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By C. G. Dunn, W. W. Martin

Various X ray diffraction techniques have been developed for determining orientations of crystals. Transmission and back-reflection Laue methods1,2,3 in particular have been found to be very useful. In many applications, however, a need still exists for a simplification of the details of analysis so that orientations can be determined more rapidly, and operators with little or no training in crystallography can be easily instructed in the necessary procedures. Filling this need in part is a system of analysis that has been in use for several years in the Metals Section of the Laboratory, General Electric Co., Pittsfield, Mass. This method is an extension of the one of Majima and Togino,4 which uses fifty-five indexed standard transmission Lauemams to cover the entire range of crystal orientations. Greater rapidity in determining the orientation of a crystal is achieved by using a slightly larger set of standard Lauegrams and by systematizing the steps followed. This system of analysis in its present form constitutes the subject matter of the present paper. Although the method uses transmission X ray Laue patterns, many of the operations are applicable to the analysis of back-reflection Lauegrams. Very briefly the method is based on the fact that Bragg and azimuthal angles for two or more identified Laue diffraction spots completely determine the orientation of a crystal, and standard Laue photographs of the Majima and Togino type provide a ready means for identifying Laue spots. After a preliminary analysis, which consists of identifying the pattern in terms of one of the standard patterns, a correctly marked Lauegram is selected to carry out the final analysis. From measured Bragg and azimuthal angles, the normals (poles) of the reflecting planes for the selected Laue spots are located on a stereographic projection and the orient.at.ion of the crystal determined (in terms of the positions of cube poles for example) by a few additional operations on a \I1ulff net. The actual procedure followed will be illustrated in detail after a brief treatment or X ray problems connected with orientation determinations and a discussion of useful principles and techniques. X Ray and Plotting Problems RELATIVE ADVANTAGES OF BACK-REFLECTION AND TRANSMISSION LAUE DIFFRACTION METHODS In the back-reflection Laue method it is easy to make a direct measurement of the tilt of a zone of Laue spots. Specimen thickness also is generally no problem. In the case of thin specimens, however, the time of exposure is less for the transmission Laue method. For example, with silicon ferrite this advantage holds up to a thickness of 0.6 mm if tungsten radiation at 40 kv peak voltage is used. Furthermore, the transmission photograph represents the entire thickness of the crystal whereas the back-reflection picture represents only a surface layer (a thickness of less than 0.05 mm for silicon ferrite), the deeper layers supplying a negligible amount to the diffraction pattern. If a crystal is distorted, the transmission Laue pattern gives more information particularly with regard to the orientation of the crystal and the overall amount of spread in orientation. Finally a 3 1/4 X 4 in. film suffices for transmission photograms; a 4 X 5 in. or 5 X 7 in. film with a central hole is needed for back-reflection patterns Not only is the manipulation of film easier because of the size of film, but

Jan 1, 1950

By W. A. Backofen, D. Lee

Titanium of commercial purity (RC-70) and two all-a (hcp) alloys (4Al-1/4O2 and 5Al-2.5Sn) were tested in sheet form under conditions of combined-stress loading. Plane-strain compression and plane -strain tension together with uniaxial compression and tension provided information for constructing the tension-tension and compression-compression quadvants of the plane-stress yield locus. Coordinate values were obtained from yield-strength measurements; the slope at each point was based on a plastic-strain ratio. The materials represented a wide range of plastic anisotropy, which was reflected in correspondingly varied departures of the yield locus from that for iso-tropy. The tension-tension and compression-compression quadrants were not identical. If the uniaxial yield strength is insensitive to direction and the sign of the applied stress, the Mises or perhaps the Tresca criterion is adequate for predicting the onset of yielding under combined-stress loading. The prediction is not so straightforward in plastically anisotropic material. Depending upon details of texture, combined-stress yielding resistance may be substantially different from that indicated by the usual criteria.' However. a continuum theory of yielding under such conditions is available for reasonably direct application.2,3 Whether or not it can be used successfully must, of course, depend upon how closely the conditions of the theory are satisfied in the material in question. In its present form the theory contains no provision for Bauschinger-type effects, nor does it allow for the consequences that might grow out of deformation by more than one mechanism, e.g., by slip and twinning.4 With those limitations, more importance must be attached to experimental information about the yield surface for anisotropic materials. Studies of the yield surface have generally been confined to states of plane stress (tension-tension and tension-compression) established in thin-wall tubes by externally applied forces and internal pressure. That approach would be entirely suitable for anisotropic materials if they were available in tube form or in sufficient bulk to be machined into test specimens. 6-9 It would not serve for the textured and anisotropic sheets of much current engineering interest, unless sheet could be formed into high-pressure tubing without structural change. Since there would seem to be little chance of the latter, practical alternatives are needed. One such alternative is the subject of this paper. It involves the use of a few selected but relatively simple uniaxial loading tests on sheet specimens. The underlying idea is presented here with some details relating to application and a few results bearing on how well it has worked. THE BASIS OF TESTING Conventional Measurements. Yielding for a general case of anisotropy is described by the plane-stress locus in Fig. 1. Reference directions in the sheet specimen are related to the stress coordinates as shown with the insert: x is the rolling direction, y the transverse direction, and z the through-thickness direction. The four reference-direction yield strengths (t, tension; c, compression), Xt, Xc, Yt, and Yc, are all different. Nine paths that can be followed in uniaxial loading experiments are identified. Paths 1, 2, 3, and 4 represent conventional tension and compression testing to establish Xt, Yt,.Xc, and Yc. Useful information also comes from the ratio of the transverse plastic-strain components associated with yielding and flow in conventional tension or compression. A condition to be satisfied in locus construction is that the plastic-strain vector be everywhere normal to the locus.10 The projection of the vector which would

