Search Documents

By Morgan Walker

INTERMITTENT injection of gas in gas-lift pumping is a variation of the common practice in that the gas is' .delivered to the well for a short, period, called the "on time," followed by a period during which none is admitted, called the "off time," instead of being admitted continuously at a constant rate or by intermittent flowing in which it is occasionally delivered to the well in just sufficient quantity to flow off the accumulated head of fluid. The term "gas" is used throughout this discussion but it should be remembered that in so far as the operation is concerned either air or natural gas may be used. The equipment at the well is the same as commonly used in gas-lift installations and gas may be admitted to either the tubing or the casing. It seldom occurs that all the wells on a lease can be pumped economically by intermittent injection from one header. Wells that. have flow characteristics which permit their being handled together have their intake lines grouped at a central compressor station. Well groups may consist of two or more wells. The total volume of gas delivered by each group of compressors goes to one well at a time through a quick opening valve, each well connected to the header receiving its charge of gas in turn. Special cases may require a deviation from this arrangement. The distribution of the gas to the different wells is effected by an automatic timing device which operates quick-opening valves on the intake lines. The timing mechanism may be actuated by electricity, liquid or gas pressure. The essential factor in the operation of the timing mechanism is that as one intake valve is closed another must be opened.

Jan 1, 1928

By D. N. Stevens, D. N. Bloom, M. A. Klugman

This paper represents a portion of a larger study concerning dispersion of major, minor and trace elements in wall-rock about orebodies in part of the Colorado Mineral Belt. This phase of the work pertains to the Mg/Ca ratio variations in wall-rock adjacent to orebodies in two mines, one in the Horseshoe Mining District (Hilltop mine), and the other in the Alma Mining District (Phillips mine). The two districts lie adjacent to one another, the Horseshoe District on the south and the Alma District on the north, and are commonly referred to as one, the Alma-Horseshoe Mining District. The Alma-Horseshoe Mining District is located 80 miles southwest of Denver, Colo., (Fig. 1) in Park County, on the eastern slope of the Mosquito range. The area is typified by rugged relief with altitudes of from 10,500 to 14,000 ft. The Horseshoe District is separated from the Lead-ville Mining District by the crest of the Mosquito range, with the Leadville District to the west. The Alma-Horseshoe District has been considered as peripherial to the Leadville District, and has not received the detailed attention that the Leadville District has. These two mines were selected for this study because of their ease of access; the attitude of the workings and the drill holes to the known mineralization; the similarity of the mineralization; and because the mineralization occurs in carbonates in both mines but are of different ages. In the Hilltop mine mineralization is confined to the Leadville dolomite (Mis-issippian) and is principally lead with subordinate zinc and silver. In the Phillips mine the mineralization occurs in the Peerless dolomite and shale (Cambrian) and in the Manitou dolomite (Ordovician) and is principally zinc with subordinate lead, silver,

Jan 1, 1964

The beginning of Colorado's mining industry is linked on one side with that of the Appalachian districts and on the other side with that of California, because the first discoveries were made by pioneers when on their way from Georgia to California. In the summer of 1849 a party of seven Georgians were taking a herd of thoroughbred horses across the continent to the Pacific coast. They reached Camp Lyon, on the Arkansas river, in October, and, meeting James Dempsey, a Government guide, they were persuaded by him that it was too late to cross the mountains that season. His advice was accepted, and, moving northward, they established a winter camp at the junction of Cherry creek with the Platte, close to the site of the city of Denver. Upon a sand-bar, lining the south side of the river, they built two cabins. During the closing months of 1849 they prospected the alluvial banks of Cherry creek, but they made no attempt to penetrate into the mountain canyons, near them, westward, for fear of the Indians.* They found gold in several places along Cherry creek, and particularly at a spot 16 miles up the stream from their camp. From the feathers of the wild geese that they shot they obtained quills in which they placed their gold-dust. This party of Georgians consisted of Dr. L. J. Russell and his brother, Green Russell, also A. T. Lloyd, G. W. Kiker, and P. H. Clark. Early in 1850 they crossed the main range of the Rocky Mountains by the Bridger pass and went on to California. They mined near Downieville, and were successful. Occasionally mention was made by them of the gold

Jan 1, 1932

By C. E. Lawall, C. T. Holland

When this study was started very little information was available regarding the physical characteristics of West Virginia coals. This was particularly true of friability and of crushing strengths of the coals. Some work had been done on specific gravity studies but the available information applied to only a few coals. In practically all coal-classification systems thus far proposed, the predominating factors of classifications are largely based on the chemical characteristics of the coal. Recently some investigators have pointed out that a classification based on chemical and physical properties of coal might serve to place coals in their proper group better than classification systems based on chemical properties alone. One of the aims of this study was to obtain some data on the physical properties of coal which might help to classify West Virginia coals. The Samples Fig. I shows where the coal samples were obtained and the approximate division line of the high and low-volatile coal fields of the state. The relative geological horizons of the coal beds tested are shown in Fig. 2. In securing samples of coal for friability and specific gravity tests, it was the aim always to obtain a representative sample of lump coal produced at a particular mine. The samples were obtained from the loading boom, the railroad car, or the mine cars. Wherever it was possible to secure samples from a loading boom they were taken from it by taking shovelfuls of coal from the loading boom every minute until about 200 lb. of coal was obtained. If there were no loading booms used at the mines sampled, the coal sample was taken from a loaded railroad car by systematically taking shovelfuls of coal at definite intervals over the whole top of the car until about 200 Ib. of coal was obtained.

Jan 1, 1932

By F. C. Blickensderfer

During 1914 and 1915, extensive experiments were conducted at the concentrator of The Detroit Copper Mining Co. of Arizona, at Morenci, Ariz., in order to test the relative grinding efficiencies of the Marathon, the Chilean and the Hardinge Mills. The Marathon mill used in the experiments was the first one ever built by the manufacturers; and since this was the first time that it was given a thorough tryout, with such remarkable results, the data have been assembled in the following paper and submitted in the interests of the milling profession. Description of the Flow Sheet The concentrator flow sheet, as it existed at the time of these tests, is shown in Fig. 1. The total mill feed, as delivered from the crushing plant, had as its maximum size 1-in. material. This material was screened dry on Cole Zig-Zag screens fitted with 4-mesh openings; the oversize going to one pair of coarse rolls, the undersize to five Wilfley tables fitted with Butchart National riffles. The coarse-roll product was elevated and screened wet on Zig-Zag screens fitted with 4-mesh openings, the undersize going to the five tables mentioned above and the oversize to one pair of fine rolls. The fine-roll product was elevated and screened together with the coarse-roll product; so that, before any concentrating operations took place, the total mill feed was reduced to pass a 4-mesh screen. After concentration on Butchart-riffled tables, the tailings from these tables passed to de-sliming hoppers, where two products were formed, viz., an overflow and a discharge. The discharge constituted the feed for the Chilean mills. It contained 16 per cent. slime (—200 mesh) and 3 parts of water to 1 of solids. The concentrator was equipped regularly with four Chilean mills (Monadnock type) to the unit; but, in order to

