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By E. V. d’Invilliers

The position of these magnetic ore-mines, with reference to the county-seat, is shown in Fig. 1. They are situated on the south margin of the Great Valley, five miles south of Lebanon, and about midway between Reading on the Schuylkill River and Harrisburg on the Susquehanna River, in three prominent, contiguous hills, standing in front of the Mesozoic red shale arid sandstone hill country of northern Lancaster and southern Lebanon counties, and overlooking the great limestone plain to the north. The three hills, named respectively Big, Middle, and Grassy Hills, and separated from each other by two branches of Furnace Creek, range in a nearly east and west direction for a distance of about one mile. Their summits are respectively 870, 715, and 700 feet above tide. The railroad line opposite the base of Middle Hill is 575 feet above tide at Cornwall and 471 feet at Lebauon. These features are all distinctly illustrated in the contour-line map accompanying this paper, but are made still more clear in the relief. map, Fig. 2, comprising, as it does, about six square miles of territory contiguous to the ore-mines. This is a photographic plate of a model made by Mr. A. E. Lehman, from surveys of my own and Mr. W. F. Shunk. The model was built in layers of cardboard, each layer representing one 10-foot contour-carve line. The scale was the same as the original map, 400 feet= 1 inch, without any vertical exaggeration. The lowest contour-line is 530 feet above ocean level; the highest, 1075 feet. The center of the model was made the focusing-point of the view, which was taken—according to scale—51/2 miles north from and 24 miles above the focusing-point, looking southward. When first mined in the last century these hills exhibited smoothlyrounded surfaces, composed of soft weathered ore of great parity, beneath which lay the great mass of hard ore, constituting the whole

Jan 1, 1886

By A. J. Garnier, N. H. van Lingen

Rock downhole is known to be lesc. drillable than when brought to the surface. This must be ascribed mainly to the presence under downhole conditions of a pressure differential across already made chips, which hinders their being lifted. The pressure differential has partly a static and partly a dynamic origin. Balling-up of bits is another consequence of this pressure differerztial. Reduction in penetration rate owing to an increase in the strength of the rock is governed by the difference between the mud pressure and the pressure of the for/nation pore liquid. Rotationally symmetric geostatic stresses have 170 effect on drillability. Fracture of rock when drilling will be brille in most cases. The above is supported by laboratory drilling experiments with drag bill and roller bits an elevated mud, pore, and confining pressures On rocks differing in strength and permeability. INTRODUCTION In oil well drilling drillability of rock is found to decrease with increasing depth of the hole. Naturally deep rock will be more compacted and, therefore, harder to drill than shallow rock of the same type. However, apart from this the drillability of a sample of deep and compacted rock brought to the surface is generally much higher than in its original location downhole. In view of the economic implications of this reduction in drillability, it seems worthwhile to analyze its causes. The origin clearly has to be sought in the difference of environment. The only conceivable factors would seem to be the presence of mud under pressure, the pressure of the formation pore liquid, and the overburden of the rock. Down the borehole the rock is compressed triaxially by mud pressure and overburden. It is well known that the strength of rock is increased when confined by external pressure1. Various authors have, therefore, ascribed the difference in drillability mainly to the strengthening of rock by triaxial compression. Another factor, mentioned by Bobo and Hoch5, is that forces. including "pressure differential forces", tend to hold a dislodged particle in place. However, the conditions determining their magnitude are not clarified nor is their effect on drilling rate assessed. In the laboratory, drilling experiments on pressurized samples of rock yielded evidence that pressure differential forces holding the chips down are the major factor in reducing rate at depth. This paper describes the experiments and shows how the results enable both a qualitative and a quantitative interpretation of the factors determining the effect of chip hold-down on drilling rate. The implication with respect to balling-up and jet action is also discussed. The greatcr part of the experiments were performed in the pressure vessel shown in Fig. 1. The spacc between the sample of rock and the vesscl is divided by "0" rings into three separate chambers. The rock sample is confined laterally by pressurized oil in thc middle chamber. Penetration of oil into the sample is prevented by partly jacketing it with brass foil. Thc pressure of the drilling fluid in the hole can be adjusted via the upper chamber. Permeable rock specimens arc water saturated before being drilled. With a properly plastering mud as drilling fluid the pressure of the pore water can be adjusted independently via the lower chamber.

