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By J. S. Bowles

THE martensite transformation in lithium, dis- covered by Barrett,' has been studied extensively by X-ray techniques by Barrett and Trautz,² and Barrett and Clifton.V he present paper reports the results of an investigation into the metallographic characteristics of lithium martensite. Such an investigation has not been carried out before. The spontaneous transformation in lithium consists of a change from a body-centered cubic to a close-packed hexagonal structure with the hexagonal layers in imperfect stacking sequence." As far as is known at present, this transformation can be regarded as being crystallographically equivalent to the body-centered cubic to close-packed hexagonal transformation that occurs in zirconium,5 although stacking errors have not been reported in zirconium. From a study of the orientation relationships in zirconium, Burgers5 as proposed that the martensite transformation, b.c.c. to c.p.h., occurs by a heterogeneous shear on the system (112) [111]. The crystal-lographic principle underlying this proposal is that the configuration of atoms in the (112) plane of a b.c.c. structure is exactly the same as that in the (1010) plane of a close-packed hexagonal structure based on the same atomic radius. The pattern in 2v2 both these planes is a rectangle d X 2v2d where v3 d is the atomic diameter. Thus a close-packed hexagonal structure can be built up from a body-centered cubic structure by displacing the (112) planes relative to each other.* This mechanism leads to orientations that can be described by the relations: (110)b.c.e. // (0001)c,p.h.; [111]b.c.c. // [1120]c.p.h Observations confirm these relations. In zirconium, Burgers' measurements indicated an angle of 0" to 2" between the close-packed directions, while Barrett's measurements on lithium indicated an angle of 3". According to the Burgers' mechanism, the martensite habit plane for this transformation would be expected to be the (112)b.c.c. plane, for this plane would not be distorted by the transformation. One of the purposes of this investigation was to find out whether the observed lithium habit plane agrees with this prediction of the Burgers' mechanism. Experimental Procedure Materials: The lithium was from the same purified ingot used by Barrett and Trautz.² The Bridgman technique was used to produce single crystals. To maintain a temperature gradient in the melt, during the production of these crystals, it was necessary to use a steel mould with a wall thickness of only 0.015 in. Metallographic Techniques: Lithium specimens could be given an excellent metallographic polish by swabbing them gently with cold methyl or ethyl alcohol.? The best results were obtained with methyl alcohol saturated with the reaction product, lithium alcoholate. With higher alcohols the reaction became progressively slower and the attack became an etch pit attack rather than a polish attack. Butyl and amyl alcohols were used for macroetching. After polishing, it was necessary to remove all traces of alcohol from the specimens; otherwise, on subsequent quenching in liquid nitrogen, the alcohol froze to a glassy film. The alcohol was removed with dry benzene. The benzene in turn had to be removed before quenching, but since it does not react with lithium it could be allowed to evaporate. The specimens could then be quickly quenched before they began to tarnish. This operation could be carried out in air on all but excessively humid days when it was advisable to use an atmosphere of dry nitrogen or argon. For examinations at room temperature, the specimens could be transferred directly from the benzene bath into a bath of mineral oil. In mineral oil the specimens oxidized slowly by the diffusion of oxygen through the oil but the structure remained visible for about an hour. Lithium Martensite: Specimens prepared in the manner described above transformed spontaneously to martensite with an audible click when quenched into liquid nitrogen; i.e., M, was above the boiling point of nitrogen (77°K). The disparity between this result and the M, temperature of 71°K, found by Barrett and Trautz, is probably to be attributed to the large grain size and freedom from mechanical deformation of the specimens used in the present work. The relief effects produced by the transformation did not disappear when specimens were quenched from liquid nitrogen into mineral oil at room temperature. This permitted the microstructures to be studied at room temperature where, of course, the martensitic phase was no longer present. Typical micrographs of lithium "martensite" made at room temperature are reproduced in figs. 1, 2, and 3. As anticipated by Barrett and Trautz, the microstruc-

Jan 1, 1952

NEW MEMBERS The following list comprises the names of those persons who became members during the period Nov. 9, 1918, to Dec. 10, 1918. ALBERT, EDWARD J 235 E. Upsal St., Philadelphia, Pa. BARRY, D. A Supt. of. Mines, The Big Five Min. Co., Idaho Springs, Colo. BATON, GEORGE S., Cons. Min. Engr., Baton & Elliott, 2412 First National Bank Bldg., Pittsburgh, Pa. BAXTER, HAROLD ALEXANDER, Met. Engr. & Operation Supt., Tacony Ordnance Corpn., Philadelphia, Pa. BENNETT, JOHN S., Chem. and Min. Engr Copiap6, Chile, South America. BOWDEN, MALCOLM, Min. Engr.; Efficiency Engr., Elm Orlu Min. Co., Butte, Mont. BROWN, HAROLD C 2d Lieut., A. S. A. P. BROWN, THOMAS ALBERT, Min. and Met. Engr.; Supt., The Royal Tiger Mines Co., Breckenridge, Summit Co., Colo. BROWN, WALTER S., Investigator, New Jersey Zinc Co. (of Pa.), Palmerton, Pa. BURDETTE, R.. S Capt., Ord. Dept., U. S. A. CALLANDER, IRA C., Asst. Supt., Gauley Mountain Coal Co., Jodie P. 0., Fayette Co., W. Va. COLLINS, J. W., Foundryman, The Aluminum Casting's Co., "L" Plant, Cleveland, Ohio,

