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By Pol Duwez, Howard E. Martens

INTRODUCTION THE high temperatures encountered in the operation of jet engines have imposed most drastic requirements upon the materials used in their construction. There are two different approaches to the materials problem connected with the design of parts in contact with high temperature gases. The first approach is to use a material which will perform satisfactorily at a temperature comparable to that of the gas. The second approach is to use a cooling system so that the material is maintained at a temperature appreciably lower than that of the gas. Besides the conventional method .of cooling by means of a liquid in contact with the material exposed to a high rate of heat transfer, a less orthodox method consists of making the part to be cooled of a porous material, so that the cooling fluid can be forced through the pores. In this scheme, the temperature of the coolant moving in a direction opposite to that of the heat flow increases gradually while passing through the porous material, and the coolant forms a protective layer on the surface exposed to heat. This method, referred to as "sweat cooling," was proposed at the jet Propulsion Laboratory in September 1944, and the experimental work on the preparation of porous metals was started early in 1945. The sweat cooling method opens new lines of research in powder metallurgy by extending the field of application of porous metals. The porous metals available at the present time are used mostly as bearings or filters, and a rather limited number of alloys are made for these applications. In addition, the fabrication of bearings or filters does not require a very dose control of the permeability, while in the case of sweat cooling such a control is of primary importance. In designing a sweat cooled part it is imperative to assure a given rate of flow of coolant under certain conditions of pressure drop. Generally, the proper pressure drop and the rate of flow of coolant are deduced from heat transfer and design considerations, and consequently the permeability of the porous metal to be used is also determined. It is therefore necessary for the designer to have available a continuous series of porous metals covering a range of permeability as wide as possible. Contrary to the requirements established in the fabrication of filters, the actual pore size of a porous metal intended for sweat cooling is of little consequence, except for its effect on permeability. The tensile strength is likely to be of great importance for some applications. Since the tensile strength of a porous alloy decreases with increasing porosity, an important problem is to develop porous alloys having the maximum permeability with minimum porosity.

Jan 1, 1948

By William Rawles

THIS report is the first of a series planned by The Mineral Inquiry, organized by the American Institute of Mining and Metallurgical Engineers to make factual studies of the world's mineral resources in their political and international relations. The efforts of The Inquiry are directed toward filling certain notable gaps in available information, without duplication of many other investigations touching this field which are being conducted by other organizations, public and private. Much of its work is indirect-suggesting topics for study by other individuals and organizations and aiding in their planning and correlation. The present report is designed to show the nationality of the commercial control of mineral resources, in addition to the geographic distribution of production by nation-in other words, the overlapping of political and commercial control. Detailed accuracy cannot be claimed for the figures presented. For some companies division of ownership by nationality has been very difficult to ascertain, as many companies are reluctant to dis-close the extent of their holdings in other corporations, and the impossi-bility of tracing all stock ownership is obvious. Also there are frequent changes in the distribution of stock control. The percentages given rep resent only the participation that has been identified definitely, even though it is known in some cases with fair certainty that larger interests exist, as, for instance, American participation in foreign chromite companies. In bringing together these figures we have had the cordial co-operation of the United States Bureau of Mines, the United States Geologi-cal, Survey, the Mineral Division of the Bureau of Foreign and Domestic Commerce, the Brookings Institution, and of many specialists connected with the mineral industries covered. The report does not cover the entire list of minerals: Coal, although it is one of the most important minerals, was omitted because not more than one or two per cent of the world's production comes from mines that are owned by companies foreign to the countries where the coal is produced. Others, such as gold, are omitted because sufficient data have not yet been accumulated. It is hoped that some of these may be covered in subsequent reports. It is not our purpose here to discuss the significance of the figures sub-mitted, but only to point out the general nature of the problems on which they have a bearing.

Jan 1, 1933

By H. Bartholomew, L. Bernstein

Experirmental bonds of Ag-In SLID to gold, copper, nickel, Kovar, Dumet, nickel-plated molybdenum, ALL-Pt vrzetallizing on ceramic, and nickel-plated Mo-Ti metallizing on Al2O3 ceramic have been made and exalnined. Metallographic taper sections were made of all bonds for Phase identification. In addition, failuve temperatltres were determined for bonds heated undev shear load. Sowe bonds were subjected to electvon-probe microanalysis. The ternary SLID system Ag-In-Sn was similarly evaluated. Plwtonticrogaphs and electron-probe data are discussed and related to over-all phase diagrams , diffusion effects, and wetting properties of the liquid. The rates of diffusion andor compound formation of gold, silver, and copper in indium and In-Sn liquids are found to descend in that order, as measured by the relatzve thickness of phase layers in the interfaces of bonds made at 200to 400°C. It is concluded that the SLID process enables one to produce bonds which are stable at high temperatures, after having been fabricated at low temperatures, and meet most requirements for application in interconnection, silicon bonding, or hermetic sealing of electronic devices. In a previous publication by one of the authors, the theory and principle of SLID bonding in simple binary systems was discussed. SLID bonding can be defined as a joining process using a preform with a laminated structure: a low-melting material clad on a high-melting core. During the bonding cycle a metallurgical transformation occurs such that the resultant structure can withstand, without reliquefying, temperatures higher than those at which the bond was made. The phases present in the joint interface region are high-melting compounds or, if equilibrium is reached, solid solution. Properties will depend on the phases present, ultimately approaching those of the solid substrate core of the SLID preform. The high-temperature strength of bonds made in these systems was correlated to phase identification in the interfacial bond regions from metallographic cross sections and electron-probe X-ray microanalysis data. When complex electronic components are assembled, multiphase systems result at the interfaces of bond regions. In these ternary and higher-order systems very little diffusion and phase-diagram data is available with which to predict the nature of the interfacia1 region that will result from conventional fabrication processes and subsequent field use of the component. Clarke and Rhines did early work in the field using the Al-Mg-Zn system as a vehicle. Guy and co-workers studied a portion of the Cu-Zn-Ni system and some mathematical4 and general5 aspects of multiphase diffusion. irkald' and coworkerse recently studied diffusion in the Cu-Zn-Sn system. No diffusion studies have been made on ternary systems directly applicable to this investigation but some phase diagrams are available: g-u-n, u-i-n, and CU-i-n. I) DESCRIPTION AND EXPERIMENTAL PROCEDURE The elements Au, Ag, Cu, Ni, In, and Sn are found in alloy systems commonly used in fabricating and assembling electronic components. Reference to the appropriate phase diagrams9 and the previous publication will indicate that the binary systems of any of the two low-melting elements, indium and tin, with any one of the other four higher-melting elements are capable of producing SLID bonds. It is not unreasonable to assume that any ternary or higher-order alloy system containing at least one of the elements gold, silver, copper, or nickel and one or both of the elements indium and tin could be used to produce bonds which are stable at high temperatures, after having been fabricated at low temperatures. Indeed, the liquidus contours and isothermal sections of those phase diagrams available7 do so indicate. The parameters affecting the bond and the equipment used in evaluating the bond were described in an earlier paper.' The current investigation used the

