Search Documents

By E. J. Ripling

A NY mechanism proposed to explain hydrogen embrittlement in titanium and its alloys must, of course, be consistent with the experimental data that characterize this embrittlement. Unfortunately, however, the mechanical behavior of hydrogen-bearing titanium, at least a-ß titanium, has not been unequivocally defined. Lenning, Craighead, and Jaffee have clearly shown that hydrogen cmbrittles a-titanium by elevating its transition temperature, probably as the result of the formation of titanium hydrides. Therefore, in these alloys, hydrogen acts like at least one other interstitial contaminant, namely, nitrogen.&apos; On the other hand, ß-titanium has been shown by these same investigators to tolerate very large amounts of hydrogen without suffering severe mechanical damage.2 Mixing these two phases, however, to form the most important class of commercial alloys, the a-ß alloys, again results in severe hydrogen embrittlement, although the mechanism by which the embrittlement is produced is not of the same type as that which causes brittleness in a alloys. Ductility damage due to hydrogen increases as the strain rate is reduced in a-ß alloys, while embrittlement in a alloys increases as the strain rate is increased, since the latter is a transition temperature behavior. Steel, like the a-ß alloys, becomes more hydrogen sensitive at slow strain rates, suggesting that the mechanism producing the embrittlement in these two metals is similar. Brown and Baldwin&apos; described the hydrogen-produced ductility depression in steel as a function of testing temperature and strain rate by defining the slope of the two surfaces that produced the depression in a three dimensional chart, Fig. 1. One of these surfaces was given by the equations (de/de)4 > 0 (de/dT) < 0 [1] while the other was defined by the pair of equations Surfaces of the type given by Eq. 1 are suggested in two ways. One is an embrittlement mechanism wherein the diffusion rate of hydrogen is competitive with the rate at which the material is being deformed, as suggested by the planar pressure theory of Zapffe and his co-workers.&apos; The other is the diffusion controlled extension of Orowan&apos;s theory on delayed fracture in glass by Petch and Stables.&apos;&apos; Surfaces of the type given by Eq. 2 are also compatible with a mechanism of pressure build up in voids, according to de Kazinczy,&apos; since the solubility of hydrogen in the metal increases with testing temperature so that as the temperature is raised, the pressure in the voids is reduced. Kotfila and Erbin recently presented some data on the dependence of ductility on testing temperature and strain rate for the a-ß 3 pet Mn complex alloy at four different hydrogen levels.H Although data were presented for only three testing temperatures at three different strain rates, their results indicated that surfaces of the types defined in Eqs. 1 and 2 are produced in the alloy when the hydrogen level is sufficiently high—200 and 300 ppm—Fig. 2. Jaffee and his co-workers presented data on a number of different a-ß alloys which indicated the existence of surfaces of the type described by Eq. 2, but the ductility recovery at low temperatures as given by Eq. 1 was not found." In an attempt to aid in crystallizing this description of the ductility dependence of hydrogen-bearing a-ß alloys, tests were conducted by the author on a-ß titanium 140A with three different hydrogen contents. The tensile properties of two as-received rods and one vacuum annealed rod* were obtained over a range of testing temperatures and strain rates as shown in Figs. 3 and 4. Hydrogen analyses were made by the Battelle Memorial Institute on four pieces of the rod whose properties are shown as solid circles in Fig. 4. The hydrogen content of these pieces, taken at widely spaced intervals within the rod, were 280, 270, 289, and 270 ppm, indicating that the hydrogen content within a single as-received rod was quite uniform. One of the broken test pieces whose properties are shown in Fig. 3 as solid circles was also analyzed, and found to have a hydrogen content of 310 ppm. The analyses obtained on two of the vacuum annealed specimens were 92 and 170 ppm.

Jan 1, 1957

By Sidney Fischler

Large striation-free single crystals of bismuth have been grown from the melt by rapid freezing. Zone-refined bismuth, together with doping impurities if desired, is placed in a shallow flat-bottomed graphite boat and melted in air with a propane hand torch. The torch is then withdrawn in a manner which causes the melt to freeze direction-ally. Crystallization, which resuires only a few minutes, usually results in the formation of a single crystal even when a seed crystal is not used. Crys -tals of any desired orientation may be grown by using oriented seeds. Undoped crystals grown by this method have residual resistivity ratios greater than 200. THE growth of large single crystals of bismuth by either the Czochralski or horizontal zoning technique is not entirely satisfactory. Specifically, difficulties are encountered in producing single crystals of the required dimensions in all desired orientations, and striations caused by low-angle polysynthetic twins are frequently present in the crystals. In addition, both methods are time-consuming and require special apparatus of some complexity. A simpler method has now been developed for growing large striation-free bismuth single crystals of desired orientation in a short time. Fig. 1 shows a typical setup consisting of a rectangular graphite boat which contains zone-refined bismuth, a 1/8-in.-thick flat quartz plate which covers the entire inner bottom of the boat, and three additional quartz plates about 1/4 in. thick which are used to separate the bismuth from the graphite everywhere except at a small area at the left of the boat. The graphite boat is 1 in. high, and its sides and bottom are about 1/8 in. thick. The quartz plates should be smooth and clean. The graphite boat is heated from the right with a propane torch, as shown in Fig. 1, until the bismuth is completely melted. The melt has the shape of a triangle with a narrow neck at the apex farthest from the torch. The melt is frozen direc-tionally by gradually moving the torch toward the right, away from the boat. The bismuth in contact with the graphite, at the left end of the neck, freezes first. The freezing interface then moves down the neck into the main bulk of material, where it develops a convex shape ideal for the continuation of single-crystal growth. The interface continues to move through the melt until the entire bulk is solid. The entire procedure may be completed, in air, in a matter of minutes. The technique described almost always yields a single crystal whose basal plane is nearly perpendi,cular to the bottom of the graphite boat. In earlier experiments, in which the bottom of the melt was in direct contact with the graphite boat, single crystals were grown with basal planes parallel, perpendicular, or at some intermediate angle to the bottom of the boat. At times the orientation of the bulk of the material differed from the orientation of the material in the narrow neck. In these cases, a nucleation site initiated the growth of a differently oriented crystal, and the thermal conditions favored the new orientation over the initial one. The thermal conditions depend on a number of factors, including the heating technique, the placement, shape, and thickness of the quartz plates, the thickness of the walls and bottom of the graphite boat, and the quantity of bulk bismuth employed. All of these factors, plus the initial orientation and the presence and effectiveness of nucleation sites, will determine the orientation of the final large single-crystal slab. When a crystal of specific orientation is desired, an oriented section of a rapid-freeze crystal is shaped by spark cutting and grinding for use as a seed. To grow a doped crystal, the desired impurity is placed in the graphite boat together with the bismuth chunks and seed. Crystals doped with mercury, cadmium, lead, and selenium have been grown. The rate of freezing is so great that the distribution coefficient of any impurity approximates unity. On a gross scale, therefore, impurities should be more homogeneously distributed in rapid-freeze crystals than in Czochralski or zoned crystals. Because of the possibility of constitutional supercooling, however, it is quite possible that impurities are not homogeneously distributed on a microscopic scale in the rapid-freeze crystals. Generally the single crystal slabs which have been prepared are initially 5 to 7 mm thick. Thicker crystals may be obtained by using one of these slabs as a seed. The slab is placed in a graphite boat resting on a large aluminum block, either air- or