Jan 1, 1919

By R. E. Gilchrist, R. F. Nielsen

When a gas is displaced by mother in a porous medium, and there is a relative immobile liquid present. there is a transition zone in which the gas composition varies from essentially that of the original gas to that of the injected gas. The change in extent of this zone with time may be governed by a mixing process in the direction of flow and by the rate of establishment of gas-liquid equilibrium, if the gases are appreciably soluble in the liquid. The relative importance of these two factors in deterri~ining the growth of the transition zone is discussed, with appropriate equations and experimental data. INTRODUCTION The various solution methods of oil recovery involve a transfer of the components of a mixture, whether within a single phase or between a liquid and gas. The similarity of a natural porous medium to systems used in separation and chromatographic processes makes it seem reasonable that certain concepts used by chemical engineers in "mass transfer" might be applicable to petroleum reservoirs. One approach, the HETP (height equivalent to a theoretical plate), involves a distance or length of time over which equilibrium might be con. sidered as established. Another approach involves direct use of a rate of mixing or rate of establishment of equilibrium. The latter approach has been preferred in relation to production problems. DISCUSSION AND PROCEDURE When a gas or liquid in a porous medium is displaced by one of different composition but miscible with it there will be set up a transition zone in which the composition varies from cssentially that of the original fluid to essentially that of the injected fluid. If the fluid being displaced is a gas, and there is a relatively immobile liquid present, the transition zone may be

By S. W. Churchill

A digital computer was used to obtain an exact numerical solution for the transient behavior of the insulation and earth adjacent to an isothermal, submerged flat surface for a single set of parametric values. Comparison of the computed results with analytical solutions for limiting conditions revealed that a complete and general solution for all parametric values could be constructed from these limiting solutions. Complete and general solutions for insulated spheres and cylinders can also be constructed from limiting solutions. The procedure is illustrated in part for an insulated spherical tank. The heat flow was found to fall rapidly to a pseudo-steady state; after some time, it then decreased slowly to zero for a flat plate and cylinder, and to a low steady-state rate for a sphere. The accumulative heat flow during the initial falling-rate period may be a significant fraction of the heat flow during the entire first year. INTRODUCTION Underground storage of liquefied natural gas and liquefied petroleum gas has received considerable recent attention. The rate of heat flow from the earth to the storage cavity and the resulting temperature field in the earth are important factors in the technical and economic evaluation of such storage facilities. The objective of this paper is to indicate how complete solutions can be developed for the transient flux and temperature field for various geometrical configurations. The representative properties and dimensions, and the resulting parameters utilized in the illustrative results, are indicated in Table 1. The results presented herein are for dry earth. The effect of the latent heat of solidification of water in the soiI will be described in a subsequent paper. The temperature field in the insulation and earth is determined by energy balances, boundary conditions and initial conditions. The physical problem can be described mathematically with virtually no idealizations insofar as physical properties are known. It appears possible to solve the equations by numerical integration on a high-speed digital computer for any geometrical configuration and conditions. For complex situations, however, the computations are expensive and the results are highly specific. Analytical solutions have been developed for a few simple but important geometrical configurations and conditions, including one-dimensional heat transfer from (or to) earth at an initially uniform temperature to (or from) flat plates, spheres and circular cylinders. The solution for an insulated flat plate has also been derived but is in the form of a slowly converging series involving taklated functions. It was planned to use a computing machine to evaluate this series for a number of representative conditions. However, upon examination of the results of preliminary computations it was discovered that a complete, general and accurate solution could be developed by interpolation between several much simpler solutions for limiting conditions. This technique then was used to develop a a complete solution for an insulated sphere. The equations presented herein are derived in Carslaw and Jaegerl and other books on applied mathematics, or they are simple extensions of these previous results. Hence, all derivations are omitted. Emphasis is instead placed on the results and their

By Å. Sterten, R. Tunold, J. Brun, K. Dalatun

The subjects of this study are two corrosion phenomena familiar to operators of aluminum plants employing Soderberg anodes of the vertical type, namely the sul-fide scale formation observed at the steel contact spikes and the corrosion wear observed at the lower edge of the gas collecting skirts. The causes of these corrosion phenomena are discussed on the basis of observations from practice and measurements on commercial cells as well as on laboratory experiments. It is concluded that the anode spike corrosion is mainly due to the attack by sulfur compounds contained in the anodically evolved gases and that this attack is greatly promoted by cracks and pores in the anode carbon embedding the lower spike ends. The corrosion of the gas collecting skirt is mainly attributed to anodic dissolution of the iron during occasional contacts with the melt. ANODE SPIKE CORROSION The visible result of this corrosion is the formation of a brittle loosely adherent scale on the lower ends of the mild steel spikes maintaining the electric contact to the carbon anode, Fig. 1. The thickness of the scale usually varies from 0.5 to 3 mm, but in some cases scale of about 10 mm thickness has been observed by the writers. During the spike pulling, usually repeated every 3 weeks, part of the scale often loosens and remains in the anode carbon, resulting in contamination of the aluminum with iron. This corrosion has been attributed to reaction of the iron with gaseous sulfur compounds originating from the anode paste materials.'-4 This paste usually consists of 70 pct precalcined petroleum coke with 1.2 to 1.4 pct S and 30 pct coal tar pitch containing 0.5 to 0.7 pct S. During operation the gasification of the anode materials takes place in two ways: First, gases are evolved by the gradual baking of the anode paste added on top of the anode. Due to the sealing effected by the upper semifluid zone of the paste, these gases, containing sulfur compounds from the pitch, are forced downward through pores and cracks in the rigid anode carbon. At the temperatures of the zone of the lower ends of the spikes, 700" to 900°C, the original components of the baking gas must undergo decomposition, similar to that taking place with the production of coal gas, resulting in a gas mixture consisting mainly of H2 and CH4 with small percentages of H2S as the main sulfur compound.5 Second, the anodic reaction results in a mixture of C02 and CO, the total quantity of which is about 15 times that of the baking gas. These anode gases must carry the entire sulfur content of the carbon gasified by the anodic reaction. Henry and Holliday6 have shown that COS is the main sulfurous component of the anode gases from Soderberg cells. Soderberg anodes usually exhibit a network of vertical cracks, frequently traversing from one spike to the next one and to the outside of the block and widening downward to the working face of the anode. In addition the anode carbon, particularly that below the spike ends, is rather porous, as indicated in Fig. 1. It therefore seems probable that the anode gases also may partly penetrate to the spike ends. As the anode is partly immersed in the melt, the static pressure of this must forward this penetration. I) SPIKE SCALE COMPOSITION A great number of freshly extracted spikes from industrial plants were inspected and samples of the scale collected for examination. Before analyzing the samples a thin iron oxide layer, formed by air oxidation of the outside of the red hot scale immediately after spike pulling, was removed. On the average the Fe/S ratio found by these analyses corresponded to I the composition FeSl. oo7. X-ray powder patterns showed no oxide and otherwise agreed with those found by Haraldsen7 for iron sulfide with 50.00 at. Pct