Jan 1, 1967

By I. J. Polmear

Earlier work revealed that small amounts of silver (< 0,l at. pct) may exert a marked effect on the structure and aging characteristics of certain a1loys based on the Al-Zn-Mg system. Silver additions have now been made to other age-hardening aluminum alloys, and significant effects have been observed in the Al-Mg system and in several Al-Cu-Mg and Al-Mg-Si alloys. Similarities exist between the behavior of alloys from all four systems, notably that in each case silver stimulates nucleation of an intermediate precipitate which forms when the alloys are aged in the medium temperature range (-120" to 220°C). The mechanism for this effect is uncer- IT has been shown that small additions of silver (<0.1 at. pct) may have a marked effect on the structure and age-hardening characteristics of tain although it is suggested it originates in a preferred affinity between magnesium and silver atoms. There is evidence to indicate that the effects of silver are unique to that element. It is suggested that silver atoms interact with vacancies during low-temperature aging of Al-4 pct Cu and certain Al-Cu-Mg alloys. Silver had little effect on Al-8 pct Zn aged at any temperature. Preliminary observations indicate that silver may have interesting effects on the strength properties of Al-Mg alloys, md a the commercial Al-Cu-Mg alloys designated as 2024 and 2618. many alloys based on the Al-Zn-Mg* system.1"&apos;*

Jan 1, 1964

By John C. Sawyer

The mechanism of catastrophic oxidation of chromium and 446 stainless steel is examined. Data are presented to show that accelerated oxidation of these two materials, as caused by lead oxide, can occur in the absence of a liquid layer contrary to presently accepted theory. An alternate theory is proposed in which the rate of accelerated oxidation is a function of the rate at which lead oxide destroys the protective oxide formed on the base metal. An example of the application of the theory is given for the catastrophic oxidation of chromium in the presence of lead oxide. WHEN stainless iron-, nickel-, or cobalt-base alloys are heated in air to moderate temperatures in the presence of certain metallic oxides, oxidation will proceed at an accelerated rate. This phenomenon, often called "catastrophic oxidation", is most pronounced for the stainless steels. With these alloys the condition is so severe that large masses of oxide will form on the surface of the alloy in 1 hr or less at temperatures of 1200o to 1700oF. While a number of oxides are known to cause this effect, PbO, V2O5, and Moo3 are the most familiar, having been the subject of one or more investigations which have appeared in the literature.1-7 In presenting the results of these investigations, many of the authors have offered possible explanations to account for the more rapid rate of oxidation observed; however, the liquid layer theory as proposed by Rathenau and Meijering 2 has been the most commonly accepted mechanism. The liquid layer theory proposes that a low-melting oxide layer is formed on the surface of the alloy as the result of the interaction of the alloy oxide and the contaminating oxide. When the temperature of oxidation is above the melting point of the oxide on the surface, a liquid layer will form and oxidation will proceed at an accelerated rate. At temperatures below the melting point of the surface oxide, oxidation will proceed more slowly in the normal manner. It is argued that the rates of diffusion of oxygen and metal ions through the liquid layer are extremely rapid thereby accounting for the high rate of oxidation. Various experimental data have been presented to show that the temperature at which accelerated oxidation first becomes apparent coincides with the melting point of the eutectic oxide which would be present on the surface. Some exceptions have been observed, e.g., silver will oxidize in the presence of Moo3 at temperatures below the lowest melting eutectic; on the other hand, stainless steel will not be catastrophically oxidized at 1500oF in a molten bath of PbO and SiO2. In reviewing the various theories which have been used to explain catastrophic oxidation, Kubaschewski and Hopkins 8 favor the liquid layer theory, but note that, ".. .as experimental observations are not altogether in agreement with this theory (liquid layer theory), one should consider it a necessary but not a sufficient condition." In contemplating the liquid layer theory, it appears that sufficient evidence has not been presented to establish the theory beyond question. As a means of further clarification, a program of research was undertaken to determine in greater detail the mechanism of accelerated oxidation as caused by lead oxide. The first part of the program deals with a comparison of the oxidation of both AISI 446 stainless steel and chromium metal in the presence of lead oxide, vs the oxidation of these two materials in air alone. These comparisons are made at a number of different temperatures, most of which are below the melting point of the surface oxides. The second part of the program is concerned with a presentation of an alternate theory of accelerated oxidation exemplified by the system Cr-PbO-Air. PROCEDURE AND RESULTS Several experimental methods are commonly used to follow the progress of oxidation. One of these, the weight-gain method, was chosen for this work. This procedure requires that a specimen of the alloy be weighed, oxidized for a given period of time at an elevated temperature, and reweighed—the difference between the two weights being noted. The weight gain of the specimen represents the amount of oxygen acquired from the atmosphere to transform a portion of the specimen to oxide. In those cases where there is a tendency for the specimen or oxide to volatilize at the testing temperature, additional data must be collected so that a correction factor can be determined. This factor must be applied to the weight change in order to ascertain the actual amount of oxidation which has taken place. The specimens used for this work were 1 1/2 in.

