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By H. Groeneveld, C. A. Connally, P. J. Hoenmans, J. J. Justen, W. L. Mason

A miscible displacement pilot using a slug of LPG driven by separator gas was conducted in the Cardiurn reservoir of the Pembina field. The injection pattern was a 10-acre, inverted, isolated five-spot. Upon completion of the LPG-gar phase, an experiment was conducted using a slug of water followed by gas. Calculated performance of the pilot is compared with actual performance. Equations are developed to calculate the distribution of LPG into zones of varying permeability, to estimate the progress of the flood at different times in the various zones and to estimate gas rates after breakthrough. The analysis indicates that permeability stratification was a dominant factor in controlling oil recovery and that oil was completely displaced from the swept pore volume. The results of the pilot indicated that miscible flooding is a practical means of pressure maintenance in this reservoir. The total recovery from the pilot area was good in spite of the early breakthrough of LPG. The effects of stratification were reduced by injecting a slug of water into the partially swept reservoir. INTRODUCTION The Pembina field,' located in Alberta, is the largest oil field in Canada and one of the largest in the North American continent. The reservoir is a stratigraphic trap producing from the Cardium sand. Neither bottom water nor free gas has been found. The recovery of oil by the natural depletion mechanism has been estimated at 12.5 per cent. Pressure maintenance studies of various areas have indicated that the recovery can be increased 21/2 times by water flooding, and a large area of the field is presently under water flood. However, reservoir studies of the North Pembina area indicated that miscible flooding might be competitive with water flooding. A pilot test was conducted to evaluate the performance of a miscible flood. A 10-acre, inverted, isolated, five-spot pattern was selected for the pilot. The pattern area was large enough to minimize wellbore fracturing effects and contained sufficient oil to provide significant working numbers. The performance of each of the four producers could be evaluated individually and compared. In the event of breakthrough in one direction, the effect would be isolated from the other producers. The use of a single injector minimized the volume of LPG required, and, because of the high mobility of gas, one well was sufficient to inject the necessary daily volume to replace the high rate of production. With four producers, the test could be completed in time for results to be evaluated, additional engineering studies to be made and a unit to be formed before the reservoir pressure in the North Pembina area declined below the bubble point. The pilot was located in an area developed on staggered, 80-acre spacing. The injection well was drilled at a regular location, while the four producers were drilled 467-ft north, east, south and west of the injector. Each quadrant and its associated producer were identified according to their direction from the injector— that is, north, east, south or west. The eight surrounding producers on 80-acre spacing were shut in to isolate the pilot area and provide for reservoir pressure observation. The pilot wells were completed using permanent-type completion techniques. After coring, casing was run through the pay section and cemented. Inside 51/2-in. casing, 2 1/2-in. tubing was hung. The wells were perforated opposite the Upper Cardium sand and lightly fractured. Fracturing volumes, rates and pressures were low to minimize the extent of the fractures. The fracturing treatments average 1,000 lb of 20-40 mesh sand in 700 gal of a low fluid-loss sand-carrying agent. Feed rates and wellhead fracturing pressures averaged 5.5 bbl/min at 2,535 psig, respectively. After fracturing, the productivity index was measured in each of the five pilot wells. The average PI of the four producers was 0.41 BOPD/psig drawdown. The measured PI'S were approximately the same as PI'S calculated from core analysis data, indicating that the fracturing treatments were just sufficient to overcome

By G. T. Hahn, C. N. Reid, A. Gilbert

An evaluation is made of the possible cautsal relationship between twinning and fracture in molybdenum. For both single and poly crystalline material no instance of twin-induced fracture was observed. Instead brittle fracture was found to be slip-induced at heterogeneities. For single-crystal material, the yield stvess in compression and the fracture stvess in tensiow obey a similar angular- relationship which follows approximately a 1/cos2 ? law. where 0 is the angle between the specimen axis and the nearest (100) plane. This sittlilarity between yield and fracture behavior casts doubt on the interpretations made previoltsly that a 1/cos2? relationship supports a critical normal fracture stress criterion. SINCE fracture in brittle materials takes place at stresses an order of magnitude lower than the theoretical strength of the lattice, it must be postulated that some stress-concentrating effect is operative during the fracture-initiation process. Several mechanisms have been proposed whereby the necessary stress concentrations could be produced. The well-known Cottrell mechanism1 describes a dislocation interaction which can lead to the formation of an incipient cleavage crack on a cleavage plane, and both zener2 and stroh3 have discussed models which predict the stress concentration at the intersection of a slip band with a grain boundary. When twinning occurs twin/twin intersections and twin/grain boundary intersections represent another possible means of fracture initiation. Although still controversial, the concept of twin-induced fracture is supported by a weighty mass of evidence, which has been reviewed in a recent paper.4 This investigation was conducted in order to assess the extent to which the brittleness of molybdenum can be ascribed to this cause. Emphasis was placed on seeking direct metallographic evidence for twin-induced fracture such as arrangements of twins located at the fracture origin. EXPERIMENTAL PROGRAM Materials. The experimental materials are described inTable I. From the chemical analyses and estimates of the amounts of interstitials likely to be retained in solution,5 it is concluded that all were mul-tiphased systems. X-ray diffraction experiments6 showed no evidence of preferred orientation in the polycrystalline materials. Single crystals were grown from Molybdenum X and Y using the floating-zone technique,' in a vacuum of better than 3 x 10-5 Torr. However difficulties were experienced with Molybdenum X due to violent gas evolution from the molten zone, and additional crystals (Crystals 3 to 6) were produced by annealing for 5 1/2 hr at 2300°C under a vacuum of 10-4 Torr. Techniques. A series of tension, compression, and bend tests was conducted. Tension tests were used to demonstrate brittle behavior, to measure brittle-fracture stresses, and to provide fractured specimens for metallographic inspection. Compression tests were employed in order to obtain some ductility and a measure of the yield stress. Furthermore, it was considered that compression loading would permit the study of crack nucleation in the absence of propagation. The bend tests were conducted to facilitate identification of the fracture origin, which would be expected to be at, or near, the tensile surface of the sample. Tests were carried out at temperatures between 78" and 298°K attained by means of a liquid-nitrogen evaporator of modified Wessel design.' Specimens were fashioned by mechanical grinding, and, prior to heat treatment, were electropolished at 10 v in a 3:l mixture of ethyl alcohol and sulfuric acid, using a stainless-steel cathode. Except where noted, a strain rate of 4 pct per min was used. All specimens had a gage section 0.75 in. long; diameters of 0.10 and 0.20 in. were adopted for the tension and compression specimens, respectively, the former of single shoulder design, the ends being gripped in split collets with a conical bearing surface. In the bend test, electropolished specimens measuring 0.080 by 0.25 by 1.25 in. were deformed at 78°K by four-point loading. The bending device had spans of 1 in. between the outer and 4 in between the inner fulcra; it was stressed between the compression anvils of the Instron machine at a constant deflection rate. EXPERIMENTAL RESULTS Polycrystals. Tension and Compression Tcsls. It is apparent from Fig. 1(a) that above 170°K fracture of Molybdenum X takes place at stresses equal to or greater than the compressive yield stress whereas

