solute concentration. On the basis of this experience with tantalum, it appeared worthwhile to search also for a carbon peak in columbium in the vicinity of the nitrogen peak." In this report, the location of the carbon peak in columbium is made known and its position relative to the nitrogen peak is compared with the authors' previous findings in tantalum. In addition, a second source of internal friction seen only in specimens of very high purity is discussed. Columbium was obtained from the Fansteel Metallurgical Corp. as either 30 or 40 mil diam stock. The manufacturer gives the following analysis as typical of this material: 99.4 pct minimum Cb, 0.10 pct maximum C, 0.5 pct Ta, 0.04 pct Ti, 0.02 pct maximum Si, and 0.01 pct maximum Fe. Wires 13 in. long were degassed near 2200°C for 4 to 16 hr in an ultimate vacuum better than 2.10" mm Hg in an apparatus described previously.8 This treatment reduced the oxygen and nitrogen damping peaks observed in the as-received material to less than 0.0001 in Q-1 above background. The latter value varied somewhat from specimen to specimen in the range 0.0004 to 0.0008, depending chiefly on the wire diameter and the frequency of vibration. The high temperature evacuation left the specimen in a rather soft condition, the Vhn being reduced to 40 from a value near 121 in the as-received conGi~iur~. The gram size was about the diameter of the specimen. After outgassing, specimens were loaded with carbon by absorption of such hydracarbons as methane or ethylene. Carbon was also added to a few speclrnens by heating them at

Jan 1, 1958

By R. E. Ogilvie, E. T. Peters

The X-ray Kossel-line method has been used preaioz~sly for measuring lattice parameters to accuracies of 1 part in 100,000.5 A second application of this method is described for determining the crystallographic orientation of a randomly positioned single-crysta1 spherical volume that can he as small as 50 µ in diameter, within accuracy limits of ±1/2 deg. The theory, experimental procedure, and interpretation of Kossel-line patterns and an experimenta1 verification of the predicted orientation relationship of the cph k and fee a Cu-Si phases are presented. This orientation relationship can be described as (111)a, (00.1)k; [110]a [11.0]k. In addition, the lattice parameter of the a phase was found to he a. = 3.62154 ± 0.00014Å. The Kossel-line method when used in conjunction with electron microanalysis is shown to he capable of providing a complete chemical and structural analysis of a given crystal. THE most easily accomplished methods for determining the orientation of a single crystal are variations of the Laue X-ray diffraction method. Although certain materials can be oriented to within ± 1 deg accuracy by observation of exterior macroscopic features, such as etch pits, cleavage faces, or growth features, the Laue method is generally preferred for routine laboratory application. Both back-reflection and transmission patterns are coordinated by appropriate reference charts and are plotted in terms of reflection plane poles (normals) on a stereographic projection. Orientation is deduced by relating the pole distribution (which is fixed by the crystallographic symmetry of the specimen) to two specified external reference directions. The Laue method has several limitations. 1) Orientation can rarely be determined to better than ±l deg accuracy. Principal errors involve inaccuracy in specimen-to-film distance, measurement confidence of individual Laue spots, and stereographic plotting. 2) Because of the indeterminacy of the X-ray wave length diffracted to a given spot, it is not possible to determine supplementary crystallographic data, such as interplanar spacings and lattice parameters. 3) The method is generally limited to specimens of cubic symmetry or to specimens of high symmetry and known structure. 4) The relatively large cross-sectional area of the incident X-ray beam generally precludes the measurement of relative grain orientation in a polycrystalline material.* _____ Several of these limitations can be overcome by application of the Kossel-line method, which has been previously employed for precision lattice-parameter determinations.'-= For this method, a point source of divergent monochromatic X-rays is generated within the crystal by means of an incident electron beam. The divergent X-rays which fulfill the Bragg law are diffracted by the specimen and are recorded on a film placed either in transmission or back reflection. Analysis of the film yields a direct measurement of orientation and lattice-spacing values to an accuracy of ±1/2 pct. As analyses can be obtained from spherical specimen volumes as small as 50 µ in diameter, the method provides a means for structural analysis of second-phase or impurity precipitates within a given matrix. The primary limitation of the Kossel-line method is the requirement for an electron microanalyzer or similar apparatus capable of producing a finely focused electron beam. This paper is designed to present the theory, experimental procedure, and geometrical interpretation of Kossel patterns. The experimentally determined orientation relationship between the k and a phases occurring in the Cu-Si system and a precision measurement of the a lattice parameter are presented as a practical application of the method. THEORY OF KOSSEL LINES Divergent X-ray beam photography utilizes an effective point source of characteristic X-rays which, when diffracted from a single crystal, form numerous diffraction and absorption cones that are recorded on film.7 The cones generated from a source lying within the crystal are called Kossel lines.' Although the X-ray scattering from a divergent point source contained within a crystal is described in terms of Laue dynamical theory,9 the directions of the diffracted spectra can be ade-

Jan 1, 1965

By F. V. Lenel, G. S. Ansell, J. A. Dromsky

The growth of aluminum oxide particles in a nickel matrix was studied eve?. the temperature vange of 2140° to 2470°F. The instability of the dispersed alumina was shown to be independent of the crystal structure of the alumina. The activation energy for the growth of the dispersed alumina was found to be 84.7 1 2.0 kcal. The particle radius increased as a function of time. These results indicate that the growth is not diffusion controlled. It is believed that the rate controlling mechanism is the dissolution of the aluminum and oxygen atoms into the nickel lattice. THE strength properties of alloys which consist of a finely dispersed second phase in a metallic matrix depend upon the spacing between the second phase particles. It is therefore desirable to achieve very fine dispersions in these alloys. Furthermore, to retain the properties these very fine dispersions must be stable during fabrication and service of the alloys. The best known of the dispersion strengthened materials, the SAP type alloys, which consist of a dispersion of aluminurn oxide in aluminum, have exceptionally good stability up to the melting point of aluminum. There is evidence, however, that other dispersion strengthened alloys, even those consisting of refractory oxides in a metal matrix, may be less stable. This investigation is concerned with the stability of Ni-Al2O3 alloys in the temperature range in which these alloys are usually fabricated. The mechanical properties of Ni-Al2O3 alloys at elevated temperatures have been previously investigated by Crelnens and rant,' and Gregory and Goetzel.2 The behavior of these alloys in stress rupture tests appears to indicate that at temperatures below 1800°F they are highly stable. There is some doubt, however, as to their stability at the higher temperatures used during the conventional fabrication. Cremens and Grant, in preparing their test alloys, cousolidated, by powder metallurgical techniques, nickel powders as fine as 0.13 µ diam and alumina powders as fine as 0.018 µ. Metallo- graphic examination of the alloys following fabrication revealed that none had interparticle spacings of less than 2 µ. Considering the size of the original component powder particles, it is likely that the dispersions coarsened during fabrication. Gregory and Goetzel, in their studies of extruded alloys of 80 pct Ni—-20 pct Cr matrixes with nonmetallic dispersion, observed a definite coarsening of the alumina dispersions in alloys sintered at 2280°F as cantrasted to those sintered at 2000oF. Similar observations on the spheroidization and growth of thoria particles finely dispersed in a nickel matrix were made by D. K. Worn and S. F. Marton.3 As a result of such coarsening, much of the effort expended in the preparation of very fine powder mixtures would be lost. The mechanical properties of the alloys which had experienced coarsening would be expected to be poorer than if the original dispersions had been retained. EXPERIMENTAL PROCEDURE Ni-Al2O3 alloys were produced from powder prepared by a coprecipitation technique. Aluminum hydroxide and nickel hydroxide were coprecipitated from chloride solutions of the metals. The mixed hydroxides were calcined to form metal oxides and the nickel oxide in the mixture was selectively reduced to nickel by treating it in hydrogen. Specimens were compacted from the resultant powder which consisted of a fine, uniform mixture of aluminum oxide and nickel particles. The compacts were sintered by resistance hot pressing4, a densification technique which requires exposure times of the order of only a second or less at elevated temperatures. A conventional sintering process was not used, since the temperature required for densification would have to be in the region in which the stability of the dispersion was to be studied. A series of hot pressed specimens were treated in hydrogen at temperatures from 2140o to 2470°F (1171o to 1355oC, for times up to 120 hr. Changes in the microstructures were studied by electron microscopy using the two-stage preshadowed carbon replica method.5 In performing a lineal analysis on a series of micrographs from each specimen it was found more convenient to determine the mean free path between alumina particles rather than particle radii as the parameter of growth. Although these quantities are directly proportional for only spherical particles, the alumina particles in these alloys