Jan 1, 1963

By R. B. Falk, F. E. Luborsky

Small iron particles in mercury pow by diffusion of iron atoms through the mercury. Iron particles, with diameters about 200Å, have been stopped from gvowing in size, even up to the boiling point of mercury, by the addition of a third metal. The metal additives were classified according to whether they produced 1) negligible inhibition, 2) some inhibition, or 3) marked inhibition of the growth rate. Metal additives which resulted in a marked inhibition of iron-particle growth form a stable monolayer on the iron particles. Only those metals which may form compounds with iron showed this pronounced effect. The one exception was manganese. Metal additives which resulted in no inhibition of growth, showed no evidence for adsorption and could not form compounds with iron. None of the added metals, even those which strongly inhibited particle gvowth and thus had adsorbed on the particle surface, actually formed any compounds with iron at room temperature at the concentrations required for complete inhibition of growth. ThE stabilization of the size and shape of metallic particles dispersed in liquid metals has important technological implications in "homogeneous" nuclear reactor systems and in elongated particle magnets and is directly related to the problem of corrosion by liquid metals. Growth in these par-ticulate systems may occur by diffusion of atoms from surfaces of high energy, corresponding to a small radius of curvature, to surfaces of low energy or large radius of curvature. Thus, small particles dissolve and large particles grow to increase the average diameter of the particles. reenwood&apos; has analyzed this diffusional growth of a system of particles containing a distribution of sizes and has obtained reasonable agreement with data on the growth of uranium particles in liquid sodium and UPb3 particles in liquid lead. Greenwood concluded from his analysis that the chief factors promoting stability of dispersed particles growing by a diffusion mechanism are large mean radius, high density, and low molecular weight with low values of the solute diffusion coefficient, the solubility, and the interfacial tension. In the case of elongated particle magnets the system of interest is that of iron particles dis- persed in mercury. Previous work,2 concerned with "spherical" particles in mercury, has indicated a diffusional growth mechanism similar to that discussed by Greenwood. It has also been shown3&apos;4 that the addition of certain metals to Fe-Hg dispersions results in the formation of a monatomic layer of adsorbed atoms on the iron particles. In this paper we will show that this adsorbed layer provides the decrease in the iron-atom flux through the mercury necessary for size stabilization. The general features of the stabilization of the iron particles in mercury through the formation of this "coating", with many different metals, will be described. EXPERIMENTAL RESULTS Iron particles in mercury were prepared by elec-trodeposition,2, 4 and measurements2, 4 of coercive force HCi) and saturation induction (Is) were used to follow changes in size and composition during thermal aging in the mercury in the presence of relatively small amounts of an added metal. In Fig. 1 is plotted the aging behavior of spherical particle samples in the presence of some of the added metals to be discussed. The iron particles have an initial average diameter of 180Å. Each sample was aged for 20 min at each temperature in succession in the presence of a twofold to threefold excess of added metal over that required to form a monolayer on the iron particles. These are the same samples used previously3 to study the adsorp-