Jan 1, 1967

By S. J. Burns, W. W. Webb

The dislocation arrangements formed during unsteady propagation of cleavage fractures on (010) planes in LiF have been investigated by high-resolution etch-pit techniques and by X-ray diffraction topogvaphy. In impurity-hardened crystals, cracks propagating in any direction on (010) planes at velocities of about 103 cm per sec are usually accompanied by light plastic deformation which is enhanced wherever the crack is allowed to hesitate. Dislocations with [011] type Burger's vectors are often formed looped around the crack tip. They are dragged along with the crack tips leaving edge dislocations lying nearly parallel with the clearage surface about 5µ from it. These dislocations are easily overlooked since they may initially intersect with the cleavage surface infrequently and later may disappear from the crystals due to image forces during aging at room temperature. Although the plastic work associated with these dislocations is only comparable with the reversible (010) surface energy, an instability in crack propagation appears due to a limit of the velocity with which dislocations can be dragged along with the crack. Transition to a fully brittle nonplastic fracture mode occurs when the crack velocity exceeds the velocity with which dislocations can be moved by the stress field of the crack. TRANSITIONS between brittle and ductile modes of fracture appear to be common processes in the behavior of materials in which the flow stress is close to the stress at which brittle cracks can propagate. High-strength engineering alloys fit into this category and the transition from ductile to brittle modes of crack propagation may be associated with catastrophic failures in these important materials. The methods of theoretical mechanics indicate the geometrical conditions in which these processes may occur but the detailed mechanisms controlling the generation of dislocations in the initiation of the plastic mode and the process by which the loss of plasticity occurs on transition to the brittle mode are not yet clear. Since these mechanisms are not amenable to direct observation in engineering alloys we have chosen a semi-ductile alkali halide LiF as a model system in which to observe transition mechanisms between plastic and brittle modes of crack propagation. Cleavage cracks in this system are known to undergo a plastic-elastic transition at velocities between 103 to 104 cm per secl and dislocation configurations in LiF are amenable to study by optical-microscopic observation of dislocation etch pits2 and by X-ray diffraction topography.3 We have been able to identify a dislocation mechanism responsible for the transition from brittle to slightly plastic modes of crack propagation in LiF. A model based on this mechanism permits calculation of the critical crack-propagation velocity for this transition and indicates the importance of the stress dependence of dislocation velocity in determining the critical crack velocity. Of course, this model cannot be carried directly over to more complex materials but it is quite possible that its qualitative features are generally applicable. Dislocations intersecting fresh cleavage surfaces in LiF were reported by Gilman,4 by Forty,5 and by Wash-burn et a1.6 Such dislocations were construed to have been nucleated where the cleavage crack hesitated or stopped, near stress concentrators, such as cleavage steps, on the crystal face. The loci of such dislocations deposited On the cleavage plane in bands were called "stop lines" or "deformation zones". Gilman, Knudsen, and walshl measured the velocity of cleavage cracks in lithium fluoride by evaporating thin resistive strips on the crystal side and recording

Jan 1, 1967

By D. G. Westlake

NUMEROUS hydride habit planes in zirconium, Zircaloy-2, and Zircaloy-4 have been observed by various researchers and have been compiled by Louthan and Angerman.1 While only partial agreement exists on the identification of these habit planes, there is a growing consensus that a tensile stress applied on a polycrystalline rod during slow cooling will cause the hydride platelets to lie perpendicular to the tensile axis. walters2 agrees with Louthan and Angerman who reason that "The particular plane on which precipitation will occur is the plane that will result in the maximum relief of stress by precipitation of the low-density hydride." The experiment described below was initiated to test this hypothesis on a single crystal of zirconium. A single-crystalline tensile specimen of zone-refined zirconium, designated number 13 in a previous study,3 had a square cross section with one pair of parallel faces only 4 deg from the basal plane of the hcp structure. The tensile axis was about 86 deg from the (0001) pole, 14 deg from the (1010) pole, and 16 deg from the (1150) pole. The specimen was charged to 50 ppm H by allowing reaction with the correct amount of hydrogen at 1073°K, homogenizing at this temperature for 2 hr, and cooling in the furnace. Two- surface analysis revealed that hydride platelets had precipitated parallel to the three {1010} prism planes. The specimen was heated to 673°K, in air, for 16 hr to bring the hydrogen back into solid solution, and furnace-cooled to 573°K.* It was then stressed in ten- sion at 1.2 kg per sq mm. (This produced a resolved shear stress of 0.6 kg per sq mm for glide on the prism plane with the highest Schmid factor. A higher stress would have produced significant strain at this temperature because the critical resolved shear stress for prism-plane glide in this crystal was only 1.0 kg per sq mm at ambient temperature.) The poles of the three prism planes were at angles of 14, 46, and 74 deg to the tensile axis so that an applied stress would exert normal stresses on the three planes in the approximate ratios 10:5:1, respectively. As soon as the stress was applied, the sample was furnace-cooled to room temperature. It was ground on wet Sic paper and chemically polished in 50 parts HNO3, 50 parts water, and 7 parts HF before examination. The result of the experiment was negative; i.e., the three (1010) planes were of equal importance as hydride habit planes.* It must be admitted that the ex- periment is not entirely comparable to those performed by walters,2 Louthan,5 and Louthan and Marshall6 on polycrystalline materials. They have found that some critical stress (below the macroscopic yield stress) is required for stress-orienting the hydrides. The applied stress in the present experiment was, by necessity, smaller by an order of magnitude than the critical stresses they report for the various materials. If the critical stress was related to the microstructure, one might postulate that it should be decreased by increasing the degree of perfection of the metal lattice and by increasing the grain size. But the experiment above suggests that, for the extreme case of a zone-refined single crystal of zirconium, the critical stress lies above that which one can apply without introducing gross plastic strain as another variable.

Jan 1, 1967

By J. L. Zakin, G. K. Patterson, J. M. Rodriguez

Correlation has been obtained between drag-reducing characteristics for turbulent flow in a pipe and measurable properties of several polymer solutions. Several concentrations of high molecular weight polymethyl methacrylate in toluene, high molecular weight polyisobutylene in both toluene and cyclohexane, medium molecular weight polyisobutylene in cyclohexane and benzene and low molecular weight polystyrene in toluene were studied. Data obtained in these nonpolar solvents and literature data for more polar solvents were successfully correlated as the ratio of measured friction factor to purely viscous friction factor vs the modified Deborah number vrl/D0.2 where 71 is the first-mode relaxation time of the solution estimated by the Zimm theory. A shift factor which is a function of intrinsic viscosity 1/(4[?] - 1) allowed all the data obtained with nonpolar solvents to be correlated as a single function. For these systems, most of the data fit a single curve to within + 5 percent of the average friction factor ratio. The shift factor did not give a unique function of the data for the more polar systems. INTRODUCTION The phenomenon of drag reduction in polymer solutions was first studied by Toms1 in dilute solutions of polymethyl methacrylate in mono-chlorobenzene. The drag ratio for flow through circular tubes has been defined2 as the ratio of the pressure. drop of the solution to the pressure drop of the solvent at the same flow rate. The drag ratio is less than 1.0 for a drag-reducing fluid. Practical use of drag reduction is being made in fracturing operations in the petroleum industry.3 A more fundamental quantity is the friction factor ratio, defined as the ratio of the observed pressure drop to that predicted for a solution of the same viscosity characteristics and density at equal flow rates using the Dodge-Metzner friction factor equation4 Viscous solutions with drag ratios greater than 1.0 can have friction factor ratios less than 1.0. For practical applications, it is drag reduction which is of interest. However, for correlation the fundamental ratio is the friction factor ratio. In recent years, drag reduction has been studied extensively. Recent studies have shown that reasonable predictions of the incipience of drag reduction in polymer solutions can be made from the properties of the solutions and the flow Variables.5 However, it has not been possible to predict accurately the amount of drag reduction to be expected for a given polymer solution without any drag-reducing turbulent flow data on the same solution.6 For flow through circular tubes it is well established that there is a concentration effect and a diameter effect on drag reduction. Toms1 observed that drag reduction increases with an increase in the concentration of the solution up to an optimum concentration beyond which, due to the increased viscosity of the solutions, the drag ratio increased with an increase in concentration. Hershey 7 observed that for dilute Newtonian polymer solutions in nonpolar hydrocarbon solvents, a normal transition region may be obtained followed by a departure from the von Karman equation (on a friction factor-generalized Reynolds number chart). The point of departure is a function of diameter: the smaller tubes depart at the lower Reynolds numbers. Concentrated solutions may show no abrupt transition region, but merely an