Jan 1, 1970

By S. S. Clarke

LATE in 1942, the increasing demand for zinc, coupled with the growing shortage of miners and the knowledge that some abandoned mines would have to be reopened for prospecting and development, led to considerable thought as to the possibilities of further mechanization in order to conserve man power. The ore bodies, locally termed "sheet ground," are an important ore bed in some mines and are in the lower part of Boone formation of the Mississippian Series. Some mines contain interbedded layers of nearly pure galena and sphalerite in horizontal sheets ranging in thickness from a thin seam to several inches and is persistent over large areas. This type of ore body offered the best opportunity for study and experiments. The use of jumbo drill carriers was advanced but some doubt was expressed as to the adaptability of jumbos as then in use to mine headings as compared to post and arm mounting (Fig I). Another consideration was the introduction of truck haulage when several mines were trackless. A drill frame, built to carry 2 4-in. drifters, was mounted on a truck chassis. This design proved a failure as it was not rigid and the resiliency of the tires and springs allowed too much play. Long screw jacks were then used to hold the truck firmly in place, but placing the jacks and securing the truck required more effort and time than setting up post drills, so the truck carrier was abandoned. The use of a bull dozer underground had been under consideration for some time to be used in "brunoing" the broken dirt from behind pillars and other inaccessable places so the slusher drags could load without double dragging the dirt. A dozer with hydraulic blade lift was purchased and R. I. Tuthill, Superintendent of Section 30 Mines, suggested that a long drill arm be substituted for the blade, and the drills mounted on the cross member. The hydraulic lift would permit drilling holes at any required spacing. The "cat" was very mobile and answered the problem of a trackless mine. The initial tests showed that the "cat" would make a good carrier, but the hydraulic controls Were not positive enough to keep the drills in alignment while drilling. The next step was to build a small hoist. This hoist was attached to the power take-off on the "cat" and by means of sheaves and a double line of 3/8-in. hoist cable, a positive control for raising and lowering the boom was obtained. To avoid serious injuries to the drillers should the cable break, a safety chain was fastened to the drill boom and hooked into a slot in the back of the "cat" frame. A different power unit was obtained for the "cat," as a state law prohibited the use of internal combustion engines in mines. A custom machine shop had just placed on the market a small hoist powered by a Model A Ford engine that had been very ingeniously converted to a slide valve engine using compressed air. To use this unit on a "cat," our master mechanic de-

Jan 1, 1947

By J. W. Downey, M. V. Nevitt

The phase boundaries for the 950" isothermal sections in the ternary systems Zr-Co-0 and `Zr-Ni-0 have been determined for the composition range from 50 to 100 at. pct Zr. The two systems show very similar phase relations, having no extensive solid solution phase fields. Each contains a ternary phase. These phases are apparently isostructural, but their structure has not been determined. Some aspects of the phase relations are discussed in terms of the alloying behavior of transition metals. THE work described in this paper is the outgrowth of a recent study of the occurrence of phases having the Ti2Ni-type structure (structure-type E9,) in certain ternary systems involving Ti, Zr, or Hf with another transition metal and O.", In the Zr-CO-0 and Zr-Ni-0 systems no Ti2Ni-type phases were found to occur. However, there are several interesting aspects of the phase relations in the two systems which have significance from the point of view of the alloying behavior of transition metals. The results of this investigation may also have some importance in studies of the oxidation of Zr-Co and Zr-Ni alloys. In both the Zr-Co-0 and Zr-Ni-0 systems only the 950" isothermal sections were investigated and, as a further restriction, the study was limited to the composition range from 50 to 100 at. pct Zr. A tentative Zr-Co binary diagram has been published by Larsen, Williams, and Pehin. In the composition range pertinent to the present work they report a eutectic at 980°C and 75.9 at. pct Zr, the products of which are the terminal solid solution based on /3 Zr and the compound Zr2Co, and a eutectic at 1080" and 64.8 at. pct Zr whose products are Zr,Co and ZrCo. The solid solution based on /3 Zr is shown to decompose eutectoidally at 826°C into a Zr and Zr2Co. The limits of solubility of Co in a and /3 Zr have not been established. The structure of Zr2Co is not identified in the publication just cited. Dwight has reported that ZrCo has the CsC1-type strcture. The Zr-rich portion of the Zr-Ni diagram has been determined by Hayes, Roberson, and Paasche." The phase relations are very similar to those of the ZrCo system. A eutectic reaction whose products are /3 Zr containing 2.9 at. pct Ni and Zr,Ni occurs at 961°c, and a eutectic between Zr,Ni and ZrNi is found at 985". The solid solution based on /3 Zr decomposes eutectoidally at 808°C. The solubility of Ni in a Zr is not known accurately but is believed to be very small. Smith, Kirkpatrick, Bailey, and Williams7 have found that Zr2Ni has a tetragonal structure of the A1,Cu-type and that ZrNi is orthorhombic. Domagala and McPherson8 have published a constitution diagram for the system Zr-ZrO,. At 950" their diagram indicates that the solid solution of 0 in 0 Zr is stable from 0 to 0.5 at. pct while the phase field of 0 in a Zr extends from 6 to 29 at. pct. These solubility limits were adopted in the present study and no binary Zr-0 alloys were made. No previous data on the phase diagrams of the ternary systems are known to exist. EXPERIMENTAL PROCEDURE The experimental details involved in the preparation of alloys in this laboratory by arc melting have been described in several previous papers"3 and they will not be repeated here. Information concerning the purity of the metals used is given in Table I. Oxygen was added in the form of reagent grade ZrO,. All of the cast specimens in both alloy systems were annealed in air-atmosphere tube furnaces at 950 3' for 72 hr and water quenched. The specimens were protected from oxidation by wrapping them in Mo foil and sealing them in quartz tubes that had been evacuated at room temperature to a pressure of 1 x 10B mm of Hg. The phase boundaries were determined by metallography, and identification of the phases was accomplished primarily by X-ray diffraction methods which employed a powder camera having a diameter of 114.6 mm. The diffraction techniques which are in use in this laboratory have been previously described.' An etchant that proved satisfactory for most of the alloys consisted of 5 pct by vol of AgNO