Jan 1, 1965

By D. L. Sikarskie, B. Paul

A theory is presented for the static penetration of a single rigid wedge into brittle material. The material considered is one which exhibits both crushing and chipping phases in the penetration process. If the wedge angle and three parameters are specified, the theory predicts forces and associated penetrations during both the crushing and chipping phases. For certain ranges of the parameters agreement with the limited experimental data available is promising, except for the initial phase of the penetration process where refinements on the proposed theory are required. The theory also predicts that for certain values of the wedge angle and other known parameters, the chipping process does not occur and penetration is due entirely to crushing. When viewed in detail, the drilling of rock by percussive means involves, among other things, system dynamics, actuating forces, dynamic stress levels and the actual penetration mechanism of the bit into rock. If rock drills are to be significantly improved, a thorough understanding of the entire system would be highly desirable. In this paper, one aspect of the system, namely the mechanism of penetration of the bit, will be studied in detail. It has been found1 that for the velocities encountered in percussive drilling (of the order of 20 fps) a static analysis adequately describes the penetration process, at least for some rocks. Hence, we will only be concerned with describing analytically the static penetration of a single wedge shaped tool. It will further be assumed that the wedge is long enough to permit a two-dimensional analysis. Numerous authors have studied the static penetration of a wedge into rock. The following papers give some of this work and provide additional references to other work. Cheatham2,3 has assumed the rock to behave plastically and has obtained force-penetration equations for both Coulomb-Mohr and parabolic yield criteria. Evans and Murrell4 have studied the penetration of two types of coal and have found equations relating the penetration characteristic (P/dSc) to wedge angle for the various strength coals tested. race&apos; attempted to find a correlation between hardness as determined by indenting the material and other mechanical properties and concluded that the results of such a test are generally inconclusive. His paper however contains a very extensive bibliography on the indenting of many different materials. Gnirk6 also has a fairly complete literature review on the static penetration of rock, of which he makes use in his indexing studies. A wide range of behavior is found in the wedge penetration of different rocks under different external conditions. For example, a rock that is essentially elastic-brittle at standard pressure may become elastic-plastic at a high confining pressure.7 It has also been observed4. " that some rocks will merely be crushed and indented by a wedge, whereas others will crack and form chips, furthermore the existence or non-existence of chips depends in great measure on the geometry of the indentor, type of rock and the depth of penetration. Hence, to attempt a theory which embraces all possible behavior is not practical at present. We, therefore, confine our attention to predicting the force-penetration characteristic and volume removal behavior for a type of rock which exhibits both crushing and chipping phases in the penetration process. Such behavior is characteristic of the harder rocks such as granites and represents a more difficult problem than the behavior of softer rocks such as coal and certain sandstones. In this preliminary study an attempt will be made to describe only the essential features of this complex penetration process. A qualitative description of these essential features is obtained with the aid of Fig. 1.1, where a wedge of vertex angle 28 is shown at some intermediate stage of the penetration process. As the wedge advances, the rock is fragmented (i.e. crushed) in some local region surrounding the wedge, the shape of this region being unknown. Simultaneous with the fragmentation in this local region, essentially elastic stresses are assumed to be building up in the surrounding rock. When a certain penetration level, di+1, is reached, the stresses along some surface are sufficient to cause failure and a chip is thus formed. The process now repeats, i.e., a crushing phase followed by the formation of a chip.

Jan 1, 1965

By R. H. Maes, R. E. de Strycker, J. J. Jacobs

A simple method, based upow density measurements, has been perfected in order to determine the critical temperature of liquid miscibility gaps. Applied to the Pb-Cu system, it yielded a value of 980"C for the critical tempevature of the miscibility gap in this system. In the Pb-Cu-As system, the immiscibility between lead and coppev evolves towards a miscibility gap be-tween lead and copper arsenides; the liquation temperatures were determined joy different avsenic contents and values up to 1065°C were found. Combining these results with those given by the clzemical analysis of liquid samples taken at dqj.event temperatures, it has been possible to make a cowzplete determination of the liquid miscibility gap in this ternary system. In the Pb-Cu-Fe-As system, a ternary miscibility gap has been discovered; for some compositions, the molten and slowly cooled alloys separated into three layers: an Fe-As alloy, a Cu-As alloy, and a lead-rich layer. ThE Pb-Cu-Fe-As system is of interest for the study of the behavior of speiss or arsenical alloys in the lead blast furnace. The speisses produced in lead metallurgy generally have complex compositions, but they may be considered, in a simplified way, as a mixture of copper, iron, nickel, and cobalt arsenides; there is often an excess of metallic constituants, and lead is also present. Speisses liquate from lead at various temperatures according to their composition; according to industrial experience, iron speisses are already distinctly separated from lead bullion at temperatures of the order of 1100°C, whereas copper speisses only begin to liquate from lead at much lower temperatures, near their solidification points, which were never accurately determined until now. Besides these considerations of practical interest, the present study is of importance for the knowledge of liquid miscibility gaps. Particularly the data of the literature about critical temperatures of liquation are incomplete and present sometimes considerable disparities; a method has been perfected in the present work to determine these temperatures with accuracy. MISCIBILITY GAP OF THE LEAD-COPPER SYSTEM Review of Literature. The generally admitted outline of the Pb-Cu equilibrium diagram is shown by Fig. 1. This figure, with its numerical indications, is taken from the review of Hansen and ~nderko;&apos; compositions are given in weight percent, as will be the case in the remainder of this paper. Between 36 and 87 pct Pb, Pb-Cu alloys separate into two layers, in the liquid state. The critical temperature of liquation has been much disputed in earlier studies, where values up to 1500°C were claimed; but a reliable study of ~ohnen,&apos; based upon differential thermal analysis, gave a value of 990°C; this rather low value was verified by a recent study of Schiirmann and Kaune? based upon thermodynamical properties of the Pb-Cu alloys, which predicted a value of 1002°C. Liquid miscibility gaps were also studied by means of chemical analyses. Satisfactory results were generally obtained with the method consisting in taking samples of the two liquid layers in equilibrium and submitting them to chemical analysis; however, this method becomes delicate near the critical temperature. With this technique, Friedrich and waehlerts ob-