Jan 1, 1967

By J. H. Palm

Experiments have shown12 that the formula -S = St - (St - Se')e ne [1] expresses very well the relation between the true stress S and the true form monotonic deformation of plastic metals in single tension and compression (Fig 1). St (asymptotic or final stress), Se' (threshold stress) and ne (characteristic or specific strain) are constants, the significance of which immediately follows from Fig 1. Only at the initial stages of deformation the experimental curve runs below the theoretical curve. This may be due, at least partly, to the uneven stress distribution caused by the anisotropy of the crystals. It is now obvious that a similar formula H = Ht- (Ht- He)e ne' [2] might be valid for the relation between the hardness and the strain. The hardness can be only a function of the strain and must therefore be independent on the hydrostatic tension present in the neck during straining. Hence, if Eq 2 is correct it will also represent the relation between hardness and strain in the middle section of the neck, the strain of which is generally assumed to be almost or fully homogeneous. For aluminum, copper, and several copper alloys it is indeed observed that Eq 2 agrees very well with the experimental Vickers hardness-strain relation from zero strain on up to the strain at fracture. Moreover nc' is, for several metals fairly equal to nc- Hence H = C1 + C2S [3] C1 and C2 are constants during uniform uniaxial tension. Eq 3 was also obtained for copper by Voce.3 According to Hencky's theory4 the ratio of the Brinell or Vickers hardness and the yieldstress (expressed in kg per mm2) of a non-strainhardening metal is approximately equal to 2.8. If Eq 1, 2 and 3 are correct the ratio Ht/S, must accordingly be the same for all plastic metals and be equal to 2.8. It is now experimentally established that this ratio varies between 2.7 and 3.1 with a mean value of 2.9. Taking into account that Ht and especially St can be obtained only by extrapolation from the range of rather low strains, the agreement with the theory is very satisfactory. References 1. J. H. Palm: Appl. Sci. Res. Al (1948) 198. 2. E. Voce: Jnl. Inst. of Met. (1948) 74, II, 537. 3. E. Voce: Metal Treatment (1948) 15, 53. 4. H. Hencky: Zlsch. f. angew. Math. and Mech. (1923) 3, 241.

Jan 1, 1950

By R. E. Reed-Hill, W. H. Hartt

High-purity magnesium single crystals, oriented with basal plane parallel to stress axis, were deformed in tension at room temperature so as to form second-order (1011)- (1012) twins. Investigation by optical and replica electron microscopy revealed that these twins can support very large internal strains approaching 1000 pct and that fracture follows as a direct consequence of these strains. At low strains the basal plane within (10i1)- (10i2) twins rotates with increasing deformation until it becomes approximately parallel to the twin boundaries. This rotation is not consistent with deformation by basal slip in the double twin, but may be rationalized in terms of nonbasal slip. Observations also indicate that basal slip in the primary (1011) twin, prior to second-order twinning, may contribute to this rotation. Deformation markings that are difficult to interpret in terms of slip on previously observed systems were noted in the twin bands. A possible means of rationalizing these markings is to relate them to grain boundary shear. The fracture associated with these twins is thought to initiate by formation of voids at the twin boundaries. These grow into microcracks, with ultimate fracture occurring by tearing of the interconnecting regions. THE significance of deformation and fracture along second-order (10i1)-(10i2) twins in polycrystalline magnesium1-' and many magnesium alloys4 was discussed in a previous paper.5 It was noted that this double twinning mode effectively controls the room-temperature deformation and fracture of magnesium single crystals, oriented with the basal plane parallel to the tensile axis. Reed-Hill and Robertson6 observed that, although these crystals may fracture at a macroscopic strain of less than 1 pct, shear strains up to 1000 pct could occur within (1011)- (10i2) twin bands. They concluded that the fracture was "ductile" and resulted from concentrating the deformation into a very small volume. The present research was undertaken to obtain a better understanding of the deformation and fracture associated with these twins. EXPERIMENTAL TECHNIQUES The experimental procedure involved straining high-purity magnesium single crystals in tension at room temperature. Specimens were oriented so that the applied stress direction was within 2 deg of the [1010] crystal axis. The specimen cross section was rectangular with faces closely parallel to (1210) and (0002). The specimen preparation and testing procedures have been described in detail elsewhere.5 Second-order (1011)- (1072) twins nucleated during straining were studied using optical and replica electron microscopy techniques. For the latter, cellulose acetate-carbon double replicas were employed. Observations were made on the (1210) crystal surface, which is the plane of shear for twins of this type. In addition to the single crystals, some investigations were carried out on longitudinal polycrystalline specimens. These were obtained from high-purity magnesium plate with a texture in which the basal planes of the grains tended to be nearly parallel to the specimen axis. All. polycrystalline specimens were annealed prior to testing in order to produce a coarse grain structure, permitting Laue back-reflection X-ray photographs to be obtained from grains of interest. EXPERIMENTAL RESULTS a) Deformation Within ( 101 l )-( 1072) Twins. Fig. 1 is an optical micrograph demonstrating the extent of the plastic deformation that can occur in (1011)-(10i2) twin bands. Notice at the lower left corner of the photograph the large displacement of the dark band (a polishing step) where it crosses the twins. The shear strain in the twin band at this position was computed to be 700 pct, using the ratio of step width to twin thickness. It was previously shown5 that (1011)- (1012) twin bands are often composed of small, separate twins aligned close to the macroscopic habit plane. Fig. 2 is an electron micrograph from near the upper end of the twin band in Fig. 1. This illustrates that the twin habit often does not coincide exactly with the macroscopic band habit, but may be inclined at a slightly smaller angle to the matrix basal plane. Figs. 3 through 6 are electron micrographs of the twin band in Fig. 2 illustrating successive steps in the development of deformation inside the band. A twin band such

