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By W. A. Stickney, J. W. Town

Tests demonstrated that aqueous solutions of sodiu~n sulfide would dissolve over 95 pct of the cinnabal- in 5 pct Hg flrotation concentrates and 60 to 90 pct of the cinnabav in low-grade ol-es. Double-leach tests showed that potassiutn sulfide would dissolve 90 pct of the avsenic sulfide in Hg-As satnples and only 1 to 10 pct of the Hg; the sodium sulfide would then dissolve 80 to 90 pct of the rernaining Hg and 30 to 70 pct of the vemaining As. 1 HIS paper describes the extraction, by alkaline sulfide solutions, of mercury from cinnabar-bear ing ores and flotation concentrates. Leaching time and solution temperature had minor effects on mercury extraction when compared to the effects of reagent quantity and ore type and grade. Cinnabar, although readily soluble in sodium sulfide, was only partially soluble in potassium sulfide and essentially insoluble in ammonium sulfide. A double leach process was developed to extract selectively the arsenic from Hg-As products with K2S and to treat the residue with sodium sulfide-hydroxide solutions to extract the mercury.' Mercury was then recovered from the pregnant solutions either by precipitation with metallic aluminum or by electrodeposition. The solubility of HgS in alkaline sulfide solutions was first reported by Volhard in 1878.' The first reported commercial application was made in 1915 when Thornhill applied the method to the treatment of an amalgamation tailing.3 Mullhollen, in 1911. described the chemical reaction of HgS with Na,S and alkaline hydrate to be:4 Most investigators have reported using a 4-pct sodium sulfide and a 1-pct sodium hydroxide soluti~n.3-4 An increase in temperature or excessive dilution of the solution causes the reaction to be reversed. Church Holmes described a similar process used for the extraction of antimony from silver concentrates at the Sunshine Mining Company's mill at Kellogg. Idaho.5 More recently. Erspamer and Wells reported the results of leaching studies on a Hg-Sb ore from Alaska.6 EXPERIMENTAL PROCEDURE Leaching tests were made on samples of mercury ore from eight deposits in Oregon, Washington, and Idaho. In addition. tests were made on flotation concentrates from seven of the samples as well as a flotation concentrate from an operating mill. Since previous investigators reported reagent requirements greatly in excess of that required by the equation HgS + Na2S = HgS . Na2S, initial studies were made on chemically pure HgS. Results showed that the reaction actually required about 1 g of Na,S per g of HgS in an aqueous solution rather than 0.338 g of Na,S as indicated by the equation. To check further the matter of reagent requirements, Na2S and HgS were mixed together in equal mole ratios. Water was then added, a drop at a time, until the components reacted; X-ray diffraction analyses showed no HgS nor Na,S present. The yellow hydrate compound which formed would not redissolve until about three additional mole equivalents of Na,S were added to the solution. This indicated that a stable solution would require at least 4 moles of Na2S for each mole of HgS. While these implications may be valid for pure HgS, the possibility of entirely different results caused by the existence of other elements present in ores and concentrates was recognized. Consequently, experiments of the same order were made on samples of a flotation concentrate from a producing mill. The concentrate, containing about 18 pct

Jan 1, 1962

By C. S. Barrett

The torsional after-effect in polycrystalline cadmium is interrupted by an abnormal twisting when the film is removed by etching. This is accounted for by the pile-up of dislocations beneath anodic or other oxide films during the original twisting, and the relief of the residual stresses from these by elastic and plastic processes during etching. WHEN an iron or zinc wire coated with oxide is plastically twisted and then released, it untwists according to a law of equal increments in equal intervals of log time, but if the oxide is suddenly removed by etching this normal after-effect is interrupted by a transient, and thereafter the untwisting is slower.' Etching may even reverse the strain temporarily. Related effects caused by hydrated oxide films on cadmium in creep tests have been observed by Phillips and Thompson:' definite transient strains caused by the removal of the film from single crystals of cadmium were observed; however, negligible effects were obtained with polycrystalline wires. Pronounced effects of both transient and steady nature were observed by Phillips and Thompson when anodic films 10-4 cm thick were etched off, and weak effects, transient only, occurred with hydrated oxide films 3x10 cm thick such as were formed by standing in water for 10 min. It was estimated that 10.' cm films, formed in water in 1 or 2 min, could be detected in the apparatus of Phillips and Thompson. Since abnormal after-effects were observed by the torsion method equally well with polycrystalline and monocrystalline iron and zinc wires,' it was anticipated that the torsion method would also disclose them in polycrystalline cadmium, in spite of the fact that analogous effects had not been observed in tensile creep tests on polycrystalline zinc by Pickus and Parker3 r on polycrystalline cadmium in tensile studies by Andrade and Randall." Because of its high sensitivity, particularly its sensitivity to surface conditions, the torsion method was thought to merit further study as a method of investigating thin films of various types and the action of various reagents on these films. Materials and Methods Cadmium wires were cold-drawn to a diameter of 0.039 in. from a stock that was determined spectro-scopically to contain 0.05 pct Bi, 0.005 pct Cu, 0.008 pct Pb, and 0.007 pct Si. The wires were cut to about 17/8 in. lengths and were annealed 1 hr in air at 290°C, after which the grain size was such that there were 5 to 7 grains across a diameter. The wires were then mounted as indicated in the insert in Fig. 1, one end being fastened with laboratory wax into a small diameter brass tube and the other end being waxed to a thin glass fiber which supported, on its upper end, a small galvanometer mirror. The wires were waxed with care to avoid straining or annealing, for it was found that the after-effect rate was sensitive to prior cold work in a wire. The assembly was then clamped above a beaker into which the reagent could be poured. The lower end of the wire was held in the hand while the wire was twisted 180"; it was then released immediately. Since precision timing of these manipulations was not attempted, the different wires cannot be compared as to exact rates of untwisting. The length subjected to plastic strain was 15/8 in. The surface coatings studied were those produced by the annealing in air, by anodic treatment of