Jan 1, 1964

By O. G. Paasche, Yuan-Shou Shen

THERE is a disagreement among the various authors about the exact manner of transformation of ZrCr2. Rostokerl and others2 stated that ZrCr2 had a C-14 (MgZn2) type of structure below 1000°C and a C-15 (MgCu2) type of structure at temperatures above 1000°C. Alisova3 and others4 reached the opposite conclusion and stated that the transformation temperature is close to the melting point of ZrCr2. A literature survey shows that various investigators3&apos;= who homogenized the specimens at a temperature higher than 1000°C have concluded that ZrCr2 had the C-15 structure at room temperature. Meanwhile, Jordan et al.4 reached similar conclusions without annealing the specimen. Other investigators1,2,6,7 who X-rayed the specimens in the as-cast condition without annealing reached different conclusions. The investigation reported herein was conducted with the aim of exploring the exact manner of transformation of ZrCr2 by various heat treatment tests. The alloys for this examination were prepared from iodide-reduced zirconium crystal bars, 99.9 pct purity, and electrolytic chromium, 99.9 pct purity. They were melted in a nonconsumable electrode arc furnace with water-cooled copper crucible in a helium atmosphere. The melting loss of each alloy was less than 1.5 pct by weight. Chemical analysis of a randomly selected specimen indicated that there was a very close agreement between calculated and analyzed compositions. Before being heat-treated each specimen was encapsulated in a vycor or quartz tube inside which an argon atmosphere was maintained at a pressure of lower than 1 atm. In determining the crystal structure of each specimen with a Debye-Scherrer camera, the standard procedure8 for X-ray quality analysis (Hanawalt method) was followed. The different series of heat treatment tests in this investigation are tabulated in Tables I and 11. The tests in Series I, specimens from 1-1 to 1-9, which were similar to Rostoker&apos;s experiment1 indicated that the transformation temperature seemed to fall between 870° and 900°C and that the crystal structure of ZrCr2 at lower temperature seemed to be of the C-14 type. However, once the compound is transformed to C-15 type, it is impossible to reverse the transformation back to the C-14 type by first heating the specimen above 900°C and then annealing it slowly below 900°C as shown in Experiments II-1 to II-3. Thus, it appears that the specimen of ZrCr2 will transform from C-14 to C-15 structure when heated above 900°C but will not transform from C-15 to C-14 when annealed slowly passing 900° C even after the extremely slow cooling process such as indicated in the experiment of Specimen II-3. As a valid transformation temperature is a temperature at which the transformation is reversible, therefore the temperature 900°C (or other temperature close to 900°C) is not the transformation temperature for ZrCr2 and the C-14 structure is not the stable structure of ZrCr2 at lower temperatures. The C-14 structure is retained at room temperature because the transformation to C-15 structure is very sluggish and the fast cooling after melting does not allow enough time for the transformation to take place. Additional energy is required to alter the metastable condition of the C-14 structure. The sluggishness of this transformation was again demonstrated through another series of experiments. Four specimens with C-14 structure were taken. Then they were annealed at 900°C but each specimen was soaked for a different period of time, Table 11. X-ray diffraction patterns of this group indicated that the C-14 structure gradually disappeared as the soaking period was lengthened. The figures listed under the column "C-14 Structure, pct" were estimated from the intensity of the d = 2.330 line of the diffraction pattern corresponding to the structure. Notice that the intensity of this line became weaker for longer soaking periods. To determine the transformation temperature of ZrCr2, specimens with C-14 structure (as-cast condition) were annealed at 1300°, 1400°, 1500°, 1550°, and 1600°C, respectively. A final specimen was first heat-treated to 1500°C in order to transform it to C-15 structure, then heat-treated at 1600°C again. From the X-ray analyses of this series of tests, Specimen Nos. III-1 to III-6, it is evident that a transition from C-15 structure at lower temperatures to the C-14 structure occurs at some temperature between 1550° and 1600°C.

Jan 1, 1969

By H. R. Crawford, P. B. Crawford, A. F. Tragesser

A method has been developed to calculate wellbore temperatures during mud circulation and the actual cementing operation to aid in the design of cement slurries. The method agrees within 10F with previously measured values. The calculation technique provides temperatures, as functions of time, at varying depths in both the casing and annulus. The technique also provides this information if a relatively cool cement slurry is pumped into the well immediately following circulation of hot mud. Circulating bottom hole temperatures of brine and a bentonite mud were measured. INTRODUCTION As wells are drilled deeper, greater demands are being made on all phases of the industry, and new technology has been developed to provide satisfactory well completions. However, little or no work has been conducted on accurately determining bottom-hole, static and circulating temperatures.. In designing a cement slurry, such factors as density, fluid loss control, viscosity, deterioration from ternperature, compressive strength and pumping time must be considered. Individual well conditions often make it necessary to include still other factors. Pumping time is a primary consideration and, as wells are drilled deeper, encountering higher bottom-hole temperatures, this property becomes even more important. Cement slurries must be designed with sufficient pumping time to provide safe placement in the well; however, the slurry cannot be overly retarded as this will prevent the development of satisfactory compressive strength. The pumping time of a specific cement is currently obtained by subjecting the cement to simulated conditions of temperature and pressure. A reasonably accurate bottom-hole pressure may be obtained by considering hydrostatic heads of fluids, friction pressure and wellhead pressures. However, accurately determining bottom-hole temperatures is much more difficult. Bottom-hole static temperatures are estimated by considering several sources of information, including logging temperatures, published temperature gradient maps and field experience. This information is usually questionable due to disagreement of data from the various sources. Temperature gradient maps were constructed based on temperatures recorded many years ago while running bottom-hole pressure tests. These thermal gradients then represent an average of well conditions and cannot always apply to a specific well. Also, logging temperatures may be affected by the time since fluid was last circulated, rate of penetration, circulating rate and many other factors. Therefore, even though logging temperatures are available, the question still exists as to the correction factor that should be applied to obtain an accurate static temperature. After obtaining static bottom-hole temperature, it is then necessary to relate this to circulating temperatures actually encountered by the cement slurry. This is accomplished by selecting a test schedule from the API RP-10B corresponding to the estimated well conditions.&apos; The API-recommended practice for testing oilwell cement provides testing schedules for various well depths and conditions. These schedules are intended to simulate down-hole conditions during cementing. They provide a rate at which both temperature and pressure are increased until the estimated circulating conditions are reached. These testing schedules represent circulating temperatures for an average well and, although there is flexibility in choosing the test schedule that most accurately simulates the temperature of an individual well, it still is not possible to consider all the well conditions that will affect the bottom-hole temperature. Many factors affect cement temperatures; for example, the length of time a well has remained static prior to running casing and cementing, the circulation time, the temperature of fluids used in cementing, fluid density and flow properties of fluids. The pumping time for a typical retarded cement could vary from 2 to 4 hours with a 10F change in testing temperature. Variations in pumping time are the most critical in highly retarded cements used in deep, hot wells; yet, predicting bottom-hole circulating temperatures is more difficult in these wells. This work was conducted to develop a means of calculating circulating temperatures as a function of well depth, casing and hole size, pumping rate and time, fluid and reservoir physical properties and thermal status of the well. PREVIOUS WORK In 1941 Farris reported on a study to develop information leading to a more practical laboratory evaluation of oilfield cementing mixtures and performance.&apos; It was then recognized that the pressure factor was being neglected,

By V. Koump, T. F. Perzak, R. G. Olsson

The rate of dissolution of recrystallized alumina in molten iron oxide in equilibrium with iron at 1450°C was studied by rotating alumina disks in the melt. The samples were rotated from 1 to 12 min at speeds from 54 to 3270 rpm. It is concluded from the experimental results that the process is limited by diffusion in the liquid boundary layer. The interdiffusion coefficient for the dissolution of alumina disks in pure molten iron oxide was estimated to be about 3 xl0-5= sq cm per sec. THE erosion of refractory materials is of considerable importance in iron- and steelmaking processes. Since many refractories contain alumina, this investigation was undertaken to develop a further understanding of the erosion mechanism. Shurygin et al.1-3 have investigated the rate of dissolution of alumina in molten silicates and molten fluorides by rotating alumina disks in various melts. The rates were reported to be limited by diffusion in the liquid boundary layer. In the present investigation the rate of dissolution of alumina in molten iron oxide in equilibrium with iron was studied in a similar manner; alumina disks were rotated at different speeds in molten iron oxide at 1450°C. The relationship between the rate of dissolution of the flat surface of the disk and the speed of rotation was used to determine whether the process is limited by diffusion in the liquid boundary layer or by some chemical reaction mechanism. APPARATUS AND EXPERIMENTAL PROCEDURE A schematic diagram of the experimental apparatus is shown in Fig. 1. The iron oxide melts were contained in 2-in.-diam high-purity iron crucibles that were centered in a cylindrical graphite susceptor. Prior to each sequence of experiments, sufficient reagent-grade ferric oxide was added to a crucible to form approximately 350 g of molten iron oxide; the depth of the melt was about 11/2 in. To prevent excessive erosion of the crucible, about 70 pct of the required iron was added to the crucible in the form of iron powder. The mixture was melted in an atmosphere of purified argon by induction heating with a 250-kc generator. After the oxide was molten, a Pt/Pt-10 pct Rh thermocouple at the exterior of the crucible was calibrated against an iron-sheathed thermocouple immersed in the melt and was thereafter used to regulate the melt temperature. All experiments were conducted at 1450°C with a probable error of 10°C. The alumina disks were constructed from pure re-crystallized alumina having a density of 3.7 g per cu cm (porosity approximately 6 pct). For each disk sample a 1-cm-diam alumina cylinder of known length (0.9 to 1.0 cm) was cemented into a ground seat at the end of a 1.27-cm-OD centerless ground alumina tube. The ends of the tube and cylinder were flush and served as the disk surface. When mounted in the apparatus the alumina tube was guided by tungsten carbide bearings and rotated by a variable-speed motor. The bearing and drive assembly could be raised and lowered in order to allow the sample to be quickly moved in and out of the melt. At the start of an experiment, a sample was preheated for several minutes directly above the melt, and then immersed in the middle of the melt to a depth of a in. and rotated for 1 to 12 min at speeds of 64 to 3270 rpm. The speed of rotation was measured by a hand tachometer. At the start and completion of each experiment, a melt sample was withdrawn with a cold copper rod and the alumina content was determined by chemical analysis. During individual experiments the alumina concentration of the melt changed by the order of 0.25 wt pct. Average melt compositions for individual experiments were used in the subsequent computations and were in the range between 0.12 and 4.3 wt pct alumina. The extent of dissolution was determined by sectioning the end of the sample and measuring the final length of the alumina cylinder with a microscope. Since the exposed end of the cylinder, i.e., the disk surface, was etched in a manner characteristic of this geometry, a number of measurements were made on