By C. A. Bruch, W. E. McHugh, R. W. Hockenbury

Commercially pure molybdenum specimens were irradiated in the Materials Testing Reactor for an estimated exposure of 1.9 to 5.9x10 20 thermal nvt. Prior to irradiation, the material was ductile in the tension test; whereas after irradiation, it was brittle. The results of tension tests conducted at various temperatures revealed that the transition temperature for this material had been increased from —30' to +70°C as a result of the radiation exposure. From metallographic studies, it is concluded that the embrittlement is due to submicroscopic changes which raise the flow stress-temperature curve of the material. PRIOR to ten years ago, the bombardment of metals with high speed heavy particles such as protons, deuterons, a particles, and neutrons was primarily of academic interest. These particles existed in extremely small numbers in a few laboratories and only under very special conditions. With the introduction of the nuclear power reactors, these particles are now available in enormous quantities and it has become important from a practical standpoint to learn how well metals withstand severe reactor radiations. The principal radiations of concern in reactors are neutrons and y rays, since the heavier charged particles exist only in limited regions of reactors. Seitz,' in a theoretical approach, has shown that all of these high energy particles, except y rays, have enough energy to knock atoms from lattice sites and that these knocked-on atoms have enough energy to cause additional displacements of neighboring atoms. The result of this process would be a lattice containing many imperfections in the form of interstitial atoms and vacancies. Hence, it would be expected that the properties of metals can be changed by reactor radiations, and such has been found to be the case. Only a limited amount of data have been published in the open literature,2,3 and these data have been for relatively low integrated-neutron-flux exposures. In general, one effect of reactor radiations on the mechanical properties of metals has been to increase strength and decrease ductility. It is now possible to irradiate specimens in the Materials Testing reactor at much higher fluxes and to higher integrated exposures than previously possible.' The data which are being presented here describe the effects of a relatively high exposure on some mechanical properties of molybdenum. Experimental Procedures Material: The molybdenum used in this investigation was obtained from the Climax Molybdenum Co. The two heats from which the material originated were melted under a vacuum of 5 to 30 microns and cast in 7 in. diam water-cooled copper molds. One ingot contained 0.061 pct C and the other 0.066 pct. Both ingots were hot worked by extrusion and rolling into % in. diam rods, after which they were annealed and straightened. The rods were subsequently heated to 1100°C in a hydrogen atmosphere and swaged to ½ in. diam rods at the Research Laboratory of the General Electric Co. The resulting material had an average hardness of 264 VPN, or 99.2 Rb, and had an average of 5000 grains per sq ml in the transverse section, Test Specimens: Substandard-size tensile specimens were prepared according to the specifications shown in Fig. 1. These specimens were used also for hardness measurements. The metallographic specimens were small disks which were polished, etched, and photographed at X1000 prior to irradiation. The field studied was

Jan 1, 1956

By E. A. Elevatorski

Wollastonite, named after William H. Wollaston, an English chemist, is a calcium metasilicate, CaSiO3. It has a short history as an industrial mineral. The earliest production of wollastonite is reported to be from a deposit near Code Siding, located north of Randsburg, Calif. At this locality small tonnages of wollastonite were quarried during 1933-34 and 1938-41, and processed into mineral wool. This operation was largely experimental and virtually no United States production was again reported until the 1950s when a large deposit near Willsboro, N.Y., was developed by the Cabot Corp. A processing plant was placed on-stream in 1953, with nearly continuous production to date, and it is currently operated by Interpace Corp. Since 1958, wollastonite deposits in the Little and Big Maria Mountains of Riverside County, and in the Panamint Range of Inyo County, both in California, have operated intermittently for production of both ornamental and commercial wollastonite. During 1973, the United States was the major producing country, furnishing about 75% of the world's output. Recent production also has come from Finland, Mexico, and Kenya. Small amounts have been shipped intermittently from India, USSR, New Zealand, Republic of the Sudan, Republic of South Africa, and South-West Africa. The principal use of wollastonite is in the manufacture of ceramics, especially wall and floor tiles. Other uses are for paints, fillers, adhesives, reinforcing agents, plastics, fluxes, and glazes. Mineralogy Pure wollastonite, CaSiO3, has the composi¬tion of 48.3 % CaO and 51.7 % SiO,. However, it is seldom found in the pure state due to the ease with which it takes into solution the metasilicates of manganese, magnesium, iron, and strontium. Predominantly, wollastonite occurs as a contact metamorphic deposit forming between limestones and igneous rocks. Commonly associated minerals are garnet, diopside, epidote, calcite, and quartz. It has a specific gravity of 2.8 to 3.0, and hardness of 4.5 to 5 on Mobs' scale. When pure, it has a brilliant white color, but with impurities it may be grayish or brownish. Luster is vitreous to pearly. Melting point of wollastonite is about 1540° C. Wollastonite occurs in coarse-bladed masses, rarely showing good crystal form. It is usually acicular or fibrous, even in the smallest of particles. The most unique property of crushed and ground wollastonite is its cleavage. Fragments of crushed wollastonite tend to be needle-shaped, imparting a high strength, and this property is the basis for many of its uses. The fiber lengths are commonly in the ratio of 7 or 8 to 1, length to diameter. Some crystals of wollastonite fluoresce under short-wave or longwave ultraviolet light, or both; colors ranging from yellow-orange to pink-orange. Specimens may also show phosphorescence. The brightness of wollastonite is a property of considerable importance to the paint industry. Material, 99% pure, with a size of -325 mesh, has a General Electric reflectance rating of 92 to 96%. Chemically, wollastonite is inert and this property makes it useful as a filler and reinforcing agent. There are two polymorphs of calcium silicates: wollastonite, a low temperature form, and pseudowollastonite, a high temperature form. Inversion of wollastonite to pseudowollastonite occurs at about 1120° C, resulting in an increase in the coefficient of expansion and a color change. Pure white wollastonite, on inversion, may change to a cream tint, or various shades of red or brown. The color change is thought to be due to the presence of iron and strontium.