Jan 1, 1954

By C. P. Besse, N. F. LeBlanc

Described in this paper are oceanographic data which should be considered by an offshore design engineer and methods for developing a design wave height from the oceanographic data. The selection of a design wave is predicated on contemplated waves which might affect the site throughout the life of the structure. Selection of a design wave height may be based on (1) arbitrarily established recurrence Frequencies of hurricanes affecting the structure (predicted wave heights are associated with the expected variations of forces resulting from these waves) and (2) a risk-type evaluation wherein all possible storms affecting the area are considered (anticipated wave heights are associated with both investment plus risk costs and expected variations of forces). It is shown how the following oceanographic predictions are integrated into design considerations: (1) a classification of storm intensity which considers all recorded storms which affected the design area, (2) the recurrence interval of storms of a given- intensity (this interval is dependent on the extent of the geographical area considered in the design problem) and (3) a forecast of all wave heights which might affect the area (geometry of a structure often necessitates consideration of waves from a multiplicity of directions). The authors believe that the described techniques can result in selecting an adequate and reasonable design wave. INTRODUCTION Since the inception of offshore operations in the Gulf of Mexico, engineers engaged in designing suuctural facilities have been plagued with the problem of selecting an adequate and reasonable design wave. In the development of any offshore structure it is mandatory that the engineer evaluate the ability of the structure to withstand the ocean waves to which it will be subjected. Selecting such design waves quite naturally necessitates a coalition of the oceanographer and the design engineer. The oceanographer must provide a detailed knowledge of scientific principles which govern the behavior of waters in the Gulf of Mexico. He should also have an adequate knowledge of the manner in which design waves are utilized by the engineer. Although the design engineer's primary responsibility is applying the oceanographer's specialized knowledge in the creation of real structures, it is important that he possess some knowledge of related oceanographic principles to reasonably evaluate and apply the recommendations of the oceanographer. In the Gulf of Mexico it is the hurricane wind waves which generally govern the design of an offshore facility. The oceanographer must therefore develop techniques for predicting the heights, periods and frequency of all hurricane waves which might affect a particular structure. From this mass of oceanographic data, the design engineer must select the design waves which will apply to his particular design. PAST STUDIES ON FREQUENCY AND AMPLITUDE OF HURRICANE WIND WAVES Past oceanographic studies on the frequency of hurricane wave heights in the Gulf of Mexico have been devoted largely to predicting the recurrence interval of hurricanes which will generate maximum significant waves of given heights. The maximum significant wave height is the average height of the highest one third of the waves in that portion of the storm producing maximum wave heights. Since these waves occur over a relatively small portion of the storm (Fig. 1) and since the paths of hurricanes vary considerably (Fig. 2), the recurrence frequency of such heights is largely a function of the extent of the geographical area considered. In establishing the frequency of maximum significant waves for the entire American Gulf coast, the major fully developed hurricanes which affected the coastline from BrOwnsville, Tex., to Key West, Fla., were considered (Fig. 2). These Storms were selected and described by C. L-

By R. N. Thurston, E. A. Nesbitt, G. Y. Chin

A general relalionship between the amount of glide shear (due to slip) and the macroscopic shape change has been developed. Since the deformation can be large, finite strain analysis is employed. In this treatnzent, the shape change is expressed by a deformation gradient matrix, F - [axi/aXj1, where X and x refer to the initial and final positions of a particle. This matrix is readily evaluated in terms of the amounl of glide shear, a , and the direction cosines of the slip-plane normal and the slip direction for a single active slip system. In the case of deformation from slip on several systems, the product F(a) of the several deformation gradient matrices is first calculated. Then, by assuming that the final configuration is reached by a long series of small shears of magnitude a, occurring more or less alternately in the several slip sys- IN problems of crystal plasticity, it is often necessary to relate the amount of glide in the operating slip systems to the macroscopic strain components. This relationship is relatively simple when only small strains are considered. In this case the separate strain contributions from several slip systems are additive. Such a procedure is incorrect in the case of large plastic deformation. To the authors' knowledge, no general treatment of the latter problem has appeared in the literature. In specific instances, Mark, Polanyi, and schmid3 have derived the relationship between glide and axial elongation during tensile pulling of a single crystal when only one slip system is active. A similar relationship for duplex slip was worked out by v. Goler and Sachs.4 Taylor and Elam5 have also studied in detail problems of large plastic deformation. Due to uncertainties as to the slip systems at the time, terns, the final deformation gradient matrix can be obtained. Mathematically, the problem is reduced to finding the limit of F(dN as N — m while a — 0 in such a way that the Product Na = 0, a ftnite conslant designatzng the accunzulated amount of shear. It turns out that this limit is simply , where F, is the matrix whose elements are the coefficients of a in the F(a) matix. Application is made of the present treatment to a fcc crystal of Perlrzalloy compressed on the (110) plane and contrained to elongate zn the [112] directzon. In addition, it is shown that one rnay readily obtain from the general analyszs the well-known formulas relating elongation, amount of glide shear, and amount of lattice rotation for crystals deforming by single and double slip under tension. however, they were mainly concerned with proving that slip occurs on {111}(110) systems in fcc metals. The present general treatment arose out of recent investigations of magnetic anisotropy in cold-worked Fe-Ni alloys6 as well as some related problems in strength anisotropy of single crystals.' Detailed application of the general analysis will be provided in the case of a Permalloy single crystal compressed on the (110) plane and constrained to elongate only in the [112] direction. It will be seen that, as expected, the small strain approximation leads to significant errors after moderate straining. It will also be shown that the present general treatment yields the formulas of Mark et al. and of v. Goler and Sachs in those specific cases. GENERAL CONSIDERATIONS Consider a homogeneous deformation in which a material point initially at moves to both positions being referred to the same set of Cartesian axes. Such a deformation can be

Jan 1, 1967

By E. P. Abrahamson, S. L. Lopata

A general relalionship between the amount of glide shear (due to slip) and the macroscopic shape change has been developed. Since the deformation can be large, finite strain analysis is employed. In this treatnzent, the shape change is expressed by a deforlnation gradient matrix, F - [axi/aXj/ where X and x refer lo the initial and final positions of a particle. This matrix is readily evaluated in termas of the amounl of glide shear, a , and the direction cosines of the slip-plane normal and the slip direction for a single active slip systena. In the case of deforma tion from slip on several systems, the product F(a) of the several deformation gradient matrices is first calculated. Then, by assuming that the final configuration is reached by a long series of small shears of magnitude a, occurring more or less alternately in the several slip sys- IN problems of crystal plasticity, it is often necessary to relate the amount of glide in the operating slip systems to the macroscopic strain components. This relationship is relatively simple when only small strains are considered. In this case the separate strain contributions from several slip systems are additive. Such a procedure is incorrect in the case of large plastic deformation. To the authors' knowledge, no general treatment of the latter problem has appeared in the literature. In specific instances, Mark, Polanyi, and schmid3 have derived the relationship between glide and axial elongation during tensile pulling of a single crystal when only one slip system is active. A similar relationship for duplex slip was worked out by v. Göler and Sachs.4 Taylor and Elam1 have also studied in detail problems of large plastic deformation. Due to uncertainties as to the slip systems at the time, ! terns, the final deformation gradient matrix can be obtained. Mathematically, the problem is reduced to finding the limit of F(aN as N — m while a — 0 in such a way that the Product Na = a finite conslant designatzng the accumulated amount of shear. It turns out that this limit is simply eafF1, where F, is the matrix whose elements are the coefficients of a in the F(a) twatrzx. Application is made of the present treatment to a fcc crystal of Permalloy compressed on the (110) plane and contrained to elongate in the [112] directzon. In addition, it is shown that one rnay readily obtain from the general analyszs the well-known formulas relating elongation, amount of glide shear, and amount of lattice rotation for crystals deforming by single and double slip under tension. however, they were mainly concerned with proving that slip occurs on {111}(110) systems in fcc metals. The present general treatment arose out of recent investigations of magnetic anisotropy in cold-worked Fe-Ni alloys6 as well as some related problems in strength anisotropy of single crystals.' Detailed application of the general analysis will be provided in the case of a Permalloy single crystal compressed on the (110) plane and constrained to elongate only in the [112] direction. It will be seen that, as expected, the small strain approximation leads to significant errors after moderate straining. It will also be shown that the present general treatment yields the formulas of Mark et al. and of v. Goler and Sachs in those specific cases. GENERAL CONSIDERATIONS Consider a homogeneous deformation in which a material point initially at moves to (xl,xz,x3), both positions being referred to the same set of Cartesian axes. Such a deformation can be