Jan 1, 1969

By B. H. Caudle, A. B. Dyes

In a recent publication&apos; it was shown that wells with a free surface in a homogeneous gravity-drainage reservoir have a hyperbolic decline with index n &apos;. This paper reports efforts to extend this theory to practical field cases. As a first step in making the extension, a large number of lease production curves from gravity-drainage fields in the stripper stage were fitted with the general hyperbolic decline equation, q/q, = 1/(1 + y) y and n in each case being determined by statistical methods. It was found that this equation is reliable for decline curve extrapolation for periods up to at least 25 years. The resrilts obtained were disturbing in the sense that values for the index, n, were often smaller than one, which fact leads to the prediction that a well would produce all infinite amount of oil in infinite time. A theoretical investigation based on the results of the previous paper provides a suitable explanation for these low values of n. It suggests that these are obtained if production occurs from two or more layers of diflerent permeabilitv and thicktress or two lnyers having different skin effects. It is predicted that as the more permeable layers are exhausted, the index of decline, n, will rise from a value smaller than 2, to n = 2. A practical method is given for fitting the general hyperbolic decline Eq. 6 to production decline data. When this method is used, the index of decline observed in gravity-drainage reservoirs in the stripper stage when wing only data over early years 01 decline usually agrees quite closely with that determined using data over 25 years or more. This indicates that the hyperbolic decline equation is reliable for predicting behavior of such reservoirs. A method is also suggested for taking into account any changes in the value of the index INTRODUCTION As a result of both experimental work with models and theoretical deductions, it was recently concluded&apos; that wells having a free surface in a homogeneous reservoir producing by pure gravity drainage should have a hyperbolic decline with n = 2. The present report details efforts to extend this work on ideal homogeneous reservoirs to actual field cases. It should be remembered that both the previous work and the present extension are based on the assumption that the reservoir is at very low pressure, and hence, nearly all the potential available for driving fluid to the well is the gravity head. Attempts to apply conclusions of this study to reservoirs which have considerable pressures supplementing the gravity drive may lead to difficulty. WELLS WITH NO FREE SURFACE In a dipping gravity-drainage reservoir such as that shown in Fig. 1, some of the down-dip wells have no free surface. The previous study&apos; showed that the up-dip wells, which have a free surface, should follow Eq. 1. The following approximate analysis, of a type which was justified experimentally for up-dip wells in the work just mentioned, leads to results for the down-dip wells.

By G. W. Neighbor, M. S. Abrahams

The actice slip plane in InSb is found to be of the {111} type by using the method of two-tmce anulysis. Measurements of the rotation of the tensile axis with increasing plastic shear strain indicate that the tnacroscopic slip direction is a <110>. Recrystal-lization is first observed at a sheav strain of about 70 pct for a crystal deformed at a temperatzcre of 473° C and at a shear strain rate of 2.90 times 10- %ecc-I; the dislocation density corresponding to this value of the strain is 1 x 108 em-&apos;. At a shear strain of about 170 pct, the recrystallization process is complete. This appears to be the fimt time re-crystallization has heen observed in a semiconductor. V HE operation of the {Hi} glide plane in Ge was first noted by Gallagher.&apos; This finding was confirmed by Treuting who also extended the result to Si. Using X-ray techniques, Treuting3 demonstrated that the macroscopic slip direction in Ge was of the <110> type. The choice of the (111) <110> slip system in Ge is exected since in the diamond-type structure the {Ill7 plane is the plane of highest density and the <110> direction is the direction of closest packing.4 There is little reason to expect that the glide elements of InSb are different from those of Ge, since the zinc-blende structure characteristic of InSb is quite similar to the structure of Ge. With regard to InSb, Allen5 has indicated that the glide plane is the {111} type. There appears to be no quantitative determination of the glide direction, however. The present work was undertaken to establish both the glide plane and glide direction in InSb and also to see if recrystallization occurs at large shear strains. EXPERIMENTAL PROCEDURE Mechanical Deformation. The single crystals of InSb used in this investigation were deformed in uni-axial tension. The tensile axis of the specimens was [145] and the periphery was bounded by (111) and (321) planes as shown in Fig. 1. The metallographi-cally polished specimens were deformed in a dynamic atmosphere of He, using an Instron machine. The tensile tests were performed at temperatures in the range of 300°C to the melting point (525OC) and at shear strain-rates varying between 1.94 X 104 sec-&apos; to 1.94 x 103 sec-&apos;. Additional details concerning the preparation and deformation of the specimens have been presented previously.&apos; Analysis of the Glide Elements. The traces of the active slip plane in both (111) and (321) planes were analyzed according to the usual method of two-trace analysis." The macroscopic slip direction was determined by measuring the rotation of the tensile axis as a function of plastic strain,8 These measurements were made from Laue back-reflection X-ray photographs. The X-ray beam was perpendicular to the (111) plane. A given crystal was used for only one measurement of the rotation of the tensile axis after it had been strained to a predetermined value. Therefore, care was taken at all times to insure that the X-ray beam was parallel to the [Ill ] direction of all the samples. RESULTS AND DISCUSSION Glide Plane. After 3.4 pct plastic shear-strain, at a temperature of 369°C and a shear strain-rate of 1.94 x 104 sec-&apos;, the slip traces on the (111) and (3.21) planes appear as in Figs. 2(a) and 2(b), respectively. The results of an analysis of these glide markings demonstrate that the slip plane is (111). The operation of this plane is expected on the basis of the orientation of the specimen, since glide on this plane and in the [I011 direction will result in