Jan 1, 1962

By W. C. Ellis, M. E. Fine

WHEN certain binary Fe-Ni alloys are worked cold and then stabilized by a stress-relief anneal, their Young's moduli are nearly invariant over a substantial temperature range determined by composition and work-anneal history.' In the present investigation addition of molybdenum to such alloys (besides increasing the yield strength) was found to diminish the sensitivity of the thermal coefficient of Young's modulus to composition changes. Such alloys are of interest for applications requiring 1 a metal with controlled, low temperature coefficient of Young's modulus and substantial magnetic permeability.' Young's modulus ordinarily decreases with rising temperature. In ferromagnetic alloys a change in modulus on heating due to loss of ferromagnetism modifies the .temperature dependence. Far below the Curie temperature ferromagnetism alters the modulus in two ways: a—Addition of the energy of magnetization, a function of interatomic distance, changes the relation between interatomic energy and distance.' This lowers the modulus in the Fe-Ni and Fe-Ni-Mo alloys of this investigation. b—An applied stress changes the ferromagnetic domain arrangement so that linear magnetostriction contributes to the strain and further lowers the modulUs.~ Since in the alloys of this investigation both effects lower the modulus, loss of ferromagnetism on heating changes the slope of the modulus-temperature curve in a positive direction. With increasing temperature the slope of the modulus-temperature curve, due to progressive loss of ferromagnetism, becomes less negative, goes through zero (the modulus is minimum), becomes positive, and resumes its normal negative value above the Curie temperature.' The increase in modulus and decrease in curvature in the modulus-temperature curve occurring on work hardening have been attributed to a reduction in effect (b).'. Vn the absence of ferromagnetism work hardening through introduction of residual strains would be expected to decrease the modulus.' Alloy Preparation—Test Methods Four series of alloys having nominal molybdenum contents of 5, 7, 9, and 10 pct and containing 47 to 58 pct Fe (the analyzed compositions are given in Table I) were melted in air and cast as bars weighing from 2 to 6 lb. The charge consisted of electrolytic nickel, Armco iron, molybdenum (99+ pct) or ferromolybdenum (62 pct Mo), manganese, and aluminum as a deoxidizer. The bars were swaged cold to the desired diameters with intermediate anneals. Alloys in this composition range are reported" to be face-centered cubic and ferromagnetic at room temperature. The Curie temperatures" decrease with molybdenum. For example, keeping the Ni/Fe ratio at unity and increasing the molybdenum from 0 to 17 pct decreases the Curie tempera- ture from approximately 500°C to room temperature. (The 17 pct Mo alloy requires quenching to retain a single phase.~) Young's modulus at a series of temperatures from —50" to 100°C was determined by a dynamic method previously described.', " In this method Young's modulus is calculated from the resonant frequency (in this case approximately 50,000 cycles per second) of a rod sample (approximately 2 in. long) vibrating longitudinally in the first mode. The densities of the samples at room temperature, Table I, were determined by the standard method of weighing in air and reweighing in redistilled bromobenzene. The linear coefficients of expansion, Table I, needed to calculate the sample length and density at the temperature of measurement, were obtained by observing the change in length (22" to 100°C) of an 8 in. gage distance with a two microscope cathetometer. The values of density and coefficient of expansion in the few samples checked are independent of treatment to the accuracy reported. This was assumed to be general for all the samples. The modulus values are accurate to &0.05x101' dynes per sq cm. However, values of thermal change in modulus are accurate to 20.005~ 10" dynes per sq cm. Young's Modulus Values at 25°C: The modulus values of O', 5, and 10 pct Mo alloys at 25 °C are plotted in Fig. 1 as functions of nickel. After working cold, the samples were either recrystallized at 950" to 1000°C or given a low temperature stabilizing anneal at 400° C. The 400°C anneal does not reduce the hardness nor cause recrystallization. Annealing at 400 °C does

Jan 1, 1952

By C. B. Daellenbach, T. D. Tiemann

The prevention of quartz flotation by chelation of activating ions with the tetrasodium salt of EDTA in oleic acid flotation was investigated using a small vacuum flotation cell. It was found that flotation of quartz could be entirely suppressed in the presence of activating ions during complete flotation of hematite. It was further shown that the concentration of activating ions liberated from Wisconsin Gogebic taconite is sufficient to float quartz in an oleic acid system. Large tonnages of low grade siliceous iron ores occur on the iron ranges of the Lake Superior district. Although the lower grade nonmagnetic types are currently not of economic importance, they constitute a large future potential source of iron and steel, and hence have been the subject for extensive research on their beneficiation. TACONITES Of particular interest to the Dept. of Minerals and Metals Engineering at the University of Wisconsin have been the low grade nonmagnetic taconites of the Wisconsin Gogebic Range. These occur as five members of the iron formation and have been described in detail elsewhere.' Essentially they consist of an intimate mixture of hematite, goethite and quartz, with minor amounts of siderite, magnetite, iron containing silicates, and manganese minerals. A typical mean analysis based on the Bureau of Mines analyses' is: SiO2-52.50%, Fe2O3-43.40%, (Fe-30.38%), A12O3-1.15%, P2O5-0.07%, Mn2O3-0.29%, CaO-0.07%, (S-0.03%), Loss on ignition-2.00%. The quartz in these ores is of a microcrystalline variety,2 and so intimately associated with the iron oxide minerals that extremely fine grinding, in the order of minus 325 mesh, is required for liberation. During grinding and subsequent flotation, the concentration of quartz activating ions in the pulp becomes sufficiently high to activate quartz which floats along with the iron minerals in the presence of oleic acid. cook3 and co-workers concluded that oxidized iron ores released appreciable quantities of quartz-activating ions in water pulps. This is shown to be true in the present investigation. As a matter of fact, application of this phenomenon, as well as the intentional addition of activating calcium ions, is made use of in the an ionic flotation of silica from the iron oxide minerals.4,5 In view of the above, it becomes almost impossible to effect a sharp separation of quartz in the oleic acid flotation of the iron oxide minerals from these taconites, even with satisfactory liberation. This fact led to the current investigation to find a means of linking the activating ions with a chelating agent. The aim was to establish basic data for the quartz-water-oleic acid system which might not only be applicable to iron ores, but also to the flotation of other siliceous ores as well. EXPERIMENTAL METHODS Vacuum Flotation Cell: The vacuum flotation used in this investigation is similar to that used by previous investigators.6,7,8 The cell consisted of a 100 ml graduated cylinder with the lip removed. Vacuum from a water aspirator was applied to the cell through a trap; suitable tubing, and a rubber stopper inserted in the top of the cell. Constant absolute pressure was maintained by means of a manometer connected to the trap with a tee. Flotation: The required amount of activating ion solution was pipetted into a beaker. This was then diluted to 99.0 ml with water having a pH below that desired for the test, or so adjusted with HC1. After addition of 1.0 ml of oleic acid solution (2.0 mg oleic acid per ml), the pH was readjusted to the desired value with NaOH, and the entire solution transferred to the cell containing 1.5 gm of quartz. After vigorously agitating the cell and its contents for 2 min, the vacuum was applied to the cell and the estimated degree of flotation observed for a 3-min period. Activating ions studied included Ca++, Mg++, Al+++, and Fe+ + + added as chloride solutions. In addition,