By R. A. Ritter, J. P. Batycky

A numerical technique has been developed to permit estimating the pressure gradient associated with laminar flow of thixotropic liquids through long Pipelines. For this purpose the pipeline is divided into a number of radial and longitudinal increments within which rheological properties of the fluid may be considered as constant at any time. Then, provided only that the fluid flow curve is defined at every duration of shear, it is possible to predict the instantaneous pressure gradient at any cross-section along the pipeline for each desired flow rate and pipe size. The technique consists of an iterative integration of shear rate to establish the appropriate value of the wall shear stress at each location. Consistency of fluid in the increment is determined by the flow history of that increment, while the radial flow associated with variations in velocity profile is accounted for by adjusting the width and radial position of the increment. A number of pressure profiles, computed at each of several flow rates, provide a convenient basis for pipeline design and pump selection. INTRODUCTION In recent years considerable attention has been given to predicting pressure drop associated with the isothermal laminar flow of time-independent non-Newtonian fluids in pipes and annuli. 6,11,12 The approach generally has been to develop analytical relationships between flow rate and pressure drop based on simple constitutive models which hopefully provide an approximate description of the rheological properties of the fluid. Analytical solutions are highly desirable since the influence of all pertinent parameters can be readily determined. Unfortunately, however, this approach is restricted to simple flow geometries and- frequently leads to erroneous results due to inadequacies in the model. In certain cases a solution may be obtained through applying appropriate numerical techniques.l,14 For example, a digital computer program is available for predicting the velocity profile and pressure drop encountered by any Newtonian or time-independent non-Newtonian fluid flowing under laminar conditions in a cylindrical pipe or annulus.l In this paper the consistency behavior of the fluid need only be described in terms of basic rheological data. Analyzing flow systems involving fluids with time-dependent rheological characteristics is considerably more complicated since substantial changes in consistency may occur because of sustained shear action. This sensitivity to shear frequently persists for several hours. Consequently, variations in pressure drop and/or flow rate resulting from the aging process and addition of unsheared or partially sheared fluid to the system must be considered for purposes of pipeline design. This paper outlines a numerical method for predicting the transient and steady-state laminar flow behavior of a thixotropic liquid in a pipeline of arbitrary length (i.e., at a specified constant flow rate, the instantaneous pressure gradient may be determined at any time after start up and at any location along the pipeline). Several such pressure gradient profiles, computed at several flow rates, may be combined to produce a complete portrait of the system response. This flow portrait provides a reasonable basis for pipeline design and for selecting a suitable pump characteristic. TIME-DEPENDENT RHEOLOGICAL BEHAVIOR The most familiar time-dependent rheological properties are those exhibited by thixotropic liquids. Many of these materials, particularly thixotropic crude oils, generally display an apparent yield stress in that a finite pressure gradient is required to initiate flow. Then, under the influence of sustained shear at a constant shear rate, the consistency systematically decreases to some final limiting value. The shear stress obtained at

By J. L. Ratliff, R. I. Jaffee, H. R. Ogden, D. J. Maykuth

An experimental program was carried out to improve the low-temperature ductjlity of tungsten through the combined use of dispersed oxides for grain-size control and Groups VII and VIII metal additions for matrix softening. Separate binary -alloy studies showed that increasing amounts of finely dispersed thoria or zirconia and critical amounts of rhenium, osmium, and iridium were effective in significantly reducing the ductile-to -brittle bend transition temperature of tungsten as well as in increasing its re crystallization temper - LIKE most of the high-melting point bcc metals, tungsten undergoes a marked transition from ductile to brittle behavior with decreasing temperatures. The material factors which influence the transition temperature include grain size, metal purity, and the presence of dispersed phases. ature and decreasing its recrystallized grain size. The beneficial effects shown by each of these additions on transition temperature and grain size in binary combinations appeared to be additive in tevnary combinations. Thus, combined additions of 2 to 8 vol pct thoria with 5 pct Re, 0.87 pct Os, or 0.30 pct Ir were effective in lowering the transition temperature of tungsten pom 2000 to 75"C, increasing its re crystallization temperature from 1600 to 1800°C, and decreasing its recrystallized gain size from 1175 to 6800 grains per sq mm. The mechanism by which these factors are operable in tungsten is not yet clearly understood. Nevertheless, it has become increasingly apparent that the interrelationship of grain size and matrix composition governs to a large extent the degree of low-temperature ductility which can be achieved. Though it is not possible to accurately assess the relative importance of these variables, the available evidence suggests that, at low temperatures, grain-size effects predominate over purity effects. This was indicated in recent studies by Allen, Maykuth, and Jaffee,' involving single crystals containing 20 to 40 ppm 0 and 1 to 30 ppm C. Thus, while each