Jan 1, 1969

By J. H. Swisher

HIGH-temperature anneals in wet hydrogen are widely used in research as a technique for decarburiz-ing Fe-C alloys. This technique is also used commercially in processing low-carbon steels for magnetic applications. The results of some calculations to be presented show that for temperatures up to 1100°C there is a critical temperature for maximizing the decarburization rate of alloys containing about 0.1 pct C. The low-temperature portion of the iron side of the Fe-C phase diagram1 is reproduced in Fig. 1. Preliminary calculations show that the critical temperature is between 727" and 905°C. In this range, the alloy contains two phases during heat treatment, ferrite and austenite. Decarburization therefore occurs by a process termed "diffusion accompanied by a phase change".' For intermediate stages of decarburization, the carbon concentration profile in a flat plate is shown in Fig. 2. The carbon content increases linearly from zero at the surface to a value, cB , at the two-phase boundary. cB is known from the phase diagram, see Fig. 1. At the interface where x = 5, the carbon content increases to ci, the initial concentration in the alloy. smith3 has derived an equation which describes the rate of movement of the boundary. With a rearrange-

Jan 1, 1969

By A. R. Troiano, R. F. Hehemann

The influence of partial decomposition to high temperature bainite on reaction kinetics at a lower temperature has been studied in two alloy steels. Reaction at the lower temperature is retarded by the prior treatment, and the extent of decomposition may be reduced. Interpretation of these results is based on a mechanism involving a limitation in the nucleation and growth of bainite plates. OF the major transformations in steel, the characteristics and general behavior of the bainite reaction are probably the least understood and appreciated. Limitations of space preclude a critical evaluation of the present status of the bainite transformation in this presentation; however, such a treatment will shortly appear elsewhere. Only the salient features pertinent to the present investigation will be introduced briefly here. Although the reaction curve for the formation of bainite is similar to that for a nucleation and growth process, other kinetic features are more in keeping with the martensitic mode of transformation. A definite temperature exists above which austenite will not transform to bainite.1-5 his temperature, which has been designated B., is determined by the composition of the austenite. Unlike other nucleation and growth processes, the amount of austenite transformed to .bainite is a function of reaction temperature. The extent of decomposition increases from 0 at H. to 100 pct at some lower temperature.' , This lower temperature will be designated B1 and appears to be relatively insensitive to austenite composition.% 5 The similarity in the effect of reaction temperature on the bainite and martensite transformations serves to emphasize the close connection between these two decomposition processes. Austenite decomposition in the bainite range proceeds without partition of the alloying elements.8-11 Partition of carbon has been proposed" primarily on the basis that partial transformation to bainite lowers M, and increases the amount of austenite retained at room temperature. Carbon enrichment resuslting from such partition has been employed to explain the influence of reaction temperature on the extent of decomposition.'" It should be noted, however, that no enrichment has been detected experimentally in high carbon steels.1,14,15 Lattice-parameter measurements of retained austenite in steels containing 0.3 to 0.4 pct C have indicated carbon enrichment, 3,10-18 although the split indicative of a high carbon martensite has not been reported. Carbon enrichment, if it does occur, must be highly localized around the bainite plates. Therefore, carbon enrichment does not account for the influence of temperature on the progress of the bainite reaction."' Thermal history is known to influence the martensite transformation through stabilization.20,21 No similar phenomenon in the bainite transformation has been reported. Materials and Procedure Two triple-alloy steels were chosen for this investigation. Their compositions were as given in Table I. These steels were chosen because the pearlite reaction did not interfere with the bainite reaction. Steel K was received in the cast condition and forged from 2 in. square bars to 1/2 x1 3/4 in. plates. The 4340 was received as 11/4 in. hot-rolled rounds. Both steels were homogenized in vacuum for one week at 2300°F in order to minimize segregation. A quenching dilatometer similar to that described by Flinn, Cook, and Fellows" was employed for the kinetic measurements. Dimensional changes were detected by a differential transformer coupled with a high speed recorder. The dilatometer was mounted so that it could be transferred to any one of three furnaces: a nitrogen-atmosphere austenitiz-ing furnace and two salt-bath furnaces for isothermal transformation. Dilatometer specimens were 1/32 x 1/4x 1/2 in. with a gage length of 1.4 in. All specimens were nickel plated in order to minimize decarburization during austenitizing. The austenitiz-ing conditions consisted of 10 min at the temperatures given above. Austenitizing temperatures were controlled to 210°F and transformation temperatures to ±3ºF. The precision of the dimensional measurements was estimated to be ± 5 x105 in. per in. Results and Discussion Isothermal Transformation: The characteristics of the isothermal bainite reaction will be described

Jan 1, 1955

"Only the coal of the Sunnyside district in Carbon County is coked. Two companies, the Utah Fuel Company and the Columbia Steel Corporation, operating adjoining mines at Sunnyside and Columbia, produce coke.The Utah Fuel Company has 819 beehive coke ovens at Sunnyside. Only 350 ovens are being operated at the present time, largely because the smelters which had been using Sunnyside coke changed their equipment from blast furnaces to reverberatory furnaces.The Columbia Steel Corporation operates a plant of 33 Koppers type by-product coke ovens at Ironton, near Provo. About 800 tons of coal are used per day to produce coke for the company's iron blast furnace plant pie ovens produce about 400 tons of coke, 12,000,000 cubic feet of gas, 9,000 gallons of coal tar, 3,000 gallons of benzol and several tons of ammonia sulphate per day.The total coke production for the State from the earliest output of record in 1880 until June 30, 1924, is 5,947,714 tons."