Jan 1, 1962

By J. T. Lapsley, I. I. Cornet, A. E. Flanigan, R. Hultgren, J. E. Dorn

With the greatly expanding use of magnesium during the war, it appeared necessary to the War Metallurgy Committee that procedures of heat treating common magnesium casting alloys be investigated systematically. Accordingly, the National Defense Research Committee of the Office of Scientific Research and Development placed a research contract with the University of California, supervised by the War Metallurgy Committee. Research on the solution heat treatment and aging of magnesium alloy AZg2, (formerly ASTM No. 17; also known as Dowmetal C or AM26o) which was done under "Restricted" Project NRC-2I, in 1942-1943, and reported in detail in a final report to the OSRD on September 3, 1943, is the basis of this paper, which has been released for publication by the OSRD. American experience with magnesium alloys prior to World War II was limited, although some of these alloys had been in commercial use since 1909. Requirements of the aircraft industry for light alloys made it imperative to understand and utilize the advantageous strength-weight properties of magnesium castings. Magnesium casting alloys are commonly solution heat- treated to increase their ultimate tensile strength and ductility; this treatment may be followed by an age hardening at elevated temperatures, which raises the tensile yield strength but diminishes the ductility. In the present investigation, AZp was studied to determine the effects on properties of various solution heat treatment and aging schedules. The principles of heat treatment of magnesium alloys are the same as for other non-ferrous alloys. The alloy studied, AZp, has a nominal composition of 9 pct aluminum, 2 pct zinc, and minor constituents. Fig I shows, by the projection of isothermal sections on the concentration plane, the solubility surface1 of the ternary magnesium-rich solid solutions as determined by X ray 'analysis. From this figure it may be seen that above approximately 375°C (707OF), AZp alloy is a single phase alloy at equilibrium; below this temperature the equilibrium condition is a two phase alloy. The beta constituent which precipitates from the solid solution exerts a hardening effect. Because the zinc concentration is only 2 pct, ancl because of its solubility relations, AZqz resembles a binary alloy of magnesium with aluminum. The solubility curve1 of Fig 2 shows how rapidly the solubility of the beta phase decreases with decreasing temperatures. Unlike many aluminum alloys, magnesium alloys do not age harden at room temperatures; precipitation hardening of AZ² alloy is performed at about 177°C (350°F). At this temperature about 3 pct aluminum

Jan 1, 1949

By J. T. Lapsley, I. I. Cornet, A. E. Flanigan, R. Hultgren, J. E. Dorn

WITH the greatly expanding use of magnesium during the war, it appeared necessary to the War Metallurgy Committee that procedures of heat treating common magnesium casting alloys be investigated systematically. Accordingly, the National Defense Research Committee of the Office of Scientific Research and Development placed a research contract with the University of California, supervised by the War Metallurgy Committee. Research on the solution heat treatment and aging of magnesium alloy AZ92, (formerly ASTM No. 17; also known as Dowmetal C or AM260) which was done under "Restricted" Project NRC-21, in 1942-1943, and reported, in detail in a final report to the OSRD on September 3, 1943, is the basis of this paper, which has been released for publication by the OSRD. American experience with magnesium alloys prior to World War II was limited, although some of these alloys had been in commercial use since 1909. Requirements of the aircraft industry for light alloys made it imperative to understand and utilize the advantageous strength-weight properties of magnesium castings. Magnesium casting alloys are commonly solution heat treated to increase their ultimate tensile strength and ductility; this treatment may be followed by an age hardening at elevated temperatures, which raises the tensile yield strength but diminishes the ductility. In the present investigation, AZ92 was studied to determine the effects on properties of various solution heat treatment and aging schedules. The principles of heat treatment of magnesium alloys are the same as for other non-ferrous alloys. The alloy studied, AZ92, has a nominal composition of 9 pct aluminum, 2 pct zinc, and minor constituents. Fig I shows, by the projection of isothermal sections on the concentration plane, the solubility surface1 of the ternary magnesium-rich solid solutions as determined by X ray analysis. From this figure it may be seen that above approximately 375°C (707°F), AZ92 alloy is a single phase alloy at equilibrium; below this temperature the equilibrium condition is a. two phase alloy. The beta constituent which precipitates from the solid solution exerts a hardening effect. Because the zinc concentration is only 2 pct, and because of its solubility relations, AZ92 resembles a binary alloy of magnesium with aluminum. The solubility curve1 of Fig 2 shows how rapidly the solubility of the beta phase decreases with decreasing temperatures. Unlike many aluminum alloys, magnesium alloys do not age harden at room temperatures; precipitation hardening of AZ92 alloy is performed at about 177°C (350°F). At this temperature about 3 pct aluminum