Jan 1, 1964

By Norman B. Pilling, George P. Halliwell

The effects of lead and tin up to maximum contents of about 0.1 per cent. each, in the presence of oxygen between 0.04 and 0.30 per cent., have been studied. Tin is retained efficiently in the oxidized condition, whereas lead is not. An approximate relation between the lead-oxygen and tin-oxygen contents is deduced which, if observed, permits the presence of these elements without detrimental effect. Excess of lead beyond this limit affects ductility and rolling properties adversely and an excess of tin correspondingly diminishes the conductivity. It is shown that copper carrying a combined lead and tin content of several hundredths per cent., yet fully equal in quality to electrolytic wirebar, may be produced under commercial conditions. Tin is not normally found in more than bare traces in current grades of American electrolytic copper, and lead is rarely found to more than a few thousandths per cent. In the electrolytic refining of copper these metals do not present great difficulty in separation and their presence in any considerable amount in copper that has been refined by this method need not be tolerated. On the other hand, the effect of these elements is not entirely academic, as a considerable tonnage of scrap copper liberally contaminated with them has been reclaimed by the fire-refining process into metal with a residual lead and tin content of several hundredths per cent., yet having excellent working properties and fully equal in quality to current grades of electrolytic wirebar. Such copper has been found entirely suitable for use in the manufacture of electrical apparatus. In the absence of oxygen, tin is known to depress the conductivity rapidly, and lead to impair ductility and the rolling properties, yet the effects of these metals on commercial grades of copper—i. e., copper containing from 0.03 to 0.10 per cent. of oxygen—apparently has not received systematic attention. Occasional reference is made to the fact that in the presence of oxygen lead may be tolerated, yet Addicks says,1 in regard to overpoling: "The unstable condition can be controlled by adding certain substances, notably lead, in small quantities, but as mere

Jan 1, 1926

By M. A. Orehoski, J. M. Hodge

It is now generally conceded that if a steel is to develop optimum physical properties in the conventionally quenched and tempered condition, the microstruc- ture after quenching should consist wholly of martensite; and that transformation of austenite to intermediate-temperature products during quenching will have a deleterious effect upon the physical properties, particularly the toughness, of the tempered material. This implies that, in order to obtain the optimum physical properties from a given steel, its "harden-ability " should be high enough to suppress all transformation except that to martensite under the heating and quenching conditions to which it will be subjected, and this premise is probably the most important reason for hardenability control. However, the criterion of hardenability that at present is most widely accepted and used is based on a microstructure of 50 per cent martensite. This criterion, which was suggested by Grossmann,' does have considerable justification. It may be readily measured by a fracture or etch test and, since it represents the point at which the hardness is changing most rapidly with variations in cooling rate, it may be determined empirically from a hardness-depth curve by determining the point of steepest slope (the inflection point). Furthermore, in plain carbon steels or in steels of low hardenability, the difference in hardenability expressed in

Jan 1, 1947

By J. M. Hodge, M. A. Orehoski

It is now generally conceded that if a steel is to develop optimum physical properties in the conventionally quenched and tempered condition, the microstruc- ture after quenching should consist wholly of martensite; and that transformation of austenite to intermediate-temperature products during quenching will have a deleterious effect upon the physical properties, particularly the toughness, of the tempered material. This implies that, in order to obtain the optimum physical properties from a given steel, its "harden-ability " should be high enough to suppress all transformation except that to martensite under the heating and quenching conditions to which it will be subjected, and this premise is probably the most important reason for hardenability control. However, the criterion of hardenability that at present is most widely accepted and used is based on a microstructure of 50 per cent martensite. This criterion, which was suggested by Grossmann,' does have considerable justification. It may be readily measured by a fracture or etch test and, since it represents the point at which the hardness is changing most rapidly with variations in cooling rate, it may be determined empirically from a hardness-depth curve by determining the point of steepest slope (the inflection point). Furthermore, in plain carbon steels or in steels of low hardenability, the difference in hardenability expressed in