Jan 1, 1917

By C. D. Hall, F. E. Dollarhide

Conventional static tests of fluid-loss agents do not realistically simulate conditions in a fracturing treatment. The dynamic tests reported here show that fluid-loss volume is better represented as proportional to time, rather than as the square root of time. This leads to a different equation for fracture area. The leak-off rate increases with increasing shear rate at the fracture wall, but appears to approach a limiting value. Pressure effects are minor. Spurt loss ordinarily is not affected by the flow velocity in the fracture and is inversely proportional to concentration of agent. The filter cake, once it is well established, is resistant to damage by the flow of plain fracturing liquid (without fluid-loss agent). The latter two findings indicate that a treatment employing a high-concentration spearhead followed by plain fluid can offer a more economical treatment under suitable conditions. INTRODUCTION The successful design of hydraulic fracturing treatments depends on accurate knowledge of the fluid-loss properties of the fracturing fluid. Howard and Fast,' in giving the basic equation relating fracture area to fluid and treating parameters, described three mechanisms which might control the rate of fluid leak-off from the fracture. One mechanism usually is dominant in a given well treatment. For each mechanism, the leak-off velocity is inversely proportional to the square root of time, and the proportionality constant is designated as the fracturing-fluid coefficient. For the wall-building type of fluid-loss agent, the coefficient is determined by a filtration test in a pressure cell, usually with a rock wafer or core as the filter medium. In these static tests, the cumulative volume generally is proportional to the square root of time, after an initial spurt volume. The static-fluid-loss test is not representative of the con,-&tions under which a fluid-loss agent performs in a fratturing treatment. The marked difference between the dynamic- and static-fluid-loss behavior of drilling fluids reported in the literature2,3 implies that dynamic testing is also needed with fracturing fluids. We have therefore undertaken a study of the dynamic-fluid-loss behavior of fracturing fluids. The testing apparatus has also afforded opportunity to evaluate the resistance of the filter cake to removal or damage by flowing fluid containing no fluid-loss agent, with and without sand. The results of these studies offer a means for more accurate evaluation of fluid-loss agent performance, and point the way to a "spearhead" fracturing technique which may offer more economical treatment for some wells. EXPERIMENTAL METHODS The dynamic-fluid-loss testing method is applicable to any type of wall-building fracturing fluid. The present study aimed first at finding what phenomena are involved, and therefore has been limited in the number of materials tested. All of the results specifically reported herein are for kerosene containing a commercial solid fluid-loss agent, which is commonly used at 50 lb/1,000 gal of oil. Another agent in liquid form, used usually at 20 ga1/1,000 gal oil, has shown all the same phenomena in dynamic tests, and generally the same level of fluid-loss control as the solid agent. The dynamic-fluid-loss core cell used in all tests is shown in Fig. 1. The fracture was simulated by the an-nulus between a 2.03 in. OD sandstone core and the surrounding pipe. Annulus widths of 0.234 and 0.117 in. were used, and the core was 3.5 in. long. The annular geometry provides a uniform fluid velocity and a well-defined shear rate over the entire filtering surface, and permits a large filter area (144 sq cm) in a reasonably compact cell. The leak-off fluid passed into a 0.5 in. diameter axial hole in the core. A hollow steel rod through this hole was threaded into a rounded "streamliner" upstream of the core, and into a mounting stud downstream. The streamliner and the stud had the same outside diameter as the core. In all tests except those where sand was circulated, the mounting stud had protruding rings which constricted the annulus, to minimize any tendency for channeling of the fluid to the side exit port. The ends of the core were sealed by Neoprene, steel and Teflon washers. The leak-off fluid was conducted from the hollow rod to an exit tube, through a metering valve (a fine-pitched needle valve) and a quick-opening toggle valve in series, and into graduated cylinders for volume measurement; Two separate circulating systems were used in the experimental program. The extensive initial testing was done at 50 to 150 Psi. The fluid was circulated by a variable speed Moyno pump, and the flow rate was read by a rota-meter flow meter. The filtration Pressure was supplied by holding a back-pressure with a throttling valve. The discharge streamcould be diverted into any of four sections

Jan 1, 1965

By R. G. Davies

The activation energy for steady state creep above -500°C is observed to be independent of the applied stress although it varies from -67 kcal per mole at 2 at. pct Si to -100 kcal per mole at 11 at. pct Si due to changes in crystallographic order. The magnitude of the activation energy, by comparison with Fe-A1 alloys, indicates FeSi type of order in certain alloys. X-ray results confirmed the presence of FeSi type of order. It is proposed that dislocation climb is the rate controlling mechanism for all the alloys. It has been demonstrated that when a diffusion mechanism is the rate controlling process, the formation of a superlattice in brass,1 Fe3A1,2 Ni3Fe,3-5 and Feco6 1) increases the creep resistance, and 2) increases the activation energy for steady state creep. Furthermore, a study of creep in Fe-15 to 20 at. pct A1 alloys7 has revealed that as the alloy composition approaches the long-range order field, there is an increase in the activation energy for steady state creep which is thought to be due to an increase in short range order. Fe-A1 and Fe-Si alloys are similar in that they both form the DO3 superlattice in which aluminum or silicon atoms have only iron atoms as first and second nearest neighbors. There are, however, two important differences between the alloy systems: 1) The superlattice formation at -350°C commences at -10 at. pct si8 as compared to -20 at. pct Al,9 and 2) Fe-A1 alloys form a FeAl (B2 type) super-lattice where aluminum atoms have all iron first nearest neighbors even at 22 at. pct Al, but so far no similar FeSi superlattice has been observed. With the similarity between Fe-A1 and Fe-Si alloys in mind, alloys of iron with 2 to 11 at. pct Si were examined for variations with composition of the activation energy for steady state creep and of creep strength. The temperature range of greatest interest was above 1/2 TM (TM is the absolute melting temperature) where it is usually observed that diffusion is the rate controlling process. A subsidiary X-ray investigation of the Fe-Si system was undertaken in an attempt to define the position of the order-disorder boundary as a function of cooling rate. EXPERIMENTAL DETAILS a) Creep. Specimens whose gage length was 1.5 in. and with a cross-section 0.04 by 0.08 in. were strained in tension by a lever-arm arrangement, and the load was adjusted between each creep test to maintain constant stress. The apparatus and mode of operation have been fully described in a previous publication.7 As each test produced a creep strain of 0.25 pct, the variation in stress during the test was negligible. Creep strain was measured at the end of one of the alloy steel grips by a displacement transducer with the out-of-balance potential being recorded on a variable speed recorder. The full-scale deflection of the recorder could be varied in steps to give limits of sensitivity of between 0.1 and 0.001 pct creep strain. The alloys, Table I, were made available by the Metallurgical Department, National Physical Laboratory (N.P.L.), england,10 and by the Research Department, General Electric Co. (G.E.), Schenectady, N.Y. They were hot worked at -850°C, warm worked at 550° to 650°C, and recrystallized in vacuum at -750°C to give a grain diameter of -0.1 mm. All the alloys had a very low impurity content; those from the N.P.L., for which a complete analysis is available,&apos;&apos; show carbon less than 0.026 pct, manganese less than 0.006 pct, and oxygen plus nitrogen less than 0.0024 pct. b) X-ray Procedure. A General Electric XRD-5 X-ray set with a focussing lithium fluoride mono-chromator in the diffracted beam, and a pulse height analyzer to eliminate harmonic wavelengths of the cobalt radiation, was used to investigate the structure of several very fine grained (grain diameter <.01 mm) Fe-Si alloys after the following heat treatments: 1) Quenched from 700°C, 2) slow cooled from 650°C (-40°C per hr), and 3) very slowly cooled from 400° to 100°C (10°C per hr with a 24 hr anneal every 100°C). The method of obtaining the diffraction pattern over the range of 20 from 15 to 45 deg was to count for at least 100 sec every l/3 deg with a slit subtending 1 deg in 20 at the focus; the probable counting error was less than 2 pct.