Jan 1, 1962

By R. M. Waterstrat

The phases occurring in the V-Mn system were studied by means of X-yay diffraction and metallo-paphic techniques, using are-melted alloy specimens annealed in the temperature range 800° to 1150°C and quenched. The bcc solid solution extends at 1250°C all the way from vanadium to 6-manganese. Below 1050°C the a-phase is formed, and the terminal a-manganese phase is stabilized up to about 900°C by vanadium in solid solution. IN the only previous general survey of the V-Mn system Cornelius, Bungardt and Schiedtl reported the existence of three intermediate phases corresponding to the approximate compositions VMn,, VMn, and V5Mn. The phase VMn8 has recently been identified as a o phase2 but the alloy VMn was found to have a bcc structure2 corresponding apparently to the vanadium solid solution rather than to the large cubic unit cell reported by Cornelius et al. 1 Subsequent work by Rostoker and Yamamoto3 has shown that the vanadium-base bcc solid solution extends to at least 15 pct Mn at 900°C. An alloy corresponding to the composition VMn, was examined by Elliott,4 who reported that the as-cast sample as well as samples annealed at 1200o and 1300°C had bcc structures, but that annealing at 1000°, 800") and 600°C produced two phases. One of these phases was apparently the bcc solid solution and the other resembled the o phase structure. Hellawell and Hume-Rothery5 established the phase relationships in manganese-rich alloys above 1000°C, and showed that the o phase in this system is replaced by the 6 Mn (bcc) solid solution at temperatures above 1050°C. These results suggest that a continuous bcc solid solution may exist above 1050°C between vanadium and 6 Mn. The present investigation was undertaken in order to develop more complete information in regard to this system. EXPERIMENTAL METHODS The alloys used in the present work were prepared by arc-melting electrolytic manganese having a minimum purity of 99.9 pct and vanadium lumps with a purity of 99.7 pct. The major impurities present in these metals were carbon, nitrogen, and oxygen and this would account for the small percentage of nonmetallic inclusions observed metal-lographically. The arc-melting was at first performed under a helium atmosphere and it was necessary to keep the melting times as short as possible in order to minimize the loss of manganese by vaporization. It was later found that the evaporation of manganese was considerably reduced when the melting was done under argon atmosphere. The final composition of each alloy was calculated by assuming that the total weight loss during melting was due to evaporation of manganese. Compositions which were calculated in this manner agreed reasonably well with the results of chemical analysis, as shown in Table I. Spectrographic analysis revealed the presence of contamination by tungsten, but in no case was the percentage of tungsten greater then 0.4 at. pct. The specimens were in each case broken in half and the fractured section was examined visually and microscopically for evidence of inhomogeneity. Each specimen was homogenized at temperatures near l100°C, as shown in Table I. After this treatment most specimens consisted of large columnar grains of the bcc vanadium solid solution. The etchant used in most of the metallographic work consisted of 20 pct nitric acid, 20 pct hydro-flouric acid, and 60 pct glycerine. It was found that this etchant would clearly delineate the phases present in these alloys although it does not produce any striking contrast between the phases. For certain manganese-rich alloys, a 1 pct aqueous solution of nitric acid was used. This etchant gave a brown color to the a-manganese phase, whereas the o phase was virtually unattacked and appeared very light as shown in Fig. 1. The etchants used by Cornelius et a1.l were found to produce spurious effects in some of these alloys. In particular, the vanadium-rich alloys etched in hot sulfuric acid often appeared to consist of two phases when both X-ray diffraction and etching with the glycerine-acid mixture indicated the presence of single phase bcc solid solution. A few percent of what appears to be an oxide or nitride phase was found at the grain boundaries and in the interior of the grains, especially in the vanadium-rich alloys. All alloys were annealed in sealed silica tubes containing 1 atm of pure argon and these tubes were then quenched in cold water. Although some manganese loss occurred during annealing, the loss seemed to be confined to the surface of the speci-

Jan 1, 1962

By A. V. Bradshaw, R. I. L. Guthrie

The behavior of large bubbles in the size range 4 to 25 cm3, rising through molten silver, has been studied. It was found that rising velocities were equivalent to those in aqueous systems of low viscosity. Mass transfer coefficients for oxygen bubbles dissolving in silver were found to be 0.036 ± 0.007 cm sec-1, being close to those predicted for transfer through the front surface of the spherical cap bubble only. It is suggested that the surface active nature of oxygen in silver could account for the relatively low coefficients obtained. MANY metallurgical processes involve interactions between gas bubbles and liquids. Examples include the removal of carbon monoxide in Open Hearth Steelmak-ing, the removal of sulfur by blowing air through copper matte during converting, and the removal of hydrogen from steel during vacuum degassing or inert gas flushing. The steps involved in such refining processes include; transport of the dissolved species to the bubble interface, adsorption and chemical reaction of the species at the interface, desorption of product molecules from the interface, and transport of product gas into the bulk gas phase of the bubble. It has been concluded1 that all the interfacial steps involved proceed so rapidly at steelmaking temperatures that transport of the solutes, present in the metal, become the important rate controlling factors provided nucleation phenomena are not restrictive. The O-Ag system was chosen for the investigation into gas bubble-molten metal interactions due to the relatively high solubility of oxygen that enables rates of oxygen transfer to be measured from changes in bubble volume. Other advantages of this system include the absence of a stable oxide phase at an oxygen pressure of 1 atm and the relatively low melting point of the metal which permits the use of a metallic container, providing that it is resistant to oxidation. In those metallurgical processes where bubbles have an important influence, bubble volumes are usually greater than 5 cm3. For this reason the present study relates specifically to single large bubbles of oxygen rising in silver. These bubbles adopt the characteristic spherical cap shape similar to that shown in Fig. 1 for a 30 cc bubble rising in water. After an initial investigation to determine the velocities of inert (nitrogen) bubbles rising in molten silver, experiments were carried out with oxygen and the rates of mass transfer between the oxygen bubbles and the silver were measured. EXPERIMENTAL Apparatus. The apparatus, Fig. 2, for containing molten silver, was constructed from "Nimonic 75" Alloy (75 pet Ni, 20 pet Cr, 5 pet Fe, Mn) and provided for the release of single bubbles from an hemispherical cup, situated at the bottom of the column. The cup was turned by translating the rotation of the drive shaft through 90 deg. This was accomplished by the use of a bevelled gear system, and a smooth drive was provided by the lubricating action of the silver on the gears. Since reliable high temperature seals at 1000°C were found to be impracticable, the filling and drive shaft tubes were extended outside the 3.5 kw resistance wire tube furnace, where connections were made using easily accessible O-ring seals. The apparatus remained gas tight to the atmosphere at pressure differentials far in excess of those used. The filling tube was connected via a small bore tube to the differential pressure transducer. Gas could be bubbled into the inverted cup from two i-in. tubes which passed down the inside of the column to preheat the gas. The temperature of the silver was maintained at 1020°C during all experiments. Measurement of Bubble Volume. In order to calculate mass transfer rates, it was necessary to obtain a continuous record of the bubble&apos;s volume during its passage through the column of molten silver. The method adopted for measuring the bubble volume involved closing off the top gas space to the atmosphere prior to each experiment, and recording the variation in gage pressure of this space during the formation and rise of the bubble. Since any change in bubble volume results in an equal change in top space volume, Boyles Gas Law may be applied (for isothermal con-

Jan 1, 1970

By A. L. Labbe

SOON after the Cottrell treater was placed in operation in the Murray plant in 1918 to treat combined lead sinter and Wedge roaster smoke, it was noticed that the power flowing through the treater did not remain constant. This was indicated by the varying milliamperes and also by the total amount of power consumed by the rectifiers. At times, for then unknown reasons, the treater current fluctuated through a wide range of from 40 to 300 milliamperes. Fortunately, these variations in power did not affect the treater&apos;s overall recovery, as this installation consisted of three independent units in series, a feature which made the Murray Cottrell an outstanding installation over many years of operation. Water conditioning of the smoke as a means of improving recovery was already known, but was not adaptable to this installation for reasons of excessive corrosion, which had been the case with other treaters using water as a conditioner exclusively. Tests conducted on the smoke had proved conclusively that water vapor played no part in the fluctuation of power, and the same was true of the SO2 contents and dust burden of the smoke. Neither did temperature variations of from 150 to 400°F have any effect on the power. Finally it was noticed that the variation in power taken by the treater, and the efficiency of recovery, had some definite relation to the number of Wedge roasters operating, and particularly to the sulphur contents of the charge. This observation soon led to the discovery that free sulphuric acid was the conditioner for Murray smoke and that variations in acid contents of smoke accounted for fluctuations of treater power. Analysis of recovered dust revealed that only a few hundredths of a per cent free acid was necessary to maintain a very efficient recovery. Once this knowledge was available, the roaster charge was adjusted as to sulphur contents to produce the necessary acid conditioner. For a number of years this practice was followed, but with improvements in the field of flotation, excess pyrites were eliminated from the smelting picture, so with a change in metallurgical practice we were confronted with the problem of providing the deficiency in acid by some other means. The situation was aggravated and our problem of acid conditioning made more difficult by the increase in the lead contents of concentrates roasted resulting from better flotation methods. The first step towards introducing acid vapors into the smoke stream by accessory means was accomplished by boiling sulphuric acid in cast iron pots placed in an open fire box. Acid fumes evaporated from the pots together with the combustion gases from the fire box were discharged into the flue through a cast iron pipe. Capacity of the iron pots was quite limited because of the comparatively slow rate of evaporation and the use of lump coal as a fuel, which did not lend itself to practical temperature control. This method of firing resulted in frequent pot failures due to overheating. In spite of these incidental difficulties, encountered in any new venture, the pot evaporator demonstrated the practicability of fuming sufficient acid to make up the deficiency in that naturally evolved in the Wedge furnace operation. As time progressed, less and less high sulphur material was available for Wedge roasting, and the necessity for accessory acid conditioning reached a climax when the supply of this material was entirely eliminated. Since the cast iron pots could only evaporate a few hundred pounds of acid per day, and were costly to replace, the logical thought presented itself of spraying the acid in a heated chamber. This new type of acid evaporator was simpler and more economical to operate and had sufficient capacity to fume several thousand pounds per day. The attached fig. 1 presents the outstanding features of the new furnace. A is the vaporizing drum, which is heated by the combustion gases of coal stoker fire box, B. Acid is sprayed into the hot gas stream within the drum through a number of air-acid atomizers, C, and the acid vapor, together with combustion gases, is introduced into the smoke stream through two cast iron pipes extending to different points in the flue. The operation of this type of furnace requires control of temperature inside the drum to prevent overheating, which dissociates the acid vapor to SO, and eventually to SO2, with resulting loss of acid. The dissociation to SO, at elevated temperatures is comparatively high, and in some cases accounts for as much as 60 pct of the acid sprayed into the drum. To reduce the loss of acid, through dissociation, some Cottrell plants, where the rate of acid sprayed reaches at times 9 ppm, use two evaporating units. Fuming the acid into the smoke stream is only