Jan 1, 1975

By E. C. Skinner

Cardox, which consists essentially of a steel tube containing carbon dioxide compressed to the liquid state, is a trade name designating a device used principally in coal mines to break down coal. The use of Cardox to break down coal was developed in the Middle Western coal fields around 1927 or 1928. At first, a large, heavy tube was used, requiring many parts. Through experimentation and technologic improvement, the weight of the tube has been brought down from 88 lb. to 20 lb. and the new, light tube makes use of larger volumes of CO2 at lower pressures than formerly was thought possible. This blasting method is extensively used today in coal mines and probably one tenth of the bituminous coal produced in the United States is prepared with Cardox. Construction and Application As used in the coal mine, Cardox looks like a piece of steel pipe 2 to 3 in. in diameter and 3 to 5 ft. long. When it is necessary to fire a shot, the tube is wired with two leg wires that are long enough so that when the tube is pushed the desired distance into the hole the wires will protrude 5 to 6 in. It has been found that No. 18 insulated copper annunciator wire is the best material for leg wires. A two-conductor rubber-covered shot-firing cable of sufficient length is then attached to the leg wires. A 2 to 6-volt current is applied, either by means of shot-firing batteries or the standard type of shot-firing generator. Upon the application of the current, the charge in the cartridge is set off. Because the breaking action upon the coal is supplied by the relatively slow physical expansion of the carbon dioxide gas rather than by the rapid chemical formation of gases as in more familiar explosives, the result is a slow heaving action, somewhat similar to that produced by black powder, rather than a shap detonating effect. Permissible powder does its work at roughly ten times the pressures in use with Cardox. The escaped gas from a Cardox shot apparently works along the planes of weakness in the coal face, such as the coal cleats, the edges of the cut, and along the top of the seam, breaking off the coal toward the free faces and heaving the mass of coal outward. Cardox tubes are manufactured in various sizes to take care of variable conditions of seam size, undercut depth, cleatage, coal hardness, and so forth. The tubes range in diameter from 2 to 3 in. (Table I) and the normal CO2 charge is from 24 to 96 oz. For any one tube size, several sizes of shear disks are available, thus excellent control of the type of shooting desired is possible. In general, local experience and experimentation dictate the size of tube chosen for any particular job. For high coal, and deeper or wider cuts, a larger size is chosen to start with, and the opposite is true for low coal, softer coal, or small cuts. After the tube size has been experimentally selected, tests are run in the mine to determine how many shot holes are needed and where the holes should be placed. The size of the disk,

Jan 1, 1944

By E. C. Skinner

Cardox, which consists essentially of a steel tube containing carbon dioxide compressed to the liquid state, is a trade name designating a device used principally in coal mines to break down coal. The use of Cardox to break down coal was developed in the Middle Western coal fields around 1927 or 1928. At first, a large, heavy tube was used, requiring many parts. Through experimentation and technologic improvement, the weight of the tube has been brought down from 88 lb. to 20 lb. and the new, light tube makes use of larger volumes of CO2 at lower pressures than formerly was thought possible. This blasting method is extensively used today in coal mines and probably one tenth of the bituminous coal produced in the United States is prepared with Cardox. Construction and Application As used in the coal mine, Cardox looks like a piece of steel pipe 2 to 3 in. in diameter and 3 to 5 ft. long. When it is necessary to fire a shot, the tube is wired with two leg wires that are long enough so that when the tube is pushed the desired distance into the hole the wires will protrude 5 to 6 in. It has been found that No. 18 insulated copper annunciator wire is the best material for leg wires. A two-conductor rubber-covered shot-firing cable of sufficient length is then attached to the leg wires. A 2 to 6-volt current is applied, either by means of shot-firing batteries or the standard type of shot-firing generator. Upon the application of the current, the charge in the cartridge is set off. Because the breaking action upon the coal is supplied by the relatively slow physical expansion of the carbon dioxide gas rather than by the rapid chemical formation of gases as in more familiar explosives, the result is a slow heaving action, somewhat similar to that produced by black powder, rather than a shap detonating effect. Permissible powder does its work at roughly ten times the pressures in use with Cardox. The escaped gas from a Cardox shot apparently works along the planes of weakness in the coal face, such as the coal cleats, the edges of the cut, and along the top of the seam, breaking off the coal toward the free faces and heaving the mass of coal outward. Cardox tubes are manufactured in various sizes to take care of variable conditions of seam size, undercut depth, cleatage, coal hardness, and so forth. The tubes range in diameter from 2 to 3 in. (Table I) and the normal CO2 charge is from 24 to 96 oz. For any one tube size, several sizes of shear disks are available, thus excellent control of the type of shooting desired is possible. In general, local experience and experimentation dictate the size of tube chosen for any particular job. For high coal, and deeper or wider cuts, a larger size is chosen to start with, and the opposite is true for low coal, softer coal, or small cuts. After the tube size has been experimentally selected, tests are run in the mine to determine how many shot holes are needed and where the holes should be placed. The size of the disk,

Jan 1, 1944

By A. D. Schwope, L. R. Jackson, A. M. Hall

IN the course of an extensive investigation of the comparative properties of several types of copper, data were obtained on the cold working and subsequent recovery and recrystallization of the copper. In this paper, relations between the degree of cold work and the course of recrystallization are developed. In particular, it is shown that the course of recrystallization over long periods of time and at low temperatures can be predicted with reasonable accuracy from short-time high-temperature recrystallization data. MATERIALS USED IN THE INVESTIGATION Three types of copper were used in this investigation; these were tough pitch, OFHC, and OFHC copper with 15 oz per ton of silver added identified herein as Lots C, D, and F, respectively. The material was supplied by the American Metal Co. Ltd., and processed by the Rome Cable Co. The analyses of the three types of copper are given in Table 1. Originally, it started as wire bars 4 X 4 X 54 in. These bars were heated to 871°C in air and hot rolled to 3/8 in. rod. This rod was cold drawn to 0.204 in. rod (70 pct reduction) in three passes and annealed at 427°C for 4 hrs. A portion of the 3/8 in. rod was saved for more detailed cold-drawing experiments. Starting with the annealed 0.204 in. rod, hard-drawn wires were produced by cold drawing to 0.081 in. wire in 7 passes (84.4 pct reduction in area). Half-hard wires were made by cold drawing the 0.204 in. rod to 0.102 in. wire, (75 pct reduction) in 6 passes and annealing at 427°C for 4 hrs. This annealed wire was then drawn to 0.081 in. wire, (37.1 pct reduction) in 2 passes. Square rod was made by hot rolling wire bars to 9/16 in, rod in 8 passes. This rod was then cold drawn to 0.445 in.-diam rod and annealed for 3 1/2 hrs at 538°C. It was then cold drawn to 0.325 in. rod and again annealed at 538°C for 3 1/2 hrs. Hard square rods were then made by cold rolling this material to 0.257 X 0.257 in. rod, a reduction in area of 20.5 pct. Half-hard square rods were made by cold rolling to 0.268 X 0.268 in. square, annealing for 332 hrs at 538°C and cold rolling to 0.257 X 0.257 in., a reduction of 8.5 pct. Soft wire or rod was made from all types of copper by annealing the hard-drawn material at 343°C for 3 hrs in steam. The grain size of the annealed 0.081 in. wire varied slightly among the three lots: from 0.015-0.025 mm in the silver-bearing OFHC to 0.035-0.045 mm in the OFHC copper. The annealed 0.257 in. wires were more uniform in grain size, averaging about 0.035-0.045 mm. Representative structures are shown in Fig I. EFFECT OF COLD DRAWING AND TENSILE STRAIN ON STRENGTH OF COPPER Conventional tensile tests were performed on the three types of copper in the various degrees of hardness described above. Results of these tests are tabulated in Table 2. The results in Table 2 are