Jan 1, 1925

By R. F. Hehemann, D. J. Cometto, G. L. Houze

The w transformation in the Zr-Nb system was studied using X-ray diffraction, dilatometric, re-sistornetric, hardness, and metallographic techniques. w forms in a diffusionless, completely reaersible manner on quenching and in a diffusion-controlled manner- on aging. The temperature at which w begins to form on quenching was determined as a Junction of composition and was found to decrease with increasing solute content. w formed bv aging establishes a metastable equilibrium with an enriched ß phase. The ß/w + ß transus has been determined for this metastable equilibrium and employed to rationalize retrogression and reversion phenomena observed in these alloys. The decomposition mechanism is discussed in terms of a gradual or continuous transformation from ß to the w state. BETA- stabilizing alloying elements lower the MS temperature of the martensitic bcc (ß) to hcp (a') transformation in zirconium and titanium alloys. In certain titanium alloys, a lower symmetry modification of the martensitic structure, termed (a"), also has been reported.1,2 These martensitic transformations are suppressed when the amount of the ß-stabilizing element exceeds a critical level. However, the high-temperature ß phase cannot be quenched to room temperature without change. At alloy contents just above the critical level, |ß trans-forms to the w phase when cooled rapidly.3-6 The amount of w in quenched alloys is reduced by increasing alloy content, and this phase virtually disappears above a critical composition.5 In addition to the transformation during cooling, w also can be formed by aging ß at temperatures below approximately 500°C, and significant alloy enrichment of retained 13 accompanies the isothermal transformation.1-8 The structure of LC is closely related to that of the ß from which it forms.9-15 The bcc (ß) lattice can be generated using a hexagonal cell with three atoms located at (000) and ±(2/3, 1/3, 1/3). This cell has an axial ratio of 0.612 and is oriented with respect to the cubic cell such that (0001)H (111)C and [1120]H [101]C. Consequently, there are four possible orientations of the hexagonal cell, associated with the four (111) planes of the bee lattice. Formally w can be obtained from 0 by allowing the two atoms at the ±(2/3, 1/3, 1/3) positions to approach the coordinates ±(2/3, 1/3, 1/2) and there are four equally probable orientations of w for each 0 grain. In titanium alloys w retains the cubic axial ratio of 0.612" and hence also can be indexed as triply cubic, but this is not the case for aged Zr-Nb alloys where w is clearly hexagonal with an axial ratio of 0.622." The lattice parameters and atomic positions of w depend on thermal history and alloy content. The atomic positions (000), ±(2/3, 1/3, 1/2) provide reasonable agreement between calculated and observed diffracted intensities for w in the fully aged condition.10,14 In the quenched condition, on the other hand, the atoms appear to be displaced from the central plane12,13,15 and assume positions ±(2/3, 1/3, Z = 0.42-0.48) rather than ±(2/3, 1/3, 1/2). This results in a structure with trigonal rather than hexagonal symmetry. The readily detectable parameter and axial-ratio changes of w in Zr-Nb alloys make this system especially attractive for studying the structural changes that occur in the formation and aging of w. Particular attention in this investigation has been directed to the relationship between the structure of w in the quenched and aged conditions, and to certain aspects of the reaction kinetics. MATERIALS AND PROCEDURE Zr-5, 12, 17.5, and 25 pct Nb alloys were prepared by a double arc-melting practice, encased in stainless-steel cans and hot-rolled to 1/8-in.-thick sheet.* Charged weights have been employed for the niobium contents and interstitial analyses are reported in Table I. Dilatometric, X-ray diffraction (filtered CuK, and MoK, radiation), metallographic, and hardness techniques have been employed to follow the transformations during isothermal and quench-aging heat treatments. In quenched specimens, the electrical resistivity also was studied. Betatizing was conducted for 1 hr at 900°C. With the exception of the isothermal dilatometric studies, samples were

Jan 1, 1965

By H. C. Rogers

A phenomenological study of the failure of polycry stalline ductile metals at room temperature was carried out using light and electron microscopy. Tensile fractures as well as sections of partially fractured bars of OFHC copper in particular were examined. The initiation and growth of the central crack in the neck of a tensile specimen occurs by void formation. After the formation of the central crack the f'racture may be completed in either of two ways: by further void formation or by an "allernating slip" mechanism. The first leads to a "cup-cone" failure; the second, to a "double-cup" failure. In the past decade or decade and a half there has been a great deal of emphasis on the solution of the problem of the brittle fracture of metals, particularly those which normally exhibit considerable ductility such as steel. Since the problem of the fracture of metals after large plastic strains has less immediate commercial or defense significance, there has been considerably less effort expended in describing the details of the phenomenology and determining the mechanism of this type of fracture. The present research was undertaken to increase our knowledge in this area. The problem of ductile fracture has not been neglected completely, however. Ludwik1 first found by sectioning a necked but unbroken tensile specimen of aluminum that fracture began with a large internal crack which appeared to have started in the center of the neck. Examination of the fracture indicated that the crack had propagated radially with increasing deformation until a point was reached at which the path of the fracture suddenly left this transverse plane and proceeded at approximately 45 deg to the stress axis until the surface was reached. This gives rise to the commonly observed cup-cone tensile fracture. When MacGregor2 was attempting to demonstrate the linearity of the true stress-true strain curve from necking until fracture, he found that copper was anomalous in that the stress dropped off markedly from the straight line value before fracture occurred. Radiography indicated that in the copper an internal crack was formed long before the final fracture, the stress decreasing during the growth of this crack. One of the most significant advances in the understanding of ductile fracture was the result of work by Parker, Flanigan, and Davis.3 By the use of etch-pit orientations they were able to demonstrate conclusively that the fracture surface at the bottom of the cup, although on a gross scale normal to the tensile axis, did not consist of cleavage facets as had been previously supposed by many investigators. Recently, Forscher4 has shown evidence of porosity near the tensile fracture of hydrogenated zirconium which he attributes to hydride decomposition. The workers at the Titanium Metallurgical Laboratory5 have also shown evidence of porosity in a number of the commonly used metals after heavy deformation. Many metals have relatively low ductility during creep tests at high temperature. The fractures are intercrystalline, resulting from the nucleation and growth of grain boundary voids. The work in this area has been recently reviewed by Davies and Dennison.6 It is possible that some of the observations and conclusions may have a bearing on the present study? especially since at least two studies7,' have been extended down to room temperature and below using magnesium alloys. However, since magnesium does exhibit low-temperature cleavage, these results may not be pertinent to the present one. The use of the electron microscope as an aid to the study of fractures has been extensively exploited by Crussard and coworkers.9 The examination of direct carbon replicas of the fractures of a large number of metals and alloys showed that the bulk of the fracture surface was covered with cup-like indentations of the order of 1 to 2 µ in size. These frequently had a directionality by which Crussard claims to be able to tell the direction of the crack propagation. With this rather disconnected background of information, this investigation was undertaken in the hope of presenting a unified picture of the initiation and propagation of a fracture in a ductile metal. To this end all of the techniques previously used were employed simultaneously so that there might be a good correlation of the data obtained by different techniques. EXPERIMENTAL PROCEDURE The metal which was chosen as the starting material for this investigation was OFHC copper. Of the dozen or so materials considered, it best fulfilled the requirements of commercial availability in large sizes, good ductility, relatively high melting point compared with room temperature and

