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By John B. Knaebel

The annual Jackling Award Lecture, sponsored by the Society of Mining Engineers of AIME, is intended to honor and commemorate one of the greatest mine-makers of all time, an innovator whose vision and determination made ore out of waste rock. He ushered in a new philosophy of mining which revolutionized some segments of the industry and notably affected our economy generally. To have been asked to given the Jackling Lecture this year is an honor that I scarcely deserve, but which I have accepted in humble gratitude. Although it is true that "many a mine has made the man," it is also true that while ore bodies are found, mines are made. That statement has been hammered into my skull many times by a man whose own career exemplifies its truth, a man who has been first and last a steadfast friend, who for a while was my boss and mentor, and who as one of this country's great mining geologists has laid his hand to the making of many a mine. I will doubly cherish the memory of the honor bestowed upon me because the presentation of the Jackling Award was made by this man. After the Lord made Dick Hunt he threw away the mold. Metals and other mineral products are essential ingredients of the social economy of the modern world. No well informed and thoughtful person can doubt that they will be indispensable to mankind as long as the fabric of civilization as we know it today can be held together. They are mined from ore bodies, which, unlike some other kinds of bodies, are physically inorganic and wholly incapable of reproducing themselves. Since in time they become mined out, an endless procession of new producers must be made by someone if the fundamental needs, to say nothing of the not so fundamental whims of an affluent society, are to be gratified. In saying that these mines must be made, I mean exactly that. It is my purpose to comment upon some of the problems we must cope with in making them. I shall try to avoid dwelling needlessly upon aspects of new mine development with which most people in this business are thoroughly familiar, and will attempt to discuss in a useful way some of the newer and more significant features of the economic climate in which these endeavors must be conducted. To begin with, an Almighty force put the ore into the ground, and selected the places in which to put it. Mankind was not consulted. Beyond that stage, however, mankind must find the ore deposits and contrive to extract their contents and put them into useful service. In doing this, he must use tools and techniques devised by man to overcome obstacles created in part by nature, but also in part by man himself. And most importantly of all, in making the tools and applying the techniques, he must use people. In the whole process the most pliable ingredient at hand-and also the most unpredictable-is people. Mines play out, and men die. As time goes on, we will have new mines, and they must be made by new men. The Environmental Dilemma It is perhaps timely to touch first upon the newest and in many ways the most potent consideration affecting the feasibility of developing new mines-the problems related to protection of the environment and to the

Jan 1, 1973

By R. L. Perrine

This paper presents results obtained by consider ation of unstable miscible displacement within a porous medium as a quasi-turbulent displacement process, that is, a process characterized by fluctuations around a simpler, mean flow behavior. The instability usually results from the miscible displacing fluid having a viscosity lower than that of the displaced fluid. Such instability is evidenced in suitable experiments by the presence of viscous fingering. A unified theoretical relationship applicable to both stable and unstable displacement is derived. Parameters in the theory are obtained by empirical means. Although the theory describes an idealized conceptual model, agreement with experiment is quite reasonable up to moderately adverse viscosity ratios, such as 17:1. INTRODUCTION Many experimental studies of miscible displacement in porous media have been conducted with prediction of reservoir displacement efficiency as the ultimate objective. Most such studies have utilized lower displacing than displaced fluid viscosity, scaled to potential reservoir fluid pairs. Theoretical approaches have been largely limited to unit viscosity ratios, however, in spite of the necessity for an understanding of the mechanism of the unstable, adverse viscosity ratio displacement process. An obvious reason is the difficulty of describing in mathematical form the viscous fingering characteristic of these conditions. Observation of experiments conducted with dyed fluids in transparent systems suggests an analogy between unstable miscible displacement and turbulent flow in a pipe. In both cases there are fluctuations around a simpler, mean flow behavior. An important difference is that flow behavior of interest in the porous medium is entirely transient, contained within a transition zone between displacing and displaced fluids. Transient behavior complicates description in that coefficients in the equations become variables rather than constants. In the study reported here, the analogy with turbulent flow has been used in treating unstable miscible displacement as a quasi-turbulent displacement process. The purpose has been to derive, even if restricted to an idealized conceptual model, a unified theoretical relationship applicable to both stable and unstable displacement. A relationship meeting these specifications up to moderately adverse viscosity ratios, such as 17:1, is presented. One fluctuation parameter in the theory and dispersion coefficients are obtained by empirical means. The idealized theory is compared with experimental results. UNSTABLE MISCIBLE DISPLACEMENT AS QUASI-TURBULENT DARCY FLOW The miscible displacement of one fluid by another within a porous medium is considered in this study. Flow conditions are such that Darcy's law is applicable. It is further assumed that, by the stability criterion proposed by perrine,1 initial flow conditions lie well within the regime of instability. Thus substantial viscous fingering is assured. We wish to show how this flow regime can be represented as quasi-turbulent. That is, the Reynolds number for the established flow conditions is below that for inertial or turbulent flow within a porous medium, and lies within the regime for which Darcy's law is valid. Yet flow behavior can be described as the combination of some relatively simple average result, and characteristic fluctuations that are superimposed on the simpler behavior. Stated another way, flow behavior includes the movement of layers of fluids with different velocities past or over one another. Such descriptions are characteristic of turbulent pipe flow. The fluctuations in the present case are a direct consequence of the local viscous fingering which accompanies the unstable displacement process. Should displacement become stable, fluctuations would die out. It is of particular importance to note that fluctuations such as these can interact in a way that contributes to material transport by the basic flow. The basic transport equations required for engineering calculations must be modified to reflect this fact.