Jan 1, 1951

By H. O. McLeod, J. M. Campbell

This paper presents the results of the data obtained in the first stage of a long-range study at high pressures of the system, vapor-hydrate-water rich liquid-hydrocarbon rich liquid. The data presented are for the three-phase systems in which no hydrocarbon liquid exists. Tests were performed on 10 gases at pressures from 1,000 to 10,000 psia. One of these was substantially pure methane, and the remainder were binary mixtures of methane with ethane, propane, iso-butane and normal butane. Several conclusions may be drawn from the data. 1. Contrary to previous extrapolations, the hydrocarbon mixtures tested form straight lines in the range of 6,000 to 10,000 psia which are parallel to the curves for pure methane, when the log of pressure is plotted vs hydrate formation temperature. 2. The hydrate formation temperature may be predicted accurately at pressures from 6,000 to 10,000 psia by using a modified form of the Clapeyron equation. The total hydrate curve may be predicted by using the vapor-solid equilibrium constants of Carson and Katz&apos; to 4,000 psia and joining the two segments with a smooth continuous curve between 4,000 and 6,000 psia. 3. The use of gas specific gravity as a parameter in hydrate correlations is unsatisfactory at elevated pressures. 4. The hydrate crystal lattice is pressure sensitive at elevated pressures. INTRODUCTION Prior to 1950 many studies had been made of the hydrate forming conditions for typical natural gases to pressures of 4,000 psia.""&apos;"&apos;"" Most of these attempted to correlate the log of system pressure vs hydrate formation temperature, with gas specific gravity as a parameter. One of the more promising correlations was made by Katz, et al, which utilized vapor-solid equilibrium constants. The only published data above 4,000 psia are those of Kobayashi and Katz7 for pure methane to a pressure of 11,240 psia. In the intervening years, most published charts for the high-pressure range have represented nothing more than extrapolations of the low-pressure data, with the methane line serving as a general guide. The reliability of these charts has become increasingly doubtful (and critical) in our present technology as we handle more high-pressure systems. The portion of our high-pressure hydrate research program reported here was designed to: (1) investigate the reliability of existing charts; (2) obtain actual data on gas mixtures to 10,000 psia; and (3.) develop a simple hydrate correlation that was more reliable than those which simply used specific gravity as a parameter. Binary mixtures of methane and ethane, propane normal butane, or iso-butane were injected into a high-pressure visual cell containing an excess of distilled water. Hydrates were formed and then melted to observe the decomposition temperature of the hydrates at pressures from 1,000 to 10,000 psia. EQUIPMENT The equipment consisted of a Jerguson 10,000-lb high-pressure visual cell, a 10,000-1b high-pressure blind cell and a Ruska 25,000-1b pressure mercury pump. The visual cell was placed in a constant-temperature water bath controlled by a refrigeration unit and an electric filament heater. A Beckman GC-2 gas chromatograph was used in analyzing the gas mixtures after each run was completed. EXPERIMENTAL PROCEDURE After evacuating the gas system, the heavier hydrocarbon was injected into the high-pressure mixing cell to that pressure necessary to give the desired composition. This cell then was pressured to 1,100 to 1,200 psia by methane from a high-pressure cylinder. The mixing cell holding the gas contained a steel flapper plate and was shaken intermittently over a period of 15 minutes. After mixing, the valve to the high-pressure visual cell containing excess distilled water was opened, and the gas mixture was allowed to flow into the cell. The temperature in the water bath was lowered 10" to 15&apos;F below the estimated hydrate decomposition point. As a first check, the temperature was increased at a rate of 1°F every six minutes to find the approximate point of decomposition. It was again lowered 1.5° to 5°F to form hydrates. The temperature was raised to within l° of the estimated decomposition point and then increased 0.2F every 10 to 15 minutes until the hydrates decomposed. This procedure was repeated at various pressures to obtain 7 to 13 points for each mixture between 1,000 and 10,000 psia. After completion of the hydrate decomposition tests, the gas mixture composition was analyzed with a calibrated gas chromatograph. These gas analyses have an estimated error of ± .1 per cent.

By S. G. Proctor, D. L. Baaso, W. V. Conner

A double crucible system was developed for I-g scale plutonium reductions. The equipment consists of an inner MgO crucible, an outer MgO crucible, and a stainless steel pressure vessel. The reduction charge is PIaced in the inner crucible and the annulus between the inner and outer crucibles is filled with a mixture of calcium and iodine. The exothermic reaction between the calcium and iodine in the annulus supplies the heat required for complete reaction of the reduction charge ana&apos; good metal coalescence. Metal yields of 80 to 85 ,pct were obtained from I-g scale reductions and yields as high as 97.5 pct were obtained from 0.5-g scale reductions. The system was used to reduce charges containing as little as 0.1 g of Pu resulting in metal yields up to 90 Pct. THIN metal foils of highly enriched plutonium isotopes are used as targets for cross section and other measurements. The isotopes are separated in the calutrons at Oak Ridge and are very expensive to prepare. Often, only 1 or 2 g of material are available thereby emphasizing the need for a method of preparing these small quantities of metal. Procedures for 1-g scale plutonium reductions have been described, but these procedures require elaborate or expensive equipment. For example, the procedure described by Anselin et al.1 requires an inert atmosphere glovebox and induction heating equipment. The procedure described by Baker2 also uses induction heating equipment to obtain the recommended heating rates and liner temperatures. Baker&apos;s procedure also requires high purity PuF4 with very little PuO2 present. The procedure described in this paper resulted from a search for a simple and inexpensive method for making 1-g scale plutonium reductions. EXPERIMENTAL Equipment. The equipment required for the double crucible system consists of a pressure vessel, two MgO crucibles, a MgO crucible lid, and a resistance-heated, vertical crucible furnace. The crucibles were slip cast from high purity MgO and were supplied by the Coors Porcelain Co. and the Norton Co. The pressure vessel and lid were fabricated from 316 stainless steel. Materials. The PuF4 used for this study was obtained from the P1utoniu:m Metal Production Department at Rocky Flats. The PuF4 was prepared in a con- tinuous hydrofluorinator by reacting Pu02 with HF at 650°C. The isotopic composition of the plutonium was approximately 93 pct 239PU, 6 pct 240PU, and 0.5 pct 241PU. Some of this PuF, contained less than 1 pct Pu02 (Batch No. 6 and 7), while other batches contained up to 15 pct Pu02 (Batch No. 3). The chemical analyses of all the PuF4 used for this study is given in Table I. The calcium was 99 pct pure, AEC grade, and only that fraction which would pass through a 20 mesh screen was used. The I2 was USP grade re-sublimed I2 which was ground before being used. Procedure. Various charge compositions and methods for loading the double crucible system were tested. The optimum conditions for 1-g scale reductions are described below. The double crucible system was loaded as shown in Fig. 1. The outer crucible was placed in the pressure vessel and the annulus between them was filled with MgO sand. The inner crucible was placed in the outer crucible, supported by a layer of MgO sand in the bottom of the outer crucible. A mixture of 5.2 g of Ca and 25 g of I2 was placed in the annulas between the two crucibles. The upper portion of the annulas was filled with MgO sand. Next, a layer of a mixture of calcium and I, equal to 20 wt pct of the calcium and I2 used in the main charge was placed in the bottom of the inner crucible. The PuF, to be reduced was mixed with a 30 pct excess of calcium and 1 mole of I2 per mole of Pu and this mixture was placed in the inner crucible. The charge was topped with a layer of a mixture of calcium and I2 equal to 20 wt pct of the calcium and I2 used in the main charge. The pressure vessel was sealed with a flat copper gasket and was purged by alternately evacuating and filling with argon. The purge valve was closed and the vessel was placed in a vertical crucible furnace which was preheated to 950°C. The furnace was turned off after 20 min of heating and the vessel allowed to cool. Experience has shown that. this amount of heating is sufficient to assure complete reaction of the charge. RESULTS AND DISCUSSION The double crucible system has been used to produce plutonium metal on a 0.1- to 1-g scale. Reduction