Jan 1, 1947

SEC. 1. The membership of the Institute shall comprise seven classes, namely: 1. Members; 2. Honorary Members; 3. Senior Members; 4. Associates; 5. Junior Members; 6. Rocky Mountain Members; 7. Junior Foreign Affiliates. All shall be equally entitled to the privileges of membership, excepting that no member of any class whose residence shall be outside of the United States of America, except those residing in Mexico and Canada, shall be entitled to vote. SEC. 2. Members: A person to be eligible for election or transfer into the class of Member must be at least 27 years of age and must have had at least six years' employ¬ment in the practice of engineering, mining, geology, metallurgy, or chemistry, during at least three years of which he must have held positions of responsibility in one or more of these fields. This shall not apply to Rocky Mountain members. SEC. 3. Honorary Members: Each candidate for Honorary Membership must be recommended by at least ten Members and must be elected by ballot at a meeting of the Board of Directors by the unanimous vote of the Directors present; provided, how¬ever, that the number of Honorary Members shall at no time exceed twenty. Such election shall automatically act as a transfer of such member from any other class of membership to which he may belong. SEC. 4. -Senior Members: The Directors may in their discretion place on the Honor Roll of Senior Members and transfer from any other class of membership other than Honorary Membership the name of any Member or Associate who has reached the age of seventy years and who has paid dues for thirty years, and for good reasons in special cases may vary the age or number of years. SEC. 5. Junior Members: A person to be eligible for election to the class of Junior Member must be qualified through education or experience to hold a subordinate position in engineering, mining, metallurgical, geological or chemical work. He shall be not more than 33 years of age at the time of election. SEC. 6. Associates: A person eligible for election or transfer into the class of Associate shall be one, who, in the opinion of the Committee on Admissions and the Board of Directors, is suitable for such election or transfer by reason of his interest in or connection with mining, geology, metallurgy, or chemistry. SEC. 7. Rocky Mountain Members: Membership in a special class of life member¬ships without annual or other dues and without right to receive the printed literature distributed to other members or the right and privilege of wearing the official badge of technical members, is given to all members of the Rocky Mountain Club at the time of consolidation. These special Rocky Mountain memberships shall be claimable by any member of the Rocky Mountain Club, not already a member of the Institute, availing himself of the privilege within ninety days after mailing of notice to him, by the Secretary of the Institute. SEC. 8. Junior Foreign Afliate: Upon appropriate action by the Board of Direc¬tors and with its approval, members in good standing of any technical society having its headquarters outside of the United States of America, Canada, or Mexico,

Jan 1, 1937

By F. E. Streigert, J. H. Mallery, A. H. Geisler

The existing data on precipitation in Cu-Be alloys have been reconciled with some new data to present a more complete and rational analysis of the atomic rearrangements involved in the process. Two different coherent precipitates were proposed. Each is responsible for a type of diffuse X-ray diffraction effect. THE nature of the atomic rearrangement involved in the precipitation of the compound y-Cu-Be from supersaturated solid solutions of the copper-base alloys has been the subject of two extensive investigations.'-" he precipitate in this alloy system has a rather simple structure; at equilibrium it has a body-centered cubic unit cell with copper atoms on the corner sites and beryllium atoms on the center site. While a simple mechanism would be expected to obtain for the precipitation process, those which have been proposed do not adequately explain all of the reported data. The purpose of the present report is to rationalize some of the previous observations and to contribute new data on the subject obtained by the usual techniques of structure analyses. The Cu-Be system of alloys represents one of the most extensively investigated precipitation systems, so a detailed review of the pertinent literature on structure is in order. Striations: Since one of the important features of the mechanism—the habit plane—is determined by microscopic examination, microstructural features that have been observed should be considered at this time in order to assess their role in the precipitation process in Cu-Be alloys. One of the most prominent features appears as regularly oriented markings or stripings (ripples or striations) across the crystals on a suitably polished and etched sample of the aged alloy. These have been observed and interpreted by many workers. Masing and Dah14 observed striations in both the a and /3 phases of a 2.5 pct Be alloy, and attributed them to the beginning of the precipitation of very fine y crystals along favored planes of the a matrix and to the local decomposition of the £ crystals into a plus y. These markings were apparent in both the a and /3 phases of the micrographs given by Smith5 for a 2.75 pct Be alloy, while Moore" detected them in the a matrix phase using polarized light. Desch7 attributed the markings to fine precipitate particles lined up on a family of matrix planes like Masing and Dahl had done, and he identified the planes as those of the (111) family, the slip planes of the matrix. The striations were presumably the structure which Northcott- eferred to as "martensitic." He found that this structure was present in samples that had been aged under widely different conditions with hardness values ranging from 160 to 400 VHN. Later Guinier and Jacquet' also found striations in both the a matrix and secondary /3 phase. Their analyses showed that the striations were along (110) planes of the matrix phase, apparently in contradiction of the results of Desch. They found that the striations were developed during aging on the surface of a specimen that had been polished prior to aging.' On this basis they concluded that the striations represented mechanical deformation by slip along (110) planes of the matrix resulting from precipitation stresses.'" Similar but fainter striations were observed in the same manner on an A1-Cu alloy sample, in which the change in volume was much less than for Cu-Be alloys and

Jan 1, 1953

By R. H. Hiltz

Three titanium alloys, known to provide a mechanically unstable p structure after quenching, were selected as material for a study of the Origin and nature of stress-induced transformation. Data from hardness and tensile tests were used to study the initiation of marten site and the effect of the transfornation and transformation products on the mechanical properties of B titanium, particularly the stress-strain relationships. With the exception that the materials investigated showed a natural aging reaction in the P condition, their behavior is analogous to that of' stress-transforming stainless steels. The transformation starts after initial yielding and is essentially complete just prior to .maximum loading. Once initiated, the martensite itself is capable of propagating the transformation. The deforn-ing materials show low strength, high ductility, and a high strain-hardening exponent. In metallic systems which undergo a diffusionless transformation on cooling, it is possible to produce a similar transformatiorl by plastic deformation under special conditions. Dewez' in delineating Ms temperatures for the titanium-rich end of most titanium binary systems:, established the point of intersection of Ms with room temperature for those particular alloys. Alloy compositions which would quench to mechanically unstable p were established from these data. The first investigation into the possibility of stress-induced transformation occurring in titanium met with little success. It was concluded that such a transformation in titanium was possible only under very special conditions. Continued effort in this direction, however, showed certain Ti-Mo binaries to be highly susceptible to this phenomenon. This fact was utilized by Liu to study the habit planes of martensite and by Machlin and wenig5 to study the effects of stress and plastic deformation on the P to a prime transformation in titanium alloys. The first investigation of the mechanical properties of these alloys was conducted by Holden and his coworkers at Battelle Memorial Institute.' The potentiality of the alloys i3s sheet-forming material was recognized, and the data reappeared in a survey of this type of titanium material.? The initial optimism was tempered by the difficulties in obtaining homogeneous 2i-Mo alloys in commercial-size ingots. The increasing need for good sheet-forming material kept investigation in this direction active, resulting in the development of a commercial alloy of this type.' This study was undertaken to obtain a clearer picture of the mechanism of stress-induced transformation in titanium. MATERIALS AND TEST PROCEDURE Three alloys were selected for this investigation: one a Ti-Mo alloy; another a Ti-Cr-Mo alloy, known to deform martensitically; and the third, a Ti-V alloy with an Ms below room temperature and an Md above room temperature, but known to contain w phase in the quenched condition. The compositions of these alloys are given in Table I. These alloys were cast as 5-lb experimental ingots, employing skull melting. The ingots were hot-forged into 1-in.-round bars. Two 72-in.-thick slices were cut from each bar for hardness testing; the remainder was cut into 0.252-in. tensile blanks. The tensile blanks and one hardness disc from each bar were heated in salt to a temperature 100°F above the 0 transus, held 2/2 hr, and brine quenched. X-ray analysis showed the pieces to be 100 pct 8.* *Omega cannot be distinguished from p by the technique employed. Tensile blanks were machined to size after this treatment. TO prevent transformation due to machining stresses, the shoulders and gage length were ground. True stress-true strain curves were