Jan 1, 1961

By J. R. Jargon, H. K. van Poollen

Non-Newtonian fluids may be injected into a reservoir during secondary recovery operations. The non-Newtonian fluid used in this work is a power-law type of fluid that is, the viscosity of the fluid decreases as the flow rate or rate of shear increases. This paper presents equations for steady-state linear and radial flow of such fluids, transient behavior results from a finite difference model of a radial system, and transient behavior results from a field test. The equations that describe the flow of a non-Newtonian fluid are non-linear and are solved numerically. Finite difference solutions are presented as curves of dimensionless pressure drop at the wellbore us dimensionless time for a constant injection rate. Solutions were obtained for 5-percent, 10-percent and 100-percent PV of a non-Newtonian fluid for injection rates of 1, 10, 100 and 1000 cc/sec and for a 5 percent PV of non-Newtonian fluid located at T = Tw' 3, 10, 20, 50 and 100 ft for a flow rate of 1 cc/sec. The buildup curves do not exhibit a stmight-line portion as is the case for Newtonian flow through porous media. Correlations also are shown for the productivity index vs rate for the computer model study and the field tests. INTRODUCTION During various secondary recovery operations non-Newtonian fluids are injected. Such fluids, in general, include thickened water and gelled fluids. The term "non-Newtonian" implies that the viscosity is not only dependent upon temperature and pressure, but also on the rate of shear that is applied to the fluid. For example, water, which is a Newtonian fluid, will have essentially the same viscosity no matter what rate of shear is applied. A pseudo-plastic fluid exhibits a decreasing viscosity when higher rates of shear are applied; a dilated fluid has an increasing viscosity with increasing rates of shear. The objective of studies performed and described in this report is to obtain relationships and mathematical and empirical descriptions of the flow of non-Newtonian fluids through porous media. Simple mathematical relationships, computer studies that include the unsteady-state behavior of such fluids, and field tests are used. PREVIOUS LABORATORY STUDIES Laboratory studies have been performed by several investigators. The fluids used in one investigationl were surfactant-stabilized dispersions of water in hydrocarbons. Its Fig. 6 is reproduced as our Fig. 1. This effective-viscosity vs frontal-velocity diagram shows that this fluid is of the power-law type over a fairly large range of frontal velocities. Gogartyl developed an equation for the effective viscosity as a function of shear rate that reduces to Eq. 2 for frontal velocities greater than about 2 ft/D. The fluid characteristics used in the present study are similar to those reported by Gogarty. Christopher and Middleman2 performed experiments with power-law fluids flowing through pipes packed with glass beads. They used a modified Blake-Koteny

Jan 1, 1970

By F. H. Wilson, E. W. Palmer

G. L. Bailey (British Non-Ferrous Metals Research Association, London, England)—I was glad to see this further work on the structure and properties of the Cu-Ni-Fe alloys of low nickel content in which we have been so interested in the United Kingdom since before the war." The authors' data on the variation in solubility of iron in both 5 and 10 pct Ni alloys with temperature is based on more detailed observations than we have made but is generally in line with our own work. The difficulty in these matters is in the identification of the structures obtained. The precipitation of the second phase which we presume to be the a' solid solution, richer in nickel and iron, from the a solution with which it is in equilibrium occurs in such a fine state of division that its resolution by ordinary micro-graphic methods is impossible in the early stages. We have published some electron micrographs which illustrate the effects observed more clearly but we agree with the authors when they say that dark staining characteristics when the alloys are subjected to sea-water corrosion are more sensitive in detecting the presence of the a' phase than is the microscope. After taking this point into consideration, however, I confess to some difficulty in understanding Fig. 4. I can only assume that the precipitates visible at such a low magnification as 150 diam in these relatively coarse forms are produced from iron-rich areas of the initial solid solution due to inadequate homogenization. If this is so I find it difficult to accept the authors' contention that such local variations in iron concentration as are capable of giving the heavy precipitates shown in Fig. 4 do not in fact materially affect the solvus surfaces of their equilibrium diagrams. I share the authors' enthusiasm about the resistance to corrosion of these alloys by seawater but I am not fully in agreement with them on the vexing question of the best iron content for these alloys. In Fig. 3 of ref. 8, we give many results showing that the resistance of these alloys to impingement attack in the Jet Test increases as the iron content rises to 2 pct. With higher iron contents pitting at shielded areas is likely to occur. Admittedly the addition of 0.75 pct Fe produces a large improvement in corrosion resistance compared with a pure 90/10 Cu-Ni alloy, but the further improvement on raising the iron to 2 pct is substantial. We recommend from 1.5 to 2 pct Fe as the desirable range from the point of view of corrosion resistance. While it is true that British practice generally uses a hard tube, the corrosion resistance of the 10 pct Ni alloys has in our work been little affected by the structural condition. Although the annealed alloys blacken in seawater, this effect seems to be superficial and not to be accompanied by any real decrease in

Jan 1, 1953

By T. W. Smoot

Bentonites are composed chiefly of montmorillonites which are clay minerals generally distinguished from other clay minerals by their surface activity and extremely fine particle sizes. The uses of bentonite are many and varied, depending principally on one or more of the following characteristics: 1) Thixotropy, 2) Dispersion, 3) Plasticity, 4) Swelling, 5) Surface area, activity, ab- and adsorbtive properties. Some of the major uses of bentonite are briefly reviewed as are the mineralogical and physical properties. The uses of X-ray diffraction, differential thermal analysis and other tests for physical properties are reviewed as to their uses in evaluation. In the past 40 years, the industrial uses of bentonite have grown to include many applications. The physical and chemical properties of bentonites are such that they can be used as catalysts, adsorbants, plasticizers, binders, drilling muds, and pelletizers. Not all bentonites, however, can be used for all of these purposes and some are probably of little or no value for any use. The purpose of this paper is to briefly review the mineralogical, physical, and chemical characteristics and variations of bentonites. The determination of these characteristics and their uses in industry are also reviewed so that these determinations can be used as guides for evaluation. DEFINITION OF BENTONITE The word bentonite, as used in this paper, refers to a clay material, usually (but not necessarily) the result of the in-place-weathering of volcanic ash. Its mineralogical composition is dominated by the clay mineral montmorillonite and generally carries other clay minerals and silica minerals as the major 'impurities'. For this discussion, such a definition is justified since most of the economic uses of bentonite~ depend upon the physical and chemical characteristics of the montmorillonite minerals involved. The terms montmorillonite and bentonite have been used almost interchangeably in industry although the latter is much more commonly used. Therefore, where montmorillonite is discussed, it should be understood that the major feature of a bentonite is the subject. STRUCTIJRE OF MONTMORILLONlTE The economic importance of bentonites is generally based on one or more of the following characteristics of montmorillonite: 1) high-thixotropy and easy dispersion in water, 2) high-plasticity, 3) swelling properties, 4) large surface area/unit mass, 5) high surface activity, and 6) adsorptive properties. To understand these properties, an appreciation of the structural scheme of montmorillonite is helpful. Montmorillonite IS one of the three-layer clay minerals composed essentially of hydrous aluminum silicates. The essential building blocks of these layered silicate:; are: 1) silica tetrahedron composed of a silicon atom surrounded by four oxygen atoms, and 2) aluminun (iron, magnesium, etc.) octahedral unit that, ideally, has one of these atoms surrounded by six hydroxal units. Fig. 1 is a diagrammatic sketch of the montmorillonite structure. In the montmorillonite structure, the tetrahedron are pointed inward and attached to the octahedral units so that the O= atoms replace the hydroxal units where they are bonded; hence two-thirds of the hydroxal units