Jan 1, 1947

By C. S. Matthews, H. M. Girner, C. D. Williams, M. J. Edwardson, H. R. Parkison

Quantitative interpretation of electric logs requires knowledge of formation temperature. In this paper, methods are developed for computing changes in formation temperature caused by circulation of mud during drilling operations. The basis of the method is the mathematical solution of the differential equation of heat conduction. The solution of this equation is presented in a series of graphs. These graphs are used to determine formation temperature disturbance at various radii for arbitrary mud circulation histories. Example comparisons with field results show reasonable agreement. It is concluded that, in general, the temperature disturbances caused by circulating mud are small beyond 10 ft from the wellbore but are quite significant near the wellbore. INTRODUCTION Quantitative interpretation of electric logs requires knowledge of the formation temperature in order to establish the resistivity of the formation water with accuracy. To determine the formation temperature, the temperature disturbances produced by circulating drilling mud must be evaluated. -The objective of this investigation was to develop a method for the numerical determination of these temperature disturbances at any distance from the wellbore as a function of time. The first step was to solve the differential equation describing the temperature behavior in the formation during and after mud circulation. This solution was used to calculate a series of curves relating temperature disturbance to shut-in time for various circulating periods. An "exact" method for computing formation temperatures with the use of these temperature-disturbance plots is described, and the results of an application of this method to a well in Montana are presented. Procedures for approximating formation temperatures are also discussed. BASIC EQUATIONS AND ASSUMPTIONS With the assumptions that (1) cylindrical symmetry exists, with the borehole as the axis, (2) heat flow is 1. The formation can be treated as though it is radially infinite and homogeneous in extent, as regards heat flow. 2. The presence of mud cake can be disregarded. 3. The effect of heat generated by bit action is negligible. 4. After mud circulation ceases, the rate of radial heat flow at the wellbore is negligible and is assumed zero for purposes of calculating temperature change with time. These assumptions are believed to be reasonable because of the agreement achieved between calculated and measured borehole temperature in a well, to be discussed later. Further, the agreement at the wellbore indicates that temperatures in the formation were also approximately correct. since it is the formation temperatures that determine wellbore temperature after ceasing circulations SOLUTION OF THE DIFFERENTIAL EQUATION AND ITS APPLICATION The solution of Eq. 2 for the temperature behavior during and after a single period of mud circulation is

By B. Post, F. W. Glaser

Three borides of zirconium have been reported: ZrB,l ZrB2,2,3 and ZrB12, 4 The phase relationships, ranges of stability, and some physical properties of these compounds are described. THE zirconium borides differ considerably from one another, particularly with respect to their relative stabilities at high temperatures in the presence of carbon.4,5 ZrB2 is extremelv stable at all temperatures (to the melting point) in contact with carbon. Both ZrB and ZrB12 tend to decompose at elevated temperatures when carbon is present in excess of about 0.5 pct. ZrB forms ZrB2 and ZrC, and ZrB12 decomposes to form ZrB2 and a boron-carbide phase. In the ZrB12 preparation, boron present in excess of that needed to form ZrB2 reacts with carbon, if present, to form carbides, in preference to forming the higher boride. Similar instability of borides, other than diborides of the transition metals of the 4th and 5th groups of the periodic table, has been reported.5,6 special precautions had, therefore, to be taken to minimize carbon contamination in the preparation of ZrB and ZrB12. Reactions below 1400°C were carried out by heating powdered mixtures of zirconium hydride and boron in quartz tubes in a hydrogen atmosphere. For temperatures above 1400°C, reactions were carried out in graphite molds. Carbon contamination from the molds was minimized by employing extremely rapid heating rates to reduce carbon diffusion into the body of the specimen, and by subsequent "shaving" of sample surfaces that had been in contact with the graphite dies. A heating rate of about 2000°C per min was obtained by heating the dies (by direct conduction), using a current of 40,000 amp. Although this type of heating cycle apparently allowed little time for "homogenization," close approximations to equilibrium conditions were usually obtained as a result of the very strongly exothermic reactions between zirconium hydride and boron at elevated temperatures. Reproducibility of results in repeated runs under varying conditions was taken as a reasonable indication that equilibrium conditions had been effectively achieved. Quenching was accomplished by rapidly immersing the entire die, containing the reaction products, in water. Specimens could also be cooled at controlled rates by adjustment of the rate of flow of water through the water-cooled copper electrodes in contact with the graphite dies.