Jan 1, 1960

By E. J. Efsic, C. M. Wayman

The fcc to bcc y to a transformation in an Fe-24.5 at. pct Pt alloy was studied and found to be martensitic. Large martensite plates are formed in this alloy which are ideal for crystallographic studies; the Ms temperature is about -5°C. Ordering of the parent phase prior to transformation to martensite causes a pronounced lowering of the Ms temperature. For the first time in the case of alloys of iron, measurements of the habit plune, orientation relationship, and magnitude and direction of the shape stvain have been obtained on preselected martensite plates. It is shown that by measuring the initial and displaced (as a result of the shape strain) positions of three or more non-parallel fiducial scratches on one reference surfice the maguitude and direction of the shape strain and the state of distovtion at the parent-mavtensite interface can be determined with a reasonable degree of accuvacy. The agreement between the phenomenolog-cal crystallographic theory of martensitic transformations and the experimental results habit plune, orientation relationship, and nlagnitude and direction of the shape strain) is essentially perfect, and is thought to be the best reported to date for iron alloys. THE comparatively few investigations on binary Fe-Pt alloys have been concerned with magnetic properties, ordered phases, and the equilibrium diagram itself. Various investigations have shown that the equilibrium diagram is of the open y field type; in addition there is a peritectic transformation, a minimum melting point, and three ordered phases: Fe3Pt, FePt, and FePt3. The results of a number of investigators1-6 have shown that the temperature dependence of the y to a (and reverse) transformation in iron-rich alloys is reasonably well-established. The Fe3Pt phase was first noted by Kussman and Rittberg,7 and its existence was later confirmed by Franklin et al,8 Steijn and Donahoe,9 and Berkowitz et al.10 The FePt phase was reported by Isaac and amman' and later confirmed by a host of others.3,6,7,12,13 Kussman and Rittberg7 noted the existence of the superlattice FePt3. Prior to 1956, the 1 to a transformation in iron-rich Fe-Pt alloys was termed "irreversible". In 1957 Klokholm and Donahoe" termed this transformation martensitic, presumably because of its occurrence at temperatures well below those where diffusion normally takes place. This observation served as a point of departure for the present investigation, and in the very early stages of the work to be discussed it was readily verified that the y to a transformation in iron-rich Fe-Pt alloys is indeed martensitic. Quite large martensite plates exhibiting the well-known surface tilt were observed. These plates are ideal in the sense that the boundaries between the parent and product phases were usually very straight. It thus appeared that Fe-Pt martensites would be ideal for crystallographic studies, and this was the purpose of the present investigation. Details of the crystallography of martensitic transformations (particularly for iron alloys) and the phe-nomenological crystallographic theories15-l7 have been discussed many times and thus will require very little elaboration here. The end result of these theories, given in matrix form, is

Jan 1, 1968

By D. I. Bardos, Paul A. Beck

The 1100°c isothermal, sections of the V-Fe-Si, V-Co-Si, and V-Ni-Si systems were studied at silicon contents between 0 and 40 at. pct. The single-phase fields in the 30 and 50 at. pct Mn sections of the 1100°C tetrahedron of the V-'Wn-Fe-Si system were also determined. The direction of extension of the elongated O, R, D, and a-Mn Phase fields suggests that the stability of these phases is determined largely by the electron concentration. The peculiar shape of the above ternary a phase fields, as well as of the a phase fields in the V-;Mn-Si and Cr-MTZ-Si systems is accounted for on the assumption that in manganese-(iron-, coball-, nickel-) rich alloys silicon atoms do not exclude d electrons , while in alloys low in the above components they do. This assumption also leads to the expected electronconcentrations for q and ß-Mn phases formed by transition metals with aluminurn. In earlier work it was shown that in ternary alloy systems of transition elements the a a and certain other phases with complex structure, such as the p and R phases,3 form elongated phase fields extending in a direction corresponding to a constant electron concentration. Recently similar behavior was observed in manganese-rich ternary ß-Mn solid solutions with another transition element and with silicon, germanium, or tin,4 confirming earlier indications3 that ß-Mn is also an electron phase. The alloying behavior of silicon in these manganese-rich alloys sug- gested that the four outside electrons of silicon participate in the bonding and that the predominantly d-like bonding electrons are not excluded from the silicon sites. Presumably, some of the normally unoccupied 3d states associated with the silicon atoms are in these alloys lowered to the Fermi level;4 thus silicon behaves as if it, too, were a transition element. The manganese-rich boundary of the a phase in the V-Mn-Si system,4 as well as in the Cr-Mn-Si sys-tem,4,5 extends approximately parallel to the ß-Mn phase boundary, corresponding to the transition ele-mentlike behavior of silicon described above. However, at lower manganese contents (i.e., in the high-vanadium or -chromium region) the a phase field in both ternary systems mentioned extends in a direction roughly perpendicular to that of its manganese-rich boundary, a phenomenon that is apparently inconsistent with the behavior of silicon in the manganese-rich alloys. A similar, although somewhat less pronounced, extension of the a phase occurs also in the high-vanadium (low-iron) region of the V-Fe-Si system at 1175C, but reportedly not in the V-Co-Si or V-Ni-Si systems at the same temperature.5 The present work was undertaken to investigate in more detail the extension of the a phase (as well as of other electron phases that may be present) in the last three systems and to elucidate this phenomenon, which has not been satisfactorily explained previously. EXPERIMENTAL PROCEDURES Most of the alloys were melted in an arc furnace, using a water-cooled copper crucible under helium atmosphere. The transition metals were of better than 99.9 pct purity and the silicon used was 99.98 pct pure. Alloys containing less than 20 pct V were induction-melted in recrystallized alumina crucibles under argon atmosphere.