Jan 1, 1962

By J. R. Bell, W. J. Childs, J. H. Bucher, G. A. Wilber

A technique for determining recrystallization times as short as 0.10 sec was developed utilizing the "Gleeble", a commercially available testing system designed for the study of short-time, high-temperaLure themal and mechanical processes. The procedure consisted of heating a small tensile specimen to a given temperature of hot deformation, loading to a given reduction in area, unloading, delaying various intervals at temperature, and then reloading- to failure. The magnitude of the ultimate load obtained upon reloading decreased with delay lime as recrys-lallization proceeded. The technique was applied to austenite recrystallization in AISI 1010 and AISI 1010 uith 0.02 pct Cb steels. For each steel the reduction in area given the specimen on the first pull was mainlairred at 30 ± 5 pct and recrystallization times deterntined at various temperatures. The results indicaled a significantly slower rate of recrystallization for the columbium-modified composition, suggested the presence of- a recovery stage in the softening process , and indicated a greatly increased softening rate at a temperatuve where significant allotropic transformation to a partially ferritic Structure could occur. In recent years increasing attention has been paid to the fact that the process of recrystallization of austenite deformed at elevated temperatures is far from instantaneous at many practical hot-working temperatures.1-3 This realization has given rise to such terms as hot cold-working1 or warm-working,2 These terms generally describe processes where the recrystallization rate at the temperature of deformation is slow enough to have an appreciable effect on mechanical properties despite a relatively high deformation ternperature. The mechanical properties of interest can be either the properties at the deformation temperature as in hot-workability studies4 or the room-temperature properties after cooling as in the many recent studies of various thermomechanical processes172 where heat treatment and deformation are intentionally combined to give a unique set of room-temperature properties. Because of this interest in processes where the austenite recrystallization kinetics can be an important variable, the development of quantitative methods of following the course of short-time, high-temperature recrystallization has received increasing attention.l,3,5 The experimental methods to date have, in general, relied upon rapidly deforming the austenite, holding at temperature for various brief intervals, quenching as G.A.WILBER and W. J. CHILDS, Members AIME,are Research-Fellow and Professor, respectively, Rensselaer Polytechnic Institute, Troy, N. Y. J. R. BELL and J. H. BUCHER, Member AIME, are Research Engineer and Research Supervisor, respectively, Graham Research Laboratory, Jones & Laughlin Steel Co., Pittsburgh, Pa. Manuscript submitted March 13, 1968. IMD. rapidly as possible, and then using room-temperature measurements to follow the recrystallization process. Although such methods can be successfully applied to certain alloy steels, the existence of the allotropic transformation during cooling of plain-carbon or low-alloy steels tends to obscure the results. Thus, such room-temperature measurements as hardness and X-ray line widths do not correlate well with the extent of austenite recrystallization before quenching,5 and results based on room-temperature microstruc-tural observations are dependent upon the success in correlating the observed structure with the prior aus-tenitic grain structure.1,3,5 The purpose of the present work was to develop a quantitative method for the determination of short-time, high-temperature recrystallization rates, based on measurements made at the temperature of deformation. EXPERIMENTAL TECHNIQUE The basic technique consisted of heating a small tensile specimen to a given temperature of hot deformation, loading to a given reduction in area, unloading, delaying various intervals at temperature, and then reloading to failure. The data were obtained in the form of traces of load and elongation as a function of time. Due to the high deformation temperature, the strain hardening introduced during initial loading was progressively annealed out with holding time after unloading and the loads obtained upon reloading decreased as this softening proceeded. Although the value of the second load at any Consistent point On the load-elongation curve could have been used as a measure of the degree of softening, the most convenient to use was the ultimate load. The softening indicated by the decrease in the second ultimate load with time is essentially a process of annealing of cold-worked material at a high deformation temperature. Although some recovery grain growth may contribute to such a softening process, it is generally considered that the major softening which must take place to achieve complete removal of substantial Strain hardening will occur by the formation of new, stress-free grains. As the results of this work indicate that essentially complete removal of strain hardening did in fact occur. the primary softening process will be attributed to recrystallization, and specific reference made where it appears that other mechanisms may be contributing to the total observed softening. It would, of course, be of interest to attempt to correlate the results of this work with the actual austenite fraction recrystallized as determined by other techniques. This was not attempted in the present work because it would have required running a large number of additional specimens and, as discussed previously, there is limited assurance that the results would accurately reflect the prior austenite fraction recrys-

Jan 1, 1969

By L. E. Tanner

A study was made of the transformation kinetics of the commercial titanium-base alloys, Ti-2.5Al-16V and Ti-4A1-3Mo-1V, using two different heat treatment cycles: 1) step-quenching to aging temperatures from a ß solution anneal and 2) water-quenching from an a+ß solution anneal followed by reheating to aging temperatures. Metallographic examination and hardness testing provided the major portion of the data, while electrical resistivity, dynamic elastic modulus, and X-ray diffraction techniques were also used to obtain critical or confirnzatory data. TTT diagrams were constructed from these results. The sheet alloys, Ti-2.5Al-16V and Ti-4Al-3Mo-lV, are of current interest to the Department of Defense. As an aid to the planning of their respective heat treatments, a program was carried out to determine the transformation kinetics of these alloys using ß and a + ß reference states. This paper presents the TTT diagrams constructed from data obtained by metallographic examination and hardness testing, as well as by X-ray diffraction, electrical resistivity, and dynamic elastic modulus determinations. EXPERIMENTAL PROCEDURE Materials—Sheets of the alloys were obtained from the producers in the "mill anneal'' condition and were tested prior to shipment to assure their meeting mill specifications. Chemical analyses were also supplied and are presented in Table I. Heat Treatment—The majority of the data presented in this study were obtained from hardness testing and metallographic examination. Sample coupons were cut from the alloy sheets and heat treated on a mass scale. The annealing cycles were operated in the following manner: 1) The ß solution anneals were carried out in a horizontal tube furnace with the bare samples protected by a dynamic helium atmosphere. Following this treatment, the samples were quickly transferred to lead, solder, or oil-bath furnaces for aging. After prescribed times at the various reaction temperatures they were water-quenched. 2) The second cycle simulated commercial practice, which requires the alloys be quenched to room temperature following the a + ß sotution anneal. This anneal was carried out in a muffle furnace with a dynamic helium atmosphere. The specimens were then reheated from room temperature to the reaction temperatures in lead, solder, or oil-bath furnaces, followed by water-quenching after prescribed times at these temperatures. Information concerning the critical temperatures of the alloys was obtained from the literature and solution treatments were based on these values. A summary of the solution annealing treatments is found in Tables II and 111. The aging temperatures for the (a - ß alloys ranged in 50°C intervals from 200°C (392oF) to approximately 50°C below the ß/a + ß transus in the first cycle and approximately 50°C below the a + ß solution temperature in the second. The transformations in these alloys were known to be rather rapid, and thus aging times started at 0.5 min and went to a maximum of 7000 min. Hardness Testing and Metallographic Preparation— Heat-treated samples were sheared and mounted in Bakelite with their cut edges exposed. After a flat surface was ground, the hardness was determined in Vpn on an Armstrong-Vickers machine utilizing a 20-kg load. Three to four impressions were made on each specimen. After this procedure, the hardness impressions were removed and the edge surfaces of the specimens were prepared for metallographic examination in the usual manner. The etchant was 20 pct HF and 20 pct HNO3 in glycerine. Electrical Resistivity—It has been shown that the isothermal reactions in titanium-base alloys can be followed by measuring the changes in e1etrical resistivity attending these transformations. As mentioned earlier, the reaction kinetics of a-ß alloys are rather rapid; thus, it would be expected that resistivity data obtained from these alloys would be more significant if measurements were made continuously at the respective aging temperatures. To accomplish this end a special apparatus, based on the current-potential principle, was assembled. Its design was derived from earlier devices constructed by Colner and zmeska1 4 and Levinson.5 The set-up provided for the heat treatment zof machined specimens (approximately 2 x l0-4in. in cross section by 3 in. in length) and the determination of their relative