Jan 1, 1964

By G. C. Rauch, T. H. Courtney, J. Wulff

The superconducting behavior of Nb-Ti alloys containing 40 and higher wt pct of Nb and variable oxygen content was studied as a function of thermomechanical processing. Critical current density (Jc) us applied transverse magnetic field (H) data were obtained for 0.010 in. diam wires at 4.2° Kin steady fields up to 120 kOe. The precipitate responsible for increase in Jc in each of the alloys studied was related to niobium and oxygen content. In alloys containing 40 to 60 wt pct Nb with less than about 1600 ppm of oxygen, w precipitation in the range 350° to 450°C appears to be responsible for the maximum value of Jc observed. In alloys of higher oxygen content (-2500 ppm), the oxygen-enriched a phase which precipitates in the temperature range 450° to 550°C is a more effective fluxoid pinner. Maximum Jc in low oxygen (350 ppm) 80 wt pct Nb alloys was achieved by 600°C aging heat treatments. Measurements of resistive critical field also reflect the changes in composition accompanying precipitation in these alloys. It is well known that the superconducting critical current density, Jc, of solid solution superconducting alloys can be increased by precipitation heat treatments.1-4 The precipitate, to serve as an effective superconzlucting fluxoid pinner and increase Jc, must be of proper size and spacing.' Its electronic properties must also be different from those of the matrix.6 As in age-hardenable systems, optimum temperature and time of aging are dictated by the composition and history of the particular solid-solution alloy under consideration. The alloy diagram of Ti-Nb7-9 shown in Fig. 1 is typical of ß-stabilized Ti alloys. Precipitation to achieve high Jc in Ti-Nb and Ti-Ta alloys is usually carried out in what is believed to be a two-phase a + ß region.1,2,4,10 Previous work on Ti-Nb4,11 and on other ß-stabilized systems such as Ti-V, Ti-Mo,12-14 has shown, however, that the first precipitate to appear is not necessarily (Y. Under specific conditions the meta-stable w phase is precipitated instead. A further complication arises in such alloys from the strong influence of contained oxygen on the phase stability and on the kinetics of precipitation.15,16 The superconducting critical current density of Nb-Ti alloy containing less than 36 wt pct Nb has been studied by Vetrano and Boom1 and Kramer and Rhodes.4 The most effective aging temperature was found to be about 425°C; fluxoid pinning was attributed to the precipita- tion of a by Vetrano and Boom and of w by Kramer and Rhodes. Although the latter workers used recrystal-lized material and the former cold-worked, studies carried out in our laboratory lead us to believe that the type of precipitate responsible for increased Jc is unaffected by prior cold-work.17 That other types of precipitates may also be effective, especially in higher oxygen-containing alloys, has been demonstrated by Reuter, et al.2 in this laboratory with the aid of electron microscopy. Reuter et al. found that at an aging temperature of 600°C, fee-TiO is precipitated, resulting in high values of Jc. Even higher values of Jc were observed after heat treatment at 500°C, but the precipitate could not be indexed as fee-TiO. Whether the precipitate was oxygen-rich w or a could not be clearly distinguished. The work nevertheless indicates that appreciable amounts of oxygen alter the nature of the phases precipitated at different temperatures, at least in alloys containing more than 36 wt pct Nb. Since such compositions are more readily worked into suitable wire than lesser niobium content alloys, it was decided to study the response to aging of the alloy compositions reported in the present paper with the hope of clarifying the effect of composition on the precipitation processes responsible for high values of Jc. I. EXPERIMENTAL Nb-Ti alloys containing 40, 50, 60, and 80 wt pct Nb were prepared by are-melting 35 g ingots of iodide crystal bar titanium and electron-beam melted niobium.

Jan 1, 1969

By J. C. Wilhoit, J. B. Cheatham

Although an oil well is a long cylindrical hole with an irregular bottom, it appears likely that the nature of the stress concentration at the bottom of the hole can be ascertained from an analysis of the stresses around a short cylindrical cavity with rounded comers and smooth bottom. Such a cavity is studied primarily because it leads more readily to a solution to the problem by the use of stress functions. In this paper the stress distribution around a short cylindrical cavity subjected to bit loading, overburden and drilling-fluid pressures is determined by means of an analytical solution which approximately satisfies the boundary conditions of the problem. From this solution the stresses at the corner of the hole are calculated to be about 35 per cent lower than comparable results obtained by photoelastic and relaxation analyses. This difference is apparently due to the large radius of curvature at the corner of the cavity in the present analysis. Since good agreement is obtained between the results of this analysis and the stresses calculated for a similar loading on a semi-infinite elastic solid, it is concluded that the bit action in the region near the center of the hole is not appreciably affected by the presence of the sides of the hole. INTRODUCTION Much has been written concerning drilling "under down-hole conditions" and pertaining to the stress distribution in the rock at the bottom of an oil well.1-5 For example, it is known that identical rocks can be drilled more rapidly at the surface than under subsurface conditions of pressure and stress.6 Information on the behavior of rocks under loading can be obtained from triaxial test data.7-9 From such tests it is found that rocks exhibit brittle failure when the confining pressure and pore pressure are equal, but the mode of failure may change to ductile as the difference between the confining pressure and the pore pressure is increased. Brittle failure implies that there is very little permanent deformation before fracture, whereas ductile failure indicates that permanent deformation takes place before fracture. Some rocks are ductile at differential pressures of 5,000 psi , but other rocks are brittle even at differential pressures of more than 50,000 psi. Cuttings embedded in mud at the bottom of the hole may act as a plastic mass which the bit teeth must penetrate in order to attack the virgin rock below. 10 During drilling, the mud pressure acts as the confing pressure, and in many cases the difference between the mud pressure and the formation pore pressure IS sufficiently low so that the rock most likely falls in a brittle manner. Very little penetratiori by the bit teeth may be required for brittle failure of the rock. For these brittle materials, the elastic stress distribution is of practical interest in determicing the possible effects of a given loading on failure. The penetration of bit teeth into ductile or plastic rock has been analyzed previously ll,12 and will not be considered further in the present work. During actual drilling, many teeth act at various points ovel an irregular hole bottom. In the present analysis a solution is obtained for an idealized problem of determining the elastic stress distribution caused by only one tooth acting alone at the center of a smooth hole bottom. It obviously is not proposed that any driller should put a special one-tooth bit on bottom, but it is hoped that extensions of the simpler problems can eventually lead to a better understanding of more complex actual drilling phenomena. To obtain a solution by the use of stress functions, a short cylindrical cavity with rounded corners and smooth bottom is studied. The solution to this problem should give an insight into the nature of the stress concentration at the bottom of an oil well which is in reality a long cylindrical hole with an irregularr bottom. After selection of the proper curvilinear co-ordinate system, the stress functions are expressed as series of Legendre polynomials. The coefficient of each term in the series is then determined to satisfy the boundary conditions at a finite number of discrete points on the boundary of the cavity with least square error. Results are compared with bottom-hole stresses obtained by photoelastic and relaxation analyses.