Jan 1, 1967

By S. C. Sircar, D. R. Wiles

The rate has been studied of the hydrogen reduction of nickel ions in acetate -buffered solutions, using a nickel catalyst. At temperatures between 130°and 160°C, the rate is found to be proportional to the catalyst area, the nickel ion concentration, and the hydrogen pressure. The effect of hydrogen ions is to reduce the observed rate, likely by causing the reverse reaction (dissolution) to proceed. The experimental activation energy was found to be 25 kcal per mole. THE precipitation of metals and lower oxides from solution by hydrogen reduction has been quite actively investigated during the past few years.1 In the case of nickel, the kinetics of the precipitation have been studied in ammonia solution2 but extension of this study to acid solutions has not attracted much interest, largely because of the unfavorable thermodynamics at all but the lowest acidities. For example, Schaufelberger foundg that only 10 pct reduction of nickel could be effected in sulphate solutions. One kinetic study of the reduction of Ni++ has been reported by Knacke, Pawlek, and Sussmuth,4 in which the order of the reaction was found to be one in nickel ion and one-half in hydrogen pressure. No mention was made as to the part played by the nickel catalyst, or by the pH of the solution. Moreover, the activation energy reported (4 .12 kcal per mole) is so low as to lead to some question as to whether diffusion might be playing a major role in determining the kinetics. We have recently completed a study5 of the kinetics of the hydrogen precipitation of nickel from acetate-buffered acidic solutions whose results confirm and amplify those of Knacke et al. EXPERIMENTAL PROCEDURES The work was done in a high-pressure autoclavee by the usual technique of removing small samples of the solution for analysis. Nickel analysis was done colorimetrically by the dimethylglyoxime method.' The catalyst used was carbonyl nickel powder (Standard A powder of the Mond Nickel Co., Ltd.) which was found to be composed of spheres of less than 350 mesh. The surface area of the powder was determined in this laboratory by the B. E. T. method to be 1.85 sq m (m2) per g. The catalyst was added to the solution only after the proper temperature and pressure had been reached. Agitation of the baffled system was by a turbine-type impeller, rotated at speeds between 400 and 750 rpm. The reduction rate was found to be independent of the stirring rate, so it was assumed that diffusion in the liquid does not limit the rate. In typical experiments, the nickel concentration was unchanged during an induction period of variable length. The concentration then decreased sharply, and gradually levelled off to approach the final equilibrium value. In some cases, notably at high stirring rates, the nickel concentration was found to increase considerably during the induction period, and then to decrease. In all cases, the rate was determined from the slope of the tangent to the rate curve as the concentration first decreased below the initial value. This is seen in Fig. 1. The rate measurements were found to be reproducible to within about 5 and to be quite independent of the nature or duration of the induction period. The range of conditions used is given in Table I.

Jan 1, 1961

By Lars Rossemyr, Terkel Rosenqvist

The equilibrium Mn + 2 IiCl = MnCl2(g) + H2 has been studied at 1090°C for pure manganese and for Mn-Si alloys. For this reoction a standard free energy of Fo1363, = - 19,700 i 300 col ioas found. Combined with other thermal data, the standard free energy of the reaction Mn(s, ß) + Cl2 = MnCl2(g) was derived: Fo=- 56,140 + 9.78 T log T - 40.10 T For the formation of the various manganese silicides the following free energies were obtained. 5 Mn + 3 Si - M1z5Si3 ?F°1363 = -41,500 i 1500 cal Mn + Si - h.lnSi ?F°1363= -10,000 i 500 cnl Mn + 2 Si = MnSiz ?F°1363= -11,500 ± 500 cal In the search for a method suitable for the study of thermodynamic properties of certain alloy systems, it was found that the heterogeneous equilibrium between the alloy and a gas mixture consisting of hydrogen chloride, hydrogen, and volatile metal chlorides has considerable advantages. In addition to giving thermodynamic information about the alloys, such measurements may be used to derive thermodynamic data for the volatile metal chloride itself. The present paper gives some results which were obtained for manganese silicides and manganous chloride. Similar studies are presently being carried out for the iron silicides. The phase diagram of the Mn-Si system is rather well known,' but no information exists to the authors' knowledge about the thermodynamic properties of manganese silicides or about the chemical activities in molten Mn-Si alloys. At elevated temperature, manganese may react with hydrogen chloride according to the equation: Mn(s, ß ) + 2 HC1 = MnC1, (g ) + H2 [ 1] with the equilibrium constant The value of K, may be determined from measurements of the gas composition in equilibrium with pure manganese. When this value is used in connection with measurements on manganese alloys, the chemical activity of manganese in the alloy may be derived. Initial calculations showed that around 1000" to 1200°C the equilibrium constant K1 has a value of a few thousand, i.e., the reaction is shifted far to the right. In order to study this rather unbalanced gas composition, a flow method was used in two slightly different modifications. For all experiments with manganese silicides and for some experiments with pure manganese the apparatus shown on Fig. 1 was used. A bed of the metal or alloy was placed in a tube furnace. The refractory tube had an internal diameter of 14 mm and the bed was 75 mm long. Electrolytic manganese of high purity and silicon, 99.95 pct Si, were used as raw materials. All alloys were prepared by melting weighed amounts of the elements in a high frequency induction furnace in argon atmosphere. The melting was done in a silica crucible which was placed inside a graphite shell. The alloys were crushed to - 10 +40 mesh and analyzed chemically. Also, for the pure elements, the same size fraction was used. A controlled gas mixture of purified hydrogen and from 3 to 12 pct HC1 was passed through the bed where it reacted with manganese according to Eq. [I]. The volatile manganese chloride was collected in a condenser placed in the cooler part of the furnace tube. The gas mixture before and after the furnace was sampled and analyzed by absorption of

Jan 1, 1962

By R. DiBella, H. Biloni, G. F. Bolling

A study of the detailed relationship between solidification substructure and microsegregation in dilute tin alloys has been continued. New observations reveal that depressions form at the solid-liquid interface not only as an initial form of segregation but throughout the range of increasing constitutional supercooling. In order to show the importance of this node formation it proves convenient to make two simplifications. Although substructure changes continuously as constitutional- supercooling increases,&apos; the following discrete stages may be used for discussion: 1) the planar solid-liquid interface; 2) ordered nodes, or depressions at the interface;&apos; 3) two-dimensional or elongated cells;3 4) regular or hexagonal cells;4 5) distorted cells, or hexagonal cells with branches;4 6) dendrites. It is also convenient to ignore observations made near grain boundaries which interact with the cellular substructure.2-4 And, it should be clearly pointed out that all singular depressions at the solid-liquid interface will be called nodes. Heretofore, the term "node" has been used without clear justification to describe both the truly isolated depression and the junction of cell walls appearing in a later stage of segregation. This usage is continued because it seems proved by the observations. Standard systematic growth of tin crystals was undertaken over a broad range of constitutional supercooling using various added solutes. Except for one or two chance observations, this research failed to reveal any appreciable range of growth conditions over which a disordered array of isolated nodes exists. Even though &apos;some finite range of random instability must exist, it is now concluded that stage one passes very rapidly to stage two. Next, after confirming that stage two passes to stage three by the elongation of isolated nodes: a reexamina-tion of stages three, four, and five followed. Standard metallography showed that pockets of higher solute concentration (for ko < 1) became more numerous in cell walls the more severe the growth conditions. These high-concentration regions must arise from depressions; however, in stage three their segregation is not distinguishable from the segregation of the nodes which join cell walls together. It seems, therefore, that the presence of nodes is causal in the formation of new cell walls, and that the mechanism by which stage three passes to stage four is initiated by nodes. If a similar reasoning is followed, the separation of stage four from stage five appears to be just a convenience that has arisen from the decanted interface observations made in the past. The present observations again show that more severe growth conditions relate directly to an increase in the number of nodes in the hexagonal cell walls. These nodes apparently serve as the origin for lines of segregation extending within the body of a cell, and can thus be assumed to represent the earliest form of branching,* as shown here in Fig. 1. The examination of stage six involved decanted-interface structures as in Fig. 2, which are similar to those considered by Rutter and Chalmers,5 Plaskett and winegard,6 and Hellawell and Herbert,7 and where distinct protrusions are each surrounded by a depressed boundary. The corresponding microsegregation behind the interface, but not far enough behind to allow a smearing of the segregation because of diffusion,4 is shown in Fig. 3. Both the shape of the decanted interface and the segregation pattern behind it