Jan 1, 1968

By D. R. Uhlmann, T. P. Seward, O. Chalmers

A magnetic field has been used to damp out liquid convection during the solidification of castings. Ingots of Al-2 pct Cu were poured from various sperheats (10" to 65°C) into a room-temperature graphite mold located between the poles of an electromagnet. The effect on ingot structure of varying the strength and time of application of the magnetic field was investigated, and the results are interpreted in light 01- current theories of ingot structure. In the solidification of alloy systems, thermal gradients in the melt and differences in density of the constituents are expected to produce convective motion in the liquid. The thermal convection caused by temperature gradients should be present during the freezing of both alloys and pure materials, while the solute convection caused by density differences should be of concern only for alloys. The possible influence of this natural convection on the structure of castings has been discussed by Ruddle,&apos; Chalmers, and Tiller. More recently, an experimental investigation of such effects was carried out by Cole and Bolling. These workers took stationary molds containing lead- and tin-based alloys and. using a brine solution. quenched them from various superheats above their respective liquidus temperatures. When grids were placed across the mold diameter to suppress convection, the grid region was later observed to contain a larger percentage of columnar grains than other parts of the casting or than similar castings poured without the grids present. From thermocouple measurements, they found the temperature in the grid region to be higher than that in the region outside the grids throughout the freezing process, and concluded that convection should be an important mode of heat transfer in normal castings. In addition to such convection arising from thermal gradients or differences in density of the constituents, a further source of convective motion, which we shall call "momentum convection", may be particularly important in small (laboratory-size) castings. By this term we refer to the convective motion produced in the initial pouring process, a motion which is damped out as the linear momentum and angular momentum of the liquid are dissipated. Besides these sources of convection, which are present to some extent in all castings. external influences may also be effective in producing convective motion. Examples of such "forced" convection which have been discussed in the literature include stirring the melt, shaking the mold, rotating the mold, and so forth The purpose of the present study was to examine the effect of convection on the structure of castings poured into a cold mold. In this work, a magnetic field was employed to damp out the convective motions within the liquid, and hence control the extent of con- The effect of a magnetic field on convective currents in a conducting liquid is to induce eddy currents which oppose and damp out the convective motion. vection. The technique of magnetic damping was previously used by Utech and F1emings" in their study of convective effects on impurity banding and other substructure in alloy crystals grown unidirectionally from the melt. In their investigation, fields as small as 200 gauss were found to be effective in damping the convection currents which influence the banding. I) EXPERIMENTAL TECHNIQUE In the present study, alloys of A1-2 wt pct Cu were poured into a room-temperature graphite mold. The mold, in the shape of a circular cylinder, 2 in. ID, 2-3/4 in. OD, 5 in. inside height, and 5-1/2 in. outside

Jan 1, 1967

By Jr. Ensign C. O., J. W. Trammell

The two principal methods of room-and-pillar mining practiced at White Pine make it important to predict variations in the thickness and rock types of a stratum called the upper sandstone. In full column mining, the nearly barren upper sandstone occurs between two ore horizons (called upper shale and parting shale, respectively), and is mined as a part of the ore column. In general, full column mining is not practiced where excessive upper sandstone thickness causes full column ore grade to be substantially lower than parting shale ore. In parting shale mining, the upper sandstone forms the mine roof. Difficulty in predicting upper sandstone character arose because the extent and directions of its variations were not apparent in the property&apos;s diamond drilling, in which holes are spaced on 1000-ft centers. A study of sedimentary features, undertaken to improve the predictability of trends in the upper sandstone, led to an interpretation of the sedimentary environment in which the sandstone was deposited. Ripple marks, mud cracks, and cross-bedding, as well as other, less well known features, such as channel casts, flute casts, and current crescent casts, were mapped or recorded. These data, coupled with the knowledge of regional facies changes gained from studying drill core, show that the upper sandstone was deposited by a series of streams flowing northeastward over the underlying parting shale. Deposition of the lowermost bed of the upper sandstone, near the ancient shoreline, was locally preceded by erosion of the parting shale, and the greatest thicknesses of sandstone are found in channels scoured out of the parting shale. Awareness of the rather strong linear trends in the upper sandstone makes it possible to project continuous areas of thick upper sandstone through apparently isolated "highs" in the upper sandstone thickness contour map, which is based on drill-hole information. Since local exceptions to general stratigraphic trends exist, a method was also needed for esti- mating, in detail, the thickness and degree of shali-ness of upper sandstone forming the roof in active parting shale mining areas. Because locally the parting shale was thinned by erosion during the deposition of upper sandstone, parting shale thickness is inversely proportional to upper sandstone thickness. Utilizing information taken from short drill holes into the roof, curves were constructed for the correlation of parting shale thickness (measurable in the mine) with the upper sandstone total thickness, as well as the thickness of its basal member. The White Pine mine is situated in Ontonagon County, approximately 6 miles south of Lake Superior, in the Upper Peninsula of Michigan (Fig. 1). The orebody mined by White Pine Copper Co. occurs in the lowermost 20 to 25 ft of the Nonesuch formation, a series of middle- to upper-Keweenawan shales, siltstones, and sandstones. This paper discusses the solution of problems in ore grade control and mine planning arising from the presence of an essentially barren sandstone stratum within the orebody. THE OREBODY AND MINING TYPES A brief description of the orebody and mining methods is necessary to show how the upper sandstone affects mine planning and grade control. In the mine, the cupriferous zone of the Nonesuch shale is divided into three major portions, as shown in Fig. 2. Lying conformably on top of the Copper Harbor formation (or "lower sandstone") is the lower part of the orebody, the parting shale. The parting shale is overlain by the upper sandstone. On top of the upper sandstone, the basal portion (8 ft) of the upper shale is quite similar to the parting shale. Both the upper shale and parting shale are divided into a number of smaller stratigraphic units, each having a characteristic copper content, as shown by the histogram of Fig. 2. Other authors1&apos;3 have dealt in detail with the stratigraphy and mineralogy of the orebody. However, in this article, we are concerned with the major subdivisions only, and it will suffice to note the low copper content of the upper sandstone, shown in Fig. 2. At present, two types of room-and-pillar openings are created in the orebody. In one type, called parting