By P. R. Gage, R. W. Bartlett

At high temperatues and reduced oxygen pressuves, molybdenum silicicles oxidize to form SiO(g) vathev than a passivating SiO2 film. This is a sevious problem for low-pressure applications of sili-cide-coated rejvactory tnetals. As oxiclalion pvoceeds, the higher oxidntion potential of silicon initially suppresses oxidation of molybdenum and silicon is depleted near the surface. Eventually, silicon depletion causes secondary passivation (SiO2) mid an active to passiue transition in oxidntion behavior at moderately low pressures. At very low pressures, silicon depletion continues until steady-state oxidation of both silicon and rnolyhdenum can proceed at rates propovtional to their initial stoicliiorrzetvies. Ohsel-vations of 'oxidatiou-rate belzcriliov at temperatures from 1101° to 1700°C and oxygen pressures from 10-7 to 1 at in were made using thermogravimetn'c, X-ray diffraction, and elect?*on-rtric?ro/,vobe techniqrres. The experimental results are in good agreement with calculated lirnits for the various types of behavior that were based on tlievtnochernicnl atzd dffusion rate data for the Mo-Si system. A protective oxide film normally forms during the oxidation of elemental silicon. However, at elevated temperatures and reduced pressures this film is not present and oxidation is rapid. wagnerl has shown that this behavior occurs because SiO(g) is more stable than SiO2. The behavior of the silicides at high temperatures is similar but more complicated. Molybdenum and several silicide compounds are involved and the chemical activity of silicon, or the equivalent silicon vapor pressure, can vary considerably in these phases. Also, diffusion in the silicides resulting from reactions in which a higher silicide is reduced to a lower silicide, or molybdenum containing dissolved silicon, must be considered. EXPERIMENTAL AND ANALYTICAL METHODS Isothermal oxidation experiments were conducted using hot-pressed MoSi2 and Mo5Si3 wafers. Each experiment was conducted at a fixed oxygen pressure by controlling the flow of oxygen in and out of the furnace chamber. Most of the experiments were conducted in a molybdenum-wound tube furnace. Special construction features, including use of a high-purity dense A12O3 refractory tube, permitted operation in high vacuum to 1500°C. Higher temperatures could be obtained at higher working pressures and with a sacrifice of tube lifetime. Temperature was measured with three Pt 6 pct Rh/Pt 30 pct Rh thermocouples and controlled with a proportional controller. Weight changes were monitored with a quartz spring balance. The few experiments, at higher temperatures and low pressures, were conducted in a cold wall vacuum chamber using induction heating with direct susceptance of the sample. Because of the variable levitation effect of the induction coil on the sample, weight changes could not be monitored continuously in these experiments. All samples representing both experimental methods were weighed before and after oxidiation and examined with a microscope and analyzed by X-ray diffraction. The analytical treatment assumes equilibrium at interfaces and uses existing thermodynamic data. For convenience, equilibria are expressed as chemical potentials or as vapor pressures of the gases involved. The resulting differences in equilibrium vapor pressure between adjacent interfaces and between the surface and external atmosphere drive diffusion processes in the solid and gas boundary film, respectively. These diffusion processes govern the rate of oxidation. A bibliography of the sources of thermochemical data for the reactions pertinent to the Mo-Si-O system is given in Table I. Reactions are listed with the sources of free-energy data referenced. Since all of these reactions involve one or more gas species, usually oxygen, a plot of the product-gas chemical potential, -RT In Pg, vs absolute temperature for each reaction is given in Fig. 1. These curves can be used to determine equilibrium vapor pressures or to compare the stability of different compounds in the various gases: O2, Si, SiO, and Moo3,. The data of Table I and Fig. 1 were used to determine the analytical results presented in this paper. EXPERIMENTAL RESULTS The weight changes for three MoSi, samples oxidized at 1500°C are plotted against time in Fig. 2. The treatments were identical except for the indicated variance of oxygen pressure and each curve represents one of three characteristic types of behavior that were observed. The upper curve is an example of passive oxidation, which is defined as

Jan 1, 1965

By M. A. Grossman

THE hardenability of most steels can be predicted within 10 to 15 per cent provided the complete chemical composition is known, including "incidental" elements; and provided the as quenched grain size is known; and provided, finally, that the composition and heating temperatures for hardening are such as to result in austenite free from carbide particles. In method proposed herein, the steel is considered as having a base hardenability due to its carbon content alone (the hardenability of a "pure steel" of the given carbon content, without any other elements), and this base hardenability is multiplied by a multiplying factor for each chemical element present. After multiplying all these together, the final product is the hardenability. Grain size may be taken into account either in the base hardenability or after the multiplication. Hardenability is stated in terms of "ideal critical diameter"; namely, the diameter of bar, in inches, that will just harden all the way through (absence of unhardened core) in an "ideal" quench, and the calculation may also be related to the Jominy test. The data bring out certain features rather clearly. For example: I. It is quite useless to attempt to predict hardenability unless all elements, including "incidentals," are known; thus an "incidental" chromium content of o.-2o per cent increases the hardenability by about 50 per cent. 2. The relative effectiveness of different alloys is sometimes unexpected; thus molybdenum when calculated in this way appears to be of the same order of effectiveness as manganese, rather than much more powerful as would appear to be common experience; observe that an increase from no manganese to 0.20 per cent Mn provides a multiplying factor of 1.67, and an increase from no molybdenum to 0.2o per cent Mo provides a factor of 1.63, an increase of over 6o per cent in each case. However, if to a steel containing 0.50 per cent Mn there is added "20 points of manganese," the factor is raised from 2.65 (for 0.50 per cent Mn) to 3.35 (for 0.70 per cent Mn), or an increase of only 26 per cent. Thus the first small addition of an element has a much more powerful percentage effect than an equal further addition when some is already present, and in most cases the effect of molybdenum is considered in relation to a steel in which molybdenum is absent. 3. If two elements are equally effective, a greater harden ability will be obtained by using, for example, o.5 per cent of each than by using 1.0 per cent of either of them alone. 4. A knowledge of the as-quenched grain size is essential for precise work, since a difference of only one grain size number (say No. 7 instead of No. 6) makes a difference of almost 10 per cent in hardenability (in these units); this, however, does not apply to certain steels of high hardenability. It should be emphasized that in chromium steels (over 0.30 per cent Cr) and chrome-molybdenum and chrome-vanadium steels, undissolved carbides are likely to be present in the steel as quenched, and that in such cases the charts can indicate only a maximum possible hardenability, whereas the extent of hardening actually obtained may be much less. Thus tests on a number of chrome-molybdenum steels have indicated a degree of hardening amounting to only 50 to 65 per cent of the maximum possible, and in chromium steels from full hardening (100 per cent) down to as low as 70 per cent. On the other hand, when the amount of such elements is small (Cr under 0.30 per cent, Mo up to 0.25 per cent in the absence of Cr, and V up to 0.04 per cent), the charts provide reasonable approximations. The precise figures on the charts are suggested as tentative, subject to some modification as more data accumulate, but the fundamental concept appears to be supported by tests made on a wide variety of steels, a few of the correlations being shown in Fig. I.