Jan 1, 1947

By William H. Morris

At the Glen Summit meeting, the question of controlling the silicon in pig-iron WAS raised; and as this has been deemed by most furnace-men for years past a matter of special difficulty, an account of our experience in the matter may be of interest. When running on mill-iron, some years ago, we aimed to keep

Jan 1, 1893

By Charles W. Dickson

CONTENTS. [ ] I.THE RELATION OF NICKEL TO PYRRHOTITE. Introduction. The Sudbury district is to-day one of the two great sources of nickel in the world. The peculiar geological relations of

Jan 1, 1913

By C. W. Haworth, M. G. Hall

THE diffusion of copper in silver has been studied by Sawatzky and Jaumot1 and Cahoon and Youdelis.2 Sawatzky and Jaumot measured tracer diffusion in single crystal specimens but the 12.9 h half-life of CU64 limited the time of the diffusion anneal to about 50 hr, and in consequence the minimum temperature of the investigation was 716°C. Cahoon and Youdelis employed polycrystalline diffusion couples, and measured the interdiffusion coefficient over the range 0 to 2 at. pct Cu. These experiments were performed at temperatures down to 630°C and a discontinuity in the Arrhenius plot was observed at -700°C. The high D-values at lower temperatures were attributed to the effect of grain boundary diffusion. The present experiments were undertaken to investigate this anomaly and study the low temperature region of the diffusion of copper in silver. The method employed single crystal samples and a plated source of copper. Silver single crystals 2½ by 4 in. sq were grown from zone-refined silver bar (>99.99 pct Ag) by the Bridgeman technique, and $ in. thick slices were cut by spark machining. The specimens were electro-polished, and a 2 µm layer of copper was electroplated on to the planar surface at one end of each specimen using a potentiostatic voltage source. The following solution was used for plating: CuSO4 . 5H2O 125 gm H2SO4 (S.G. 1.8) 53 ml H2O 1000 ml The diffusion couples were sealed in capsules, in an argon atmosphere, and were annealed at either 701, 597, or 498°C. The temperature measurements were made to an accuracy of *l°C After the anneal, the penetration of copper into the silver was determined using an electron probe micro-analyzer. 100 sec scaler measurements of the intensity of Cu Ka! X-radiation were performed at 2 to 5 µm intervals along a line normal to the source, with a beam potential of 25 kv. Allowances were made for the background radiation, and the dead-time of the counting system in the normal way, and the intensity ratios were corrected to weight concentration by the method of Duncumb & de-Casa.3 For the analysis of the diffusion profiles, it was first observed that the copper layer did not completely dissolve in the host crystal for the specimens annealed at 498 and 591 whereas for the specimen annealed at 701°C, a single phase solid solution was formed. For the 498" and 597°C specimens, the concentration of copper in the silver at the interface was therefore constant throughout the anneal and the analysis for a

Jan 1, 1970

Sec. 1. The membership of the Institute shall comprise six classes, namely: 1. Members; 2. Honorary Members; 3. Senior Members; 4. Associates; 5. Junior Members; 6. Rocky Mountain Members. All shall be equally entitled to the privileges of membership, excepting that no member of any class whom residence shall be outside of the United States of America, except those residing in Mexico and Canada, shall be entitled to vote. SEC. 2. Members: A person to be eligible for election or transfer into the claw of Member must be at least 27 years of age and must have had at least six years' employment in the practice of engineering, mining, geology, metallurgy, or chemistry, during at least three years of which he must have held positions of responsibility in one or more of these fields. This shall not apply to Rocky Mountain members.