Jan 1, 1964

By F. V. Lenel

Resistance sintering under pressure is a method of hot pressing in which a powder compact is subjected to pressure and simultaneously heated by passing a low voltage high amperage current through it. Equipment for this process is described. Its basic characteristics such as resistance requirements for powders and compacts, temperature distribution in compacts, and gas reactions during resistance sintering are discussed. Examples of the sintering process in compacts made of a single metal or an alloy powder and in compacts made of mixtures of powders are presented. Potential commercial applications of the process are evaluated. THE usual sequence of operations in commercial powder metallurgy is to compact metal powders in a die at room temperature and sinter the compact subsequently without applying pressure while the compact is in the sintering furnaces. In hot pressing, on the other hand, the compacting and the sintering are combined. The loose powder, or in many cases a cold formed compact, is inserted into the die which is held at the hot pressing temperature, the compact is left in the die until it attains its temperature, and then pressure is applied and maintained for a given length of time. The die may be heated by surrounding it with a suitable furnace, by high frequency induction heating, or by passing current through the die. In hot pressing, compacts of high density and good mechanical properties can be produced in a relatively short time. One of the principal difficulties in hot pressing is the lack of suitable die materials which will have adequate strength at the hot pressing temperature. This difficulty can be overcome at least partially by heating only the material to be hot pressed without heating the die directly. This can be done by passing a low voltage high amperage current through the material and simultaneously subjecting it to pressure. Either loose powders or compacts can be hot pressed by this method in which the desired temperature is produced by the power dissipated in the metal powder. This method of hot pressing has been termed "electrical resistance sintering under pressure." The most significant differences between this operation and conventional hot pressing are: 1—The sintering times are very short, usually a fraction of a second and at most a few seconds. 2—The powder and die are initially cold. 3—Heat is generated within the powder itself and not conducted in from the die. 4— The pressure used is high. 5—Cooling following sintering is rapid, amounting to a quench. Resistance sintering of metal powders under pressure has been suggested repeatedly. In 1933, G. F. Taylor&apos; proposed an apparatus which consisted of an insulating tube made of glass, or a ceramic filled with the powder to be pressed, and plungers above and below the powder. The powder was to be heated by passing an electric current through it and pressure was then applied. The apparatus was intended principally for hot pressing cemented carbides. Although the principle of resistance sintering under pressure is clearly shown in this patent, few details are given. A sintering period of a second or a fraction of a second, which is terminated by the operation of an inertia switch, is mentioned but no values for current density are mentioned. Low pressures, such as atmospheric pressure or a somewhat higher pressure exerted through a lever, are suggested. W. D. Jones&apos; discussed electrical resistance sintering under pressure and proposed using resistance welding apparatus for this purpose. In a patent issued in 1944, G. D. Cremera described a method of electrical resistance sintering under pressure using resistance welding apparatus which was to be applied mostly to nonferrous metals such as copper, brass, bronze, and aluminum. Current densities of 400,000 amp per sq in., sintering times of 1 and 2 cycles of 60 cycle current, and pressures of 5 to 10 tsi are proposed. Cremer suggested using an

Jan 1, 1956

By J. T. Williams

The equilibria in the V-Zr alloy system were investigated by solidus temperature determinations, thermal analysis, dilatometry, electrical resistance measurements, microscopic examination, and X-ray diffraction analysis. There is a eutectic reaction at 1230°C between a compound, V2Zr, and a solid solution containing 10 pct V in ß zirconium. V2Zr decomposes at 1300°C into liquid and a solid solution containing about 10 pct Zr in vanadium. The eutectic composition is probably about 30 pct V. A eutectoid reaction between V2 Zr and a zirconium takes place at 777°C at a very high rate. The eutectoid composition is 5 wt pct V. The limit of solubility of zirconium in vanadium was estimated to be 5 pct at 600°C. No attempt was made to determine the liquidus for the system. THE recent availability of large quantities of high purity zirconium has stimulated the study of zirconium binary systems. The equilibrium diagram for the V-Zr system has received little attention, however. Wallbauml appears to have made the first report concerning the equilibria in these alloys. He reported the existence of a compound, V2Zr, having the MgZn2 ((214) type of structure with a, = 5.277 kX and c° = 8.647 kX. Anderson, Hayes, Roberson, and Kroll2 made a survey of some potentially useful zirconium binary alloys and found that zirconium probably dissolves a small amount of vanadium. They reported the probable existence of a compound between the two elements and suggested that the zirconium-rich solid solution undergoes a eutectic reaction with this compound. Pfeil," in a critical review of the existing information, estimated that the solubility of vanadium in zirconium is less than 4.7 pct and probably less than 1.8 pct. Rostoker and Yamamoto&apos; proposed a partial diagram for the V-Zr system in a survey paper on vanadium binary alloys. Their diagram indicates the compound, V,Zr, a eutectic reaction at 1360°C, a peritectic reaction at 1740°C, and a limit of solubility of zirconium in vanadium of about 3 pct. They obtained no information on the equilibria in the zirconium-rich alloys. In view of the potential utility of the V-Zr alloys and the incomplete knowledge concerning the equilibria in the system, an attempt was made to establish the constitutional diagram. Preparation of the Alloys Raw Materials: The vanadium for making up these alloys came from the Electro Metallurgical Corp. Zirconium came from two sources. In the beginning of the investigation, sponge zirconium from the Bureau of Mines was used in making some of the alloys. Later, iodide metal made at the Westinghouse Atomic Power Development Laboratories became available. This material was used in the preparation of all the dilatometric and resistance specimens and about two-thirds of the solidus temperature specimens. A typical manufacturer&apos;s analysis of the vanadium is shown in Table I. No other analysis of the vanadium was made. The metal contained a dispersed second phase and did not have a sharp melting point. Typical results of spectrographic analysis of the Westinghouse zirconium are shown in Table 11. These data indicate a very high purity. The Bureau of Mines sponge metal was probably less pure but had good ductility. Melting: All of the alloys used in the investigation were made by melting pieces of vanadium and zirconium together in a dc electric arc furnace similar to those of Geach and Summers-Smith, craighead, Simmons, and Eastwood," and others. Melting was done in an atmosphere of helium scavenged of residual air by the preliminary melting of a separate charge of zirconium. Each ingot was turned over and melted at least three more times before removal from the furnace to aid in the attainment of homogeneity. Alloys prepared for use in the investigation are listed with the results of solidus determinations in Table III with the exception of the following compositions upon which no solidus determinations were made: 0.29, 0.54, 4.57, and 5.55 pct V. Analysis: The weight of each ingot made from iodide zirconium was within 0.1 g of the total weight of the initial charge, about 90 g. Since each component of each charge was weighed to the nearest 20 mg for amounts less than 10 g and to the nearest 0.1 g otherwise, the gross composition of an ingot could be calculated accurately. Chemical analysis for the vanadium content of several alloys agreed