Jan 1, 1967

By David B. Reger

The value of carbon ratios in determining the boundaries of possible oil deposits appears to have passed the hypothetical stage. The theory that the ratio of fixed carbon in pure coals is an invariable index of incipient metamorphism in both surface and underground rocks and that it may be applied in defining the limits of petroleum, advanced by David White,' has been received with keen interest by many petroleum geologists. Detailed isocarb maps have been prepared of the Pennsylvanian area of North Texas and Eastern Oklahoma by M. L. Fuller.² A similar map of the coal-bearing area of West Virginia is given here. An isocarb3 is a line showing an equal fixed-carbon percentage, pure coal basis; the term has been proposed by David White to supersede a less expressive nomenclature. The term carbon ratio is applied to the pereentage of carbon in pure coal after water and ash have been eliminated. As a comparatively small number of analyses have been made on this basis, it is usually necessary to compute the ratio by dividing the fixed carbon of the proximate analysis by the sum of the fixed carbon and volatile matter of the same analysis. Many thousands of proximate analyses have been made by the West Virginia Geological Survey, covering nearly every county in which coal is found. Numerous others have been made by the U. S. Geological Survey and the U. S. Bureau of Mines, but as uniformity of results is best secured by adhering to one set of analyses, the tests of the West Virginia Survey have been cxclusively used.

Jan 1, 1921

By H. T. Green, R. M. Brick

True stress-strain data on alloys of pure iron with up to 2.4 pct Al were obtained in the temperature range +100° to —185°C. Alumi-num was found to reduce yield and flow stresses of iron at low temperatures but to have little or no effect on ductility. The effects of temperature and composition on strain hardening are discussed. SEVERAL independent studies of the behavior of high purity iron binary alloys at low temperatures are now in progress in attempts to evaluate systematically the variables affecting the low temperature brittleness of ferritic steels. This paper reports the results of one such investigation in which the tensile properties of aluminum and aluminum plus silicon ferrites were measured from 100" to —192°C. True stress-natural strain data have been obtained in order to evaluate as many as possible of the parameters which describe the behavior of the materials involved. In comparable studies at the National Physical Laboratory in England, iron and iron alloys of high purity have been produced' and tested at subat-mospheric temperatures.' True stress-natural strain curves were obtained there also. The purest iron contained 0.0025 pct C and 0.001 pct O and N. Even this, as normalized at 950°C following hot rolling, showed little ductility at -196°C. The grain size was ASTM No. 3, and the room-temperature yield strength was 17,800 psi (which seems too high for pure iron). Some of the NPL irons contained considerably more oxygen and demonstrated intergran-ular fracture at —196°C. The authors2 carefully differentiated between intergranular fractures associated with excessive oxygen content and transcrys-talline cleavage with little ductility encountered at —196°C in the purer material. The cleavage stress was half again as great as that associated with inter-granular fracture. Test Material, Preparation, and Procedures Of a number of Fe-A1 alloys produced, eight were considered to be sufficiently pure for testing. Partial chemical analyses (Table I), low observed yield points, and high ductilities indicate these alloys to be comparatively pure for vacuum-melted irons of sizable ingots, 5 Ib or more. To produce the binary Fe-A1 alloys, electrolytic iron was melted in air, cast into slabs, and rolled to strips 0.010 in. thick. These strips, joined into a continuous ribbon and wound into 2 1/2 in. diameter spools, were subjected for four weeks to a moving atmosphere of purified dry hydrogen in a stainless-steel tube at 1050" to 1150°C. Charges of these spools were melted in beryllia crucibles under good vacuums (1 micron), and aluminum (99.97 pct Al) was added to the melts. Compositions of these alloys are recorded in Table I. The ingots were hot forged and then cold rolled at least 65 pct to 3/8 in. rods which were vacuum annealed to the desired grain size, approximately ASTM No. 4, prior to machining into tensile test bars. All tensile specimens had gage sections 1 in. long, with a fillet of 1.5 in. radius to the shoulder. Gage diameters were 0.250 in, except for a few rods where additional cold work required use of a 0.200 in. gage section. After machining, 0.002 in. was removed from the gage diameter using 240, 400, and 600-grit metallo-graphic papers. The final polish with 600 grit left the fine scratches running in the longitudinal direction. By this means, surface metal strained during machining was removed. A few specimens heat treated after machining were similarly reduced 0.004 in. to remove any material affected chemically by the atmosphere during heat treatments, as is discussed in a later section. Tensile tests of the eight alloys at constant temperatures from +100° to —185°C were performed in apparatus which has been described." The essentials include a double-walled insulated metal vessel which contained the liquid heat-transfer medium surrounding the test specimen. A constant temperature was maintained by means of a pyrometer which regulated the pressure of dry air driving liquid air through a copper coil. Temperature variation was less than ±2°C during a specific test. For axial straining, two lengths of case-hardened chain, terminating in simple shackles, loaded the specimen through threaded grips. The lower grip bar passed through a hole in the bottom of the test vessel to which it was joined by a thin-walled

Jan 1, 1955

By T. S. Fuller

The work of D. J. McAdam, Jr.1,2 at the U. S. Naval Engineering Experiment Station, Annapolis, Md., on what has been called by him "corrosion-fatigue" has focussed the attention of the engineering profession in general, and of those interested especially in the study of corrosion in particular, on the importance of making tests which will indicate as nearly as possible what can be expected from a given material in service. Pitting is unquestionably the chief agent operating so tremendously to reduce fatigue limits, but in the light of previous work3, 4, 5, 6, 7, 8 it is fair to assume that absorbed hydrogen resulting from the action of the corroding medium upon steel may, under certain conditions, have a perceptible influence. The experiments described in this paper deal more precisely with stressless corrosion, or "prior-corrosion fatigue,"9 and have been carried out to test such an assumption, to segregate, if possible, the effect of pitting from that of absorbed hydrogen, and although very preliminary in nature they indicate results which are striking.

Jan 1, 1929

By T. S. Fuller

The work of D. J. McAdam, Jr.1,2 at the U. S. Naval Engineering Experiment Station, Annapolis, Md., on what has been called by him "corrosion-fatigue" has focussed the attention of the engineering profession in general, and of those interested especially in the study of corrosion in particular, on the importance of making tests which will indicate as nearly as possible what can be expected from a given material in service. Pitting is unquestionably the chief agent operating so tremendously to reduce fatigue limits, but in the light of previous work3, 4, 5, 6, 7, 8 it is fair to assume that absorbed hydrogen resulting from the action of the corroding medium upon steel may, under certain conditions, have a perceptible influence. The experiments described in this paper deal more precisely with stressless corrosion, or "prior-corrosion fatigue,"9 and have been carried out to test such an assumption, to segregate, if possible, the effect of pitting from that of absorbed hydrogen, and although very preliminary in nature they indicate results which are striking.