Jan 1, 1963

By N. W. Lord

The process of molten-zone refining is analyzed for long ingots and many zone passages. Formulas are derived which give the resultant impurity distribution in terms of finite series. A comparison with the approximate procedure of Hamming is given. HE physical principles and applications of an extremely physicalprinciple efficient form of metallurgical refinement has been described by Pfann. The purpose of the present paper is to describe a method of analyzing exactly the particular program used which enables the segregation effect to be predicted for any number of molten-zone passages in a long ingot. The method is applied to the particular case of refinement of an ingot whose impurity initially is uniformly distributed throughout its length. A number of molten zones of equal length are passed through the ingot effecting a radical redistribution of impurity. Pfann has indicated an approximate method, due to R. W. Hamming, of calculating the resultant concentration after each successive zone pass for a particular value of the segregation constant defined in his paper. Here a solution will be presented in terms of the number of zone passes and the segregation constant. The expression, though cumbersome, is exact and susceptible to ordinary numerical computation procedures. The results of a similar computation using the procedure of Hamming are presented in a table together with the exact results of the present method. The discrepancy in terms of absolute concentrations is tabulated for the first eight zone-lengths. To establish the notation (which follows that of Pfann1 as closely as possible) and physical basis of the analytical equations, the physical model and principal assumptions may be reviewed. An alloy of two elements, where there is formed a continuous range of solid solutions, usually does not melt as a simple compound. Rather, a temperature is reached where the solid solution is in heterogeneous phase equilibrium with a liquid solution of different composition. The temperature dependence of these equilibrium compositions forms part of the phase diagram. For very small concentrations of a solute B in a solvent A, this usually takes the form of Fig. 1. Sometimes the solidus and liquidus slope upward. This corresponds to a segregation constant (defined below) which is greater than unity. The segregation constant is now defined as k = Cn(x)/CnI(x) [1] where C,,(x) is the impurity concentration in the solid ingot at distance x during the nth passing of a molten zone and Cnl (x) is the impurity concentration of the liquid zone from which the solid at dis- tance x is formed (see Fig. 4 of ref. 1.) C (x) remains the same after passage of the zone. The constant k may be either greater or less than unity in general. Purification in the former case is effected only in a finite ingot and in the portion that is melted last. For k less than unity purification is effected even in an infinite ingot. The method which follows gives, in the former case, the successive increases in impurity concentration and, in the latter case, the successive decreases in concentration. The general case of impurity redistribution will be considered first, and purification will be discussed later on. The analysis rests on the following assumption: The movement of the zone is too rapid to allow appreciable atomic rearrangement in the solid sections and too slow to disturb the uniform impurity distribution in the liquid zone characteristic of equilibrium. Hence, the composition in the solid at the left solidifying interface will be determined by Eq. 1 while the impurity concentration of the liquid zone will be uniform throughout its length. The reasoning which follows closely parallels that of Appendix 11 in Pfann&apos;s paper. It is reviewed here for the case of the nth zone pass in order to make clear the meaning of an operator essential to the present method. Fig. 4 of ref. 1 shows the movement of a molten zone of length 1 in an ingot of total length d. Each Cn(x) can be determined from the condition that the amount of solute added to the zone during an incremental advance, dx, is due to the melting in of a solid portion C(x)dx and the freezing out of kCnl(x), that is d I —r- CnL (x) dx = Cn-1 (x+l)- kCnL (x) dx or, in terms of Cn(x) d k k —— C,(x) +—c.(x) =— Cn-1 (x + l). [2] dx l l This, of course, is derived from the main assumption, the fact that 1 is constant, and that the total impurity content previously present up to x + 1 is constant . A correction has to be made for the region (d — nl) < x < d. This is due to the zone length changing during the passage of the solidifying interface beyond x = d — 1. Since the general solution would be too complicated otherwise, only the region 0 < x < d — nl is considered. The general solution of Eq. 2 is

Jan 1, 1954

By M. C. Fuerstenau, D. W. Fuerstenau

Studies of the collection of alkali and ammonium halides utilizing vacuum flotation techniques and contact angle measurements show that ionic size controls the flotation of techniquesthese halides with amine salts measurementsas collector. Contact angles of air bubbles on sylvite in saturated brines were withaminemeasured salts asascollector.a function of such variables as collector addition, length of collector chain, and pH of the brine. No contact occurs between halite and an air bubble in brines containing dodecylammonium acetate as collector. LONG-CHAINED aliphatic amine salts have been used for the separation of sylvite (KCl) from halite (NaCl) by flotation.1,2 It is puzzling how these two minerals, which are so similar chemically and crystallographically, can be separated by this method. Gaudin" has postulated that the difference in floatability of halite and sylvite with salts of primary amines depends on ionic size: In the case of amine flotation, the cation would attach itself to the chloride. I have a speculation there, which I cannot prove, that the ammonium group, that is the —NH3 group in the amine, floats potassium chloride because the dimensions of this grour, as it has been measured in other compounds is almost identically the dimensions of the potassium ion, quite different from the sodium ion, and so it fits where potassium had been, in place of it and not attached to it. Apparently, because an aminium ion (RNH3+) is much larger than a sodium ion, it cannot fit into the lattice of halite. Taggart also has speculated that ionic size may control the floatability of sylvite.4 The object of this experimental investigation has been to test this hypothesis and to study what controls the adsorption of cationic collectors at the surface of sylvite. Since collection is to be approached from the viewpoint of ionic size, the ionic radii that are of interest in this work are presented in Table I. The values of the ionic radii of the ions listed in Table I, except NH4+, are those given by Pauling." Several different values for the radius of the ammonium ion have been given, but that of Goldschmidt6 seems to be preferred. The radius of the charged head of a dodecylammonium ion is assumed to be the same as that for the ammonium ion. Little experimental work has been reported in the technical literature concerning the separation of sylvite from halite by flotation. Guyer and Perren studied the separation by flotation of 50 pct binary mixtures of NaCl, KC1, NH,Cl, NaNO3, KNO3, K2SO4, and Na,SO, using either oleic acid or a sodium sul-fonate as collector.&apos; It is possible to measure floatability under actual flotation conditions where all three phases, air- water-mineral, are present by vacuum flotation tests and contact angle measurements.9 Both of these techniques were used in the experimental approach in this paper. Experimental Method and Materials The vacuum flotation tests were run with glass-stoppered pyrex graduated cylinders. Twenty-five ml graduates were used to test the floatability of all salts studied except rubidium and cesium salts. For each test distilled water containing the desired collector concentration was saturated with the salt to be floated. Sufficient salt (—48 mesh) was added to leave about 2 ml of solids in the bottom of the graduate. After the graduate had been agitated several minutes to saturate the solution with air, a vacuum was applied. If the salt were floatable in the collector solution, the gas bubbles attached themselves to the particles, and the particles floated to the surface. In determining the floatability of the expensive Rb and Cs halides, the experiments were run in 10 ml graduates with about 11/2 ml of collector solution initially. Contact angles were measured in the usual manner except that the solutions had to be previously saturated with the mineral to avoid dissolution of the crystal. Solutions for studying contact angles were made by adding the desired amount of collector to a saturated brine, giving the collector concentration in molarity. The mixture was agitated until dissolution of the collector was complete, with the exception of those concentrations greater than about millimolar. At these high concentrations complete dissolution of the collector was impossible. The face of the mineral to be tested was a freshly cleaved crystal of halite or sylvite. The mineral was placed in the brine and conditioned with collector for at least 15 min, which was found to be long enough to obtain a maximum value for the contact angle. The temperature remained constant during each experiment. The experiments were run at 24°C ±2°C. For contact angle measurements, a crystal of halite from Carlsbad, N. M., was used. Several samples of sylvite were used in this work: a crystal of sylvite from Stassfurt, Germany; a crystal from Carlsbad, N. M.; and a crystal of chemically pure potassium chloride. Saturated brines were made from reagent grade chemicals and distilled water.