By Avnulf Muan, P. V. Riboud

The effect of Cr2O3 on melting relations of iron oxide at oxygen pressures slightly above those prevailing in contact with metallic iron has been determined. Liquidus and solidus temperatures of wüstite are raised by the addition of Cr2O3, from 1385°C for pure iron oxide to 1420°C at the peri-ledie point where wüstite, iron chromite, liquid, and gas coexist in equilibrium under the atmospheric conditions used in the present investigation (CO2/H2 = 1). It has been shown in a previous investigation1 that the liquidus and solidus temperatures of iron oxide are drastically increased by the addition of Cr2O3 under the relatively oxidizing conditions prevailing in air. This is to be expected in view of the easy substitution of cr3+ for Fe3+ in crystalline oxide structures which are stable at high oxygen pressures, as demonstrated for instance by the completeness of the solid solution series Fe2O3-Cr2O3 at high temperatures (>1000°C). The stable phase of iron oxide at oxygen pressures approximating those prevailing when metallic iron is present is wüstite. This phase, although commonly represented by the simplified formula FeO for sake of convenience, in reality has a defect structure in which a part of the iron ions are present as Fe3+. It is to be expected that at least a limited amount of cr3 ions will substitute for Fe in this structure when Cr2O3 is added to wüstite. Depending on the ease of this substitution in comparison with that of cr3 substitution for Fe in iron oxide-Cr2O3 liquids under the same atmos- pheric conditions, the liquidus and solidus temperatures of wüstite may increase or decrease upon addition of Cr2O3. The methods of establishing these melting relations and the results obtained are presented in the following. I) EXPERIMENTAL METHOD 1) General Procedure. The quenching technique was used in this study. Oxide mixtures were heated at chosen constant temperatures in an atmosphere consisting of equal parts of CO2 and H2 until equilibrium was established among gas and condensed phases. The samples were then quenched in mercury to room temperature, and the phases present were determined by microscopic and X-ray investigation. 2) Preparation of the Mixtures. Starting mixtures were prepared by weighing and thoroughly mixing Fe2O3 and Cr2O3, or wüstite and Cr2O3. The Fe2O3 and Cr2O3 used were "Baker Analyzed" reagent grades, whereas wüstite of composition was synthesized by reducing Fe203 at 1325°C in an atmosphere consisting of equal parts of CG andHz. In order to avoid contact with crucibles, the equilibration runs were made with small pellets suspended from thin (-0.004-in.) platinum wires. Mixtures made from FOs and Cr203 were readily pelletized at a pressure of 500 atm in a 9.5-mm die, whereas mixtures made from wüstite and Cr203 when treated identically were too weak for handling. In order to obtain a stronger product, these pellets were sintered by heating for 2 min at -1200°C in the same atmosphere as was subsequently used for the equilibration runs (CG/H2 = 1:1). During the sintering, these pellets were wrapped in thin (0.0004-in.) platinum foil, which was stripped off before the equilibration runs. In order to reduce the time needed for equilibrium to be attained, preliminary equilibration of relatively large pellets of each mixture was carried out at -1200°C for 4 to 12 hr. Smaller pellets

Jan 1, 1964

By Kehsin Kuo

DIE steel with 1.47 pct C, 0.42 pct Mn, and 8.22 pct W contains Fe,C and WC in the annealed state and WC embedded in a martensitic matrix in the hardened state (quenched from 800°C). The presence of a large amount of undissolved angular WC carbide particles makes this steel extremely wear-resistant. However, the WC carbide begins to dissolve in the austenite when the heating temperature is raised above 1000°C, but this solution process is still not complete at 1250°C although partial melting has already taken place at the austenite grain boundaries. At 1325 C, this steel contains austenite and liquid metal but no WC carbide, Fig. 1. On quenching into brine, the liquid metal solidifies into a constituent which etches differently from the austenite-martensite matrix. This constituent appears lighter than the matrix at the 1250°C stage when etched with 2 pct nital but darker in Fig. 1 when etched with Murakami's reagent. Such a constituent has also been found in burnt tungsten high speed steels. This constituent has been isolated from the matrix by anodic dissolution in a 5 pct HC1 electrolyte and was found by the X-ray powder method to be iso-morphous with &manganese, Table I. The agreement in line intensity between these two phases indicates that there is no ordering in this new phase. The calculated cube edge of this phase is 6.395A, somewhat larger than that of ß-manganese (6.30A). Besides this phase, the anodic residue also contains free carbon from the matrix and amorphous WO,, although the X-ray powder pattern of the residue shows the presence of only this phase. This makes it impossible to determine directly the composition of this phase from the residue. Attempts to prepare an alloy containing only this phase have hitherto been fruitless. The following experiments serve to show that this phase can only be obtained by drastic quenching from the liquid state and that its tungsten and carbon contents must be lower than those in the Fe,W,C carbide. On air cooling from 1325oC, the liquid metal solidifies into a eutectic consisting of austenite and the Fe3W3C carbide. The nature of this dark-etching Fe3W3C has also been ascertained by X-rays. This eutectic becomes coarser on furnace cooling. The solidification of a specimen cooled for 5 min from 1325" to 1200°C and water quenched is still not complete and the fine eutectic is formed during the quenching. Since this new phase can be replaced by a eutectic of austenite and the Fe,W,C carbide, its composition must lie somewhere between the latter two phases. The special features of this phase are: 1—it can only form directly from the liquid state by drastic quenching and 2—it is a ternary phase of iron, tungsten, and carbon. The ß-manganese structure of this phase is also interesting. In the Periodic Table, iron and tungsten lie on opposite sides of manganese; a certain combination of these two elements, in the presence of carbon, may have average properties similar to those of manganese. It has already been established that the x phase, a ternary one of iron, chromium, and molybdenum, has the a-manganese structure."' A similar phase has also been found in the ternary system Fe-Cr-W and in the following binary systems: Re-Cb, Re-Ta, Re-Mo, and Re-W.6 References 1 G. D. Preston: Philosovhical Magazine 11928) 5. 7th Series, p. 1207. - K. W. Andres: Nature (1949) 164, p. 1015. 35. G. McMullin, S. F. Reiter, and D. G. Eberling: Trans. ASM 11954) 46, p. 799. ' J. S. Kasuer: Acta Metallurgica (19541 2, p. 456. • H. J. Goldschmidt: Trans. ASM (19541 46, p. 807. P. Greenfield and P. A. Beck: Private communication.

Jan 1, 1957

By T. E. Leontis

ONLY a limited amount of information has been published on the effect of thorium as an alloying ingredient in magnesium. McDonald13'2 showed that the addition of thorium in amounts up to 3 pct increases significantly both the ductility and the strength of rolled magnesium sheet. A recent publication by Sauerwald" indicates that thorium contributes strength and creep resistance at elevated temperatures to magnesium in both the cast and wrought states. The observations of Sauerwald are confirmed in this paper, which presents the results of a systematic study of the effects of compositional variation on the tensile properties, creep resistance, conductivity, density, and metallography of Mg-Th alloys containing up to 50 pct Th both in the sand-cast and in the extruded states. The effect of the addition of cerium to Mg-Th alloys is also reported. Experimental Methods The alloys studied in this investigation were pre.pared in small laboratory melts according to the melting practice described by Nelson4 as the crucible method. Electrolytic magnesium was used as the starting material to which the thorium* was introduced in metallic form. It was observed that the behavior of thorium in molten magnesium with respect to reaction with protective fluxes is similar to that of rare-earth metals. In order to obtain the highest possible alloying efficiency, it is necessary to use non-MgC1, fluxes and to apply all the precautions described by Marande6 for melting, alloying, and casting magnesium-rare-earth alloys. .Under these conditions it is possible to obtain an alloying efficiency in the range of 85 to 95 pct. The Mg-Th-Ce alloys were made with cerium of the highest purity available, not with the commercial product "misch-metal" which contains approximately 50 pct Ce. The total rare-earth content of this material was 97.1 pct, of which 0.3 pct was rare-earth metals other than cerium; the major impurities comprising the balance were iron and magnesium. The thorium and cerium contents of each alloy were determined chemically. Test bars 6% in. long with a 2% in. long reduced section of % in. diam were cast in sand molds using a four-baf pattern. In most cases, a 3 in. diam extrusion billet 10 in. long was cast in a steel mold from the same melt used for making the sand-cast test bars. For the higher thorium alloys, it was necessary to remelt the foundry scrap from the test bars in order to make the extrusion billets. The billets were scalped to a diameter of 2 15/16 in. and the ends faced to a length of 9% in. as limited by the size of the extrusion container. The alloys were extruded into % in. diam rod on a 500-ton direct-extrusion press. The details of the extrusion process are as follows: Billet preheat 925°F (2 to 16 hr) Container temperature 900 °F Die temperature 900°F Extrusion speed 10 ft per min Reduction ratio 36:1 Percent reduction 97.3 Alloys containing up to 10.2 pct Th were preheated at 925°F for 2 hr; all alloys with higher thorium contents were preheated for 16 hr. In order to establish the effect of extrusion temperature, billets of