Jan 1, 1970

By R. P. Dolan, W. M. Shedd, B. L. Buchanan

Analytical expressions are derived from empirical data relating the basic physical device parameters to the radiation dose. To put in perspective and justify the approach taken, the overall problem and methods of attempted solution of radiation damage to transistors in a missile or satellite environment is briefly examined. The main emphasis is on a new criteria for neutron radiation hardness and the device design implications resulting from using this criteria. Contrary to other derived criteria, the derived bipolar hardness relation depends mainly on the initial lifetime rather than the base transit time. It is shown, that contrary to previously accepted beliefs, the inherent attainable radiation tolerance of JFET&apos;s is at least a factor of ten greater than the radiation tolerance ofbipolars. Experimental results on JFET&apos;s designed for maximum neutron radiation tolerance are presented. The accuracy, utility, and range of possible validity of various published lifetime and carrier removal damage constants are discussed. It is well known that transistors are subject to damage due to radiation at much lower radiation levels than those which would cause structural damage to a missile or satellite. The three principal categories of radiation damage to transistors are displacement, ionization, and thermomechanical. The main cause of displacement damage is fast neutrons, and the main effects which permanently degrade transistor operation are minority carrier lifetime reduction and carrier removal. A second-order effect is mobility reduction. The main cause of ionization damage is gamma radiation which results in permanent effects due to charge build-up in the oxide and transient effects due to photo-current generation. The transient effects are usually handled by using circumvention techniques in the electronics system. The extent of the transistor degradation due to displacement and ionization damage depends strongly on the transistor type, whereas the extent of thermomechanical damage does not depend on the transistor type, but depends mainly on packaging and metallization which are common to all transistor types. The standard transistor types are categorized in Fig. 1 and typical cross-sections are shown in Fig. 2. The original reason for the designation of the two classes of transistors was that the unipolar operation depends on only one carrier (majority) and the bipolar operation depends on both minority and majority car- riers. The bipolar is then seriously degraded by minority carrier lifetime reduction, whereas the unipolar or field effect transistor is not. One of the two principle classes of field effect transistors, the junction field effect transistor (JFET), depends strongly on carrier removal. However, the insulated gate (MIS) is little affected by either carrier removal or lifetime changes, but is virtually eliminated as a radiation resistant device because of permanent charge build-up in the insulator caused by ionizing radiation. The bipolar and JFET suffer only minor effects due to charge build-up in the oxide passivation when compared1 to the MOS. This is because the active region of the MOS is adjacent to or part of the oxide layer. It is then the bipolar and JFET which are potentially resistant to permanent radiation damage. DERIVATION OF RADIATION TOLERANCE RELATIONS Criteria. The bipolar is by far the most widely used transistor type. The usual criteria for hardening bipo-

Jan 1, 1970

By H. E. Mauck

The many factors that control bumping must be carefully studied for each coal seam where bumps occur, and specifications known to exclude bumping should be incorporated in the mining plans. This calls for complete knowledge of the seam&apos;s characteristics and its adjacent strata, and in many instances these characteristics are not revealed until the seam is actually mined. Pressure and shock bumps, the two general types, occur jointly and separately. In this discussion no differentiation will be made. Whether pressure or shock, they are treated as bumps, and both must be eliminated. Bumps in mines have occurred in several places throughout the coal fields of the world. A study of many of these occurrences indicates that geologic characteristics, development planning, and mining procedure have contributed. But more specifically, there are conditions usually associated with bumps: thickness of cover, strong strata directly on or above the seam, a tough floor or bottom not subject to heaving, mountainous terrain, stressed and steeply pitching beds, and the proximity of faults and other geologic structures. Mine planning should incorporate these known factors (not necessarily in order of importance): 1) Main panel entries should be limited to those absolutely necessary to ventilate and serve the mine. This reduces the span over which stresses may be set up that will later throw excessive pressures on barrier and chain pillars when they are being removed. 2) Barrier pillars should be as wide as practicable so that they will be strong enough to carry the loads thrown on them when final mining is being carried out. 3) Pillars should never be fully recovered on both sides of a main entry development if the barrier and chain pillars are to be removed later. The excessive pressures placed on the main chain and pillar barriers by arching of the gob areas can result in bumping when these barriers are being removed. 4) Full seam extraction is better accomplished by driving to the mine boundary and then retreat-drawing all pillars. If there are natural boundaries in the mine—such as faults, want areas, and valleys —retreat should be started there. 5) Pillars should be uniform in size and shape. The entire development of the mine should call for uniform blocks with entries driven parallel and perpendicular. Only angle break-throughs should be driven when necessary for haulage, etc. 6) For better distribution of rock stresses and reduction of carrying loads per unit area, both chain and barrier pillars should be developed with the maximum dimensions. 7) Pillars should be open-ended when recovered. If they are oblong, the short side should be mined first. Both sides of a block should not be mined simultaneously, but under no circumstance should the lifts be cut together. 8) Pillar sprags should not be left in mining. If they are not recoverable, they should be rendered incapable of carrying loads. 9) Pillar lines should be as short as practicable. (Three or four blocks are adequate). Experience has shown that rooms should be driven up and retreated immediately. The longer a room stands, the more unfavorable the mining conditions. This contributes to bumping. 10) Pillars should not be split in abutment zones (high stress areas lying close to mined out areas) and if slabbing is necessary, it should be open-ended. 11) Pillars should be recovered in a straight line. Irregular pillar lines will allow excessive pressures thrown on the jutting points. Experience has shown that the lead end of the pillar line can be slightly in advance. 12) Pillar lines should be extracted as rapidly as possible. This appears to lessen pressures on the line and render abutment zones less hazardous. 13) Extraction planning should call for large, continuous robbed out areas. Robbing out an area too narrow to get a major fall of the strata above the seam tends to throw excessive pressures on a pillar line. 14) Timbering in pillar areas should be adequate but not excessive. Too heavy timbering or cribbing is likely to retard roof falls and throw excessive weight on the pillar line. 15) Experience has shown that when pillar lines have retreated 800 to 1000 ft from the solid, bumps can occur. Because this distance may vary in different seams, impact stresses should be studied for each individual condition. In any event, extra precautions should be taken against bumps in this area. This list of controlling factors may or may not be complete. It probably is not, but it covers most of the problem&apos;s significant aspects. The question is whether or not bumping can be eliminated. The answer is that bumping can be minimized and possibly eliminated if these and other established factors are thoughtfully considered and incorporated in the mining and extraction plans. If a mine has already been developed or the pattern set so that little change can be made, then it will be necessary to adjust to the most nearly practicable system that can incorporate the known factors.