Jan 1, 1957

By J. R. O’Connor

Stimulated emission from Nd+3 in yttrium uanadate fYVOJ is reported. Single crystals of YVO4:Nd, obtained from Linde Col-p., have improved substantially in the last several months. Pulsed thresholds of YVO, laser rods are now approximately 2 to 3 5, cowparable to tliose for YAG:Nd. Yttrium uanadate crystalli~es irz a space group similar to zircon. All rare-earth vanadates have this structure. Rare-earth ions sucll as Nd+3 which substitutes for Y+3, aye situated irz a strong tetragonal crystal field which lacks inversion symmetry. This condition increases the p&apos;robability of the parity-forbidden f — f transitions. Yttrium anadate has strong absorption bands beyond -1000A. These are clue to Y-O, V-O charge transfer and molecular transitions. Under 2537 and 3660A irradiation pure YVO, fluroresces u bright yel-LOLO. This fli&orescence is corrzpletely quenched in YV04:Ncl crystals. This and other evidence of energy trut~sjer from the lattice is repovted. Optzcul atz 1user pvope 1-ties o! YI&apos;U4:E[t are brieJy described. THIS paper describes some of the optical and laser properties of Nd+3- and Eu+3-doped yttrium orthovana-date (YVO4). It reports laser action for the first time in this low-symmetry host. For some time we have pursued a research program concerning laser hosts&apos; wherein the rare-earth (RE) ion is situated at a site of low crystal symmetry so as to increase the probability of radiative transitions. Single crystals of doped and undoped YV04 are grown from iridium crucibles in an oxyhydrogen gas-fired furnace by a modified Czochralski technique.&apos; This material crystallizes in a D4li tetragonal space group of the zircon (ZrSiO,) type.3 All RE vanadates have the same structure and form solid solutions with YVO4. Therefore, it will be possible to investigate a variety of cross-pumped laser systems, as in the case of yttrium aluminum garnet (YAG).4 At present, ~d&apos;~-doped YAG is one of the most efficient solid-state lasers.5 Accordingly, most of the material to follow will compare the properties of YV04 to those of YAG. Fig. 1 shows the relative transmission of YAG and two types of YVO4. "Pure" YVO4 has a normal absorption edge and is colorless. A second type has a broad absorption peaking near 4200A and is yellow. Rubin and Van uitert6 suggest the yellow material is slightly reduced. Samples of each type are being investigated by electron spin resonance,7 but these studies are so far inconclusive. The pulsed laser threshold is much lower in the yellow material than in the colorless. Therefore, the absorption at 3500 to 5000.4 transfers energy to the Ndi3 ions. Photons of wave length between 2000 and 4500A cause undoped YVO4 to fluoresce at 4800, 5250, 5460, 5540, and 5750A. This emission, previously reported by Brixner and Abramsom,8 is partially quenched by EU&apos;~ and completely quenched by Ndt3 at room temperature due to energy transfer from the lattice to the RE ions. At low temperatures, some lattice fluorescence is present. This implies that the energy-transfer process is in part phonon-assisted. Although the YVO, single crystals used in this work were prepared from Y2O3 containing less than 0.01 pct rare-earth impurities, there is aopossibility that emission lines between 4800 and 5750A are due to dysprosium, terbium, and so forth. However, these lines are not observed in other compounds, prepared from Y2O3, such as YP04, Y2MoO6, and so forth. Furthermore, extensive absorption measurements on our "pure" YV04 single crystals between 0.4 and 5.0 failed to reveal any characteristic rare-earth lines. Fig. 2 compares the absorption spectra of Ndt3-doped YAG and YVO4from 0.6 to 1.0 . The Ndt3 absorptions are labeled according to free ion, R-S coupling. These term designationsQ are appropriate for YAG:Nd. They appear to be inappropriate for YV04:Nd. In YVO, neodymium must substitute for yttrium. The yttrium site is situated In a strong tetragonal field, where point symmetry is (42m) or possibly lower.1° However, the reduced splitting of the Stark components of the YV0,:Nd spectrum implies that the NdT3 ion is in a cubic site. The only plausible explanation for this discrepancy is that the Ndt3 ion is in a low-symmetry site, lacking inversion symmetry, so that a substantial admixture of 4f and 5d wave functions occurs. In this case, R-S coupling is not valid and J is no longer a good quantum number." Consistent with this view, the 4~ metastable level of YV04:Nd has an oscillator strength larger and a

Jan 1, 1968

By J. D. Defilippi, E. M. Gilbert, K. G. Brickner

FCC chromium-manganese-nitrogen (Cr-Mn-N) steels differ from most other fcc materials in that these steels undergo a ductile-to-brittle transition. Transformation to martensite is considered to be responsible for this behavior in some metastable Cr-Mn-N steels. However, very stable Cr-Mn-N steels also exhibit a ductile-to-brittle transition. The results of this study indicate that deformation faulting is the probable cause of the brittle behavior of stable Cr-Mn-N steels. Deformation faulting accounts for the ductile behavior of these steels in a tension test at -320°F and brittle behavior in an impact test at -320°F. Deformation faulting also accounts for the toPological features observed on the fracture surfaces of impact specimens of these steels. FACE- centered- cubic chromium-manganese-nitrogen (Cr-Mn-N) steels differ from most other fcc materials in that these steels undergo a ductile-to-brittle transition. Many Cr-Mn-N steels transform to martensite during deformation,l-5 and several investigatorsl-3 have suggested that the brittle behavior of these steels is caused by martensite formation. However, very stable Cr-Mn-N steels also exhibit brittle behavior. Schaller and Zackeyl reported that a very stable Cr-Mn-N steel (less than 3 pct martensite formed at -320°F) exhibited a transition temperature higher than that for steels in which large volume fractions of martensite formed during testing. The explanation given by Schaller and Zackey for this observation was that in the very stable steel the martensite, because of its higher interstitial content, was more brittle than that formed in their other steels. This explanation was questioned by Tisinai and samans4 and Baldwin.6 Moreover, because the toughness of stainless martensite at cryogenic temperatures is generally very low, this explanation does not account for Thompson&apos;s7 observation that small additions of nickel (1 to 3 pct) greatly improve the toughness of high nitrogen (0.35 pct) Cr-Mn-N steels. The present paper summarizes the results of an investigation of the low-temperature brittleness in very stable Cr-Mn-N steels. The importance of the mode of deformation on the toughness of these steels is discussed. Table I. Compositions of the Steels Invertigated, Pet Steel C Mn P S Si Ni Cr N - A 0.09 14.70 0.018 0.011 0.47 0.22 18.40 0.54 B 0.12 14.90 0.001 0.008 0.48 0.14 17.80 0.38 C 0.12 14.95 0.004 0.005 0.62 3.95 18.43 0.38 MATERIALS AND EXPERIMENTAL WORK The compositions of the steels investigated are shown in Table I. Steels A and B had compositions within the limits of a proprietary Cr-Mn-N stainless steel,* whereas Steel C was similar in composition to the proprietary steel except for its 3.95 pct Ni content. All steels were hot-rolled to 1/2-in. thick plate. The plates were subsequently annealed for 1 hr at 2000°F and water-quenched. Standard longitudinal and transverse Charpy V-notch impact specimens were machined from the annealed plates. Duplicate longitudinal and transverse impact specimens were tested at 212", 80°, 32", 0°, -100°,-160°,-200°,-256", and -320°F. Longitudinal tension-test specimens were also machined from the plates and tested at a crosshead speed of 0.05 in. per min at the aforementioned temperatures. The fractured impact and tension-test specimens of all three steels were examined to determine whether martensite had formed during testing. Magnetic, X-ray, electron-diffraction, and electron-microscopy techniques were used to detect the presence of martensite in the highly deformed areas of these specimens. Metallographic examination of highly deformed areas of impact and tension-test specimens revealed the presence of dark-etching bands, such as those shown in Fig. 1. These bands were observed only in deformed samples and were thought to be associated with the low-temperature brittleness of the Cr-Mn-N steels. Accordingly, a sample 1 in. wide by 3 in. long was cut from the 1/2-in.-thick plate of Steel C. This sample was surface-ground to a in. and then cold-rolled 60 pct at -320°F. Thin foils were prepared from the cold-rolled sample and examined in a JEM electron microscope. Brightfield, dark-field, and selected-area diffraction techniques were used to determine the cause of the dark-etching bands. Fractographic experiments were also performed. Impact specimens Of Steels A, B, and C were broken at -320oF, and the fracture surfaces of these specimens were immediately shadowed with carbon. The carbon replicas were examined in a Siemens electron microscope, and attempts were made to correlate the topological features of the fracture surfaces with the deformation mechanisms that could be occurring during an impact test of these steels.