Jan 1, 1968

By R. Simon, L. Yarborough

This paper presents a correlation for predicting the critical pressures of gas-solvent-reservoir oil systems of the type encountered in gas and solvent injection processes. Experimental data from 14 systems were used to develop the correlation which expresses critical pressure as a function of composition, C7+ mol weight, and C7+ characterization factor. Critical pressures predicted by the correlation for the 14 gas-solvent-reservoir oil systems have an average deviation of 2.8 per cent of the experimental values. The correlation war used to study a wide variety of other hydrocarbon mixtures reported in the literature whose critical conditions are in the range of interest in oil recovery work (100 to 280F and 2,000 to 6,000 psia). Critical pressures predicted by the correlation had an average deviation, expressed as per cent of experimental value, of 5 per cent for 20 miscellaneous mixtures from the literature, and 6.4 per cent for I7 methane-butane-decane mixtures reported by Sage and Lucey. INTRODUCTION Predicting critical points of hydrocarbon mixtures is an important part of phase behavior calculations. If the critical point can be predicted for a specific mixture, the separation between the bubble-point and dew-point regions will be defined, and physical properties of the mixture can be calculated by using the principle of corresponding states. Numerous correlations for predicting critical points of hydrocarbon mixtures have been published. Six of these correlationsl-8 were for light hydrocarbons, and two correlations7,8 ere for high boiling-point petroleum distillates. Organick9 And Eilerts" presented correlations which combined the published data from both light and heavy mixtures, and these correlations became the only available methods for predicting critical pressures for wide boiling range mixtures. In the course of the authors&apos; work on secondary recovery by gas and solvent injection, experimental data were obtained for 14 mixtures of dry gas, ethane-propane-butane solvents, and various reservoir oils. The critical pressures and critical temperatures of these gas-solvent-reservoir oil systems were in the range of 1,900 to 4,800 psia and 180 to 265F. The critical pressures predicted for the 14 mixtures by the correlations of Organick and Eilerts had an average deviation from the measured values of about 13 per cent. The authors felt that it was necessary to develop a better correlation for engineering calculations and the resultant correlation, which is presented on Figs. 1 and 2, predicts the critical pressures of the 14 mixtures with an average deviation of 2.8 per cent. The validity of the correlation was tested by applying the correlation to other hydrocarbon mixtures reported in the literature. This subject is discussed later in the paper. The authors feel that there is also a need for an improved critical temperature correlation for gas-solvent-reservoir oil systems. However, developing a critical temperature correlation is a significant study in itself and is not included in this paper. EXPERIMENTAL DATA The compositions and critical points obtained by the authors for 14 gas-solvent-reservoir oil mixtures are shown on Table 1. The methane concentrations in these mixtures varied from 20 to 63 mol per cent. The intermediate hydrocarbon group, consisting of ethane, propane, and butane, varied from 20 to 64 mol per cent, and the concentration of the C,, fraction from 10 to 20 mol per cent. The physical properties of the C,, fraction covered the range from light, paraffinic, nonasphaltic (Mixtures 5, 6, 7, 8, 9 and 10) to heavy, aromatic, and asphaltic (Mixtures 13 and 14).

By W. L. Kerlie, J. H. Richards

The variation in hydrogen content of basic open-hearth steels during refining and the effect of changing from steam to air atomization were studied. The water content of the furnace atmosphere, the slag basicity, and the degree of oxidation of the slag-metal system largely control the hydrogen in the steel. SMALL amounts of hydrogen have an adverse effect on the ductility of steel and a wide range of hydrogen contents has an adverse effect on the internal soundness of both ferrous and nonferrous castings. This has been of major concern to the metal industry, both manufacturers and fabricators, for many years. As one of the results of continuing investigations of the hydrogen problem, much has been learned about methods of reducing the hydrogen content of liquid metals below the relatively high levels that are normally associated with such phenomena as rising ingots and excessive porosity in sand casting. However, it was not until a reliable and simple method for sampling and analyzing liquid metal was developed that much progress was possible in quantitatively evaluating the effect of steelmaking practice on hydrogen contents at levels below 6 to 8 ppm. Throughout all the work described in the following presentation the Geffner pintube technique was used to obtain samples of liquid metal1 * and the tin-fusion method, developed at Massachusetts Institute of Technology,2 was used to analyze the samples. Before attempting a systematic experimental program for minimizing the hydrogen content of open-hearth steels, it was necessary to determine the variation in hydrogen in the open-hearth bath during the normal processing of a heat. The logs of ten heats are illustrated in Figs. 1 through 7. The number at each point on the graphs refers to the number of pintube samples analyzed for hydrogen. An increase in the hydrogen content of the liquid metal during the early period of refining is apparent in all heats. There is evidence of a decrease in hydrogen later in the refining period, see Fig. 2, for example. Therefore, the existence of a maximum hydrogen content at some intermediate time during the refining period is anticipated. To test for the existence of a maximum, the data illustrated in Figs. 1 to 5 were grouped together so that they could be examined on a collectively comparable basis. The grouping was effected, as illustrated in Fig. 8, by superimposing Figs. 1 to 5 so that the apparent hydrogen maxima coincided at a fixed position, indicated by the zero on the time axis. Also, by vertical adjustment until the graphs appeared to the eye to be superimposed, the hydrogen scale was made to show the change in hydrogen content. That is, a standard amount was subtracted from the hydrogen values for each heat. Therefore, Fig. 8 represents a generalized presentation of the data and, for this reason, cannot be used to predict the actual hydrogen content of the bath during refining, but rather to show the probable rate of change in hydrogen during refining. Statistical examination of the grouped data in Fi.g. 8 shows that the apparent. increase and suhsequent decrease in hydrogen are real and confirms the existence of a maximum in general. The average increase In hydrogen during the early period of refining is between 0.8 and 1.3 ppm per hr, and the

Jan 1, 1958