Jan 1, 1942

By V. G. Leak, R. A. Swalin

The dilhsion of silver and trace concentrations of tin in liquid silver has been rrzeasured in the temperature range from about 975° to 1350°C. The difBsion dala. fil lhe following equations: fov self-diffusion of silver The ratio of DSn to DAg is found to be about 1.34. The higher diffusivity of tin is interpreted in terms of the coullombic repulsion which results from the fact that tin dissolved in silver has a valence of +3 relative to silver. THE phenomenon of diffusion has been well-described theoretically for hard-sphere gases and solids and found to be in good agreement with experimental data for certain gases and solids. Liquid-diffusion phenomena have not been well described theoretically and there is generally a lack of good experimental data. In this investigation self-diffusion and solute diffusion in liquid silver were studied. Silver was chosen as a solvent for two main reasons. First, silver is a noble metal and the atoms are considered to have a spherically symmetric charge field; hence the liquid may be considered to be a random array of approximate hard spheres. Mercury, gallium, and other lower-melting metals were eliminated from consideration since it appears possible that in their liquid states there is some residual long-range order and directional bonding. Second, silver behaves as a monovalent solvent and, while cadmium, indium, tin, and antimony all have nearly the same atomic size, they have chemical valences relative to that of silver of +1, +2, + 3, and +4, respectively. Slifkin, Lazarus, and coworkers1-5 investigated the diffusion of these solutes in solid silver and found that their diffusion rates increased with an increase of the excess valence of the solute. In the solid state the diffusion-rate increase was calculated to be due to a change in local modulus of the solvent caused by the excess valence of the solute.1"3 For the present investigation it was planned to determine the effect of valence upon solute diffusion in liquid silver. It was deduced that a coulombic repulsion between solute and solvent might be responsible for larger volume fluctuations in the vicinity of the solute thereby enhancing the diffusion of the solute atoms. Tin was the first solute investigated and was chosen for experimental convenience. If an excess-valence diffusion effect exists in the liquid state, the solute tin with its excess valence of + 3 might show a large enough effect to distinguish it from the self-diffusion of silver in silver. The silver self-diffusion data were obviously required as a base line for comparison. In addition silver self-diffusion was investigated with a view toward examining the data in connection with the Sutherland- Einstein: Coheen-Turnbull,7 and swalin8 theories of diffusion in liquids. I) EXPERIMENTAL TECHNIQUES The diffusion coefficients for tin and silver in liquid silver were determined in separate experiments using the capillary-reservoir method of Anderson and saddingtono but following closely the experimental techniques outlined by Ma and Swa1in. The radioactive alloy bath was prepared by plating either tin-113 or silver-110 m isotope* onto 99.999 *Isotopes obtained from Oak Ridge National Laboratory. pct Ag rods.* The silver and isotope were melted *Silver obtained from Cominco Co. and mixed in a graphite crucible under a purified argon atmosphere, then outgassed for capillary filling. Some of the fused-silica capillaries were filled individually as reported by Ma and Swalin, but most were filled in a different manner. A long piece of capillary tubing was sealed off on one end and held in the system so that the open end was just above the surface of the molten alloy. The system was evacuated, the open end submerged into the alloy, and argon was admitted into the system up to atmospheric pressure. The molten alloy was forced up the capillary several inches where it solidified and was subsequently cut into segments of the proper length for diffusion annealing. The apparatus for diffusion annealing was the same as that used for capillary filling except that a Pt—Pt, 10 pct Rh therinocouple was placed in the solvent bath in order to accurately measure the temperature during the run. A diagram of the dif-

Jan 1, 1964

By Richard A. Queeney, Lance G. Peterson

ESTIMATES of the stacking fault free energy of copper reported in the literature show an extensive divergence of results. Based on measurements of dislocation node radii, Thornton et al.7 find the lower limit of stacking fault energy to be 60 ergs per sq cm: Jossang and coworkers4 estimate the same value should be less than 40 ergs per sq cm. Assuming stacking fault free energy to be twice the coherent twin boundary free energy, Valenzuela8 reports 72 ergs per cm, while Inman and Khan2 cite a value of 24 ergs per sq cm. which can no doubt be accounted for by almost complete absence of interstitials in solution. However, the minute concentration was sufficient to permit observable recovery of dislocation damping. Assuming a dislocation density of 1010 CM -2 and an average loop length of 0.5 , a concentration of interstitials as low as 0.01 at. ppm is sufficient to reduce dislocation damping by more than a factor of ten. The concentration of interstitial solute atoms in Ferrovac E-0.15 pct Ti was thus estimated at 0.01 at. ppm < (C + N) < 2 at. ppm. The absence at elevated temperature of strengthening due to interstitials, Fig. 1, was therefore anticipated in the titanium alloy. The temperature dependence of the flow stress below 300°K was unaffected by addition of either titanium&apos; or zirconium.&apos; Observed, however, was an initial difference in the absolute value of the flow stress below 320oK, Fig. 1. Grain size may account for the difference; the grain diameter of Ferrovac E was 30 to 45 µ and that of Ferrovac E-0.15 pct Ti 65 to 90 µ. Although second phase particles in the form of titanium carbides and nitrides precipitated in the Ferrovac E-0.15 pct Ti alloy, the spacing was too large to produce any strengthening effect. The mechanical properties measured, therefore, were inherent to the bcc lattice. 1W. C. Leslie and R. J. Sober: Trans. ASm, 1967, vol. 60, pp. 99-111. 2 H. Conrad: Journal of Metals, 1964, vol. 16, pp. 582-88. 3 A. S. Keh and W. C. Leslie: Materials Science Research. H. H. Stadelmaier and W.W.Austin, eds., vol. 1.pp. 208.50, Plenum Press, New York, 1963. 4A. S. Keh, Y. Nakada, and W. C. Leslie: Dislocation Dymmics, Rosenfield, Hahn, Bement, Jr., and Jaffe, eds., pp. 381-406, McCraw-Hill Book Co., New York, 1968. &apos;W. J. Bratina, J. T. McCrath, and H. E. Rosinger: Can. Met. Quart., in press. 6 W. J. Bratina, J. T. McGrath, and D. Mills: Suppl. Trans. Japan Inst. Metals. 1968, vol. 9, pp. 436-43. &apos;H. E. Rosinger and C. B. Craig: Can. Met. Quart., In press. %. Mills. J. T. McGrath. and W. J. Bratina: ScriptaMet,, 1968, vol. 2, pp. 311-13. The purpose of this note is to present the results of a direct measurement of the stacking fault free energy in copper. The present method avoids both the question of interaction forces of nodal dislocations and the assumption that the coherent twin boundary free energy is one-half that of a stacking fault. Fig. 1 shows a transmission electron micrograph of a stacking fault intersecting a high angle grain boundary in copper: Fig. 2 is a schematic representation of the intersecting interfaces. To produce the specimen, pure copper was rolled to 0.002 in. thickness and annealed for several hours at 875°C. These foils were strained in tension to l04 µ in. per in., and rean-nealed to 8 min. Recrystallization and/or grain growth were not observed during the second anneal, indicating that intersections such as Fig. 1 are equilibrium configurations. Many such stacking faults were observed, but few were noted as intersecting grain boundaries in well-illuminated areas. The intersection of Fig. 1 was well away from the specimen edge. Since similar defects have not been observed in fully recrystal-lized and annealed copper, it is felt that this defect is due to the tensile straining. Electron transmission