Jan 1, 1929

By T. A. O’Hara

SINCE May 1948, when tungsten carbide bits were introduced at the Flin Flon mine, they have been popular with the miners because of their fast drilling speed and low gage loss. The high cost of commercial carbide bits and tipped drill steel, however, prevented their use except for the hardest rock. In an effort to extend the use of tungsten carbide on a basis economically competitive with detachable steel bits, experimental work was begun in 1950 to test the feasibility of making tungsten carbide tipped drill steel in the mine drill steel shop. This work showed that tipped drill steel could be made locally at less than half the cost of the commercial product. The performance of the local tipped drill steel was comparable to that obtained with commercial carbide bits and tipped drill steel and the cost per foot drilled was much lower. Local tipped drill steel was adopted for all mine drilling in November 1951. Since then drilling costs per foot have been sharply reduced and footage drilled per manshift has increased markedly. Experience at Flin Flon has shown that production of satisfactory carbide tipped drill steel is not difficult and that highly skilled labor and costly equipment are not required. As long as wise selection of brazing materials is made and certain simple precautions are rigidly maintained, there is no reason why small mines with relatively unskilled labor cannot produce a satisfactory product. The following description outlines the technique used at Flin Flon for making carbide tipped drill steel and discusses characteristics of the brazing process that make special precautions necessary. Drill steel is forged to four-wing shape in a conventional steel sharpening forge. Standard steel dies are modified to minimize forging cracks around the central waterhole and to forge a blunt bithead on the steel. The steel is preheated to 1500°F and held at this temperature for at least 2 min. When the temperature has equalized throughout the steel section, the drill steel is transferred to the forging furnace and heated rapidly with a reducing flame up to 2000°F. This two-stage method of heating minimizes the grain growth and decarburization of the steel while ensuring that the steel temperature does not vary greatly throughout the forging zone. After forging the steel is allowed to cool in air to about 1600°F before being annealed in a bath of vermiculite. Despite the high hardenability of the 3 pct Ni-Cr-Mo drill steel used, this simple treatment anneals the drill steel sufficiently for milling. The forged and annealed drill steel is slotted on a plain horizontal milling machine that is equipped with a quick opening chuck and a slot depth stop. The full depth of the slot is milled in a single pass of the 3-in. milling cutter which is fed at 33/4 in. per min across the crown of each bit wing. The slots are cut to a width of 0.342 to 0.344 in. Maintenance of this slot width is necessary to ensure that the optimum brazing clearance of 0.002 in. will result after assembling of shims and carbide in the slot. Prior to March 1953, when the milling machine was installed, drill steel was slotted on a small manually fed ¾ hp milling attachment mounted on the bed of a lathe. Over 16,000 drill steels were slotted on this unit, and in view of its small size and low cost it gave excellent service. Brazing of Tipped Steel Drill steel that has been milled and cleaned in carbon tetrachloride is mounted in a rotating cradle holding six drill steels, the length of which may be from 2 to 12 ft. The slots in the drill steel, the shims, and the tungsten carbide inserts are thoroughly fluxed with a fluoride flux and assembled as shown in Fig. 1. Fig. 2 shows the brazing equipment in use. As the ring burner is lowered over the bithead a spring valve opens the gas lines, and the gas mixture, preset to give a slightly reducing flame, is fed to the ring burner where it is lit from a pilot flame. The ring burner heats the drill steel over a zone about 1 to 2 in. below the bithead, which becomes heated by conduction through the steel. By this means the bithead is heated rapidly and evenly, and contamination of the brazing joint with soot from the flame is avoided. The bithead is heated to the melting temperature of the brazing alloy within 1 min. This rapid heating minimizes the disadvantage of a non-eutectic brazing alloy. The brazing alloy, a nickel-bearing quaternary alloy, is placed at the bottom of the slot below the carbide insert, as shown in Fig. 1. As the brazing alloy melts it is drawn by displacement by the carbide and by capillary action into all parts of the joint to displace liquid flux from metal surfaces. As soon as the brazing alloy melts, each insert in turn is wiped by being moved back and forth along the slot. This action assists wetting of the carbide by the brazing alloy and assists in displacing molten flux from the joint. After continuous heating for about 75 sec, when the bithead has reached a temperature of about 1500°F, the ring burner is raised and the gas supply is shut off automatically by the spring valve. As soon as heating is stopped a hand press is placed on the bithead and the inserts are squeezed down firmly. This action minimizes the clearance between the bottom of the insert and the slot. Correctly brazed steel should maintain a clearance at the bottom of the slot of 0.001 to 0.002 in. After six steels have been brazed they are removed from the cradle and allowed to cool in air. As soon as each drill steel is cool it is dressed on a grinding wheel to remove excess flux and braze and is ground to the gage appropriate to the length of the drill steel.