Jan 1, 1944

By Lloyd R. Jackson

IN the past 15 years there has been a great deal of interest in the fundamentals of plastic flow and rupture in metals and a number of papers have presented substantial advances toward a fundamental insight into the mechanism of these two processes. This paper is an appraisal of the present position with respect to the following: I. The relative importance of shear and hydrostatic stresses in controlling the plastic flow of metals. 2. The relative importance of shear and hydrostatic stresses in determining the amount of flow before rupture in metals. 3. The use of generalid strains in determining the effect of plastic flow on the work-hardening of metals. FUNDAMENTALS INVOLVED IN PLASTIC FLOW In attempting to devise a quantitative description of plastic flow in metals, it has been customary to assume that metals are homogeneous and isotropic. The fact that neither of these assumptions is correct complicates the experimental appraisal of relations developed. Never- the less, as will be shown, metals are close enough to being isotropic and homogeneous so that considerable insight into the mechanism of plastic flow can be obtained. The assumptions of .isotropy and homogeneity allow both the elastic and plastic behavior of metals to be described in terms of 12 variables-six being stress variables, and six being strain variables. ) From this viewpoint, it should be possible to represent the stress-strain relations in the plastic range for metals by a surface in 12 dimensions. One boundary of this surface would be the surface of transition from elastic to plastic behavior; the other would be the surface of rupture. These two surfaces would be joined by the surface of plastic flow. The shape of this composite surface would depend on the test temperature and speed of straining. This paper is concerned only with results obtained at constant temperature and constant speed of straining. While such a representation is conceivable, the demonstration of such a surface would present almost insurmountable obstacles. The most fruitful simplification of this concept has been the recognition that the behavior of metals in the plastic region is primarily a result of shear stresses and strains. Based on this viewpoint, it is possible to utilize stress and strain invariants that are independent of hydrostatic components to define an effective stress and strain. In effect, this amounts to reducing the task of representing stresses and strains in 12 dimensions to one of representation in two dimensions. The two invariants (one for stresses . and one for strains) that have received most attention are the following: [ ]

Jan 1, 1946

By P. S. Kotva

Dislocation substructures in superalloy matrices of varyzng co)npositions have been studied. In general, it has been found that the alloys can be classified into ''high", ''medium", and "low" stacking fault energy classes based on the type of dislocation substructure observed in the matrix and that the substructure can be correlated to the stacking fault energy. The effect of different types of dislocation substructure and dislocation reactions on the intragranulur precipitation of carbide phases has been studied. In a Ni-Cv-Mo-Fe matrix, precipitation of MC carbides in association with stacking faults has been observed. In most superalloys, solid-solution strengthening and precipitation hardening are the chief mechanisms employed to achieve strength. The latter contribution to strength is usually achieved by the precipitation of / in certain wrought alloys. Insufficient attention has been given to the problem of obtaining strength in su-peralloys by controlling precipitation of carbides on imperfections within the matrix. The present work was undertaken to investigate the dislocation substructure in various superalloy matrices, to study the effect of such substructure on subsequent precipitation of carbides in the matrix, and to investigate whether certain modes of precipitation of carbide phases found in austenitic stainless steels2"4'6 would occur in nickel-base alloy matrices with dislocation substructures of the same type as those found in austenitic steels. 1) EXPERIMENTAL TECHNIQUES Five-pound heats of the various alloy compositions reported here were vacuum-cast. The ingots were given light deformation by rolling to break up the as-cast structure and then homogenized for 24 hr. HASTELLOY alloy X (nominal composition: Ni-2OCr-17Fe-8Mo-0.05C) was homogenized at 2150°F and In-cone1 625 (nominal composition: Ni-20Cr-5Fe-8Mo-3.5Cb-0.05C) was homogenized at 2280°F. Fabrication of 0.004-in. sheet was achieved by cold rolling with intermediate annealing treatments being carried out at the same temperature as those used for homogeniza-tion. Each solution anneal was followed by quenching. The aim of this procedure was to redissolve as much of the primary carbide phase as possible. Samples of the 0.004-in. sheet were cut and encapsulated in quartz capsules and then heat-treated in the tube furnaces. Thin foils were prepared using an ethanol-10 pct perchloric acid bath at 32°F and at a voltage of 22 v. A "window" technique was employed. Observations were made on a JEM-7 electron microscope operating at 100 kv. 2) EXPERIMENTAL RESULTS a) Types of Dislocation Substructure. Fig. 1 shows a schematic correlation between stacking fault energy, SFE, and the type of dislocation substructure observed in various matrices of nickel- and cobalt-base alloys. A precise quantitative determination of stacking fault energy is not implied in the figure but the correlation between stacking fault energy and the type of dislocation substructure obtained allows alloys to be divided into three classes in analogy with the classification employed by Swann and ~uttin~' for binary alloys of copper. Class I alloys are associated with a "high" SFE and show a cellular substructure of dislocations as typified by the micrograph of a thin foil of pure nickel deformed 4 pct at room temperature in Fig. 2. With decreasing SFE the tendency toward cell formation is lessened and dislocations tend to be arranged in coplanar groupings. Examples of this class of alloys with "medium" SFE are provided by the mi-crostructure of solution-heat-treated, quenched, and deformed thin foils of HASTELLOY alloy X, "Waspaloy" (prior to any aging), and Inconel 625. Fig. 3 shows a thin-foil micrograph of an alloy of Inconel 625 composition, solution-heat-treated, quenched, and deformed 5 pct at room temperatures. No evidence of any cell structure can be obtained in materials of "medium" stacking fault energy, Fig. 3, even after severe deformation. The stacking fault energy of the alloy shown in Fig. 3 is, however, not low enough to make the dissociation of dislocations visually obvious. As stacking fault energy decreases further, with successive addition of solute in the matrix, there is an increased tendency toward dissociation of dislocations and cross slip becomes progressively more difficult. Eventually, when the stacking fault energy is "low" enough, complete dissociation of dislocations is seen to occur as shown in Figs.