Jan 1, 1970

By P. E. Busby, C. Wells, M. E. Warga

Fundamental data on the rate of diffusion of boron in austenite and solubility of boron in the a and y phases of iron and steel have been obtained from deboronizing experiments and provide partial explanations for some of the phenomena observed in boron steels. The rate of diffusion of boron is about the same as carbon in austenite. The solubility of boron in austenite at normal heat-treating temperatures is less than 0.001 pct. A partial tentative Fe-B phase diagram in the important low boron concentration ranges and an equation representing the diffusion of boron in austenite are presented. A LTHOUGH a search of the literature revealed ALTHOUGHthat boron diffusion had not been studied quantitatively and systematically prior to 1948, a few qualitative observations which were made by following the diffusion of boron into iron packed in ferro-boron1-3 suggest that boron obeys normal diffusion laws and that the rate of diffusion increases with increasing temperature. The claim of Cornelius and Bollenrath1 that boron does not appear to influence the rate of carbon diffusion has recently been substantiated by Wells, Batz, and Mehl.4 Calculations based on data from the paper by Campbell and Fayv indicated that the diffusion coefficient (D) for boron in iron at 900°C is approximately 3x10 sq cm per sec; a value of 2x10-7 sq cm per sec at 1038°C has been reported by Digges et al in connection with decarburizing experiments on a commercial boron steel containing 0.43 pct C. During the period 1948 to 1950 several diffusivity constant (D) values for the diffusion of boron in the y phase of iron and steel and a number of solubility values for boron in both a and y phases were reported. It was concluded tentatively as a result of Metals Research Laboratory studies that D values for boron in y and a phases are about the same as for carbon in the comparable phasesa and that carbon up to 0.4 pct did not affect the rate of diffusion of boron in y iron above 1000°C within the limits of experimental error. Whether saturation of the y or a phases with carbon would significantly affect rates of diffusion of boron through them was then and still remains in doubt. The best estimate of Q (activation energy) was reported to be about 25,000 cal, somewhat lower than that for carbon in y iron. The solubility of boron in y iron at 1000°C was thought to be higher than 0.004 pct B but not as high as a more recent analysis of available data shows it to be. The solubility of boron in a iron at 700°C was reported to be 0.0004 pct B, but it would not have surprised the authors if the true solubility value is actually lower than this. Breaks in diffusion curves which were not understood when first obtained prior to 1950 have now been recognized as indicating solubility limits, and these are included in the present paper. Experimental Procedure Two methods were employed in an effort to determine diffusion coefficients (D) for boron in austenite. Early experiments involved the use of welded couples prepared in accordance with the procedure of Wells and Mehl,&apos; but when the results of these experiments proved to be unsatisfactory for calculating D values, studies were continued by means of deboronizing experiments. In both cases, boron analyses following the diffusion anneal were carried out spectrographically and all D values were computed using the Grube solution of Fick&apos;s law. The precision of the spectrographic method used proved to be about equal to that reported by Corliss and Scribner" —average deviation from the mean about 5 pct of the amount present when the boron content is 0.003 pct. Concentration-penetration curves obtained from spectrographic analyses of welded couples generally indicated that practically no boron was transported across the weld interface during the diffusion anneal. Furthermore the concentration-penetration

Jan 1, 1954

By G. E. Keller

In connection with the work of Committee D5, Coal and Coke, of the American Society for Testing and Materials in developing new standards and improving old standards, series of tests were made to check the variance of sample preparation and analysis of American coals. Table I in the "Tentative Method of Preparing Coal Samples for Analysis - 02013-62T" is rather complicated, and the tests were designed to develop data for the simplification of this table. An allowance of 20% of the total variance of sampling and analysis for the sample preparation and analysis steps was used. Also, instead of coals being grouped into ash ranges, they were grouped as washed and raw (unwashed) coals. As a result of these tests, a simplified Table I for Tentative Standard 02013 has been proposed. A standard test for laboratories to check their variance of sample preparation and analysis has been developed. Committee D5 on Coal and Coke of the American Society for Testing and Materials has been making a study of all standards applying to coal for the purpose of improving and bringing them up to date and for breaking some of the existing standards covering a number of tests into single standards. As an illustration, sample preparation is now partly covered in ASTM Designation D492-48, Sampling Coals Classified According to Ash Content, and partly in ASTM Designation D271-58, Laboratory Sampling and Analysis of Coal and Coke. The latter contains, in addition to parts of sample preparation, methods for determining Moisture, Ash, Volatile Matter, etc. Although revision of sampling methods and sample preparation methods are now assigned separately, it is possible that they may end up as one standard. One of the first new Tentative Standards developed as the result of these studies was ASTM Designation D2013, Tentative Method of Preparing Coal Samples for Analysis. Although this method is quite an im- provement over that in the old standards and gives very specific instructions, Table I (attached) is quite complicated and rather hard to understand. In order to take advantage of reduced quantities or weights, it is necessary to know the ash content within rather narrow limits before sample preparation and no distinction is made as to whether the coal is cleaned or uncleaned. No limits of accuracy are assigned to the sample preparation and analysis. The total over-all accuracy of sampling, preparation and ash determination as given in D492-48 has been specified for commercial sampling procedure as follows: "This procedure is intended for an accuracy such that if a large number of samples were taken from a single lot of coal, the test results in 95 out of 100 cases would fall within ? 10 per cent of the average ash content of the sample." Although Appendix 11, Method for Determining the Over-All Variance of Reduction and Analysis in the Tentative Methods for Mechanical Sampling of Coal, ASTM Designation D2234-63T, does give a method for determing the over-all accuracy of preparation and analysis expressed as variance and also the F test, it does not permit the determination of the variance at the various stages of reduction or of the ash determination. It was decided in 1961 to ask a number of laboratories to run a series of tests for the purpose of determining the variance at various stages of reduction and analysis with the view of simplifying Table I of the Tentative Standard D2013-62-T. This paper is a record of this series of tests with recommendations which resulted from the tests. At the same time and preceding these tests the International Organization for Standardization under Committee TC/27 WG 7 was studying the same problem. Tomlinson1 of England and Dr. paul2 of West Germany had made studies giving quite a bit of information. Also Schwerd and coryel13 published some information which was the main basis for the Mechanical Division provision in ASTM Tentative Standard D2013-62T, Table I. In checking with the various laboratories which offered to cooperate in these tests, none was found which used mechanical division through the entire preparation procedure. It was then decided to use riffles at all stages with the thought that if riffles gave satisfactory variances, approved mechanical