Jan 1, 1970

By C. H. Li, R. J. Stokes

This paper compares the fracture behavior of tungsten rods in three conditions: recrystallized. recovered, and wrought. Notched specimens suhjected to a 50 in.-lb impact load showed ductile-brittle transitions at 700, 4.90°, and 440°C, respectinely. The recrystallized material had an equiaxed pain structure and jracbred by simple cleavage from a grain boundary source at all temperatures up to 700°C. The wrought and recovered material had an elongated fibrous structure and at low temperatures fractured by cleavage originating from the notch. As the transition temperature was approached cleavage was preceeded by more and more intergvanular splitting which deflected the crack front into planes parallel to the tensile axis. The enhanced toughness of wrought and recovered tungsten was attributed both to its inability to initiate cleavage because no pain boundaries were suitably oriented perpendicular to the tensile stress and to its inability to maintain cleavage because of intergranular splitting ahead of the crack. It has been appreciated for a long time in a qualitative manner that the room-temperature brittleness of fully recrystallized tungsten may be alleviated by working the material at relatively low temperatures.&apos; More recently this difference in mechanical behavior between wrought and recrystallized tungsten has been examined quantitatively by measurement of the tensile properties as a function of temperature. In these experiments brittleness has been expressed in terms of ductility or reduction in cross-sectional area upon tensile fracture or in terms of the bend radius before fracture under bending.&apos; This work has shown the existence of a fairly sharp transition from brittle to ductile behavior with an increase in temperature. The ductile-brittle transition temperature for recrystallized material is approximately 200°C higher than for wrought material. An increase in strain rate, small additions of impurity,&apos; or an increase in grain size4 shift the respective transition temperatures to higher values, but the difference between them remains approximately the same at 200°C. A number of explanations for this embrittlement by recrystallization have been given. It has been blamed either on the concentration of impurity at the grain boundaries, the increase in grain size, or the change in texture which occurs upon recrystallization. The present paper examines the effect of different heat treatments on the notch-impact behavior of commercial powder-metallurgy tungsten rods. The change in the ductile-brittle transition temperature for this method of loading and the fracture mode has been related to the different mi-crostructures produced by heat treatment. EXPERIMENTAL PROCEDURE Commercial swaged powder-metallurgy tungsten rods 1-3/8 in. in length and 1/8 in. in diameter were machined to introduce a sharp V notch 0.030 in. deep. To change the microstructure from that of the as-received wrought material some of the specimens were subjected to an anneal in nitrogen either at 1300° or 1400°C for 8 hr or at 1600° or 2000°C for 1/2 hr. The notched rods were then placed in a miniature Charpy-type impact machine and struck at their midpoint (opposite the notch) with a hammer designed to deliver 50 in.-lbs of energy. The strain rate at the base of the notch was estimated to be approximately 100 sec-1 at the instant of impact. The specimens were heated in situ to the desired impact temperature. The microstructures produced by the various anneals were studied by both X-ray diffraction and metallographic techniques. Fig. 1 reproduces the microstructures observed metallographically following a 10-sec electroetch in a 10 pct KOH solution. Fig. l(a) shows the elongated fibrous grain structure of the as-received material. Following the anneal at 1300" or 1400°C the grain structure was still elongated as shown in Fig. l(b) but the etch pits delineated dense polygonized dislocation arrays within many of the grains. Occasionally a relatively dislocation-free recrystallized grain was found growing into the matrix. The anneals at 1600° and 2000°C resulted in complete recrystallization and some grain growth. The grains produced at 1600°C were still slightly elongated as shown in Fig. l(c) whereas the anneal at 2000°C produced equiaxed grains. The changes in grain size produced the expected changes in the X-ray back-reflection patterns; there was no indication either in the as-received material or the annealed material of any preferred orientation. RESULTS a) Impact Behavior. Fig. 2 reproduces the ductile-brittle transition curves measured in the manner described in the previous section. It can be seen that under these testing conditions the as-received

Jan 1, 1964

By K. G. Wikle, J. G. Klein, W. W. Beaver

Consolidation of hydride process, electrolytic, calcium reduced, and comminuted thorium powder, as well as saw chips and lathe turnings, by vacuum hot pressing and by cold pressing-vacuum sintering was studied. The mechanical properties of the consolidated material in the extruded form are compared with those of wrought castings. AT present there little little industrial use for thorium metal, although it has some important though small scale applications in electronic equipment. Despite its high inelting point—about 1750°C —a low modulus of elasticity, 11.4xl0 si at 20°C;&apos; relatively low mechanical properties coupled with a high density, 11.7 g per cu cm; and an unusually high chemical activity with normal atmospheres limit any structural applications. The metal is utilized as an alloying element principally in magnesium. Pure thorium finds utility as electrodes in gaseous discharge lamps such as the high intensity mercury lamp&apos; because its low work function and high electron emissivity provide lower starting potentials and more uniform operating characteristics than other available materials. The metal is also found in photoelectric tubes used for the measurement of the ultraviolet spectrum." Thorium metal has been used in germicidal lamps of the cold cathode type as sputtered coatings on nickel in order to provide a low work function surface and a low starting voltage. Other applications have involved the radioactive properties of thorium for the production of ionized particles." The potential value of thorium is much greater than its present use pattern because of possible utility in the field of nuclear power. Th may be converted through nuclear reaction to a fissionable element U which should be capable of acting similarly to U in the g&apos;eneration of atomic power. Thorium has been reported to be about three times as plentiful as uranium in the earth&apos;s crust, placing it in the order of abundance of lead and molybdenum." Thus, it is of interest in augmenting the potential supply of fissionable material for nuclear power. Because of its high melting point, thorium is usually produced as a powder through the calcium reduction of its oxide or thermal reduction of halides by sodium, magnesium, and calcium. It may also be produced in flake form by electrolysis of fused alkali or alkaline earth chloricles and fluorides. Therefore, powder metallurgy assumes importance in the fab- rication of thorium metal shapes. Furthermore, it is rather difficult to obtain pure thorium by melting, as the molten metal reacts readily with graphite as well as oxide, carbide, and nitride refractories. These contaminate the melt with oxides, carbides, and metallic impurities." The current investigation was undertaken to examine the fabrication of thorium by powder metallurgy methods which have been used for the commercial production of beryllium and other metals.&apos; A sparcity of data concerning the comparative cold and hot compaction of thorium powders of different derivation existed. Therefore, all commercially available types were examined along with other experimentally produced thorium powders in order to round out the comparison of consolidated thorium powders with melted reguline metal. Review of the Literature By heating a mixture of ThC1, with potassium, Berzelius made the first thorium metal as an impure powder in 1828. Improvements in the basic process, increasing thorium assay to 99 pct, were made by several investigators including Arsem," Lely and Hamberger10 and Von Bolton." Calcium reduction of Tho, to make powders was investigated by Berger," Huppertz,&apos;" Kroll," and Kuzel and Wedekind.&apos;" A thorium powder produced by this method using a CaC1, fluxing agent assayed 99.7 pct, as reported by Marden and Rentschler.&apos;" Compacted and sintered, this product was found to be ductile, and could be fabricated into wire and sheet. Improvements of the calcium reduction process were made later" wherein CaCl, was eliminated from the reaction, producing metal assaying 99.8 pct Th. Further work by Lilliendah118 howed that a coarser metal could be obtained by the substitution of ThC1, or ThOC1, for oxide with consequent advantage of stability to atmospheric reaction. Reports on the technology of thorium developed in Germany during World War II have been made by Espe."&apos; Thorium powder of 99.5 pct Th was obtained by reduction of the oxide by calcium. Screening to —200 mesh, compacting with about 20 tsi, and sintering in vacuo at 1320" to 1360°C for 3 hr resulted in a porous sinter cake. The sinter cake was sufficiently ductile to be worked into bar, wire, and sheet which could be employed as electrode materials.