Jan 1, 1969

By C. F. Nagel

THIS paper is written primarily for those who are familiar with the processes mentioned but who desire a further insight into some of the fundamental principles. It does not give a complete description of the various practices employed in the working of aluminum and its alloys, but presents a few illustrations of the essential respects in which aluminum differs from other common metals, and emphasizes the controlling factors of the processes under consideration. HEAT TREATING Certain aluminum alloys are capable of having their mechanical properties greatly altered by a heat-treatment process. The degree of this improvement may be illustrated by some typical properties of duralumin (Table 1). It is essential that the material be in the wrought condition in order to obtain maximum heat-treatment effects. [ ] In the case of duralumin, the process consists of heating the metal to a prescribed temperature and then quenching it. Immediately after quenching, the material is relatively soft though not as soft as in the annealed state. If allowed to stand at room temperature, the material changes, increasing in hardness and strength without loss of elongation, but at the expense of the plasticity. This action is very rapid at first, gradually diminishing in intensity and to all practical purposes being completed in about four days. This process of hardening is termed

Jan 1, 1928

By Kenneth Seaver

It is rather gratifying that in the field of the manufacture of refractories, in which, as a whole, foreign practice has surpassed us, the United States stands pre-eminent in the making of silica brick. In some instances European manufacturers have refused point-blank to credit the possibility of making to such specifications as are Common with us, until the actual completed shapes were shipped for inspection. But both the manufacturer and consumer in foreign lands are now frankly learning from us, and benefiting by our advanced practice in this field, as is evidenced by the orders placed here and the efforts put forth abroad to come up to the standard of American-made material. In no metallurgical operation has the employment of silica brick spread with such leaps and bounds as in the building of plants for the carbonization of coal, particularly in the byproduct coke oven. Today, in practically all types of oven, the structure above the floor level (except the facings) is of silica, while in the recuperative and in some regenerative types, its use is extended to that portion of the oven below the floor as well. By reason of the fact that there are a number of classes of refractories of higher silica content than the usual so-called fire-clay brick, there is sometimes a misunderstanding as to the class of material referred to as "silica" brick. We here refer only to a brick having a silica content of 94 per cent. or more and made usually from quartzite rock with a small percentage of lime as a binder. The term "silica" brick is sometimes rather loosely applied to brick from highly siliceous clays, or to a quartzite refractory of mingled fire-clay and quartz rock; but such products are not considered in this discussion. : Materials The rock from which practically all high-grade silica brick are made is a true quartzite or metamorphosed sandstone, variously known as

Jan 1, 1916

By A. Blake Caldwell

Northwest frontier is here defined as Alaska, the Yukon, the Northwest Territories and northern British Columbia. The mineral potential of this area is great and will unfold in the decade ahead but now the largest single producer is the Pine Point lead-zinc mine in the NWT about 1200 km north of Edmonton, Alta., close to the south shore of Great Slave Lake. The mill, initially designed for 4540 mtpd, was expanded to 7250 mtpd in 1969. This operation is described in detail in Mining and Concentrating of Lead and Zinc, Vol.1, Chapt. 17, SME/AIME in the article by W. H. R. Gibney.