Jan 1, 1962

By Juan Sodi, John F. Elliott

The standard free energy of ReS2 has been measured in the range of 1050° to 1250°K using H2/H2S mixtures and a slight variation of the method described by Hager and Elliott.1 The result is: The experimental method and apparatus were modified slightly for this study. Measurements on Cu2S were made to verify the application of the method to the work on ReS2. THE EXPERIMENTS AND RESULTS Briefly, the experimental method consisted of exposing a chip of copper or rhenium at a known temperature for 8 hr to a slowly flowing gas stream at the same temperature in which Ph2S and PH2 were known. The chip was withdrawn quickly from the hot furnace, and subsequently it was inspected for the presence of a sulfided surface. In the experiments described here, there was no ambiguity in any case as to the presence or the absence of the sulfide. At a given temperature, gas compositions for sulfidization were explored systematically until two compositions were found whose values of ?G°, Eqs. [I] and [2], were within approximately 100 cal of each other, one of which was sulfi-dizing and the other was not. These are termed the "straddle" compositions and it is assumed that the equilibrium composition lies between them. The chief modification to the apparatus, which is shown schematically in Fig. 1 of Ref. 1, was to support the metal specimen on a small alumina boat which could be moved along the reaction tube, 6 mm ID, by platinum wires. An appropriate seal at each end of the reaction tube permitted the sample to be moved from the cold end of the tube into the hot zone in 2 to 3 sec, and the sample could be withdrawn equally rapidly. Thus, it was possible essentially to quench the specimen from the reaction temperature with the reaction gas or helium flowing and without danger of breaking the reaction tube. The usual practice at the end of the experiment was to switch the gas system to the helium tank, flood the reaction chamber with helium, and pull the sample out of the hot zone. The purpose of the modification was to permit study of the sulfidization of copper without the complication of the back-reaction between the gas and the specimen as the latter cooled during slow withdrawal of it from the hot zone; this was a problem in the earlier work.&apos; A further improvement located the tip of the temperature-indieating thermocouple and the specimen precisely at the hottest part of the furnace. A carefully calibrated thermocouple, with its tip at the position of the specimen and with other conditions duplicating those of an actual experiment, showed that in the temperature range of 900° to 1122°C the temperature of the specimen differed from that of the tip of the indicating thermocouple by less than 0.5°C. The two positions were 0.5 cm apart. The reaction gas was prepared from ultrahigh-purity hydrogen (<l ppm O2, <0.5 ppm H2O) and CP grade hydrogen sulfide (99.5 pct H2S). High-purity helium (99.995 pct He) was used. All of these gases were purchased from the Matheson Co. All flow meters were recalibrated by the soap-bubble method with hydrogen, H2S, helium, and several gas compositions used during the study. These calibrations gave a linear relationship with a slope of 1.0 for the plot of log flow rate vs log pressure drop across the flow meter, in accordance with the Hagen-Poiseuille equation. The analysis of the gas was determined in the same manner as was reported previously. Good checks were obtained between the composition of the gas established by the flow-meter settings and by chemical analysis of the gas taken after the mixing bulb and ahead of the furnace. The pressures of H2S, H2, S2, and HS in the equilibrium gas at temperature were calculated from the following data :3 The pressures of the species S and S8 were negligible for the conditions of the experiments.3 There was no sign of vaporization of ReS2 either by weight loss or deposits in the reaction tube. Thus it is not possible to account for the apparent volatility of the compound reported by Juza and Biltz.2 The inlet gas composition and the calculated equilibrium ratio of PH2 S/PH2 for the "straddle" points of each experiment are shown in Table I. The specimens of metal for the experiment were small clippings of annealed copper (99.9+ pct) sheet 0.005 in. thick that was obtained from Baker and Adamson and of "high-purity" rhenium (99.9+ pct) sheet 0.005 in. thick that was purchased from Chase Brass and Copper Co. A specimen was removed from the apparatus; inspected for the presence of the sulfide, and then stored in a sealed vial. A fresh clipping was used in each measurement. The condition of the surface of each specimen after the experiment is noted in Table I.