Jan 1, 1956

By N. A. Warner

Dense cylindrical specimens of artificial hematite were reduced in hydrogen over a range 0-f total pressures between 0.1 and 1.0 atm and temperatures between 650" and 950°C. Hydrogen reduction at a total pressure of 1.0 atm in the presence of a chemically inert diluent and reduction in hydrogen/water mixtures were also studied Znter.face movement during reduction was followed using metallographic techniques. Substantiu1 partial-pressure gradients are shown to exist in both the gas film and the reduced iron layer. The results are consistent with a mixed-control mechanism involving an interaction of gaseous-difusion effects with a first-order reversible chemical reaction at the iron/wüstite interface. Over the hydrogen-pressure range of 0.1 to 1.0 atm, the rate is not directly proportional to the hydrogen pressure in the bulk gas stream. Nitrogen dilution of the reducing gas masks the true effect of hydrogen partial pressure by influencing the effective molecular dif-fusivity and gives an apparent first-order relationship. The transport of hydrogen and water vapor across the iron layer is consistent with molecular gaseous diffusion rather than a Knudsen diffusion mechanism. ALTHOUGH the literature on the reduction of iron oxides is very extensive, it is only comparatively recently that a definite quantitative model for the reduction of hematite has been proposed. This development has been largely due to the work of McKewan. An extensive literature review of previous work has recently been compiled by Themelis and Gauvin. Exclusive control by gaseous diffusion in the reduction of hematite has been recently advocated by Kawasaki and coworkers.&apos; Although the results of the latter workers are of considerable practical significance, the extreme porosity (approximately 30 pct) of their specimens precludes their application to a kinetic evaluation of the reduction of dense polycrystalline hematite. The salient features of the kinetic model used by McKewan for the gaseous reduction of dense hematite at temperatures in excess of 560°C (when FeO is stable) are: 1) The reduction is a topochemical reaction taking place through &apos;the series FeO/FesO4/FeO/ Fe; 2) The gas-solid type of reaction takes place only at the FeO/Fe interface and the internal reduction: FeO proceeds by solid state diffusion; 3) The iron layer formed is quite porous and offers negligible impedance to gaseous diffusion; 4) The rate-controlling step in the reduction is a chemical process occurring at the Fe/FeO interface. In view of the relatively fast reaction rates involved when hematite is reduced in a stream of hydrogen, it seemed possible that the simplifying assumptions of negligible diffusional resistance in both the bulk gas phase and the reduced iron layer could introduce serious errors in the interpretation of experimental reduction data. The present work was initiated to elucidate the effect of gaseous diffusion on the reduction rate and to determine its possible influence in the formulation of the reaction mechanism at the metal/oxide interface. EXPERIMENTAL Procedure and Apparatus. Dense cylindrical compacts of artificial hematite were prepared by sintering pressed compacts of reagent-grade ferric oxide, containing 99.6 pct FezO, in air at 1200°C for 80 hr. The resulting compacts had a final porosity of less than 2 pct and a height and diameter of approximately 1 cm. Variations in sintering atmosphere from air to pure oxygen, in sintering time from 20 to 80 hr, and temperature from 1150" to 1250°C had no effect on the subsequent reduction behavior. The reduction of the samples took place in a vertical tube furnace in which was placed a 1-in. transparent silica tube. The 24-in. heated length

Jan 1, 1964

By Phillip B. Bucky, R. S. Taborelli

The term "longwall mining" is best known to coal men, although modifications of the method are continually being used in other fields. Longwall mining is of interest today because it makes for greater inechanization and better supervision at the working face, and therefore lower production costs and increased safety. It is an approach to the factory system of production, the high efficiency of which as compared to that in mines is well known. There is available an accumulation of experience regarding this work but coordination of experience is lacking and this series of experiments was undertaken to help serve this purpose. Model experiments were used for reasons of safety, cost and time, and because it has been previously shown1,2,3 that a scalar inodel built of the same material as the prototype will behave in a manner similar to that of the prototype, with time effects the same, if the effective weight of the model is increased in the same proportion as its linear dimenbions are decreased. This is accomplished by substituting a centrifugal for the gravitational field; i.e., placing the model in a centrifuge and controlling the strength of the centrifugal field by varying the rate of rotation. Many factors affect the behavior of the geologic material overlying a bed, therefore results and conclusions contained in this paper apply only to the conditions presented herein. Longwall Mining What Is Longwall Mining?—Agreement should be reached as to what constitutes longwall mining and therefore what constitutes a longwall face. Let longwall mining be defined as the mining procedure in which the roof and overlying bed breaks tend to parallel the working face. In the ordinary stope or room-and-pillar mining the roof breaks tend to