Jan 1, 1971

By Leonard Logan

Production control in the petroleum industry may be discussed under three general heads: engineering, legal, and economic. Though the problems that fall under the respective heads of engineering, law, and economics are separate, they are dependent and interrelated. The present study will deal with the economic phases of production control and only incidentally with the legal phases as they affect the economics of the problem. No one of these forms of control is more important than the others. The engineers have performed their service well, and although there are many problems of engineering control yet to be solved, they are not pressing. Legal control is a problem for the legislatures and the courts to solve. The petroleum industry is such a comparatively young industry and the legal problems are of such recent origin that the courts have not yet been called on to settle all the issues and the legislatures still have before them bills concerned with the regulation of the industry. Economic control of production in the petroleum industry is a highly desirable goal but it has not yet been completely achieved. Since the first commercial well was drilled in 1859, there have been recurring periods of overproduction. (Over- production will be used here to describe a condition wherein so much crude oil is available that a satisfactory and reasonable price is not received for the product.) The production of oil in 1859 approximated only 2000 barrels. In the following year production jumped to 500,000 barrels, thus bringing into

Jan 1, 1932

By P. L. Allsman

A method of analyzing blasting action indicates that major cost savings are possible by revising practice and bringing the classical blasting formulas up to date; difficult problems such as taconite and throw-breakup can be attacked by engineering Denonation pressure — the peak pressure developed by the explosive reaction; "shatter blasting." Average detonation pressure — the average pressure over the time the reaction continues; "heave blasting." Radial presswe — the stress in a radially expanding sphere about the detonation. Corresponds to the diminishing peak pressure. Lateral stress - the stress in a circumferential direction, induced by outward movement of the rock away from the detonation. Limiting stress or Slim— the rock property at which failure will occur in the case at hand; may be tensile or compressive strength, or combinations. Recent developments in explosives technology, following the advent of nuclear explosions, have led to a rather complete understanding of their action. A survey was made for the purpose of uncovering any knowledge which might be applicable in the mining and quarrying industries, the principal users of explosives. Although no revolutionary techniques have become apparent, much basic data pertaining to blasting — fortunately not classified — have been developed by recent research, from which a good general understanding of the process can be derived. It is hoped this explanation of the scientific principles governing explosive action, together with a proposed analysis of blasting practice, will further the development of mining engineering. NATURE OF THE REACTION: The phenomenon of explosion, termed detonation, is in reality a rather complex chemical reaction, and as such is completely explainable by physical and chemical laws. Detonation is characterized by high temperature and pressure, extreme rapidity — often being complete within one microsecond — and most importantly by formation of a shock pulse which accompanies the reaction zone. As with any chemical reaction, there is a critical value of temperature and pressure below which detonation can not occur; this is termed the "critical point". Whereas an explosive substance will decompose slowly at ordinary temperature, with the formation of gases which readily dissipate; at elevated temperature the reaction is considerably speeded. If pressure is suddenly applied at some point in an explosive medium, adiabatic compression results, and the temperature is locally raised. When the temperature becomes high enough the decomposition will be so rapid that the gaseous products do not have time to dissipate, but contribute to a further build-up of pressure. This, in turn, furthers adiabatic temperature rise over a widening area, and with any heat generated by an exothermic reaction, causes a rapid increase in temperature. Both temperature and pressure are spontaneously built up in this manner until the critical point is reached, resulting in detonation. The values of pressure and temperature necessary to create detonation are established by the unusual nature of the process. If temperature alone is raised above the critical point, without sufficient increase in pressure, deflagration or "low-order detonation" occurs. In the true, or "high-order detonation", the shock pulse resulting from the violent decomposition becomes an integral part of the reaction, and itself aids in increasing the detonation