Jan 1, 1954

By R. H. Fillnow, D. J. Mac

IT has been demonstrated that alloys in systems structurally analogous to steel undergo reactions during heat treatment similar to those of steel, and yet very little work has been done on such systems, particularly from the standpoint of isothermal transformation. The literature shows that most of the isothermal work has been done on copper-alu-minum alloys.'-3 One paper has been reported on a eutectoid copper-silicon alloy,' and one on iron-nitrogen alloys.' There are no published reports of previous investigations on the isothermal transformation of a eutectoid beryllium bronze. At least two papers have been reported on isothermal age-hardening phenomena in beryllium-copper alloys of lower beryllium content.5,6 No information on eutectoid transformations in this alloy is given in Rimbach's book, "Beryllium." ' This paper reports the results of an investigation on the isothermal transformation characteristics of a eutectoid beryllium bronze. Materials Used and Procedure The eutectoid beryllium bronze used in this investigation was centrifugally cast at the Ampco Metal Co., Inc.. Milwaukee. Wis. A master beryl-lium-copier alloy, purchased from the Brush Beryllium Co., was used to make the beryllium addition to the copper. The alloy was melted under a protective carbon cover and Centrifugally cast into a graphite mold preheated to 920°F. Some trouble with segregation of beryllium in the melt was encountered but was eliminated by violent stirring during melting and pouring. Since very thin .specimens were used for the various heat treatments, only one casting of the four prepared was necessary for the complete isothermal study, thus eliminating any discrepancies caused by a change in composition. The analysis of this casting was 94.035 pct copper and 5.965 pct beryllium. A slight trace of iron was found. The castings were sawed into specimens approximately 3/32 in. in thickness. The specimens were first given an "austenitizing" heat treatment at 788°C, well within the single phase beta region, in a vertical tube furnace. The specimens were then quenched individually into an adjacent molten salt bath and held there for a definite time, after which they were quenched into water at room temperature. Temperature was held to ± 2°C in the salt bath. No homogenizing treatment was given the specimens other than the soak at 788°C. Despite the fact that this was a cast alloy, no evidence of segregation or coring was found in the microstructure of any specimen, either as cast or as heat-treated. Upon completion of the heat treatment, the speci-mens were subjected to metallographic examination

Jan 1, 1951

Discussion of the papers of Messrs. Howe and Levy, Burgess, Crowe and Rawdon, and H. M. Howe, presented at the New York Meeting, October, 1913, and printed in Bulletin No. 78, June, 1913, pp. 1075 to 1136. ALFRED STANSFIELD, Montreal, Canada :-In Professor Howe's paper on the position of Ae3, he shows its industrial importance in determining the temperature to which steel should be heated. for "grain refining." Several years ago I carried out a research on the " burning " of steel and found that while hypo-eutectoid steel should be heated to Ae3, hyper-eutectoid steel that has been very much overheated must be reheated to the curve Sa on the carbon-iron diagram, in order to redissolve the pro-eutectoid cementite and thus, to obtain recrystallization and grain-refining. This treatment was necessarily only partly successful, and I should like to ask whether the conditions of heat treatment for refining overheated high-carbon steels have been ascertained. Speaking of the " Persistence of Solidificational Segregation " Professor Howe quotes an experiment made by Mr. Stead in which a bar of 1.2 per cent. carbon steel was immersed in a large mass of molten blast-furnace slag and allowed to cool with it. The steel was heated above the solidus (curve Aa on the carbon-iron diagram), but below the liquidus (curve AB). A drop of eutectic carbon-iron alloy, rich in carbon, had liquated out and hung below the bar. This agrees exactly with my own observations on the burning of steel, but my interpretation differs somewhat from that of Professor Howe. Professor Howe considers that the fusion of a part of the metal (rich in carbon) shows that the solidificational irregularities of composition had persisted through the rolling and heat treatment of the steel, as otherwise no part of the steel would have been sufficiently fusible to melt. It does not appear to me that such an explanation is necessary. A piece of steel represented in temperature and composition by any point within the area ABa will ultimately resolve itself into a solid phase and a liquid phase in equilibrium with each other. It may appear that if the steel were absolutely uniform no fusion could take place, but the smallest nucleus of high-carbon, fusible alloy would ultimately bring about this separation into liquid and solid phases, by diffusion of the carbon from the solid to the liquid areas. A nucleus sufficient for the purpose might be provided by an area of pro-eutectoid cementite which had not become diffused during the

Jan 12, 1913

By Samuel Daniels

The Al-Cu-Ni-Mg alloy is much benefited by heat treatment and, in such condition, is preferable to the Al-Cu-Fe-Mg alloy either as cast or as heat-treated, when both are reheated to temperatures of from 400" to 600" F. and compared, cold, with respect to strength and to hardness. The main differences between the alloys do not arise until the reheating temperature exceeds 400" F., above which they gradually soften. This softening is characterized metallographically by the appearance of intragranular precipitate. The Al-Cu-Ni-Mg alloy is weakest and least hard after being reheated at 600" F., the Al-Cu-Fe-Mg alloy at about 700" F. The former starts to reharden at 700" F., the latter at 800' F. There is a tendency for both alloys to lose strength and hardness with prolongation of time at the reheating temperatures of 500" and of 700" F. The Al-Cu-Ni-Mg alloy is considerably stronger and harder than the Al-Cu-Fe-Mg alloy when both are heat-treated and reheated for long periods of time at 500" and 700" F. Although the latter can be made harder initially than the Al-Cu-Ni-Mg alloy by suitable quenching and aging, heat treatment, in general, favors the retention of this initial (strength and) hardness after reheating to a much greater degree in the Al-Cu-Ni-Mg alloy. The percentage of elongation of the two materials in any condition is very small. Certain aluminum alloys have been used for pistons and other parts operating at temperatures up to 650" F. because of their lightness, high thermal conductivity and specific heat, fair strength and hardness at elevated temperatures, excellent machineability, and moderate resistance to corrosion. In the piston, the inertia forces and the bearing loading, and with this the vibration, are decreased. The piston head functions at a lower temperature (about 200" F. below one of cast iron), so that, because of the lessened tendency toward carbonization and detonation, higher compression ratios may be used. On the other hand, the aluminum alloys do not have the wearing and bearing qualities of cast iron; also they have a large coefficient of expansion, allowance for which must be made in the design. The range in analyses of aluminum alloys, used for pistons in domestic and in foreign internal-combustion engines is shown in Table 1. There seem to be three general types of composition; one containing about 10