Jan 1, 1955

By R. G. Davies, P. Beardmore, T. L. Johnston

The flow stress of a series of Ni-Cr-A1 alloys consisting of a dispersion of y' (based on Ni3Al) in a rnatrix of nickel-base solid solution y has been measured at temperatures up to 950°C as a fwzction of the volume fraction of y'. At high temperatures the flow stress is controlled by the amount of Y' in the alloy, i.e., the higher the volume fraction of y', the greater is the flow stress. This simple relationship is not obeyed at low temperatures in so far as a peak in the flow stress-volume fraction relation occurs at about 25 pct y'. The variation in the mechanical properlies of these alloys as a function of both temperature and volume fraction of y' has been correlated with changes in distribution of both the dislocations and y'. The results are interpreted on the basis that at low temperatures the y matrix is strengthened significantly bv the presence of a hyperfine y' precipitate due to decomposition on cooling; at high temperatures the y matrix is a single phase of low strength. It is clearly recognized that the high temperature strength of most nickel-base superalloys depends upon a dispersion of the ordered fcc phase y', based on Ni3A1, in a fcc solid solution matrix y based on nickel. Although the volume fraction of y' varies widely from about 0.2 in Nimonic 80A to about 0.6 in Mar-M200, all such nickel-base alloys manifest an unusual insensi-tivity of the flow stress with respect to temperature. In Mar-M200 for example, the 0.2 pct flow stress remains essentially constant from room temperature to 750°C. The conclusion has been drawn1 that the characteristically low temperature dependence of the flow stress of y-y' nickel-base alloys is obtained when the state of dispersion of y' is such that dislocations are forced to cut through the y' particles at the onset of yielding. When the spacing between the y' particles is so large that the flow stress is controlled by dislocation bowing between particles, then the initial flow stress decreases progressively with an increase in temperature at a rate determined by changes in elastic properties. The same conclusion is inherent in the detailed, mechanistic model of the deformation process in commercial superalloys which has been developed by Copley and ear' in which the temperature independent flow stress is attributed primarily to the contribution of the antiphase boundary energy created in the y' particles during deformation. In this theory the temperature insensitivity of the flow stress is a reflection of the constant antiphase boundary energy as a function of temperature. An important microstructural parameter that is relevant to the explanations that have been suggested' to account for the temperature insensitivity of the flow stress is the volume fraction of y'. To vary the latter to any significant extent in a given commercial alloy is clearly difficult. However, it is possible in a relatively simple Ni-Al-Cr ternary system which manifests analogous microstructures in terms of the distribution of y' in y and contains specific alloys which have flow properties that depend on temperature in a manner quite similar to their more complex commercial counterparts. Hornbogen et . have studied precipitation phenomena and deformation mechanisms in such alloys but only where the y' volume fraction was small (less than 0.2) and the y' particle size varied from less than 100A up to a maximum of -1000A. In the present study, a series of alloys was prepared in which the volume percent of y' at 900°C was varied from 0 to 100 pct with the y' particle size (of the order 0.5 p) comparable to the sizes obtained in commercial superalloys. Particular attention has been given to the relationship between variations in the volume fraction and distribution of y' and the temperature dependence of the flow stress EXPERIMENTAL TECHNIQUES The Ni-Cr-Al system was selected because it is well characterized, bears a close relationship to commercial alloys, and offers the advantage of an extra degree of freedom over a binary system. In the present investigation, a series of alloys across the tie line between NisA1 and Ni3Cr (Ni3Cr is not an in-termetallic compound, the nomenclature is only used to designate the composition) were vacuum cast. The pseudobinary6 and the composition of the alloys used are shown in Fig. 1. It is important to note that the compositions of the y phase and the y' phase in the two-phase alloys was always the same. Alloy compositions were selected from the binary diagram, Fig. 1, in order that aging at 900°C would produce from 0 (100 pct y) to 100 pct y' by volume percent. (The size of the y' particles produced during the equilibrium aging treatment increased as the volume fraction of y' increased, ranging from about 0.2 p at low volume fractions up to about 0.8 p at the highest volume fraction.) The y' phase is based on the inter-metallic compound Ni,A1 which has the fcc LIZ type superlattice structure, and chromium substitutes for aluminum in the structure. The y phase is a disordered fcc solid solution. The alloys were heat treated at 1150°C for 2 hr, air cooled to room temperature, and finally annealed for 16 hr at 900°C. The rods were then centerless ground to 0.25 in. diam and cut into compression samples 0.5 in. long. The compression tests were made on an In-stron machine at a strain rate of 7 x 10"4 sec-'. A rapid heating radiant heat furnace was used which minimized the heating and temperature stabilization time to 10 min for the highest testing temperature. All the tests were stopped after 5 pct plastic strain.

Jan 1, 1970

By C. E. Birchenall, C L. McCabe

IN attempting to extend vapor pressure measurements of the type previously reported by Schadel and Birchenall1 for silver and by Schadel, Derge, and Birchenall' for silver-silicon to other systems, it was observed that the materials melted at indicated temperatures 10" to 15" below their accepted melting points. Further investigation revealed that the thermocouple readings were in error due to appreciable conduction losses along the reference thermocouple wires. If the wire diameter of the reference couple inserted into the Knudsen cell was reduced, the correction for the indicating couple changed in a manner tending to explain the melting behavior. When extrapolated to zero wire diameter from measurements with several reference thermocouples of different wire thickness, the melting point of silver then agreed with the indicated temperature at which silver chips were observed to coalesce into a sphere. Approximately the same calibration was given by observing the melting of small wires of silver or gold in the Knudsen cell connected in series with an ammeter, where the leads into the cell were very fine in order to minimize heat conduction. Unfortunately neither of these methods seemed to yield a sufficiently precise temperature calibration to match the apparent precision of the other aspects of the vapor pressure measurement. It was decided. therefore, to redetermine the vapor pressure of silver in another setup under conditions permitting precise temperature measurement. The vapor pressure of pure silver could then be used as an internal calibration of temperature in the older unit in making runs on alloys. This has been done; the present report is a correction to ref. 1. Experimental Procedure The apparatus, shown in Fig. 1, was very similar to that employed by Harteck,3 except that the orifice sizes were smaller and the residual pressure in the vacuum system was probably much lower. A small, sharp-edged hole, nearly circular in shape, was ground into the rounded end of a quartz tube. The orifice area was then measured by tracing the image at known magnification on graph paper and counting the squares enclosed. The silver specimen was sealed into the tube to make a Knudsen cell. A tantalum jacket surrounding the cell served to increase the uniformity of temperature. This assembly was placed in the bottom of a long quartz tube with an inside diameter of about 1 in., which was connected to the vacuum system through a ground joint sealed with picein wax well removed from the furnace. A thermocouple tube inserted through the top of the vacuum line reached into the tantalum jacket so that the thermocouple junction was immediately adjacent to the Knudsen cell except for the protection tube wall. A resistance furnace could be raised to cover the end of the quartz tube containing the cell in such a way that the cell was in the uniform temperature zone 13 in. from the end of the furnace. An ionization gage was included in the vacuum system in the cold lines of wide diameter, immediately beyond the ground joint. The vacuum system consisted of a mercury one-stage diffusion pump, backed by a Welch duo-seal mechanical pump. The pumps were separated from the reactor chamber by a dry ice trap. The ionization gage always read less than 10-5 mm Hg after initial outgassing and before each run was started. Each newly filled Knudsen cell was evacuated at high temperature overnight before the first weighing was made. The cell was returned to the system, heated for a measured time at constant temperature, cooled, and reweighed. The heating and cooling times were quite short since the hot furnace was raised to receive the reactor at the beginning of the run and removed again at the end. The tube heated or cooled quickly. The total mass loss was attributed entirely to effusion of silver vapor from the quartz cell, since empty quartz cells maintained constant mass through similar heating cycles. The vaporized silver condensed on the cold walls of the quartz tube extending above the furnace. Earlier studies in the induction heated unit had shown that the same vapor pressure was found for silver, whether the silver was in contact with the tantalum metal cell or with porcelain or quartz liners. The Pt-Pt-10 pct Rh thermocouple was calibrated against a secondary standard of the same material and found to agree with the published tables. Always operating in air at temperatures below 100O°C,