Jan 1, 1959

By J. W. Edington

THE Hopkinson bar has become a popular technique for the measurement of the mechanical properties of materials deformed at high strain rate. Maximum use of the equipment is made in the arrangement first used by Kolskyl in which a short compression specimen is sandwiched between two pressure bars and is loaded by a single pulse travelling through the system. The pressure bars are used both to apply the load to the specimen and as transducers to obtain continuous strain-time histories of three pulses, incident on, reflected, and transmitted by the specimen. The data measured by the pressure bars can be analyzed in terms of the stress/strain behavior of the specimen.2&apos;3 However, one of the assumptions of the analysis of the observed pulses is that the total stress and total strain do not vary significantly from point to point within the specimen at any given instant during the deformation process. Although this assumption is generally justified for very short disc-like specimens2 the situation is uncertain for larger specimens. For example, at small plastic strains (-0.01) Hauser et al.2 have some evidence of small flucations in the total stress within the crystal during deformation, even in relatively short aluminum specimens. In addition, Karnes4 has shown that the plastic strain, and by inference strain rate, is different at each end of a compression specimen tested in a Hopkinson bar, although the length of the specimen was not specified. Recently, the mechanical properties and the dislocation substructure have been investigated in single crystals of columbium5 (length 0.25 in., diam 0.19 in.), and copper6 (length 0.5 in., diam 0.5 in.) deformed at high strain rates. As part of this research program the assumption that the plastic strain is constant throughout the specimen has been checked by measuring the total dislocation density as a function of position in the specimen. Compression specimens of the same orientations and dimensions were tested as described previously5,6 sing a split Hopkinson bar. Since any discontinuity in strain distribution is most likely to arise during the initial stages of deformation the investigation was performed on specimens deformed to plastic shear strains of 0.054 (copper) at a strain rate 1.2 x l03 sec-1, and 0.06 (columbium) at a strain rate 1.5 X l03 sec-1. The orientation of the single crystals is shown in Figs. 1 and 3. Thin foils were taken parallel to the most highly stressed slip plane, i.e., (111) in copper and (011) in columbium, using conventional disc techniques. The dislocation densities were measured using first order reflections with compensation for invisible dislocations.5&apos;6 In the copper single crystals the discs were randomly distributed throughout the cross section of the specimen. However, the dislocation density obtained from each disc was plotted vs the disc positions relative to the ends of the specimen. The results for the copper specimens are shown in Fig. 1. Clearly the dislocation density is constant throughout the main portion of specimen within the experimental error. The error bars on the dislocation densities correspond to a shear strain variation of 0.015 on the basis of previous measurements% ± of the rate of increase of dislocation density with strain in copper single crystals of the same geometry. Thus within this experimental error the plastic strain can be concluded to be constant within the specimen and the assumptions used in the analysis of the stress/time curves are therefore reasonably valid. The higher measured dslocation density near the impact end and the lower dislocation density at the bar end of the copper specimen is in agreement with the results of Karnes4 who showed that this strain/time curve rose to a maximum more rapidly at the impact end compared with the bar end. Hauser et al.2 have also pointed out that at small plastic strains (-0.01) the strain at the impact end of the specimen may be greater than that at the bar end. Thin foils taken from different points within the columbium single crystals demonstrated that the dislocation density could vary significantly within the specimen, see Fig. 2. Large areas of some thin foils up to 30 µ sq contained very few dislocations, see Fig. 2(a). However, in other parts of the compression specimen dislocation configurations like those shown in Fig. 2(b) existed over large areas (-30 µ sq). As a result, when the average dislocation density in a thin foil is plotted as a function of the position of the thin foil relative to the ends of the specimen, considerable scatter is observed, see Fig. 3. In this material then, the local dislocation density, and consequently the

Jan 1, 1970

E. A. Loria (Product Metallurgical Engineer, Crucible Steel Co. of America, Pittsburgh)—In this interesting paper, our introductory work was quoted. We would like to call attention to our sequel paper on the experimental determination of oxygen in cupola-melted cast iron,20 which was not mentioned. Vacuum-fusion oxygen values (as well as hydrogen and nitrogen) were reported for nine heats of cast iron melted in the Battelle 10-in. cupola under normal operating practice and under oxidizing conditions. The oxygen analyses ranged from 12 to 68 pprn compared to the author&apos;s computed range of 10 to 80 ppm. The average amount of oxygen found in our irons was about 20 pprn and changes in the silicon content of the iron from 1.32 to 2.35 pct had no consistent effect on the oxygen content of the iron. The gas determination specimens were poured in split steel molds that produced a clean pin, 3/8 in. diam and 2 in. long. Because freezing was almost instantaneous, the pins were entirely white iron (nongraphitic). In the early stages of the investigation, the pins were transferred to a mercury-filled trap system immediately after pouring. This was done to collect gas evolved between pouring and analysis. However, it was found that during storage for 4 weeks gas evolution was negligible. Because the vacuum-fusion analysis was usually completed within 4 days of pouring, pins from later heats were not stored in the mercury-trap system. We found some evidence that cast iron picks up oxygen during long storage, because of rusting. Earlier work by the British Cast Iron Research Association has shown that cast irons may be stored for a long time without significant change in their oxygen content. The practical significance of this study (and our own) would be in the improvement of cast-iron quality. Has the author investigated this aspect and reached any conclusions on the effect of oxygen on the mechanical properties of cast iron? The second phase of our study was to determine the properties of the test bars poured simultaneously with the gas analysis specimens. We realize that there may be complicating factors attendant in this procedure.21 Results from many test specimens measuring chill depth, transverse flexure and deflection strength, spiral fluidity, and sensitivity to hardness of gray irons ranging from 12 to 68 pprn oxygen showed that the lowering of transverse strength was the only significant undesirable effect of high oxygen content. A statistical study of the chill test results21 showed that the iron containing 22 to 46 pprn oxygen had forced chill depths that were 2/32 in. below the expected value from their composition, and irons containing less than 16 ppm oxygen had forced chill depths averaging 1/32 in. greater than the expected chill depth. Higher oxygen contents, within the range of 12 to 68 pprn did not increase forced chill depth. With the wedge tests, there was a good linear relationship between carbon equivalent of the irons and their chill depth. The results indicated that oxygen contents below 50 ppm in the iron did not affect chill depth. With 50 to 70 ppm oxygen in the iron, oxygen appeared to have a slight graphitizing tendency. These results are in disagreement with the common belief in gray iron foundries that "oxidized irons" produce high chill depths. It would be appreciated if the author would comment on this subject. Gustaf Ostberg (author&apos;s reply)—In Fig. 1 the legend of line I should read 2 pct C, 1 pct Si. The author wishes to thank Mr. Loria for calling attention to his later work, which was published after the present paper was concluded. The range of oxygen contents quoted seems to agree well with the author&apos;s values. The lack of response to variations in silicon content is probably due to the fact that the oxygen content in most cases was below the saturation level. The absence of temperature dependence, even in the case of saturation, is understandable if the difficulty in formation and escape of the deoxidation products is taken into account.

Jan 1, 1960

Alfred E. Hunt, Pittsburgh, Pa. (communication to the Secretary): My experience has been similar to that described by Mr. Sperry, regarding a 50 per cent. alloy of nickel and aluminum. Moreover, the same phenomena occur with a 50 per cent. alloy of iron and aluminum, manganese and aluminum, chromium and aluminum, or tin and aluminum. It seems that there is a narrow limit of composition of these metals, within which, when they are alloyed, they crumble down into a powder after standing for any considerable time. The resulting powder does not seem to be a separation or disintegration into the elements composing the alloy, but simply a powdering down of the alloy itself. The physical properties and utility of any of these alloys in pondered form have not been carefully investigated as yet, and offer a field worthy of further investigation. Samuel Peters, Pittsburgh, Pa. (communication to the Secretary) : In the early seventies, during the early days of the open-hearth process in this country, the writer was engaged as

Jan 1, 1900

William Kelly, Vulcan, Mich.:—In the northern part of Bedford county, Pa., the county immediately south of Huntingdon, the Clinton measures appear along the eastern slope of Tussey mountain near its base. The only iron-ore bed that has been worked is in the upper part of the formation. Along the mountain this seam has a dip of more than 60°. At the outcrop the ore is soft and rich, but with depth it passes often abruptly into an almost pure limestone similar to that described by Dr. Rutledge. The soft ore extends deeper in the cross-gullies than it does under the higher gound between the depressions. The slates adjoining the rich ore have been weathered, but where in juxtaposition to lean ore and limestone they are hard. The alteration of the ore-bed is evidently the result of the conditions of drainage in quite recent geo-