Jan 1, 1953

By Henry Wurtz

Jan 1, 1880

By W. E. Howlett, R. L. Slobod

Jan 1, 1965

By E. Gybbon Spilsbury

(Pittsburg meeting, March, 1910.) THE recovery of gold and silver from their ores by means of the cyanide process has been so successful in the last few years that any radical improvement would seem impossible; yet the appliance to which I wish to call attention in this paper is really a radical departure from the methods now in general use. The most modern and approved of these, known as the all-sliming method, depends for its success on the grinding of the. ore so fine that practically 90 per cent. of it will pass through a 200-mesh screen. The slimes thus produced are then agitated and aerated in tanks of various types. The main object of this treatment is to insure such a thorough admixture of the pulp and solution that every particle of the ore is surrounded by a volume of solution sufficient to insure the dissolving of the whole of the gold- and silver-content. In addition, in order to expedite the action of the cyanide, and to oxidize such elements as would, if left in their active state, become cyanicides, air is blown in under pressure. In the Pachuca tank, this air is the active and sole means of agitation. The chief objections to all the methods off agitation used heretofore are the expense of operating and maintaining the mechanical devices for keeping the pulp in suspension, and the length of time required to obtain a fairly complete extraction. of the values. These objections are inherent to any method of agitation effecting a circulation of the pulp and the solution together, by giving the whole mass a circular movement, as in the low-tank system, or a vertical circulation, as in the Pachuca-tank system. In both methods, the particles of ore are kept traveling in the same direction as the solution, and with very little difference of speed; so that, while the whole mass is in

May 1, 1910

By G. W. Healy, L. F. Sander

was investigated by reacting thin discs of calcium oxide with gas mixtures of CO2, CO, and Son. Its value was 19,300 * 300 cal independent of temperature in this range. No solid solubility of sulfur in calcium oxide was detected within the limits of the experimental method and it is estimated to be below 0.025 pct by weight. The importance of lime in desulfurization is well-established but complete information on the pure phase equilibrium: CaO + 1/2 s2 = CaS + +02 [11 is not yet available. The goal of this work was to evaluate solid solubility of CaS in CaO and to determine the free-energy change associated with Reaction [I] at temperatures of 1400" to 1650°C. The equilibrium constant for Reaction [1] can be written: It is convenient to rewrite Eq. [2] in the form: where A = {Ps /PqJ1'2 has been referred to' as the "sulfurizing power' of a gas mixture. In this work, thin discs of CaO were suspended in a vertical tube furnace and exposed to CO + CO2 + SOz gas mixtures having known values of A. The samples were then analyzed for sulfur. As expected, X-ray diffraction confirmed that CaS was the only sulfur-bearing phase formed at the relatively low oxygen pressures used. EXPERIMENTAL PROCEDURE Reagent-grade CaCO3 was pressed in a 3/8-in.-diam pill die and prefired in air to produce CaO discs weighing between 0.004 and 0.01 g. Several discs were used to provide a suitable weight for chemical analysis while maintaining a large surface area to react with gas mixtures. These were placed in a platinum mesh basket and suspended in the gas stream in the hot zone of a vertical tube furnace. Desired gas mixtures were prepared from cp grade CO and CO2 and anhydrous grade SO2. The method of soap bubble displacement was used to calibrate capillary flow meters. While this gave excellent results with CO and Con, some problems with bubble insta- bility and soap film "drag" arose with the use of SO2 at low flow rates. Hence, frequent sampling and analysis of gas mixtures was carried out to insure proper control of the ingoing SOZ. The furnace used for gas:solid equilibration was a vertical mullite tube externally wound with 60 pct Pt-40 pct Rh wire having a diameter of 0.028 in. An inner tube of $ in. ID served as the reaction chamber having Pyrex ground joints sealed to the mullite to provide gas-tight connections at top and bottom. A Pt-Pt 10 pct Rh thermocouple was inserted into a protection tube adjacent to the sample basket to measure sample temperature during a run. Constant-temperature control to 2C was observed at any desired set point within the range of this investigation. This was accomplished by a control thermocouple imbedded in the furnace windings which served to actuate an electronic controller wired for high-low operation. The sulfur analyses of the solid samples were carried out using a stoichiometric combustion technique based on the method of Fincham and Richardson. Some analyses were done using a modified evolution method3 but these were used primarily to check the results of the combustion method. The results were in good agreement but the combustion technique of-ferred an advantage in economy of time and material. CALCULATION OF GAS EQUILIBRIA Heating a given mixture of CO + CO + SO2 to high temperatures gives rise to a large number of product species. The details of calculating the partial pressures of these products of interaction and dissociation can be found in several references4,5 and need not be repeated here. The thermodynamic data selected for the major species in the gas mixtures are shown in Table I. Equilibrium constants from these reactions were combined with oxygen, carbon, and sulfur balances and a computer program written to facilitate the calculations. Some early difficulties in reproducing experimental results were finally traced to the effect of atmospheric pressure changes. No reference to consideration of this question had been found in the