Jan 1, 1961

By M. A. Scheil, S. L. Hoyt

About three years ago eight standard steels were prepared for the cooperative investigation of methods for oxygen analysis, sponsored by the Iron and Steel Division of the American Institute of Mining and Metallurgical Engineers and by the National Bureau of Standards. A preliminary report of the results was made at the annual meeting of the A.I.M.E. in February, 1936. † This report showed the reliability of vacuum-fusion methods and further brought out certain advantages of the fractional vacuum-fusion method. The results obtained by the authors with the latter method will be discussed here, with some comments on the procedure for those who are interested in determinations of gases in steel. Some additional comments are offered on the use of the polarizing microscope for identifying the inclusions in the steel to check and amplify the findings of the fractional vacuum-fusion analysis. For details of the fractional method, see the paper by L. Reeve1. Degassing the Crucible One of the sources of error in vacuum-fusion methods has been the blank gases evolved from the graphite crucible after preliminary out-gassing. We have used 1750" C. as the outgassing temperature. This gives a base pressure at that temperature of less than 0.030 mm. after one hour and of 0.004 mm. on dropping the temperature to 1575" C. At lower extraction temperatures the base pressure is 0.001 mm. at 1050" C., 0.002 mm. at 1170" and 0.003 mm. at 1320' C. Because of the size of our induction coil we are limited to this temperature, although it 'would be advantageous to use 1950" to 2000" C. Hamilton2, after outgassing at 1850" C. for 100 min., was able to reach a base pressure of 0.001 mm. for a temperature of 1700' to 1725" C. This would give a blank corresponding to less than 0.00015 per cent 02 and 0.00005 per cent H2. Chipman and Marshall outgas at 2000" C. for one hour by using a special construction

Jan 1, 1937

By A. Taylor

NICKEL-RICH alloys hardened with small additions of titanium and aluminum and centered around that region of face-centered-cubic primary solid solution, 7, where the atomic ratio of nickel chromium is about 3 to 1, form an important series of high temperature materials because of their resistance to oxidation, high temperature strength, and excellent creep characteristics at elevated temperatures.' With suitable heat treatment it is possible to precipitate a phase based on the ordered face-centered-cubic y' phase, Ni,,Al, from the randomly ordered nickel-rich y face-centered-cubic primary solid solution. By adjusting the ratio of aluminum to titanium the precipitation of needles of the hexagonal intermetallic compound n-NiaTi may be effected. In the present investigation, the main interest was focused on the interrelationships among the y, y, and 7 phases. By using alloying elements of high purity and a suitable melting technique the presence of carbides and nitrides could be completely obviated, although it is realized that their presence is probably essential in alloys of commercial importance. As a preliminary to the study of the Ni-Cr-Ti-A1 system, it was necessary to establish the constitutional diagrams of the three contiguous systems, Ni-Cr-Al, Ni-Ti-Al, and Ni-Cr-Ti, which together form the nickel-rich corner of the Ni-Cr-Ti-A1 quaternary tetrahedron. Accounts of this work on the ternary systems have already been published by the author. A preliminary disclosure of the author's investigations on the ternary and quaternary systems has been made by Hignett2 and has since stimulated work in the same field by Nordheim and Grant." Before the detailed account of the quaternary system, a brief description will be given of the three contiguous ternary diagrams on which it is ultimately founded. Ni-Cr-Ti System The 750° and 1000°C temperature isothermals of the Ni-Cr-Ti system as determined by Taylor and Floyd' are shown in Figs. 1 and 2. The single phase fields of greatest interest are occupied by the extensive face-centered-cubic y primary solid solution of chromium and titanium in nickel, and the close-packed-hexagonal Daltonide, n-Ni3Ti, from which tie-lines radiate across the very wide y+n two-phase field. The chromium-rich primary solid solution is represented by a. As may be seen from the diagrams, the extent of the y phase field increases quite appreciably as the temperature is raised from 750" to 1000°C. Isoparametric lines for the y phase are shown in Fig. 1. They are interesting in that they are almost parallel to the y/r+n boundary. Thus the determination of the tie-line directions is entirely dependent on the exact stoichiometry of composition of the 7 phase. Ni-Ti-Al System The 750°C isothermal section of the nickel-rich portion of the system Ni-Ti-A1 as determined by Taylor and Floyd is shown in Fig. 3.' The regions of greatest interest are the extensive y primary solid solution of titanium and aluminum in nickel, the y' Berthollide phase field based on the ordered face-centered-cubic structure of Ni3Al, and the hexagonal Daltonide phase, n-Ni3Ti. As may be seen from Figs. 3 and 4 the area occupied by the y phase is substantially the same at 750°C as at 1000°C, but the area of the 7 phase is considerably greater at the higher temperature. The tie-lines radiate from n-Ni3Ti across the y+v and y' + v two-phase fields and are evenly distributed

Jan 1, 1957