Jan 1, 1965

By Michael Hoch, Walter C. Wyder

The Cr-Ni system was investigated between 900" and 1400°C and 45 to 100 wt pct Cr. The Cr-Fe system was studied between 1200" and 1385°C and 70 to 100 wt pct Cr. Powdered alloy compacts were produced from high-purity chromium, nickel, and iron poulders. The compacts were sintered in an argon atmosphere and then machined into small, 0.030-in. diam cylindrical specimens. These were then heated by induction in the high-temperature X-ray diJfractwn camem under a helium atmosphere and irradiated with copper Ka radiation. From the results of the high-temperature X-ray diffraction patterns it can be concluded that chromium exists only in the bcc form, and that no eutectoid reaction occurs in the Cr-Ni system in the temperature interval 900" to 1400°C. The X-ray diffraction patterns are in agreement with the Cr-Ni phase diagram of Bechtoldt and Vacher, and Williams. It is possible that chromium undergoes one or more allotropic transformations between room temperature and its melting point. Most of the evidence was gathered from the study of the chromium-rich ends of the Cr-Ni and Cr-Fe binary systems. All the investigations rely heavily on indirect methods: thermal analysis, electrical resistivity, and micro-graphic inspection of quenched samples. In the present work the chromium-rich side of the Cr-Ni and Cr-Fe binary systems is studied, using high-temperature X-ray diffraction techniques. Thus the coexistent structures can be observed directly at temperature. PREVIOUS INVESTIGATIONS ~ansenl shows two possible Cr-Ni phase diagrams 1) a simple eutectic [Liquid =a-Cr(bcc) + y-Ni(fcc)], 2) a eutectic [Liquid+ P-Cr(fcc) + y-Ni(fcc)] and euctectoid [p-Cr(fcc) ==a-Cr(bcc) + yNi(fcc)]. The latter is based on the work of Grant and co-workers.2 Stein and rant&apos;" placed the eutectoid composition at 1215°C and 68 wt pct Cr; their diagram is shown in Fig. l. Similarly, Sully3 also discusses the two proposed diagrams. Price and Grant,4 working on an isothermal section of the Cr-Ni- Fe ternary diagram at 1300°C, made no mention of a ß-chromium phase (fcc). The ß-chromium structure was also reported by Misencik 5 in the Cr-Nb(Cb) system. Williams 6 studied the Cr-Ni system using precipitation techniques. His work included the chromium and nickel solvus lines below 1250°C. He reported no ß structure nor any eutectoid reaction. The solvus lines which he produced were very similar to those of Jette, et al 7 below 1000°C. He concluded that Grant and coworkers2 had misinterpreted their data, and that effects which they had reported were actually due to the rapid and extensive precipitation of the nickel-rich phase. Bechtoldt and vacher8 used sintered powdered alloy compacts in an attempt to find the ß phase. They used metallographic and room-temperature X-ray diffraction techniques and also failed to detect the fcc chromium phase. They attributed the low slope of the chromium solvus at approximately 1150°C as the probable cause of the abundant precipitation obtained. They also stated that the precipitation was the probable cause for the abrupt discontinuities in the thermal analysis and electrical resistivity data which led to the conclusion that a eutectoid reaction existed. Fig. 1 also shows their diagram. Grigor&apos;ev and coworkers9 investigated the polymorphic changes of chromium in the Cr-Ni phase diagram. Using standard microstructural and thermal analysis techniques, they published a phase diagram which contained five polymorphic forms of pure chromium and four eutectoid reactions. They stated that the room-temperature bcc chromium transformed to a fcc structure at 930°C, to bcc at 1300°C, to hcp at 1650oC, and then finally to bcc at 1830°C. The eutectoid reactions were placed at 850°, 960o, 1140°, and 1220°C. Also to be noted was the fact that the uppermost phase, the bcc above 1830°C, was different from the upper phase of Grant and coworkers,2 the fcc above 1840°C. Their eutectoid reactions, temperatures and compositions fall

Jan 1, 1963

By M. C. Flemings, R. Mehrabian, G. E. Nereo

Analytic expressions were developed and applied in two preuious papers to predict effects of solidification variables on macrosegvegation. In this paper, experiments are reported to test the analyses quantitatively and qualitatively. Good quantitative agreement is obtained between predicted and measured compositions for the following types of segregation (all experiments were on laboratory-size ingots of Al-4.5 pct Cu alloy): (a) inuerse segregation, (b) composition distribution in ingots solidified with unidirectional heat and fluid flow, (c) simulated centerline segregation, and (d) segregation resulting fro, change in solidification cross section. In addition, qualitative agreement of prediction of theory with experiment is obtained for (a) segregation resulting from sudden change in heat extraction during solidification, (b) centerline segregation in a bidirectionally solidified ingot, and (c) under-riser positive segregation. IN two previous papers,&apos;&apos;2 it was shown that a variety of apparently different types of macrosegregation result from the same basic mechanism. This mechanism is the flow of solute -rich liquid to feed solidification and thermal contractions. Analytic expressions were derived and examples given of their application to predict macrosegregation. Calculations were for Al-4.5 pct Cu alloy. In this paper, we describe a series of experiments designed to test results of the foregoing calculations. All experiments were on laboratory-size ingots (weighing 10 lb or less), of A1-4.5 pct Cu alloy, nominal composition. APPARATUS AND PROCEDURE A typical mold and chilling arrangement used for unidirectional solidification is shown in Fig. 1. Before casting, this assembly was placed in an electrically heated air furnace, and heated to 680" to 700°C. When the plaster reached the desired temperature (usually 1 to 2 hr after the furnace temperature), the mold was filled through a funnel extending from the outside of the furnace into the cavity. The assembly, with liquid, was held for several minutes to allow convection to subside. The coolant, either air or water (depending on the desired solidification rate), was turned on and the ingot solidified. With the four sides and top insulated, heat removal for solidification was through the bottom surface, resulting in a fully columnar structure. Design of the casting cavity was varied, as discussed below, to produce desired macrosegregation effects. However, for castings solidified with unidirectional heat flow, the top was always open to the furnace atmosphere and temperature. The bottom sur- face rested on a chill through which the heat was extracted, and all other surfaces were insulated with plaster. In one case to be discussed, the mold design was changed to obtain bidirectional solidification. In this case, two vertical chills were used (on two faces of the mold) and the bottom was insulated. In a last case discussed, that dealing with under-riser segregation, experiments were conducted in bottom-chilled sand molds at ambient temperature. Except for this one case, all experiments were conducted in the electrically heated air furnace as described above. In many ingots, thermocouples were used to obtain continuous records of solidification. In all cases, these were chromel-alumel couples with output continuously monitored on a twelve-channel recorder. Where thermocouples were employed, these were embedded in the refractory mold walls during mold-making so their heads extended to the vertical centerline. Six to nine couples were employed, placed at intervals from the chill. METAL-MEL TING PROCEDURE Aim analysis was A1-4.5 pct Cu; alloy was prepared from high-purity aluminum (99.9 pct Al) and A1-Cu master alloy (50 pctOFHC Copper, 50 pct high-purity Al). The A1-4.5 pct Cu alloy was cast in 5- to 7-lb ingots for subsequent remelting for the macrosegregation experiments. Melting for ingot casting was as follows. After preheating the melting crucible for an hour or more, the charge was melted and superheated to a temperature of 690" to 720°C. After stirring, the melt was degassed with chlorine for 10 min. After stirring again for 1 min, a gas sample, under reduced pressure, was taken. When the sample was gas-free, the crucible