Jan 1, 1969

By D. R. Bomberger, M. Deul

An automatic microscope has been developed to measure the size and populations of macerals and minerals in fine coal particles. Differentiations are made on the basis of reflectance. This study is concerned with the size distribution of pyrite inclusions. Most fine coal cleaning processes do not remove pyrite as completely as a straight density separation. Consideration of the pyrite size distribution is suggested as a more accurate method of predicting sulfur reduction by a cleaning process. Analyses of float-sink results and Deister Table operation in terms of pyrite particle sizes are made. Other potential applications of automatic microscopy are also discussed. In the study of coal cleaning by pseudo-gravity methods (tables, cyclones, hydrotators, etc.), it is customary to express the deficiency of the process relative to float-sink measurements in some form of "percent efficiency". These have been calculated in many different ways in order to provide a means of extrapolation of test data to untested coal feeds. This works quite well with coarse coal (here defined as larger than one-eighth in.). Large, heavy pieces of refuse are rejected rather thoroughly by many cleaning processes, and the efficiencies relative to float-sink results approach 100%. In coarse coals, the difference in the sizes of fine pyrite is masked by the averaging effect of the large coal envelopes. The existence and fate of fine pyrite is often hidden in the larger numbers representing coarse (and easily removed) pyrite. There are cases where it is desirable to subject fine coal (here defined as smaller than one-eighth in.) to cleaning processes. This is true of naturally occurring fie coal streams and in cases where comminution followed by cleaning is necessary to provide the quality desired. It is in this area that the use of process efficiency relative to float sink fails. Percent efficiencies decline to very low figures, and equivalent separation densities become a matter for conjecture. The net result is that an attempt to estimate product quality from float-sink data and a process efficiency figure often results in large errors. This work represents a new approach to the problem by observing the fate of pyrite particles, grouped by sizes, in a cleaning process. Examples are given which relate pyrite size measurements to float-sink washability and to a Deister Table operation. EXPERIMENTAL WORK The automatic microscope used here has been previously described.' Briefly, it is an instrument which measures the diam of particles (as defined by a range of reflectance) exposed on a polished specimen. The instrument scans a sample briquet, measures diam of exposed pyrite inclusions and records the data by size intervals. Preparation of the briquet is critical, due to the automatic nature of the process. The principal requirement is for an optically flat surface, so that focus will be maintained as the motorized stage moves the specimen under the lens. The various substances present must be polished to the same level in spite of the wide difference in hardness between coal substance and minerals. The polishing method must not pluck out small grains, which might be rather loosely held by surrounding coal. The procedure below was successful in producing all of these results. The coal sample (8 mesh or finer) was riffled (a micro vibrating riffle is essential) to obtain 15 g. One volume of Lucite powder (thermosetting plastic mounting medium) was mixed with three volumes of coal and mixed thoroughly. A briquet was formed using a cylindrical die (1 1/4 in. diam) at 3000 lb per in.* pressure and 150°C. Preparation of the optical surface was accomplished on automatic polishing equipment, according to the following schedule: 1) Grinding: 3 stages, using successively fine Carborundum grits of 180, 320 and 600 mesh. 2) Rough polishing: 2 stages, using 3µ alumina on cast iron, followed by 1µ alumina on lead. 3) Fine polishing: 2 stages, using 0.5µ alumina on napless broadcloth over bronze, followed by 0.1µ alumina on silk filament cloth over bronze. Raw data were corrected for the fact that the chords

Jan 1, 1964

By Robert K. Waits

The resistivity, temperature coefficient of resistance, stability at 200°C, and structure of annealed Si-Cr films have been studied as a function of film cowposition. Colorimetric analyses of the films indicated 20 to 60 wt pct Cr over the resistivity range of 1 to 10-4 ohm-cm. Resistors were delinealed by photolithographic etching, contacted with aluminum, and mounted in TO-5 packages. The average temperature coefjicient of resistance, a, between 50" and 150°C is related empirically to the film resistivity, p, by a (ppm per °C) = -2900 - 1140 log p (ohm-cm). Near a resistivity of 3.5 x 10-3 ohm-ctn (30 at. pct Cr) , the temperature coefficient changes from positive to negative. The temperature at which the resistance maximum occurs increases with decreasing resistivity. The conduction mechanism between —196" and +360 C appears to have two components—an activated conduction dominating a1 lord chromium content and higher temperatures and a linearly temperature-dependent conduction dominating at high chromium content and low temperatures. Electron diffraction patterns showed the films to consist of Cr3Si and amorphous silicon. Resistors thicker than 150A typically drifted less than 0.6 pct after 5000 hr stor-age at 200°C. Resistors as thin as approxinzately 80°A that had a silicon overlayer showed similar stability. Such films had sheet resistivities of 10 kilohm per square and temperature coefficients near —600 ppm per "c. THIN-FILM resistors having sheet resistivities greater than 1 kilohm per square are required in many integrated circuit applications. A process for producing such resistors should be compatible with present circuit fabrication and packaging, including exposure to maximum temperatures of 300° to 570°C during contact alloying, die-attach, and ceramic package sealing. In addition, the resistors should have a low temperature coefficient of resistance and change less than 1 pct during storage at temperatures up to 200°C for 1000 hr or longer. Reports of work beginning in 1949 at Battelle Memorial Institute indicated that, among other materials, the silicides of chromium, Cr-Co, molybdenum, tantalum, and tungsten were worthy of further study as high-stability resistive films. Later work on Si-Cr films vacuum-deposited by flash evaporation from a tungsten boat1 led to the conclusion that Si-Cr had "excellent properties for development into high- stability ... resistive elements". More recently, Si-Cr films have been vacuum-deposited from a tantalum boat,2 from a BeO-coated tantalum boat,3 by flash evaporation from an electron-beam-heated tungsten Source,4 and by triode sputtering.5 This paper describes a study of Si-Cr films vacuum-deposited by electron-beam heating of Si-Cr alloys. Resistivity, temperature coefficient of resistance, stability at 200°C, and structure were studied as a function of film composition. EXPERIMENTAL PROCEDURE Film Deposition. Si-Cr sources were prepared by vacuum-melting pressed pellets of mixtures of powdered chromium (99.9 pct, -100 mesh, Lunex Co., Pleasant Valley, Iowa) and silicon (99.99 pct, -325 mesh, United Mineral and Chemical Co., New York). A permanent-magnet focused electron-beam gun having a water-cooled copper crucible was used to heat the pellet until complete melting occurred. The mixture was then kept molten for an additional 5 min. The alloy thus formed was used as the source material for Si-Cr film depositions using the same electron-beam gun. Pieces of semiconductor-grade silicon ingots were used as sources for silicon depositions. The films were deposited within a standard vacuum evaporator having an 18-in.-diam Pyrex bell jar, a liquid-nitrogen-cooled trap, and a Consolidated Vacuum Corp. 6-in. diffusion pump containing Dow Corning DC-705 pump fluid. The substrates were thermally oxidized* silicon wafers 2 to 2.5 cm in diam. The