Jan 1, 1969

By R. W. Armstrong, P. C. Jindal

The hardness dependence on grain size for polycrys-talline cartridge brass and a leaded brass has been measured by Brine11 and Rockwell B testing. In each case, the hardness, H, depends on the average grain diameter, 1, according to: H =Ho + kHl-1/2 where Ho and kH are experimental constants. Diamond pyramid hardness values have also been measured as a function of the indentation size and grain size to give additional information on the nature of the hardness test and the dependence of hardness on micro-structure. The hardness of polycrystalline brass depends on its grain size. Bassett and Davis&apos; demonstrated this as early as 1919 by making Brinell hardness measurements on cartridge brass. Since then, the hardness of this type of material has been measured as a function of grain size by making Rockwell,2&apos;3 Vickers,4 and Brinell5 tests. he hardness dependence on grain size has also been measured for other materials. Angus and summers6 investigated the grain size dependence of the Brinell hardness of polycrystalline copper and a Cu-4.5 pct Sn bronze. In other studies, nickel,? Armco iron,Big an Fe-0.07 pct C alloy,I0 and an 0.39 pct C-12.45 pct Cr stainless steel" have been investigated. In some of the preceding cases, the hardness results have been analyzed to show that the hardness varies with the average grain diameter, 1, according to an l-l\4, l-1/4 or I-2 dependence,11-13 The studies of the influence of grain size on hardness have not been based on any theoretical model. This may be because the hardness of a material is itself a complicated property. However, attempts have been made to correlate, experimentally and theoretically, the hardness of a material with its unidirectional stress-strain behavior.14-l6 On this basis, Hall" proposed that the polycrystal hardness dependence on grain size might follow directly from the Hall-Petch18,19 relation for the grain size dependence of the yield stress. Thus, the hardness-grain size relation was given as: H = Ho + kHl-1/2 [1] where Ho and kH were taken as experimental constants. The relation was applied to the measurements on brass,&apos; copper,6 bronze,= and Armco iron.&apos; More recently, this relation was shown by Armstrong and jindal20 to adequately describe the measurements on cartridge brass made by Bassett and Davis&apos; and Babyak and Rhines.5 In this case, the relationship was taken a step further by independently relating the values of Ho and kH to the values of oyand ky, previously reported by Armstrong, Codd, Douthwaite, and petch21 from measurements of the yield stress dependence on grain size for this type of material. In the present investigation, new Brinell and Rockwell B hardness measurements have been made as a function of grain size for a cartridge brass and a leaded brass. In addition, diamond pyramid hardness values were measured as a function of the indentation size. All these results are applied to a further analysis of the hardness dependence on grain size. MATERIALS AND EXPERIMENTAL METHODS Cartridge brass and a leaded brass were selected for this investigation for two main reasons: it was anticipated 1) that these materials could be cold-worked and recrystallized to a wide range in grain size and 2) that the results to be obtained on these typical industrial materials could be usefully compared with previous investigations. The chemical analyses of the actual materials which were employed are given in Table I. The as-received 1/2- and 3/4-in.-thick plates were given various reductions in thickness by cold rolling. The rolled material was heat-treated at various temperatures between 330" and 850°C for differing time periods from 5 min to 9 hr to achieve a variation in the average grain diameter between 0.0339 and 0.000543 cm.22 During heat treatment, the brass was protected from zinc loss by packing it in chips or foils of the same composition material. Reasonably equi-axed grain structures were obtained in each case. The metallurgical grain sizes of the specimens were determined from measurements of the average linear intercept on a random line. Annealing twin interfaces were not counted along with grain boundaries. The

Jan 1, 1970

By I. J. Lin, A. Mitzmager

This paper describes a series of experiments conducted to determine how the physicochemical properties of the suspending fluid affect the size reduction of quartz in a ball mill. Particular attention was directed toward the differences in rate of grinding and particle size distributions when either polar or non-polar fluids were the suspending liquids. Grinding was performed with water, carbon tetrachloride and methylcyclohexane as the liquids, and also in nitrogen for comparative purposes. From the results ob- tained, the authors have concluded that the presence of water greatly enhances efficiency in the grinding of quartz. Comminution is generally regarded as a purely physical process in which fragmentation of a solid is obtained by the application of suitable forces. In 1944, however, Rehbinder et al.1 pioneered the idea that the interfacial forces between the solid and the media plays an important role in comminution. Since then, a number of researches have shown conclusively that comminution is affected by surface chemical phenomena,2-5 and Gaudin went as far as to compare comminution to a chemical reaction.6 The effect of the viscosity and the density of the fluid media in grinding was also investigated, but

Jan 1, 1969

By D. J. Jr. McAdam

PART I. OUTLINE OF INVESTIGATION, DESCRIPTION OF MATERIAL AND METHODS RESULTS of investigation of corrosion-fatigue of metals at the U. S. Naval Engineering Experiment Station, Annapolis, Md., have been discussed by the writer in several recent papers.1,2,3,4 In those papers references were also given to the work of Haigh in 1917.6 Efforts by the writer to obtain a better understanding of some of the factors involved in corrosion-fatigue finally developed into a general investigation of the effect of time and cyclic stress on corrosion of metals. Corrosion fatigue of Metals As described in the papers mentioned, rotating-cantilever specimens, conically tapered so that the stress is nearly uniform over a length of about 1.5 in., were subjected to simultaneous fatigue and corrosion. For this purpose a stream of water was diagonally applied so as to sweep the specimen from the outer to the inner fillet and surround the stressed surface with water. The fresh water is a calcium carbonate water, the composition of which has been given in a previous paper.2 The salt water is Severn River water having a saline content about one-third that of sea water.

Jan 1, 1928

By David Mitchell

THE production of fluorspar is one of the smaller nonmetallic indus-tries in the United States with a capital investment-about $10,000,000. Shipments from United States mines&apos; in 1936 totaled 176,231 short tons valued at $3,111,268, an average price of $17.65 per ton. The United States produces about 40 per cent of the world&apos;s fluorspar, of which about 90 per cent comes from the Illinois and Kentucky deposits along the Ohio River.2 Further detailed information on the fluorspar industry in the United States is given by Hatmaker and Davis.28 Fluorspar is marketed according to the following relative quantities and specifications: Consump- Specifications, Per Cent Use Form tion in 1936, Per Cent CaF2 SiO2 Fe2O3 Metallurgical Gravel 79 +85 -5 or lump Ceramic,. Ground 10 +95 -3 -0,1 Acid Ground 11 +98 -1 The concentration of the coarser sizes presents relatively little diffi-culty, good recoveries being obtained with jigs and tables. However, sizes smaller than about 10 mesh cannot be successfully treated by ordinary gravity separation methods and so have been accumulating as

Jan 1, 1938

By David Mitchell

THE production of fluorspar is one of the smaller nonmetallic indus-tries in the United, States with a capital investment-about $10,000,000. Shipments from United States mines1 in 1936 totaled 176,231 short tons valued at $3,111,268, an average price of $17.65 per ton. The United States produces about 40 per cent of the world&apos;s fluorspar, of which about 90 per cent comes from the Illinois and Kentucky deposits along the Ohio River.2 Further detailed information on the fluorspar industry in the United States is given by Hatmaker and Davis.28 Fluorspar is marketed according to the following relative quantities and specifications: Consump- Specifications, Per Cent Use Form tion in 1936, Per Cent CaF2 SiO2 Fe2O3 Metallurgical Gravel 79 +85 -5 or lump Ceramic.. Ground 10 +95 -3 -0.1 Acid Ground 11 +98 -1 The concentration of the coarser sizes presents relatively little diffi-culty, good recoveries being obtained with jigs and tables. However, sizes smaller than about 10 mesh cannot be successfully treated by ordinary gravity separation methods and so have been accumulating as

Jan 1, 1938