Jan 1, 1926

By M. F. Hawkes

Austenite grain size has long been recognized by metallurgists as an important property of steels because of its influence on toughness, hardenability, ma-chinability and creep strength. Much research has been carried out on wrought steels to develop methods for disclosing austenite grain size, and to, provide knowledge of the grain-size characteristics of steels produced and heat-treated in various ways. Much less work in these fields has been done on steel castings; partly because, owing to the variable and complex microstruc-tures of cast steels, it is frequently very difficult to obtain results that are not questionable. The object of this investigation, therefore, was to establish, for a wide variety of carbon and alloy cast steels, suitable methods for disclosing austenite grain size and some knowledge of the grain-size variations to be expected. The steel foundryman is interested in austenite grain size at two different stages in the manufacture of his products; one is just after the steel is cast, and the other is after any reheating operation carried out above the critical temperature range. In both, it must always be kept in mind that the grain size being sought is actually a previous austenite grain size, and that the austenite no longer exists once the steel has cooled to room temperature. The prob-lem of the investigator is always to use evidence obtained from the structure at room temperature to deduce this previous austenite grain size, which was established the last time the steel existed at a temperature above the critical range. Since most steel castings today are at least annealed or normalized, the measurement of austenite grain size established by heat-treatment will he considered first. More than 50 commercially produced cast steels rcpresenting a wide variety of compositions and melting practices were used in the study. For each steel, grain-size determinations were made after each of three widely varying heat-treating schedules namely, one hour at I550'F, one hour at I700°F, and six hours at 1700 F. The list of steels and results obtained on them will be found in the appendix. (See page 530.) A discussion of the knowledge gained on the applicability of various test methods and on grain-size characteristics of cast steels in general follows. MEASUREMENT OF AUSTENITE GRAIN SIZE ESTABLISHED BY HEAT TREATMENT Methods Suitable foR Determining Austenite Grain Size after Heat-treatment at Ordinary Temperatures and Times It was desired, as one of the results of this investigation, to be able to recommend to steel foundries a method, or methods, of disclosing grain size best suited to accurate

Jan 1, 1948

By M. F. Hawkes

Austenite grain size has long been recognized by metallurgists as an important property of steels because of its influence on toughness, hardenability, ma-chinability and creep strength. Much research has been carried out on wrought steels to develop methods for disclosing austenite grain size, and to, provide knowledge of the grain-size characteristics of steels produced and heat-treated in various ways. Much less work in these fields has been done on steel castings; partly because, owing to the variable and complex microstruc-tures of cast steels, it is frequently very difficult to obtain results that are not questionable. The object of this investigation, therefore, was to establish, for a wide variety of carbon and alloy cast steels, suitable methods for disclosing austenite grain size and some knowledge of the grain-size variations to be expected. The steel foundryman is interested in austenite grain size at two different stages in the manufacture of his products; one is just after the steel is cast, and the other is after any reheating operation carried out above the critical temperature range. In both, it must always be kept in mind that the grain size being sought is actually a previous austenite grain size, and that the austenite no longer exists once the steel has cooled to room temperature. The prob-lem of the investigator is always to use evidence obtained from the structure at room temperature to deduce this previous austenite grain size, which was established the last time the steel existed at a temperature above the critical range. Since most steel castings today are at least annealed or normalized, the measurement of austenite grain size established by heat-treatment will he considered first. More than 50 commercially produced cast steels rcpresenting a wide variety of compositions and melting practices were used in the study. For each steel, grain-size determinations were made after each of three widely varying heat-treating schedules namely, one hour at I550'F, one hour at I700°F, and six hours at 1700 F. The list of steels and results obtained on them will be found in the appendix. (See page 530.) A discussion of the knowledge gained on the applicability of various test methods and on grain-size characteristics of cast steels in general follows. MEASUREMENT OF AUSTENITE GRAIN SIZE ESTABLISHED BY HEAT TREATMENT Methods Suitable foR Determining Austenite Grain Size after Heat-treatment at Ordinary Temperatures and Times It was desired, as one of the results of this investigation, to be able to recommend to steel foundries a method, or methods, of disclosing grain size best suited to accurate

Jan 1, 1948

By N. B. Davis

I. Introduction. II. Definition of Plasticity. III. Theories of Plasticity. A. Structure of the clay particles. (1) Fineness of grain. (2) Plate structure. (3) Interlocking particles. (4) Sponge structure of particles. B. Presence of hydrous aluminum silicates. C. Molecular attraction between particles. D. Presence of colloidal gelatinous matter. (1) Review of theory of colloids. (a) Suspension colloids. (b) Emulsion colloids. (2) Experiments. IV. The Formation of Clays. A. Residual clays. B. Transported clays. V. Summary. VI. Conclusion. I. IntRodUction While working with a number of very sticky cracking clays from western Canada the writer became interested in a study of the cause of the excessive plasticity. This led to a review of the rather bulky literature on the subject and the finding that plasticity still remains to be satisfactorily explained. Most of these investigations seem to have been carried on solely from the chemical viewpoint, the geologic aspects having been ignored. In nearly every case, the investigators have worked with the higher-grade clays of low to medium plasticity, such as the kaolins, the excessively plastic ones having generally been neglected, partly because of their limited application, yet the author believes that these are