Jan 1, 1969

By P. Pietrokowsky

lntermediate phases in the binary metal alloy systems Ti-Zn, Ti-Hg, Zr-Zn, Zr-Cd, and Zr-Hg have been investigated by X-ray powder diffraction methods. Gamma-TiBHg and Zr3Hg have a beta tungsten structure; TiHg and ZrHg are analogous to ordered AuCu (L l0 type); delta-TixHg and ZrHg, are isomorphous with ordered AuCu,, (L l2 type); ZrZnz is face-centered cubic (C 15 type); TiZna crystallizes in a C 14 type structure; and Zrll-, occurs in cubic and tetragonal modifications in which a random distribution of atoms exists. Additional information for the intermediate phase TiZn, is presented. THE physical metallurgy of titanium and zir-conium when alloyed with elements in subgroup 11-B of the periodic table has received comparatively little attention in the literature. Laves and Wallbaum' reported the existence of several intermediate phases in the system Ti-Zn. They identified TiZn as a CsCl (B 2) structure and described TiZn, as an ordered Cu,Au (L 1, type) crystal structure. Gebhardt' investigated very limited zinc-rich regions in the systems Ti-Zn and Zr-Zn. Phase relationships in these partial constitution diagrams were deduced from the results of thermal analysis and metallography, X-ray results being inconclusive. The first zinc-rich intermediate phase in either system was not identified. More recently, Anderson, Boyle, and Ramsey" have had occasion to refer to a partial Ti-Zn phase diagram (unpublished data) covering the region 0 to 10 atomic pct Ti, wherein two intermediate phases corresponding approximately to TiZn,, and TiZn,, were indicated. The metals used in this investigation and typical analyses as determined by the suppliers are given in Table I. The iodide-process zirconium which was used was not hafnium free. Experimental Methods Alloy Preparation: The relative ease with which titanium and zirconium combine with other elements places extensive restrictions on the methods by which alloying can be achieved. The direct current electric arc melting procedure was found to be unsatisfactory for preparing these binary alloys and for this reason solid state diffusion methods were utilized. The iodide process titanium and zirconium rods were filed by hand. Fines were subjected to a magnetic separation and classified as to size as either finer than 200 mesh or coarser than 200 but finer than 80 mesh. Powder mixtures of various binary combinations of titanium and zirconium with zinc and cadmium were compacted in a 12 mm diameter die and then placed in Vycor or quartz vials which were evacuated and sealed. The mercury alloys were prepared in a similar manner; the proper weight of liquid mercury being placed in a vial with minus 200 mesh titanium or zirconium. These ampoules were then heat treated at various temperatures and times depending on the characteristics of the particular alloy system. Upon completion of the heat treatment the ampoules were cooled rapidly and the specimens weighed. The loss in weight varied from 1 to 3 pct (except as noted in the text) and for this reason the calculated compositions were accepted as a fair measure of the actual alloy compositions. X-Ray Dif raction Techniques: Powder diffraction experiments were conducted at room temperature with a 143.2 mm Debye-Scherrer X-ray camera in which the film was placed in the Straumanis arrangement. When copper radiation was used, a nickel filter was employed, and for the cobalt target a filter of iron oxide was used. In both cases, the /i? filter was placed between the X-ray target and the powder specimen which was mounted on a glass fiber. The lack of information based on optical metallography placed reliance on the results obtained from the X-ray experiments for determination of

Jan 1, 1955

By B. L. Landrum, P. B. Crawford

The rise of a new laboratory model for studying tran-sient fluid flow problems, is described. The theory of he model is based on the analogy between the equa-ions which describe the flow of compressible fluids in porous media and the conduction of heal m solids. The "thermal model" is used to study the Iransient Pressure distributions during fluid displacement in five Standard well patterns. The studies indicate quantita-ively how die pressures change with lane at any poit in the reservoir. It was found that, in general, the equipressure contours are more radial about the wells early in the program than after steady-state conditions are established. In addition two studies of fractured five-spot patterns are presented to illustrate the opplicability of the model for a wide range of reservoir conditions. INTRODUCTION In reviewing the oil recovery techniques involving fluid displacement it was found that a very considerable quantity, if not all of the fluid within the pattern, mav be displaced (during a time-dependent or transient period rather than the usually applied steady-state conditions. Since the oil is displaced during a transient period, it was believed desirable to study the nature of the transients for several of the commonly used fluid displacement patterns. A literature review indicated fairly complete studies of the streamlines and isopotentials for steady-state: yet few data were available on the transient displacement process. Several simulated steady-state multiwell displacement programs were carefully mapped by Foster and Lodge on an electrical conducting sheet as early as 1875, hut the more recent work of Muskat provides the best source for steady-state phenomena having direct petroleum production applications. In order to study the transient phenomena a new lab-oratory model was developed which is based on the similarity of the equations that describe the flow of fluids and the conduction of heal. The equations describing heat flow were developed early in the nineteenth century. Regorous solutions tor most transient problems are in gcneral complex and usually limited to fairly restricted houndary conditions.2 the thermal model may be used to obtain approximate solutions for many transient problems without some of the restrictions required for analytical solution. The application of the analogy between heat conduc-. tion in solids. electrical flow and fluid flow, in a porous media has long been recognised. Techniques originally developed for the solution of heat conduction prol?lems have in many instances proved useful in ohtaining solutions to fluid flow prohlcms. Ir is not surpri5ing then that a heat conduction or thermal model might bf usefull in sludying transient fluid flow. As an example ol the problems suitable for the study with the thermal model, transient pressure distributions during fluid displacement programs are studied for the five-spot, line-drive, staggered line-diibe. seven-spot 211 two fractured patterns. A model was constructed of an iron plate of geometry similar to the reservoir patterns under study. The principle of mirror images as suggested by Bryan was employed in shaping the model. A temperature differential was obtained by soldering two sections of copper tubing to the plate at points analogous to the well location in the field. Hot and cold water was passed through the tubing. The temperature distribution, analogous to the pressure distribution in the field. was obtained a a function of time from thermocouplcs distributed over the plate.