Jan 1, 1910

By J. S. Billheimer, B. H. Sage, W. N. Lacey

A restricted ternary system made up of n-pentane, tetralin, and a purified bitumen was investigated at 70, 160, and 220 °F. Most of the experimental observations were at atmospheric pressure or at 200 psi." However, some experimental measurements were carried out at a pressure of approximately 8000 psi. It was found that the purified bitumen was precipitated from its solution or dispersion in tetralin by the addition of n-pentane and that the separation occurred at lower weight fractions of n-pentane at the lower temperatures. The bitumen-tetralin solutions show some colloidal characteristics at temperatures below 160 °F when near compositions at which the bitumen separates as a solid phase. At states remote from the phase boundaries and at temperatures above 160 °F these characteristics become less evident. Under these latter circumstances the mixtures tend to follow the behavior of true solutions, particularly in regard to the approach to heterogeneous equilibrium. An increase in pressure appears to increase the solubility of bitumen in tet-ralin-n-pentane solutions. This effect is more pronounced at temperatures above 160 °F than at lower temperatures. INTRODUCTION Asphaltic phases of plastic or solid nature have appeared in numerous instances during the recovery of petroleum from underground reservoirs. Such depositions occurring underground appear to have caused adverse production histories for particular wells or zones. Because of this field experience, it is desirable to understand the factors which influence the formation or separation of the asphaltic phases from petroleum. The problem is unusually complex because the number of true components involved is very large and the details of the phase behavior encountered are difficult to ascertain experimentally. The literature relating to asphalts, asphaltines, and bitumen is voluminous and widespread.&apos; Only those references which are directly pertinent to the work at hand are cited. The separation of an asphaltic phase, hereinafter called bitumen? from naturally occurring hydrocarbon mixtures has been the subject of several investigations.2&apos;3&apos;4&apos;5&apos;6 It has been found that as many as four phases4 may be produced from a crude oil by the solution of a natural gas and propane at a pressure of 1500 psi and a temperature of 70 °F. The separation of bitumen from such naturally occurring mixtures results in at least one liquid phase which is substantially free of high molecular weight components.³ The influence of the solution of lighter hydrocarbons on the separation of bitumen from a Santa Fe Springs crude oil has been investigated. The results indicate that in the case of the methane-crude oil system, the quantity of plastic or solid phase separated reaches a maximum between 0.14 and 0.19 weight fraction methane and then decreases until negligible at higher weight fractions of methane. Similiar behavior was encountered in the case of mixtures of ethane and crude oil. The decrease in the quantity of the solid phase with an increase in the weight fraction of the lighter component appears to result from the formation of an additional liquid phase6 in which the bitumen is relatively soluble. The formation of this additional phase probably occurs at a weight fraction of methane close to that at which the quantity of separated solid reaches a maximum. A comparison of the deposition of bitumen in the field with the separation of asphalts from lubrication oil has been made&apos; and apparently the phenomena are similar. The phase behavior of bitumen also appears to be comparable to that of coal tar."&apos; The chemical and physical characteristics of asphalts and bitumen have been the subject of extended investigations which have been reviewed in some detail by Katz.¹º The conclusion was reached that the dispersion of bitumen in a number of organic liquids was not entirely colloidal since it was impossible to isolate individual dispersed particles even with the electron microscope. However, the evidence appeared to indicate that at states close to phase boundaries the extent of the dispersion of the phases influenced the equilibrium to a greater extent than is encountered in many simpler systems. From earlier study of field samples it became apparent that the phase behavior of bitumen-hydrocarbon systems was unusually complex. It was difficult to characterize in detail the phase behavior involved in naturally occurring hydrocarbon systems, even after a relatively extended investigation. For this reason, the study of a somewhat simpler system which still behaved in a similar manner became desirable. Three major constituents were necessary as-follows: a bituminous solid, a liquid constituent which was a reasonably good solvent, and a constituent in which bitumen was largely insoluble. A sam-

Jan 1, 1949

By T. B. Massalski, H. W. King

An hcp phase may be induced by cold working the ß&apos; phase of the Ag-Zn system. This phase reverts to ß&apos; on subsequent aging. No phase change occurs on cold working the o phase, but ß&apos; is formed when the deformed alloy is subsequently aged at room temperature. It is concluded that for alloys near 50 at pct Zn the ordered bcc ß&apos; phase is the equilibrium structure at room temperature. WhEN the disordered bcc ß phase of the Ag-Zn system is cooled to temperatures below 258o to 274oC, it transforms to a complex hexagonal phase <o.1,2 The nature of the o ß=o transformation has been the subject of some discussion,2&apos;3 and the structure of o has been described in detail.&apos; The latter phase appears to be stable on aging at room temperature but decomposes following cold work. When alloys containing approximately 50 at. pct Zn are rapidly quenched from the 0 phase field, the ß ? o transformation may be suppressed; but the ß phase undergoes an ordering reaction (ß ? ß&apos;). The ß&apos; structure may also be obtained as a result of cold working and aging at room temperature.4 Kitchingman, Hall, and Buckley4 have suggested that the decomposition of (o following cold work proceeds in two stages, (o ? ß followed by ß ? ß&apos;, but did not confirm this by experiment. When the ordered &apos; phases in the systems Cu-Zn5 and Ag-Cd6 are cold worked, they become unstable and transform to a close-packed hexagonal phase (( ) indicating that when order is destroyed in a ß&apos; structure the close-packed hexagonal phase may in many cases be more stable. It thus became of interest to study more closely the effect of cold work and annealing on the stability of both the ß&apos; and o phases in a Ag-50 at. pct Zn alloy. Predetermined weights of spectroscopically-pure Ag and Zn, supplied by Johnson and Matthey, were melted and cast under 1/2 atm of He in transparent vycor tubing. The ingot was homogenized for 1 week at 630°C and quenched into iced brine. Since mechanical polishing was found to induce a phase change, sections were first polished at room temperature, sealed in tubes under 1/2 atm of He, reannealed for several days at 630o or 200°C and then quenched into iced brine. Sections of the alloy thus prepared were found to be homogeneous when examined under the microscope. The sample quenched from 630°C (ß -phase region) was pink in color, whereas the sample quenched from 200°C (o-phase region) was silver. The latter sample showed the characteristic hexagonal anisotropy when examined under polarized light. Filings of the alloy were examined at room temperature, after various heat treatments, using an RCA-Siemens Crystalloflex IV diffractometer with filtered CuKa radiation. The X-ray reflections from flat powder specimens quenched from 630o and 200°C and sieved through 230 mesh were recorded graphically at a scanning speed of 1/2 deg per min. The resultant patterns are shown in Figs. 1(a) and 1(b) and may be identified as those of the 8&apos; and <02 structures respectively. The lattice parameter of the ß&apos; phase was determined as 3.1566Å.* This value compares very well withthatto be expected for a 50 at. pct Zn alloy from the data of Owen and Edmunds? and indicates that no loss of Zn occurred during casting. In order to study the effect of cold work upon the ß&apos; and o phases, filings made at room temperature and sieved through 230 mesh were mounted immediately in the diffractometer-i.e., without a strain-relief anneal. Changes in structure on subsequent aging were followed by scanning repeatedly over the regions of the low index reflections of the ß&apos; and o structures-i.e. , 28 from 35 to 44 deg. Immediately after filing the 8&apos; specimen, additional diffraction peaks were observed in the low-index region of the pattern, as shown in Fig. 1(c). These additional peaks do not coincide with those of the o structure, Fig. l(b), but may be indexed as the (10.0), (00.2), and (10.1) reflections of an hcp phase (<) with nearly ideal axial ratio. However, this hexagonal phase appears to be very unstable since within a very short time at room temperature it reverts back to the ordered ß&apos; phase, the reversion being complete within seven hours. The 5 ? ß&apos; reversion reaction is, therefore, very similar to those already reported in Cu-Zn5 and Ag-Cd6 7&apos;alloys. The action of filing caused the deformed surface of the originally pink ingot to become silver in color, indi-cating that the ( phase possesses similar reflecting properties to the o phase. Hence, the subsequent

Jan 1, 1962