Jan 1, 1915

By Norris Johnston, Carrol M. Beeson

For many years the permeability of reservoir sands has been measured by flowing air through a cleaned and dried core sample. This differs from the true reservoir permeability in one important respect: the rock particles in the reservoir are surrounded by interstitial water, not air, and their physical shape and condition of hydration are greatly dependent thereon. Permeability as defined must be measured with a single-phase fluid. Since no means exist for removing the oil and gas from a core sample by simply flowing water through it, the sample must be cleaned and then resatu-rated with water before testing. The present discussion attempts to show that after the cleaning process a considerably different permeability is determined with salt or fresh water than is obtained with air. The postulate is made that the salt-water permeability is probably closer to the true reservoir permeability than is the measurement with air. This is discussed in relation to both physically possible and economically feasible measuremerits. Data On more than 1200 core samples are given to show the nature of the effects observed, and a plea is made for others to consider water permeability measurements as a routine necessity, eventually replacing air permeability in regions where the differences are great. Introduction Permeability is defined1 as a "a measure of the capacity of a porous medium to transmit fluids, when there is no interaction between the solids and the fluid." Darcy's law of liquid flow in porous media states: kA?P QµL/A?L [I] where Q represents milliliters of liquid of viscosity p centipoises passed per second through a porous medium of permeability k darcys and cross-sectional area A sq. cm. under a pressure gradient ?P/L atmospheres per centimeter, the flow being understood to be laminar or viscous, not turbulent. This equation defines the darcy. For the case of radial flow in a uniform sand, such as usually exists in the neighborhood of a well in an oil or gas sand, Darcy's law becomes:2 Here h is the net sand thickness in feet, Pe and Pw are pressures at the drainage radius4 and well radius, and re and rw are these radii, respectively. The productivity index3 is defined as the barrels of liquid produced per day per pound per square inch mean pressure differential between ferences and well bore, and is thus: By further definition, the specific productivity index is: or in terms of practical units, and allowing for shrinkage:

Jan 1, 1945

By L. L. Handy

Two equations are compared for calculating imbibi-tion rates of water into porous media. A linear relation was found to apply when the square of the volume of water imbibed was plotted against time for water imbibing in natural rock materials. The effect of adsorbed organic materials on surface wettability was determined from measurements of rate of imbibition, water saturation and water permeability. INTRODUCTION An understanding of the mechanisms by which water displaces oil from porous media requires an understanding of the role of capillary forces in the displacement process. The capillary forces are determined by properties of the water-oil-solid surfaces. Consequently, some measure of the surface properties of reservoir rocks is necessary if displacement behavior in oil reservoirs is to be predicted from laboratory measurements. Several experimental procedures have been proposed to assign quantitative wettability indexes to reservoir rock surfaces. The most recent of these proposals are those of Bobek, Mattox and Denekas' and Amott.' These indexes are designed to show a continuous variation from the preferential oil-wet to the preferential water-wet systems. They require measuring some property of the rock which is a function of surface wettability. The quantities are measured on unaltered core material and compared with values obtained for known oil-wet and water-wet extremes on the same material. These methods are useful but are semi-empirical in nature. They have the disadvantage that the measured quantities may be functions of other variables in addition to surface wettability. Some understanding of the behavior of water-oil-rock systems may be obtained from studies of the simpler water-air-rock systems. Surfaces of a given wettability are more easily maintained for the latter. Also, certain approximations can be made for the flow of water and air which are not justified for water-oil flow. The theoretical and experimental work discussed in this paper is part of a study of the role of capillarity in the dynamic displacement of air by water. The expec- tation is that observations for the water-air system will lead to generalizations which are applicable to water-oil systems. THEORY Imbibition is defined as that process by which one fluid displaces another from a porous medium as a result of capillary forces only. Differential equations have been derived to describe the displacement of oil by water during imbibition, but these equations have not been solved.".' With certain assumptions, however, solutions are possible for water-air imbibition.

By H. H. Smith, G. J. Danek, M. R. Achter

Nickel, Type 316 stainless steel, and copper show decreases in reversed bending fatigue strength with increases in air and oxygen pressure. The leveling off of the log N us log p curve for nickel at 300" and 816°C is discussed in terms of a model in which surface saturation gives rise to a critical pressure. A metallographic study indicated that the effect of environment on crack initiation and growth rates is a function of temperature. It is shown that small amounts of impurities in inert gases may substantially reduce the fatigue strength of nickel. 1n a previous investigation1 of high-temperature bending fatigue in air and in vacuum, reversals in the effect of environment with changes in test conditions were atrributed to a mechanism consisting of two processes acting in competition. One of them, oxidation strengthening, has been demonstrated in a straightforward manner by measuring the rupture strength of a sintered nickel-oxide bond.' The detailed operation of the other process, facilitation of crack propagation by surface absorption of oxygen, is uncertain. Thompson et UZ advanced the possibility that the formation of an oxide layer prevented rewelding of cracks during the compressive part of the cycle, while Wads worth and Hutchings4 further proposed that the reaction of oxygen with atoms at the tip of the crack helped the metal to break there, a possibility which has also been advanced to explain crack formation in creep.5 A difficulty in arriving at a mechanism for this lowering of fatigue strength in air is the inconsistency of some of the data. For example, Wadsworth and Hutchings4 found an approximately linear decrease of log cycles-to-fracture (N) with an increase of log pressure-of-air (p) for copper and aluminum at room temperature, while snowden,' although finding a similar linear dependence for aluminum, observed a stepped curve for lead. Between l0-' and 10"1 torr there is a sharp decrease in the fatigue life of lead with increasing air pressure, but plateaus, signifying no dependence on pressure, are found at higher and lower pressures. If, as discussed elsewhere: these environmental effects arise from the adsorption of gas on crack surfaces, it would be expected that as the pressure is increased a saturation point would be observed, beyond which there is no further effect. To gain further information about the form of the log N vs logp curve, it was decided in the present investigation to repeat this type of measurement at elevated temperature on nickel, copper, and Type 316 stainless steel. An additional object of this program was

Jan 1, 1963