By J. L. Jorstad

The hypereutectic Al-Si alloys 390 and A390 have wear characteristics superior to any of the more common aluminum casting alloys. This excellent wear resistance, coupled with good mechanical properties, high hardness, and low coefficient of thermal expansion, has made these alloys candidates for the all-aluminum internal combustion engine, the application for which they were primarily developed. A390 alloy is intended for sand and permanent mold casting, and requires a lower iron ccxltent (0.5 pct maximum) than its die casting companion 390 alloy (0.6 to 1.1 pct Fe). The 17 pct nominal silicon in these alloys provides sufficient quantities of the hard primary silicon phase to assure a high degree of wear resistance, yet little enough of this phase, to minimize the casting and m achining problems associated in the past with the hypereutectic alloys. The mechanical and physical properties of 390 and A390 alloys are competitive with the best of the common aluminum casting alloys with the exception of low ductility. This low ductility is not considered a detriment, since in many applications these alloys replace cast iron. Corrosion resistance is similar to the standard alloys 380 and 333. Numerous parts, of varying degrees of complexity, have been cast of alloys 390 and A390, including engine blocks and heads. The deviations from normal good handling and casting practices that are required for the alloys are those associated with the presence of the primary silicon. To obtain optimum strength and machinability, the molten alloys should be treated with phosphorus to refine the primary silicon phase. Care must be exercised to properly control the pouring temperature and rate in order to minimize primary silicon growth and segregation. AUTOMOTIVE engineers have for years recognized the advantages of lighter weight and better heat transfer that aluminum could offer as material of construction for engine blocks. These advantages alone, however, are insufficient justification for acceptance of aluminum engines by the automotive industry. To be acceptable, an aluminum engine first must pass all tests to which cast-iron engines are subjected, and, in addition, must represent no cost penalty in manufacturing. None of the conventional aluminum casting alloys have sufficient wear resistance to withstand the tests to which cast iron is subjected as a cylinder bore material, and recent attempts to manufacture aluminum engines with cast-in-place iron sleeves could not compete costwise with the equivalent cast-iron engine. The prerequisite to an acceptable aluminum engine seemed to be a casting alloy with a high degree of wear resistance that would eliminate the necessity of cast-in iron sleeves or liners. The hypereutectic A1-Si alloys 390 and A390 have met this challenge.' Combined with modified pistons and a patented cylinder bore surface finish,2 the performance of these alloys in more than 1300 cold start and cold scuff tests, over 3500 hr of dynamometer wear and endurance tests, and more than 350,000 miles of road tests leaves no doubt as to the feasibility of running on bare aluminum cylinder bores. ALLOY DEVELOPMENT The desirable characteristics of the hypereutectic A1-Si alloys have been recognized for some time. Their excellent wear resistance and low coefficient of thermal expansions have been used to advantage in the production of diesel-type pistons in Europe and to a limited extent in this country. The major contributor to the good wear resistance of these alloys is the extremely hard primary silicon phase in the microstructure. Although the percent silicon used experimentally has been much higher, most commercial alloys have ranged from about 14 to 25 pct. Generally, as the silicon content is increased, wear resistance increases, and the coefficient of thermal expansion decreases. With increases in silicon, however, machinability and cast-ability tend to suffer. The development of a satisfactory aluminum engine block material at Reynolds involved the testing and evaluation of many alloys, covering a wide range of silicon contents and numerous variations of iron, copper, magnesium, manganese, and other alloying additions. Castability, mechanical and physical prop-

Jan 1, 1969

By Henry Wurtz

HERACLITUS, a sage of antiquity, called the dark philosopher, who refused a throne, preferring a hermit's cell, propounded, twenty-four centuries since, the maxim : [ ] War (or strife) engenders all things. This, though probably intended by Heraclitus to apply especially to the internal forces of nature, is often said, with equal reason, of the affairs of men. Controversial strife, whether fortunately or unfortunately, is a crucible through which all new discoveries in science, and all technical applications of science must pass-a test which they must all endure before they can become so vitalized as to germinate, so to speak, take root in the human mind, grow tip, and overspread the earth. The greater the number and the power of the elements arrayed against such a growth, and of the influences hostile thereto, the greater should be the inherent vitality of the germ, the more strenuous, skilful, and persistent its cultivators and upholders. During the decade last past we have had, in spite of the severe stringency of the times, an active growth of this kind in progress, whose prospective importance it would now be difficult to overrate. This is the movement which has for its motive the idea that, generally speaking, fuel should be gaseous in form, and which has for its goal the introduction into general public use of gaseous products, made by cheap and rapid processes and on a gigantic scale, distributed throughout our cities and towns in distribution-systems which shall be proportionately gigantic, and sold at prices which will bring such fuel within the means of the poorest householders. Personally, for ten years past, the writer has never failed, on occasions public or private, to urge his belief that the realization of this idea, deemed by him a certainty of the future, will bring about

Jan 1, 1880

By W. Hodge, A. F. Haskins, R. M. Evans

Refractory boron compounds are shown to resist corrosion by molten zinc. Coatings were made from ferroboron and manganese boron by several methods: welding, hard facing, and pack diffusion; and techniques of coatings are very important. Mechanical failure of the diffusion coatings can be partially eliminated by applying them to type 416 chromium steels rather than carbon steels. Welded coatings made with a tungsten arc are better than those made by other welding methods. Sintered compacts of mixtures of iron and chromium borides developed strengths of about 30,000 psi. They resisted corrosion of zinc at 600°C and oxidation at higher temperatures. MOLTEN zinc rapidly attacks common metals of construction. The rate of attack is greatly amplified at temperatures over 500°C. However, there are many applications where it is necessary, or at least advantageous, to use metal construction, such as for galvanizing pots and zinc smelter equipment. Condensers and agitator linings used in continuous zinc smelting processes are eroded and corroded by violently agitated molten zinc. Pumps, tubing, and other parts to permit transfer of molten zinc would be desirable in many zinc handling operations. The development of metallic materials which are resistant to molten zinc also might permit its use as a heat-transfer medium in modern high efficiency power plants. In reviewing previous work on zinc-resistant metals, it was evident that most of the available data refer to corrosion at temperatures just above the melting point. Tungsten and high Mo-Fe alloys containing more than 80 pct Mo resist corrosion by molten zinc.', ' Imhoff states that because chromium is not attacked by zinc, chromium plating improves galvanizing pot life." Nitrided cast iron has been used in die-casting machines' where the molten alloy is 96 pct Zn and 4 pct Al. Other metals that have been suggested as of possible value for zinc resistance are columbium, tantalum, and titanium.V he corrosion of steel by zinc has been discussed by many authors,z, "2 None of these materials resisted corrosion by molten zinc at the temperatures used in this investigation, i.e., 600" to 700°C. I. Exploratory and Resistant-Coating Tests To make a study of zinc-resistant metals, exploratory tests were conducted to sort out roughly those materials which seemed most worthy of further development. This preliminary study included both solid metals and surface coatings. Exploratory Corrosion Tests in Molten Zinc—In initiating a study of this problem, numerous samples of metals and alloys were evaluated by dynamic corrosion tests made at 440 °C. Each metal was tested in its most readily available form, which, because of shape and size, was often unsuited for standard corrosion testing. The equipment used for dynamic corrosion testing of all materials, except in a few initial experiments, is shown in Fig. 1. Samples were tied to tungsten rods with tungsten wire, immersed in the zinc, and revolved at 12 rpm about the circumference of a 4 in. diam circle. Speed of travel was about 12.5 fpm. The molten zinc was protected from oxidation by an atmosphere of carbon monoxide, or, at a later period in the program, by a 9:1 mixture of nitrogen and hydrogen. When the effect of both molten metal and atmosphere was desired, the zinc level was lowered. Special high grade zinc was used in all of the corrosion tests. Corrosion losses for materials tested are shown in Tables I and 11. Whenever the materials were available in suitable form, the losses are reported in mils per year. It was recognized that weight losses on irregularly shaped specimens did not show the corrosion rate but the important consideration—poor resistance to corrosion—could be determined readily.