Jan 1, 1969

By H. R. Peiffer

Hardening of soft metals can be accomplished by dispersing finely divided hard particles in them. The dispersing of finely divided alumina in silver in the presence of oxygen yields a high strength material which is unusual in that its mechanical properties above the melting point of the continuous Ag-O alloy matrix are similar to other solids. The tensile strength is studied for two of these alloys, one of which contains 15 pct by weight alumina and the other 20 pct by weight alumina. The average fracture strengths above 960°C of these alloys were found to be 0.4 x106 dynes per sq cm and 3.8x106 dynes per sq cm respectively. The strengths appenred to be independent of temperature above 960° C. The creep behavior of the 20 pct alumina material was studied above 960°C. The initial creep rate, 6 , of this material can be represented by where s is the applied stress and E the activation energy for the process. This energy is of the order of 1.7 to 2.1 ev. HARDENING of soft metals by the addition of finely divided, hard particles has resulted in the production of materials with excellent creep resistance and with inhibited recrystallization.&apos; It has been demonstrated that such composites have useful strength up to temperatures very near the melting point of the soft metals, but one might expect that the strengthening by the dispersed particles ceases to be of importance above the melting point of the softer phase. This, however, is not so.2 In this paper the strengthening of a liquid Ag-O solution by the presence of dispersed alumina is discussed and experimental data concerning the mechanical properties of such a material at temperatures above 960°C presented. EXPERIMENTAL PROCEDURE Specimens were prepared using powder metallurgical techniques. Finely divided silver oxide was mixed in the appropriate proportions with "Linde B" alumina. For these experiments the specimens contained 15 or 20 wt pct alumina in pure silver.* These alloys were used in these experiments since they were the ones that could be handled readily enough to yield experimental data. Silver oxide was used in preference to silver because it is readily available in a very finely divided form and because the presence of excess oxygen in the silver is important to the properties2 of the material. The powders were mixed wet, using alcohol as the medium, in an ordinary food blender until a uniform mixture was obtained. They were then dried and heated to reduce the silver oxide to silver. Only a fraction of the finely divided silver remains unoxi-dized upon removal from the oven since the finest particles of silver oxidize immediately upon contact with air. The remaining fine silver served as a binder during pressing. The mixture was then pressed at 20 tons per sq in. into shoulder-grip (reduced section 1/4 in. by l/4 in. by 1 1/4 in.) and rectangular bar specimens (l/4 in. by 1/4 in. by 2 in.). The green specimens were heated very rapidly in a globar furnace to 1000°C, held at temperature untic they had reached theoretical density and then furnace cooled. The strength of the two compositions at temperatures above 960 °C was determined by pulling shoulder grip specimens inside a furnace mounted on a tensile machine. Specimens were gripped by means of wires wrapped around the shoulders of the specimen. Temperatures were measured by means of a chro-mel-alumel thermo-couple placed in the vicinity of the specimen. Control of temperature (within ±°15C) was accomplished by means of a Weston Recorder-Controller. Before loading, specimens were held at temperature for approximately 15 min in order to insure uniform specimen temperature. The low loads were measured by a special load cell designed by Baldwin Lima Hamilton. All testing was performed on a FGT testing machine of the same com-pany. Creep studies were performed in bending utilizing rectangular specimens of the 20 pct alumina alloy. Specimens supported at both ends by knife edges of inconel were placed in a globar furnace, heated quickly to temperature, and the deflection of the beam measured optically as a function of time. The weight and size of the specimen were predetermined and the maximum stress on the beam calculated from

Jan 1, 1962

By C. P. Lawhon, J. P. Simpson, W. M. Evans

Data obtained under controlled test conditions using a microbit drilling machine showed that oil emulsified in water muds may either increase or decrease the drilling rate, depending upon drilling conditions. A low-viscosity oil such as diesel fuel can give drilling rates in limestone almost equal to that of water. Data obtained for water emulsified in oil muds showed little decrease in the drilling rate in water-saturated cores as the water percentage of the mud was increased above the 5- to 10-percent range. Changes in drilling rate were found to be dependent upon the oil or water concentration of the mud and upon the type of formation drilled. Changes in static filtration on paper (API filtrate) did not correlate with filtration while the mud was circulated across rock. INTRODUCTION Oil additions to water muds have been reported to increase drilling rates, provide hole stability and improve filtration control. Eckel&apos; showed that water-base emulsion muds used in the West Texas area increased drilling rate with increasing oil concentration up to 15 percent oil by volume, but drilling rate decreased at a concentration of 20 percent by volume. Based on laboratory tests using water muds to drill shale, Cunningham and Goins&apos; reported that drilling rates increased and tendency for the bit to ballup decreased with the addition of oil. Percentage increase in drilling rate varied with the particular formation. They showed oil additions to improve drilling rates ap proximately 75 percent in Vicksburg shale and as much as 150 percent in Miocene shale. Each investigation showed an optimum oil content for the particular formation. Most data that indicated improved filtration control due to oil additions were based on static API Eltrates through paper rather than dynamic filtration through permeable rocks. Some types of dynamic test give a better representation of filtration down-hole while drilling and might be more likely to show some correlation with drilling rate. Static filtration would be important, of course, in relation to hole stability and formation damage. This laboratory&apos;s drilling tests, conducted on water-raturated Berea sandstone, indicated that improvements in drilling rate were not evident with increasing oil concentration in water-base muds. Investigation also showed similarity between oil-emulsion (water-in-oil) muds and water-emulsion (oil-in-water) muds while drilling these formations. In Lueders limestone high concentration of water-in-oil muds and high concentration of oil-in-water muds provided the same relative drilling rates. In Berea sandstone there was a large reduction in relative drilling rate with both the oil and water muds that contained low percentages of emulsified fluid. Dynamic filtration rates of water muds on rock did not always decrease with increasing oil percentages even though the static API filtration rates on paper did decrease. Data observed in laboratory drilling of limestone and sandstone indicate that improvements in field drilling operations when water- or oil-emulsion muds are wed may not be the result of increased drilling rates but of improved hole conditions. In some cases, actual drilling rates might be slower but improved hole conditions will result in less total time on the hole. DEFINITIONS Mud pressure—Pressure of drilling fluid as measured after leaving the drilling chamber. This is considered as the approximate mud pressure just past the bit and at the face of the formation. Terrastatic pressure—Pressure representing weight of overburden. Formation pressure—Pressure of formation fluid as measured at outlet of drilling chamber. This is considered as approximate pressure of fluid in the pores of the formation. Differential pressure—Difference between the mud pressure and formation pressure. Relative drilling rate, percent—Drilling rate with experimental fluid divided by drilling rate with water times 100 equals percent. LABORATORY EQUiPMENT AND TESTING PROCEDURES The drilling equipment has been described in previous publications The microbit drill is a closed system (capacity, approximately 7 gal) that can be pressurized to 15.000 psi and heated to 500F. Main components are a drilling chamber, filter-heater, rotary-drive and variable-speed cir-