Jan 1, 1954

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

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

By Raymond C. Boettner

The present work had for its purpose: 1) the identification of crack nucleation sites in AISI 4340, quenched to martensite and tempered over a range of 'temperatures; and 2) the comparison of fatigue processes in AISI 4340 with processes observed previously in pure metals From constant def1ection-bending fatigue tests, martensite boundaries were identified as the favored crack nucleation sites in quenched and tempered AISI 4340. It, also, was concluded that the fatigue processes operating- in this lous-alloy steel were similar to Processes observed in pure tnetals. ALTHOUGH much engineering data has been accumulated on the fatigue properties of quenched and tempered martensitic steels,' fatigue as a process is not as well understood in martensite as it is in pure metals.' Important features of the fatigue process, such as the identity of the nucleation sites, have not been determined in the commercially important high-strength low-alloy structural steels. The present work had for its purpose: 1) the identification of crack nucleation sites in a low-alloy steel, i.e., AISI 4340, which had been quenched to martensite and tempered over a range of temperatures; and 2) the comparison of fatigue processes in the AISI 4340 with processes observed previously in pure metals. This comparison of the fatigue processes in the different tempers was restricted to the high-strain low-cycle part of the S-N curve. Under these test conditions, previous work on a number of metals has shown that a large number of cracks are nucleated in less than 30 pct of the fatigue life.3 Furthermore, crack nucleation sites are not restricted to inclusions but are also associated with intrinsic structural characteristics of the metal. MATERIAL A 20-lb ingot of vacuum-melted AISI 4340 (for composition see Table I) was hot-rolled to 1-in.-diam rod and then cold-rolled to a 1-in.-wide strip, 0.08 in. thick. Fatigue specimens, see insert of Fig. 1, were machined from the strip with the long dimension parallel to the rolling direction. m this orientation, the stringers of 1 to 2 p inclusions present in the sheet lay parallel to the stress axis in the specimens. The specimens were austenitited at 2050°F in order to obtain a large prior austenite grain size, i.e., 2 mm, which facilitated the subsequent identification of the prior austenite boundaries. A helium atmosphere was used to minimize decarburization. After austenitiza-tion at 2050°F, the specimens were transferred to a 1450°F furnace so that specimen distortion was held to a minimum in the subsequent oil quench. Previous work4 indicated that refrigeration in liquid nitrogen prior to tempering reduced the percentage of retained austenite in the quenched specimens to less than 5 pct. Tempering was carried out in air over the temperature interval of 200°to 800°F to produce a range of mechanical properties, Table I. The preparation of the fatigue specimen was completed by grinding about 0.005 in. from each surface and electropolishing in a chrome trioxide-acetic acid solution for 30 min. Examination of etched cross sections of specimens prepared in this fashion showed the foregoing specimen preparation to be adequate for the removal of the decarburized layer present after the heat treatment. Transmission electron microscopy showed that the as-quenched microstructure of this alloy consisted of a mixture of martensite plates containing either a high density of dislocations or microtwins. Previous work5'6 indicated that in the course of oil quenching autotem-pering resulted in the formation of E carbide on the martensite and microtwin boundaries. Tempering for 2 hr at temperatures up to about 400°F resulted in further precipitation of the E carbide. Finally, at about 400°F, cementite began to replace the E carbide on the martensite and microtwin boundaries in addition to forming a Widmanstatten structure within the plate matrix. EXPERIMENTAL S-N curves were obtained using electropolished specimens cycled at 1800 cpm as cantilever beams in fully reversed bending at selected constant deflections. The deflections were translated into surface strains by means of a calibration curve obtained through the use of strain gages. An argon atmosphere was used to minimize environmental effects. To investigate the development of fatigue slip bands, the specimens of the different tempers were unidirec-tionally bent to a surface strain of 0.005 to 0.007, photographed to record the location and appearance of slip bands so introduced, and then cycled to failure

Jan 1, 1968

By J. R. Cahoon, H. W. Paxton

The mechanical properties of unidirectionally solidified A1(rich)-Mg and A1(rich)-Cu castings containing up to 15 wt pct solute have been determined with re -spect to the volume fraction of interdendritic eutectic. Pioperties were determined in the directions pumllel and Perpendicular to that of solidification; the volume fraction of eutectic was varied between the "as-cast" and equilibrizcm amounts by approperiate heat treatment following solidification. The principles of fiber strengthened composites and dispersion strengthened materials are adapted to explain the mechanical properties of these castings. It is generally accepted that castings often have inferior mechanical properties when con~pared to wrought products. However, there is little quantitative data available concerning the factors which make apparently sound castings weak and/or brittle. The relative ease and inexpensiveness of the casting process have always been attractive and, therefore, an understanding of the factors which contribute to the mechanical properties of castings would seem desirable. Such an understanding may lead to an improvement in the mechanical properties to an extent where castings would become competitive in applications where presently only wrought products are considered to have the requisite properties. Such an understanding could also improve the reliability of present cast products. Much of the recent research on castings has centered about determining the extent of segregation in cast alloys. Macrosegregation, particularly inverse segregation, has been studied in some detail 1-8 and a considerable understanding of microsegregation has been obtained.9'10 The effect of solidification rate on dendrite spacing and on the amount of interdendritic eutectic in binary alloys has been established, particularly for Al(rich)-Cu alloys.""0 However, the extension of these ideas to relate the amount of interdendritic eutectic, concentration gradients, micro-segregation, dendrite spacings, and so forth, to the rnechanical properties has been limited. Dean and spear" have related the mechanical properties of an Al-Si-Mg alloy, A356-T62, to the dendrite spacing and have shown that the mechanical properties improve with decreasing dendrite spacing. Passmore et al.12 have shown that annealing at high temperature improves the mechanical properties of Al(rich)-Cu al- loys and Archer and Kempf 13 have shown that an Al-1 pct Mg-1.75 pct Si alloy behaves in a similar manner. Ahearn and Quigley 14 have shown that high temperature homogenization also enhances the mechanical properties of an SAE 4330 steel. However, in the above investigations, no underlying reasons were suggested for the improvement in mechanical properties. The purpose of the present investigation is to relate the mechanical properties of castings to some of the solichfication variables and to derive some equations by which calculations of the mechanical properties may be attempted. In particular, the effect of the amount of interdendritic eutectic and the effect of stress direction with respect to that of solidification on the mechanical properties will be considered. The Al(rich)-Mg and Al(rich)-Cu binary alloy systems were chosen for study. The A1-Mg system was chosen because its constitutional relationships are such that large volunles of eutectic (up to 24 vol pct) may be obtained in the as-cast condition and then be completely dissolved by subsequent heat treatment at about 440°C. This allows a comprehensive study relating the mechanical properties of castings to the amount of interdendritic eutectic. Also the Al(rich)-Mg eutectic is almost a single phase 15 which should make the experimental results more amenable to theoretical interpretation and calculation. The A1-Cu system was chosen for study because of the large amount of related information available concerning segregation, dendrite spacing, and so forth. Unidirectionally solidified castings were used throughout the investigation so that the effect of solidification direction with respect to the direction of applied stress could be determined. THEORETICAL It is well known that upon solidification of binary alloy castings, the nonequilibrium amount of eutectic which forms is given by 10 where fe o is the weight fraction of eutectic, Cs is the solid solubility of solute at the eutectic temperature, k is the equilibrium partition coefficient, and C, is the average composition of the alloy. In the development of Eq. [I], it is assumed that the effects of inverse segregation and diffusion in the solid are negligible, and that no porosity is present. If the casting is homogenized at a high temperature for a long period of time, some (or all) of the eutectic is dissolved and the amount of eutectic for this "equilibrium" condition may be calculated directly from the constitutional diagram. By appropriate intermediate annealing, the