Jan 1, 1956

By R. I. Jaffee, E. S. Bartlett, D. N. Williams

The oxidation resistance and fabricability of molybdenum powder-metallurgy alloys containing up to 20 wt pct nickel in the form of a fine dispersion were studied. The effects of dispersion distribution and of amount and type (pure nickel or nickel compounds) of additive were considered. A definite improvement in oxidation resistance was observed, but the amount of improvement fell far short of requirements for an oxidation-resistant alloy. Fabricability by hot-rolling was limited. INVESTIGATIONS at Ohio State university1 and Battelle Memorial institute2 have indicated that the attainment of a truly oxidation-resistant molybdenum-base alloy for high-temperature service, with the coincident retention of high-temperature strength and toughness, will be a difficult, if not impossible, feat. Among the most promising alloy systems studied was the molybdenum-nickel system shown in Fig. 1. Alloys in this system show good oxidation resistance at temperatures well above the normal melting point of molybdenum trioxide because of the regenerative formation of protective nickel-molyb-date (NiMoO4) on the surface of the alloy during exposure at high temperatures. Unfortunately, the required alloy content to obtain oxidation resistant molybdenum-nickel alloys far exceeds the limits of the narrow terminal solid-solution boundary, and closely approaches the peri-tectic composition involving the brittle intermetal-lic compound, MoNi. Because of the peritectic reaction, fusion alloys are characteristically brittle, since the MoNi phase, being the last to form, tends to occur as a continuous grain-boundary network around each molybdenum grain. An added problem with this system is that the protective nickel-molybdate layer formed at high temperatures is unstable, undergoing one or more allotropic modifications upon cooling. The volume changes associated with this modification during cooling are such as to cause violent expulsion of the protective molybdate coating from the surface of the piece. This is, of course, quite detrimental in material exposed to cyclic temperature during service. The research reported here resulted from an attempt to overcome the aforementioned problems associated with the molybdenum-nickel alloys primarily by controlling the amount, size, and location of the MoNi phase. By using powder-metallurgy techniques of alloy preparation and sintering the blended molybdenum-nickel powder mixtures at a temperature below the peritectic reaction temperature, molybdenum-nickel alloys with a randomly dispersed MoNi phase were produced. MATERIALS The molybdenum powder used in the preparation of alloys for this study was commercial hydrogen reduced, minus 325-mesh (Tyler) material. Master alloys, by which nickel was added, consisted of elemental carbonyl nickel powder (12 average particle size), an alloy of 50 wt pct Ni-50 wt pct Cr, and compounds of MoNi, Ni,Si, and Ni2B. The nickel-chromium powder was prepared by filing a vacuum-arc-melted ingot, separating impurities magnetically, and screening to minus 200-mesh. The MoNi powder was prepared by pulverizing a hydrogen-sintered powder-metallurgy compact to minus 325-mesh material. Further sizing of this powder for study of effects of dispersion distribution was accomplished by air classification of the minus 325-mesh powder. The Ni2Si and Ni2B powders were prepared by crushing vacuum-sintered powder-metallurgy compacts to minus 200-mesh material. EXPERIMENTAL PROCEDURE Specimen Preparation—The general procedure followed in the preparation of alloys for testing is described as follows: 1) Blending. Either hand-blending of the powders in a ceramic mortar and pestle for X 1/2 or mechanical blending was used. Hand-blending seemed preferable for optimum homogeneity. 2) Compacting. Most of the blended powders were pressed to green compacts at pressures of 15 tons per sq in. in a single-action die, although some were pressed at 80 tons per sq in. A 1 X 11/4 X 1/4-in. green compact was the standard size for oxidation test material, and 7 X 1/2 X 1/4-in. compacts were used for fabricability test specimens. 3) Sintering. All green compacts were sinteredat

Jan 1, 1959

By P. Duwez, C. B. Jordan

The phenomenon of the change of volume of pressed powder compacts upon sintering is well known in the field of powder metallurgy. Depending upon the metal or metals involved and the pressure used in forming, a compact may, in the course of time of sintering at a given temperature, expand mono-tonically, contract monotonically, or first show a volume change of one sign followed by a change of the opposite sign. It is clearly desirable to have accurate knowledge of the magnitude and sign of the change in dimensions to be expected in any given case, both from the point of view of direct usefulness in the fabrication of parts by powder metallurgy, and from the longer range viewpoint of elucidating the fundamental mechanism of metallic sintering. The present study was therefore undertaken as a first step in acquiring systematic and reasonably quantitative knowledge of the change in density of metal powder compacts during sintering. For practical reasons, copper was selected as the material to be studied first, and its densification followed as a function of temperature and time of sintering in hydrogen and in vacuum. Experimental Procedure The copper powder used was that designated by the manufacturer (Metals Disintegrating Co., Elizabeth, N. J.) as MD-151. This powder was sifted through Tyler standard screens to separate the fraction having particle size range between 200 mesh and 325 mesh, and this fraction was used in all the subsequent work. Compacts weighing about 10 g were then pressed in a 1 in. diam round die, using a pressure of 20,000 psi throughout. Sintering was carried out in commercially built electric furnaces in which the resistance windings are so disposed as to produce a nearly uniform temperature along the axis of the furnace for a length of about 18 in. centrally located. In order to be able to sinter in a controlled atmosphere, a 2 in. stainless steel tubing was inserted in the furnace. Each end of the tube was cooled by a water jacket about 7 in. long, and closed with a rubber stopper. The hydrogen used for one series of specimens was purified as described in Ref. 1. For the other set, a pressure of about 0.5 mm Hg was maintained during sintering by a Welch Duo-Seal Pump. The specimens were heated on square trays made of stainless steel. In placing specimens in the trays, a thin even layer of powdered aluminum oxide was first sprinkled on the bottom of the tray. A copper guard disk about half the thickness of the specimen was then placed in the tray and covered with a second layer of alumina. The actual specimen was then set on the guard disk, and a final coat of alumina sprinkled over the specimen. This technique was evolved for sintering the specimens in such a way as to reduce the influence of unknown extraneous factors to a minimum. If the specimen is placed directly on the tray and sintered, it is found that the resulting shape is that of a frustum of a cone, rather than a section of a right circular cylinder, since friction with the tray prevents the bottom of the specimen from contracting at the same rate as the top. In the arrangement used in these experiments, the guard disk provided a support which shrank at the same rate as the specimen, and the alumina powder reduced to a minimum friction between guard disk and tray, and between specimen and guard disk. The procedure followed in sintering consisted of bringing the furnace to the required temperature, and then inserting the specimen into the central heated portion of the furnace tube in one of the two atmospheres used. At the end of the heating period, the specimen was cooled by bringing it into a portion of the furnace surrounded by a water jacket. These manipulations were carried out without opening the furnace, by means of rods which were attached to the trays and operated through a sliding seal in the rubber stopper. The progress of densification of the copper compacts was studied at 1300, 1400, 1500, 1600, 1700, and 1800°F. At each of these temperatures, a specimen was allowed to sinter for each of the following time intervals: M, 1, 2, 4, 8, 16, 32, and 64 hr. The thickness and diameter of each specimen were measured with micrometers before and after sintering, and each was weighed on an analytical balance after sintering. Results The techniques described in the preceding section were found to give satisfactory results. The specimens were not detectibly warped after sintering, and were usually of uniform diameter (that is, truly round) to within 0.001 in., a very few showing a variation in diameter of ±0.002 in. All specimens were found to have the same diameter