Jan 1, 1957

By V. L. Minear

This paper presents some of the problems encountered and solved by the construction industry during recent years while pressure grouting the foundation rock of dams. Pressure grouting has become "Standard Operational Procedure" and consequently many hundreds of thousands of bags of cement have been used for that purpose. There is nothing weird, mystical, or even particularly difficult about the basic idea. Holes are drilled in the rock and grout is pumped into them. This procedure is used to shut off the flow of water under the dam and to reduce the uplift on the dam. Like most other engineering and construction problems, the fundamental principle is simplicity itself, but some of the "minor details" are rather complex. When faced with the problem of actually forcing the grout into the rock, or keeping it there once it has been injected, an engineer frequently has need for all of the ingenuity and initiative that he has at his disposal. It is proposed to discuss these minor details at length, even at the risk of boring the reader. The difference between success and failure in a grouting program can well hinge upon some such detail that may not be fully understood. Pressure grouting is a method used to make good foundations from bad ones by plugging actual or potential passages within the rock through which water might escape and thereby diminish or even destroy the usefulness of the dam. In some special cases, pressure grouting has been used for the purpose of consolidating broken rock to increase its bearing power. Lacking this useful method, many of the dams constructed during recent years could not have been built. Most of the really excellent dam sites in this country have been utilized long ago. Consequently it becomes necessary to find methods of utilizing the less desirable sites. Pressure grouting plays a vital role in the methods developed. Engineers have learned the hard way, that the term "solid rock" is a misnomer. Solid rock does not exist to any considerable extent in nature. Practically all rock contains imperfections such as shrinkage cracks, bedding planes, joints, solution channels, folds, shear zones, crush zones, and the like. All these imperfections are weaknesses through which ground water is likely to flow. The function of rock grouting is to rectify these imperfections and make the rock as nearly "solid" as practicable. The location, direction, depth and number of grout holes to be drilled are important. Clearly, a hole must pierce fractures in order to serve a useful purpose, for grout cannot be forced into the rock itself. Were it possible to determine the pattern and extent of the cracks within the rock, it might be feasible to drill two or three strategically placed holes so as to make practicable the grouting of the whole integrated network of cracks from them. Such is not the case. While a study of the geology, surface indications, and drill cores is helpful in gaining an idea of subsurface conditions, in the final analysis, the seams must be located by trial and error methods. There can be no doubt that the closer the

Jan 1, 1949

By F. E. Rizzo, L. R. Bidwell, D. F. Frank

THE poor high-temperature oxidation resistance of the refractory metals becomes an important problem as higher operating temperatures are sought. A meaningful analysis of their oxidation behavior requires a knowledge of the phases that can appear during oxidation as well as their relative stability. Although several extensive investigations have been carried out to obtain this information for the W-0 system,1-10 serious inconsistencies still remain. First, the standard free energies of formation (AG°) determined in recent studies which used the solid electrolyte galvanic cell technique6,7,9 differ significantly from those obtained primarily from gas equilibration measurements.1,3,4,10 The disagreement is especially large for the highest oxides. These data have been used to predict the decomposition temperatures of WO2. 72 and WO2.90.1 A considerable extrapolation is required, resulting in a large uncertainty in the decomposition temperatures. Second, Bousquet and perachon3 have presented convincing evidence, ob- tained by gas equilibration measurements, in support of the long-contested existence of W30. Phillips and chang,ll on the other hand, were not able to observe this phase under similar experimental conditions. The purposes of this work, which is part of a continuing investigation of the W-Mo system, were to resolve the discrepancies between the AGO measured by gas equilibration and solid electrolyte galvanic cell techniques, to extend the latter measurements to lower temperatures, to redetermine the WO2.72 and WO2.90 decomposition temperatures, and to apply the more sensitive electromotive-force method to the question of the existence of W3O. I) EXPERIMENTAL PROCEDURE The measurement of ?G° of reaction by the use of calcia-stabilized zirconia solid electrolyte galvanic cells has been described previously.12 The method is based on the measurement of the difference in chemical potential of oxygen between an electrode containing the two oxides of interest and a reference electrode of known oxygen activity. Two different types of cell holders were used in this study. The first type of apparatus has been described earlier." It consisted of two solid electrodes separated by a disc of electrolyte. Both electrodes were exposed to the same stream of purified helium. With this design a gaseous transport of oxygen between the two electrodes is possible. The second type of apparatus, which has also been described earlier,14 effectively separated the two electrode atmospheres. It utilized a closed-end impervious, calcia-stabilized zirconia tube that was open to the atmosphere. The tube served as an electrolyte for the cell and provided a means for separating the anode compartment, which was held under a static vacuum, from an air reference cathode. The cells and reactions of interest are described in Table I. High-purity tungsten metal was obtained from the Fansteel Metallurgical Corp. and the Allied Chemical Corp. Table II contains an analysis of each material. WO3 was prepared by oxidizing the metal in air at approximately 600°C in high-purity recrystallized alumina boats and by the dehydration of reagent-grade tungstic acid in alundum boats. The two methods of preparing WO3 resulted in slightly different purities.