Jan 1, 1970

By C. D. Williams, C. E. Ells

The susceptibility of quenched and aged Zr-2.5 wt pct Nb alloy to embritt2ement during irradiation has been examined for a number of solution temperatures and aging times. Material quenched from temperatures approximately 40°C below the transus has high tensile ductility, and this ductility is insensitive to aging at 500°C or to irradiation. If, however, the material is quenched from temperatures above the transus it becomes highly susceptible to loss of ductility either from aging at 500 or from irradiation. Inter granular failure is characteristic of the materials having low ductility. The distribution of the equilibrium phase is found to control the susceptibility to embrittlement by restricting 6 grain growth during heat treatment and thus influencing crack propagation. IN zirconium, as in titanium, -stabilizing alloy additions are used to obtain high strengths via quench and age heat treatments, and the Zr-2.5 pct Nb alloy has been developed1 because of its strength advantage over the Zircaloys. Early in the development of the Zr-2.5 pct Nb alloy the problem of 13 embrittlement was appreciated, and for this reason the solution temperature was chosen below the p transus.' In the course of irradiation studies on quenched and aged Zr-2.5 wt pct Nb alloy it was found' that irradiation introduced an important aspect of p embrittlement, riz., material quenched from the phase and aged 24 hr at 500°C was severely embrittled by moderate doses of neutron irradiation. This effect had not been studied in titanium alloys. In titanium the metallurgical features leading to 0 ernbrittlement were found to be structures with: a) coarse a platelets at the grain bondaries, b) finely dispersed a uniformly distributed throughout the (0) matrix,6 c) Widmanstatten a-13 with more than 50 pct P, d) the presence of some metastable p transformation products,3 and e) large prior -phase grain size.5 Alternatively, the presence of a uniform distribution of coarse a was conducive to high ductility and a structure largely of equiaxed a was very dctile. The detailed mechanisms of the embrittlement have not been worked out for all of these conditions, although weakness at either a-matrix boundaries or prior p grain boundaries have been prominent in the eculation. It was proposed that acicular a might act as a mild notch, and low ductility has been associated with easy fracture along its boundary.' There have been two opposing suggestions for the source of the high ductility associated with equiaxed a phase. JaffeeB proposed that this a would accept a large por- tion of the oxygen, thus increasing the ductility of the matrix, whereas after study of a Zr-Nb-Cu alloy Weinstein and oltz proposed that the a phase, softer than the martensitic matrix, acted to blunt cracks formed in the matrix. In the present work we have studied the effect of neutron irradiation on the ductility, particularly the P embrittlement, of the Zr-2.5 wt pct Nb alloy. By a variation of solution temperature and aging time a variety of metallurgical conditions have been examined, and a range of resultant ductilities obtained. The ductility has been related to the material microstructure and mode of fracture. EXPERIMENTAL The alloy used in the present work came from two separate ingots fabricated into rod of 3/8 or i in. diam, Table I. For both batches the P transus temperature was approximately 890° C. Most of the heat treatments were done directly on lengths of the j} in. diam rod, after which the tensile test specimens were machined. Quenching was achieved by dropping rods from a dynamic vacuum into water, the cooling rate estimated to be 2 100°C per sec. For aging the rods were encapsulated in evacuated silica tubes. Round tensile test specimens, with gage diam and length 0.160 and 1.0 in., respectively, were used throughout and pulled at room temperature or 300°C on Instron tensile machines, at a crosshead speed of 0.05 ipm. Specimens were irradiated in the NRX and NRU reactors, in facilities described in previous publications.'0 The metallurgical conditions examined have been: All tensile test specimens were machined with axes in the axial direction of the swaged rod. Although the specimen had a degree of preferred crystallo-graphic orientation with basal plane normals both parallel with and perpendicular to the tensile axis, the material was comparatively isotropic." The techniques of thin foil examination in the electron micro-

Jan 1, 1970

By N. A. D. Parlee

The diffusion rate of hydrogen in liquid iron has been measured by a gas-liquid metal diffusion cell technique. The diffusion cell was formed by immersing an alumina tube containing hydrogen gas at 1 atm in a bath of stagnant liquid iron. The change in the composition of the melt in the cell was determined by measuring the rate of absorption of the gas in the cell. The appropriate solution to Fick's second law was used to examine the data and calculate diffusivi-ties. The absorption of hydrogen in stagnant pure liquid iron has been found to be diffusion-controlled. The results show that the chemical diffusion coefficient, D, of hydrogen in pure iron in the range of 1547" to 1726°C can be represented by the following Arrhenius relation: D(sq cnz per sec) = 3.2 x X exp(- 3300 i 1800/RT) where the uncertainty in the activation energy corresponds to the YO pct confidence level. Oxygen in the melt (above 0.015 pct 2) increased the apparent rate of absorption of hydrogen. The importance of diffusion data on liquid metals for predicting the rates of certain metallurgical processes has been recognized for a long time. Moreover, these data are much needed to test and develop theory for diffusion in liquid metals. Despite this practical and theoretical interest, however, relatively little reliable information about diffusion in liquid metals is available in the literature. This is particularly true for gas components such as hydrogen, oxygen, and nitrogen in liquid metals, where almost no data on chemical diffusion coefficients are to be found. This is probably due to a multitude of experimental difficulties particularly associated with high-temperature melts. In an effort to fill this gap in information, a research program was undertaken to study the diffusivities and rates of solution of gases in liquid metals. This paper presents the results of a study of the diffusion of hydrogen in liquid iron. EXPERIMENTAL METHOD Two methods for the study of the kinetics of dissolution of gases in liquid metals are being employed in this laboratory. Both involve the measurement of the volume of gas absorbed by the melt as a function of time and as such both avoid the uncertainties involved in chemical analyses of quenched samples for relatively small amounts of gas. In the first method, the gas dissolves in an inductively stirred melt and, in the absence of a slow surface reaction, the results are often interpreted in terms of mass transport across a liquid "boundary layer" between the homogeneous gas phase and well-stirred part of the melt. Other interpretations of the results of such experiments have also been described in the literature.1'5 In the second method a gas-liquid metal diffusion cell is used.' The gas dissolves in a cylindrical column of stagnant liquid metal and, in the absence of a slow surface reaction, the results are interpreted in terms of a non-steady-state diffusion solution to Fick's second law. The weakness of the first method is that while it gives information on the mechanism of absorption by stirred melts it yields an overall rate constant which even in the simplest cases depends on the nature and the thickness of the "mass transport layer" or "boundary layer". It yields no values of diffusion coefficients. The second method was used in this research because in many cases it is possible to determine the diffusion coefficient of the gas component in the liquid metal. In this research it has been utilized to measure diffusion coefficients of hydrogen in liquid iron. The apparatus used was essentially the same as that described by Mizikar, Grace, and par lee but certain modifications have been introduced to meet the elevated temperatures and special conditions of this research. Fig. 1 is a schematic drawing of the apparatus and Table I gives the identification of various parts in this figure. The diffusion cell, shown in detail in Fig. 2, was formed by immersing an impervious alumina tube (hereafter called absorption tube) in a bath of pure liquid iron contained in an alumina crucible. Two types of tubes were used, Morganite triangle RR and McDanel AP35. The crucible was contained in a vertical impervious alumina combustion tube (32 mm ID by 914 mm long) which was closed at both ends by water-cooled brass heads employing O-ring compression seals, Fig. 1. A protection tube enclosing a Pt, 5 pct Rh-Pt, 20 pct Rh thermocouple was introduced through the lower end of the combustion tube

Jan 1, 1968

By R. G. Agarwal, Ramey Jr. H. J., Al-Hussainy R.