Jan 1, 1950

By R. M. Waterstrat, E. C. van Reuth

BINARY- alloy phases having the A15-type crystal structure have been described as occurring at a simple and more or less invariant stoichiometric composition (A3B) which corresponds to the relative number of atoms occupying each of the two crystallographi lattice sites in this structure.1,2 It is frequently assumed, therefore, that each crystallographic site is occupied exclusively by one kind of atom. In most cases, however, there have been insufficient experimental data to establish whether atomic ordering is, in fact, complete. Recent studies have shown that binary A15-type phases are sometimes stable over an appreciable composition range3''* and, occasionally, the composition range of stability does not even include the "ideal" A3B stoichiometric composition.5-7 We have observed the existence of nonstoichiometric A15-type phases in the binary systems Cr-Pt and Cr-Os. This has not been reported in previous work on these alloy systems.1,8-11 A series of alloys, each weighing approximately 30 g, was prepared by are-melting in an Ar-He atmosphere using 99.999 pct Cr, 99.999 pct Os, and 99.99 pct Pt as starting materials. Each alloy was melted four times with a total weight loss of less than 1 pct. The stoichiometric (A3B) alloys were sealed in evacuated quartz tubes and annealed at 1200°C for periods of time ranging from 3 days to 2 months. Examination of the alloy microstructures revealed that little change had occurred over this time interval and it was therefore assumed that the microstructures were fairly representative of equilibrium conditions. No evidence of contamination was observed although there was apparently some loss of chromium which was confined to a thin layer at the surface of the specimens. The quartz tubes were quenched from the annealing temperature into cold water. X-ray diffraction and metallographic examination of the stoichiometric alloys revealed an estimated 10 to 30 pct of second phases which were tentatively identified as phases previously reported in these binary-alloy systems.8-11 A second series of alloys was prepared by mixing -325 mesh metal powders having a nominal purity of 99.9 pct and compressing these mixed powders in a cylindrical die at a pressure of 43,000 psi. These alloys, each weighing 15 g, and some of the arc-melted alloys were annealed in a high-temperature vacuum furnace heated by tantalum strips at a pressure of 10-8 Torr and were rapidly cooled by turning off the furnace power. X-ray and metallographic examination of both series of alloys served to establish the composition ranges of the A15-type phases. Although some chromium losses occurred during the vacuum annealing, they were largely confined to a thin layer on the outer surfaces of the samples. It was established that the A15 phases occur in the Cr-Pt system at 21 ± 1 at. pct Pt after 1 week at 1200°C and in the Cr-Os system at 28 ± 1 at. pctOs after 1 day at 1400°C (see Table I). We also observed that an arc-melted stoichiometric (A3B) alloy in the Cr-Ir system was single-phase (A15-type) in the "as-cast" condition in agreement with previous work.8,13 In addition we obtained a sample of the Cr-Os A15-type phase from Argonne National Laboratory. This alloy contained less than 1 pct second phase12 and was submitted to a density measurement. The density measurement yielded a value of 11.14 g per cu cm in comparison to a theoretical value of 11.25 g per cu cm calculated using the observed lattice constant (4.6806Å) of this alloy. The uncertainty in measurement was 0.1 pct but the sample may have contained some cracks or minor imperfections which could account for the low experimental value. We have also studied the atomic ordering in these phases by means of integrated line intensity measurements using thick, flat, rotating powder samples and CuK a radiation in an X-ray diffractometer. We have obtained order parameters of 0.90 for the Cr-Pt phase, 0.89 for the Cr-Ir phase, and 0.64 for the Cr-Os phase using the formula: where s is the usual Bragg and Williams order parameter, ra is the fraction of chromium atoms in A sites, and FA is the fraction of chromium atoms in the alloy. The values obtained are estimated to be accurate within ±4 pct. If the unusually small value for the order parameter of the Cr-Os A15 phase were due to the existence of lattice vacancies on the "B-atom" sites, then a density of 10.04 g per cu cm would be expected in contrast to the observed value of 11.14 g per cu cm. We, therefore. conclude that the fraction of lattice vacan-

Jan 1, 1967

By G. O. Hiers

IN 1948 in the United States, 184,300 tons of lead was fabricated as coverings for electric power and communication cables. Such covering generally is called "sheathing" for the principal lengths of the cables, while the short lead encasements of the joints are called "sleeves." For the purpose of this paper, both lead sheathing and lead sleeving can fittingly be called lead tubing. In the United States in the same year, 34,262 tons of lead was consumed in making extruded lead pipes. Such pipe is employed frequently for water or chemical solution conveyance with gage pressures up to 100 psi. Consequently, the walls are thicker in ratio to diameter than is the case with the cable sheaths. For this discussion, lead pipes can be thought of as tubes or tubing. About 19 centuries ago, the Romans produced metallic lead of about 99.9 pct purity. Evidently this was cast in flat sheets of Y8-in. thickness, for instance. Such sheet of 10-ft lengths was bent around a removable core producing a tube somewhat oval in cross-section and with a typical internal diameter of, say, 3 in. This method of production necessitated a seam or longitudinal joint evidently made by casting the seam metal adherently. The seam metal, as analyzed recently, contains about 1 ½ pct Sn. Blow pipes were said to have been employed when the joint was made. Specimens of such pipe were found both in Italy and in the British Isles. It is believed the pipes conveyed water for drinking purposes and for baths. Numerous specimens are still in existence in a good state of preservation, due to lead's usual inertness in air or water. According to Claude E. Pearson,1 there were various stages in the evolution of lead extrusion presses for making tubes: I. First press by Joseph Bramah, 1797. 2-First successful press by Thomas Burr, 1820. 3-First bridge die press, J. and C. Hanson, 1837. In 1845, Wheatston and Crooke made the first lead sheathed electric cable. Flat lead strip was wound spirally around the insulated conductor. Overlapping edges of lead were soldered. Early cables were also made by pulling insulated conductors into 50-ft lengths of lead tube. Borel in France, and Wesselau are credited with devising the first presses for extruding lead sheaths directly on cables in 1879. In the United States, Eaton, in 188o, and Roberston, in 1885, introduced new presses of improved design. Following Cunningham's original idea of making lead traps and. bends in 1873, the presses apparently were improved and made successfully soon thereafter. Between the years 1930 and 1940, two continuous extrusion presses for lead were developed abroad. One apparently was originated by Dunsheath and made by the Henley Telegraph Works Co. The other was developed by the Pirelli-General Cable Co. CURRENT PRODUCTION In addition to the several descriptive chapters on lead extrusion in Pearson's book, the Metals Handbook2 recently described pipe, sheath, and collapsible tube production. The latter is by impact extrusion.

Jan 1, 1951