Jan 1, 1968

By J. T. Williams

The equilibria in the V-Zr alloy system were investigated by solidus temperature determinations, thermal analysis, dilatometry, electrical resistance measurements, microscopic examination, and X-ray diffraction analysis. There is a eutectic reaction at 1230°C between a compound, V2Zr, and a solid solution containing 10 pct V in ß zirconium. V2Zr decomposes at 1300°C into liquid and a solid solution containing about 10 pct Zr in vanadium. The eutectic composition is probably about 30 pct V. A eutectoid reaction between V2 Zr and a zirconium takes place at 777°C at a very high rate. The eutectoid composition is 5 wt pct V. The limit of solubility of zirconium in vanadium was estimated to be 5 pct at 600°C. No attempt was made to determine the liquidus for the system. THE recent availability of large quantities of high purity zirconium has stimulated the study of zirconium binary systems. The equilibrium diagram for the V-Zr system has received little attention, however. Wallbauml appears to have made the first report concerning the equilibria in these alloys. He reported the existence of a compound, V2Zr, having the MgZn2 ((214) type of structure with a, = 5.277 kX and c° = 8.647 kX. Anderson, Hayes, Roberson, and Kroll2 made a survey of some potentially useful zirconium binary alloys and found that zirconium probably dissolves a small amount of vanadium. They reported the probable existence of a compound between the two elements and suggested that the zirconium-rich solid solution undergoes a eutectic reaction with this compound. Pfeil," in a critical review of the existing information, estimated that the solubility of vanadium in zirconium is less than 4.7 pct and probably less than 1.8 pct. Rostoker and Yamamoto&apos; proposed a partial diagram for the V-Zr system in a survey paper on vanadium binary alloys. Their diagram indicates the compound, V,Zr, a eutectic reaction at 1360°C, a peritectic reaction at 1740°C, and a limit of solubility of zirconium in vanadium of about 3 pct. They obtained no information on the equilibria in the zirconium-rich alloys. In view of the potential utility of the V-Zr alloys and the incomplete knowledge concerning the equilibria in the system, an attempt was made to establish the constitutional diagram. Preparation of the Alloys Raw Materials: The vanadium for making up these alloys came from the Electro Metallurgical Corp. Zirconium came from two sources. In the beginning of the investigation, sponge zirconium from the Bureau of Mines was used in making some of the alloys. Later, iodide metal made at the Westinghouse Atomic Power Development Laboratories became available. This material was used in the preparation of all the dilatometric and resistance specimens and about two-thirds of the solidus temperature specimens. A typical manufacturer&apos;s analysis of the vanadium is shown in Table I. No other analysis of the vanadium was made. The metal contained a dispersed second phase and did not have a sharp melting point. Typical results of spectrographic analysis of the Westinghouse zirconium are shown in Table 11. These data indicate a very high purity. The Bureau of Mines sponge metal was probably less pure but had good ductility. Melting: All of the alloys used in the investigation were made by melting pieces of vanadium and zirconium together in a dc electric arc furnace similar to those of Geach and Summers-Smith, craighead, Simmons, and Eastwood," and others. Melting was done in an atmosphere of helium scavenged of residual air by the preliminary melting of a separate charge of zirconium. Each ingot was turned over and melted at least three more times before removal from the furnace to aid in the attainment of homogeneity. Alloys prepared for use in the investigation are listed with the results of solidus determinations in Table III with the exception of the following compositions upon which no solidus determinations were made: 0.29, 0.54, 4.57, and 5.55 pct V. Analysis: The weight of each ingot made from iodide zirconium was within 0.1 g of the total weight of the initial charge, about 90 g. Since each component of each charge was weighed to the nearest 20 mg for amounts less than 10 g and to the nearest 0.1 g otherwise, the gross composition of an ingot could be calculated accurately. Chemical analysis for the vanadium content of several alloys agreed

Jan 1, 1956

By Donald B. Evans, Robert D. Pehlke

The solubility of nitrogen in liquid Fe-A1 alloys has been measured up to the solubility limit for formation of aluminum nitride using the Sieverts method. The activity coefficient of nitrogen decreases slightly with increasing aluminum content in the range of 0 to 4 wt pct Al. Based on a nitride composition, AlN, the standard free energy of formation of aluminum nitride from fhe elements dissolved in liquid iron has been determined to be: ?F" = -59,250 + 25.55 T in the range from 1600º to 1750ºC. The solubility of nitrogen in liquid iron alloys and the interaction of nitrogen with dissolved alloying elements in liquid iron have been the subject of a number of research investigations.&apos; Most of this work, however, has been reported for concentrations well below those necessary for the formation of the alloy nitride phase. Data in the concentration region near the solubility limit of the alloy nitride, particularly for systems exhibiting stable nitrides, are important in evaluating the denitrifying power of various alloying elements. They are also useful in determining the stability of a given nitride if it is to be used as a refractory to contain liquid iron alloys. In view of the importance of aluminum as a deoxidizing agent in commercial steelmaking and the fact that its nitride, AIN, is a highly stable compound and has merited some consideration as an industrial refractory, the following investigation was undertaken. The use of the Sieverts technique provided a measurement of the equilibrium nitrogen solubility in liquid Fe-A1 alloys as a function of nitrogen gas pressure up to 3.85 wt pct A1 in the temperature range of 1600º to 1750°C. The values obtained by the Sieverts method were checked by means of a quenching method in which liquid iron was equilibrated with an A1N crucible under a known partial pressure of nitrogen gas, and the solubility of A1N in liquid iron determined by chemical analysis. EXPERIMENTAL PROCEDURE The theoretical considerations involved in determining the solubility product of a solid alloy nitride phase in liquid iron by measuring the point of departure of the nitrogen gas solubility from Sieverts law have been discussed by Rao and par lee.&apos; The principal problem is to determine the variation of nitrogen solubility in an alloy as a function of the pressure of nitrogen gas over it with sufficient precision to establish the break point in the curve at the solubility limit of the alloy nitride phase. A fairly large number of data points are required to do this. A second problem is the determination of the composition of the precipitated solid nitride phase. This is necessary in order to define completely the thermodynamic relationships. The Sieverts apparatus used to make the nitrogen solubility measurements in this investigation is of essentially the same design as that described by Pehlke and E1liott.l The charge materials were Ferrovac-E high purity iron supplied by Crucible Steel Co. and 99.99+ pct pure aluminum. Recrystal-lized alumina crucibles were used, and were not attacked by the liquid alloys. The hot volume of the system which was measured for each melt ranged from 46 to 50 standard cu cm and was found to decrease linearly with decreasing pressure and with increasing temperature. The temperature coefficient of the hot volume at 1 atm pressure of argon gas was essentially constant for all experiments at a value of -6 X 10-3 cu cm per "C. The melt temperature was measured with a Leeds and Northrup disappearing filament type optical pyrometer sighted vertically downward on the center of the melt surface. The temperature scale was calibrated against the observed melting point of pure iron taken as 1536°C. The emissivity of all melts was assumed to be that of pure iron, taken as 0.43. The charge weights ranged from 110 to 140 g and the range of aluminum contents covered was from 0 to 3.85 wt pct. Aluminum additions were made as 12 to 15 wt pct A1-Fe master alloys previously prepared in the system under purified argon. The compositions of the master alloys were checked by chemical analysis and found to be in agreement with the charge analyses. Vertical cross sections of the master-alloy ingots were used as charge material for the equilibrations in order to minimize the effect of any segregation which might have occurred during solidification of the master alloys. Determinations of the solubility product of