Although it has long been realized that gas recovery from a water-drive gas reservoir may be poor because of high residual saturations under water drive, it appears that only limited infomlation on the subject has been available until recently. This study was performed to show the qiiantitative potential importance of water influx. Results indicate that gas recovery may be very low in some cases: perhaps as low as 45 per cent of the initial gas in place. Gas recovery under water drive appear to depend in an important was on: (I) the prodirction rate and manner of production; (2) the residual gas saturation; (3) aquifer propertie.); and (4) the volumetric displacement effciency of water invading the gas reservoir. The manner of estimating water-drive gas reservoir recovery can vary considerably. Examples are: the steady-state tnethorl. the Hurst modified steady-state method, and various unsteady-state methods such ac. those of van Ever-dingen-Hurst, Hurst, and Carter-Tracy. The Carter-Tracy water influx expression was used in this study. In certain cases, it appears that gas recovery can be increased significantly by controlling the production rate and manner of production. For this reason, the potential importance of water influx in particular gas reservoirs should he investigated early to permit adequtrtr planning lo optirtize the pay reserves. INTRODUCTION In recent years, the economic importance of natural gas production has become increasingly apparent. This has been evidenced by more intensive exploration efforts aimed at gas production, and exploitation of both deep, as well as low-permeability gas reservoirs. Technical developments such as deep-penetration fracturing have made development of such formations economically feasible. Unfortunately, water influx has forced abandonment of a number of gas reservoirs at extraordinarily high pressures. Although reservoir engineering methods for estimating water influx have long been available, it appears that application of these methods to the water-drive gas reservoir has been sporadic.'a Available methods for estimating water influx which can be applied to the water-drive gas reservoir problem include the steady-state method,1 the Hurst modified steady-state method and various unsteady-state methods such as those of van Everdingen-Hurst. Hurst, and Carter-Tracy. Interesting applications of these solu- tions to gas reservoir and the aquifer gas-storage problems have appeared recently.3,12,14 The experimental study of residual gas saturations under water drive by Geffen et al. in 1952 indicated that residual gas saturations could be extremely high. A value of 35 per cent of pore volume is often used in field practice when specific information is not available. The study of Geffen et al. showed that residual gas saturation might be much higher in some cases. Naar and Henderson concluded that the residual non-wetting phase saturation under imbibition should be about half of the initial non-wetting phase saturation. The Naar and Henderson result that residual gas saturation under water influx should be about half the original gas saturation is recommended as an estimate if laboratory measurements are not available. Thus, it is clear that a considerable portion of the initial gas in place might be trapped in a water-drive gas reservoir as residual gas at high pressure. A full water-drive would result in loss of residual gas trapped at initial reser.voir pressure. Consideration of transient aquifer behavior leads to the conclusion that high-rate production of water-drive gas reservoirs could result in improved gas recovery by reduction of the abandonment pressure. However, there appears to be little quantitative information on this possibility. One of the few advantages of water-drive gas production appears to be improved deliverability through water-drive support of the reservoir pressure. There may also be an advantage in higher condensate recovery caused by pressure maintenance for gas-condensate water-drive reservoirs. In view of the preceding, this study was made to assess the potential importance of water-drive in gas reservoir engineering. The Carter-Tracy approximate water-influx expression was used because this equation offers some advantages in hand-calculation which do not appear to have been generally recognized.' However, calculations were performed in the main with a high-speed digital computer to permit evaluation of the effect of water-drive under a large variety of conditions. CALCULATION METHOD Water-drive gas reservoir performance can be estimated in a manner completely analogous to oil reservoir calculations: a materials balance is written for the reservoir, and a water influx equation is written for the aquifer. Siniltaneous solution provides the cun~ulative water influx and the reservoir pressure. When reservoir performance data (gas produced and reservoir pressures) are available, it is usually possible to match performance data to determine the initial gas in place and the water influx parameters —

Jan 1, 1966

By H. D. Hoskold

As the theory and the practice of valuing mines have never been discussed in the Transactions, a paper on the subject may be acceptable, even though not exhaustive. The method here indicated is set forth, not as a model of perfection to be followed absolutely, to the exclusion of any other system based on sound scientific, commercial and equitable principles, but merely as an aid to mining engineers and financiers in estimating the value of mining-property. This paper is intended merely to point out certain financial features of the question, and not at all to go into the theories of the formation of mineral veins, extremely valuable as it must, of course, be to understand the laws of the earth's formation and of its various changes, which are the basis of such important sciences as geology and mineralogy, and highly useful as guides in determining the probable existence and extension of mineral deposits. Financial Elements of Mine- Valuation. The practice of basing the stock-exchange price of mining shares upon reports, periodically received from the mine-captain or manager, stating that in his opinion the mine is " improving " in this or that level, and is worth so much " per fathom," is, of course, absurd. Everybody experienced in metalliferous mining is aware that a mineral vein varies in thickness and also in percentage of metal in different parts of the same mine. According to the practice just indicated, a mine must have had as many values a per fathom " as there have been given reports upon it at intervals during the year,

Jan 1, 1903

By Henry M. Howe

The transformation which steel undergoes in slow cooling, from the condition of austenite whelk above the transformation rage into that of pearlite plus either ferrite or cementite below that range, is subject to great lag. This transformation is essentially the pearlitizing of the austenite. Indeed the hardenillg of steel is due to this lag, which is made so great by extremely rapid cooling that a very large part of the transformation itself is restrained, so that the rapid cooling catches the metal in an only partly transformed state, that of martensite, and under favorable contditions even in an almost wholly untransformed state, as in the cases of Maurer's austenitic manganese steel. Hadfield's austemtic. manganese steel, 25 per cent. nickel steel, and other anstenitic alloy steels, arc austenitic even after slow cooling. It has long been known that this lag increases (1) with the rapidity of cooling, (2) with the presence of certain retarding elements, notably manganese and nickel, and (3) with the temperature from which cooling starts. The present inquiry touches the last of these influences, but incidentally it throws light on the first also. This influence of higher heating ill increasing lag manifests itself in different ways which, though they may at first seem unrelated, are in fact hut different aspects of one and the same thing. Thus, among heatings to different high temperatures, all of them above the transformation range: the higher the temperature to which the steel is heated (1) the lower is the temperature at which the transformation occurs in cooling slowly down, (2) the shorter is the time occupied in the actual cooling from above the transformation range to the room Temperature under given external conditions, Whether of slow or

Jan 1, 1914