Jan 1, 1964

By Charles Luner

ALTHOUGH the kinetics of the atmospheric oxidation of tin have been studied,1-3 the kinetics in pure oxygen have not been reported. This note presents some results of the kinetics of the oxidation of tin foil in pure oxygen. The tin foil used in this investigation was 99.78 pct pure, the two major impurities being Pb (0.17 pct) and Sb (0.024 pct). The foil was degreased and air-dried prior to oxidation. The air-formed film was estimated by coulometric reduction4 of the surface oxides to be about 25A. The samples, which were in strip form and had a total geometric surface area of about 300 sq cm, were first evacuated at room temperature in an oxidation bulb to 1 X 10 -5 mm Hg and subsequently degassed for 1 hr at the temperature of the experiment. Spectroscopically pure oxygen was then introduced into the bulb. The rates of oxidation were measured by following the rate of pressure change due to oxygen consumption in a closed system, the volume of which was sufficiently large so that the pressure change during an experiment was 10 to 15 per cent of the total pressure. An oil-mercury manometer,5 with a sensitivity of 3 X x mm Hg (which corresponds approximately to 5 X 10 -7 gm of oxygen) was used to measure the pressure change. The oxidation bulb was heated by an oil-bath and controlled to within -10.3 C. The reproducibility of the rates was good, as determined by the coincidence of rates at 190oC for duplicate experiments. Fig. 1 shows the oxidation curves in the temperature range 168 to 211.5oC at 8 mm Hg pressure. The effect of pressure on the rate of oxidation at 190°C is also shown. The rate of oxidation was not markedly affected by changes in pressure in the range 4 to 9 mm Hg, except perhaps at the lowest pressure, 4 mm Hg, where after about 70 min there appeared to be a decline in the rate, as compared with the rate at the higher pressures. At the lower tem- perature, 168oC, and for short reaction times, it is difficult to distinguish between a linear relationship and a slight curvature in the oxidation rate. However, for longer times, there was a somewhat larger indication of curvature. The samples after oxidation were gold-colored in appearance. The oxide that was formed at 170°C was identified as a-SnO by reflection electron diffraction. This is in agreement with results obtained by other workers.6, 7 The rate data best fitted the logarithmic-oxidation rate law, w = K1n ln(1 + -t/to) [1] where w is the weight of oxygen consumed, t is the time, and K In, and to are constants. First estimates to were obtained according to the method of Taylor Thon8 and were then employed to obtain the best simuItaneous estimates of to and K In, that would fit the logarithmic expression. These estimates were found numerically, so that the residual variance could be minimized. The calculations were performed on an

Jan 1, 1961

By D. H. Woodard

One of the principal uses of polarized light in metallurgy is to show the granular structure of metals by contrasting reflections. This use is confined largely to anisotropic metals, such as beryllium, tin, magnesium, and the like.1 Recently, Hone and Pearson2 have extended the use of polarized light to include the isotropic metal aluminum. This was made possible by the formation of an anodic film that was optically anisotropic and that had a fixed and reproducible structural relationship to the grains in the substrata. With their method, it was possible to indicate localized textures in annealed aluminum sheet. In the course of an investigation of plastic deformation at the National Bureau of Standards, it was found that the isotropic single phase alloy Monel could be made optically anisotropic by treatment with the usual Monel contrast etchant. Experiments soon showed that dif- , ferences in orientation could be detected easily. Previous investigations involving the use of polarized light have been restricted to annealed metals. In this paper it will be shown that inhomo-geneous plastic deformation which results in lattice bending and development of deformation bands, experienced by individual grains in polycrystalline Monel, is revealed in the microstruc-ture. This interpretation of the micro-structure was confirmed by a study of slip lines formed after prior plastic strain. Materials and Procedure The material used was Monel, the chemical composition of which was C—0.17 pet, Mn—1.28 pet, Fe—1.61 pet, S—0.009 pet, Si—0.17 pet, Cu— 30.70 pet, Ni—66.04 pet. The material was divided into 2 lots which were treated as follows: Lot A was heated at 1700°F for 3 hr to produce a fully annealed, coarse grained structure. Lot B was heated at 1400°F for 1 hr to produce a close approximation of a fully annealed structure without appreciable increase in the grain size of the original material. After each of these heat treatments the material was cooled in air. Samples for metallographic examination were cut from tensile and compression test specimens and were prepared by grinding on lead-tin laps, followed by electrolytic polishing, † Etching and the development of ridges were avoided by keeping the voltage above 3.8 and by preventing contamination of the polishing solution with water. Monel, when polished and treated with most etchants, is optically isotropic. However, when given a light etch for about 4 sec with the Monel contrast solution the surface of the metal becomes optically active. The microstructure when observed under crossed nicols* appears different from that viewed by ordinary light. This is shown by a comparison of micrographs made of the same area of a plastically deformed specimen using polarized light (Fig la) and white light (Fig 16). If Monel, which has been lightly treated with the contrast etchant, is slightly deformed by compression, rectangular figures are apparent over the specimen. However, within too short a period to be photographed, these rectangles become circular, probably because of surface tension effects within the etched film. The two observations,

Jan 1, 1950

By John M. Bostwick, James C. Bradbury

Bitumen is a general term for a group of materials composed of mixtures of hydrocarbons that are fusible and soluble in carbon disulfide. Included in the group are petroleum, asphalts, asphaltites, and mineral waxes. Asphalt is a dark-colored, cementitious solid or semisolid, which generally softens on heating and resumes a more viscous or solid state when cooled. It is found in native form in fissures and pore spaces in rocks and as lakes. It is also obtained from the residues produced in petroleum refining. Asphaltites are black, naturally occurring solid bitumens. They are generally shiny and break with a conchoidal fracture. Common asphaltites are gilsonite (uintahite), grahamite, and glance pitch. Mineral waxes are solid or semisolid, relatively nonvolatile bitumens of various colors, having an unctuous feel and waxy luster, Of the native waxes, ozokerite is the best known. Montan wax is also classed as a natural wax, although it never occurs in the pure state and must be extracted from lignite or cannel coal (Ladoo, 1920). Mineral waxes may also be obtained by pyroprocessing of petroleum, lignite, and peat. Pyrobitumens are dark-colored, solid hydrocarbons that are infusible and relatively insoluble in carbon disulfide. They are so-called because they will yield liquid and gaseous hydrocarbons only when heated. Asphaltic pyrobitumens are defined as "substantially free of oxygenated bodies" and include wurtzilite (elaterite), albertite, and impsonite (Thrush, 1968). Nonasphaltic pyrobitumens contain oxygenated bodies; examples are coal and peat. The nonasphaltic type will not be discussed further. Petroleum Asphalts Asphalts refined from crude petroleum have almost completely replaced the native types. Refineries can produce asphalts in every desirable grade more economically and in greater quantity than the native types and with a variety of properties that meet the requirements of the various manufacturers using them. For all practical purposes, it is estimated that more than 95% of the asphalts consumed in the United States are refined from crude petroleum (private communication). Native Asphalts Native asphalts are dark brown or black, semisolid or solid substances that, in theory, formed by slow, natural fractionation of crude petroleum at or near the surface of the earth (Goulston, 1938). They are colloidal systems of high molecular weight components dispersed in a medium consisting of low molecular weight components (Pracy, 1954). Native asphalts may occur as lakes, as impregnations of rock (rock asphalt) or as deposits consisting of mixtures of natural asphalt and mineral matter in various proportions (Stern, 1960). Practical use of native asphalt dates to about 3000 B.C. It was mined from surface pits or lakes and used, then as now, for masonry mortar, waterproofing aqueducts and tunnels, binders for road construction, and as an article of commerce (Miller, 1938). Generally the native asphalts find little commercial use today except where they are geographically abundant, easily mined and transported, or have specific or special application. Petroleum-refined asphalt is produced by unusually economical techniques and is much less costly to use. The largest and best known of the lake deposits is that on the Island of Trinidad, where the lake covers an area of 114 acres and is 285 ft deep at its center (Stem, 1960). The lake asphalt is refined by heating to evaporate in-

Jan 1, 1983