Jan 1, 1938

By A. W. Jenike

A theory of gravity flow of bulk solids like ores, concentrates, coal is outlined and an example of design for unobstructed flow is given. Mass-flow and plug-flow patterns are defined and related to the geometry of the bin. The effect of these patterns on segregation, first in-first out flow, and caking is described. This paper is a report on the results from the AIME sponsored project "Flow of Bulk Solids" which was carried out under the author&apos;s direction at the Utah Engineering Experiment Station, University of Utah, during the period 1956-1962. The theory has been applied to the design of numerous industrial storage plants. A QUANTITATIVE METHOD OF DESIGN FOR FLOW There is hardly an industry which does not store and handle solid materials in bulk form. When the volume of the solids is large, gravity is invariably relied upon to cause the solids to flow out of storage, through channels and chemical reactors. Such materials as ore, coal, cement, flour, plastics, clay, soil, to which a general term of bulk solids is applied, flow by gravity or are expected to flow by gravity in thousands of installations and by billions of tons annually. The chemical and process industry depends on gravity flow of solids in process and out of storage. Mining relies on gravity flow in block-caving, in ore passes, as well as in storage and loading bins. Agriculture uses gravity flow in elevators, in feed plants, and in farm silos. In practice, satisfactory gravity flow is seldom attained. Such problems as doming, piping (rat-holing), lack of live capacity, segregation, deterioration in storage, harass industry. In spite of the universality of these problems, until recently, there has been no theory of solids flow, no method of predicting how a solid would behave, and how the channel, bin, or hopper should be designed. In the late nineteen forties and early fifties, the theory of plasticity came of age and simultaneously the computer became widely available. These two factors permitted the development of a theory of flow of bulk solids. The work was conducted during the years 1953-1961. In conjunction with the development of the theory a method to measure the flowability of solids was also developed1-7 and the theory was ap-plied to the design of ore pockets, storage plants,16 chemical reactors and adsorption columns. A comprehensive report of this work is given in a previous work." THEORY OF FLOW The most important question which a theory of flow has to answer is: what conditions need to be satisfied for a given solid to flow uniformly and without obstructions through a channel? First of all, the following observations are made:

Jan 1, 1967

By Wheeler Davey

WHEN metals are subjected to mechanical working, such as rolling, one of the phenomena that take place is a movement of the crystals of the metal into a system or systems of orientation which bear an angular relation to the direction of working. In rolled metal the crystals may orient themselves with respect to the rolling plane or the rolling direction, or both. Direct evidence of the existence of preferred orientations in rolled metal has been obtained by numerous investigators from X-ray patterns of the rolled metal. Most of these workers used a monochromatic pinhole method. Where the degree of preference for a single orientation is high the monochromatic pinhole pattern may be interpreted by a geometrical calculation based on the angular relationships of the planes in the crystal and the diffraction arcs on the X-ray pattern. This process may be aided by the geometrical construction of the locus sphere (Langenkugel), and reflection circles of Polanyi,1 Polanyi and Weissenberg,2 and Mark and Weissenberg,3 and of the plane polar figures (flachenpolfigur) of Franz Wever.4 R. M. Bozorth5 solves X-ray patterns from metals having a single preferred orientation by means of a series of charts which when placed over the X-ray pattern give directly certain angular relationships, between the crystal planes and the surface of the specimen.

Jan 1, 1929

By Clair Upthegrove, Elbert S. Rowland

THE widespread interest in the alloys of beryllium with copper is due principally to the fact that certain compositions show very favorable precipitation-hardening characteristics and are, in fact, the strongest nonferrous alloys known at the present time. Investigations have been conducted by Masing and Dahl and others on the physical characteristics of ternary beryllium alloys, in an attempt to produce precipitation-hardening effects with a smaller amount of beryllium than that required in the binary copper-beryllium alloys (2.0 to 2.5 per cent Be). The results of these investigations show that ternary alloys of beryllium with Cu-Sn, Cu-Al and Cu-Zn, among others, exhibit promising improvements in physical properties after precipitation-hardening treatment and, while the increases in properties obtained are not as high as those realized from the binary copper-beryllium alloys, excellent properties are: obtained with beryllium contents averaging about one-half the amount required in the binary alloys. This paper presents the results of an investigation of the equilibrium relations of certain alloys containing copper, tin and beryllium. The alpha-phase boundary of the ternary system is given for alloys containing from. 0 to 1 per cent beryllium. The equilibrium conditions beyond the alpha-phase boundary are presented for alloys containing up to 32 per cent tin and 1 per cent beryllium. The data were collected by means of

Jan 1, 1935

By Paul G. Shewmon

The migration of slightly solzrhle spherical particles through a solid under the infllrence of a temperature gradient is analylzed for the cases of various transport mechanisms. It is shown that the variation of the velocity of the particles with radius, r, depends on the dominant mechanism of matter transport around or through the inclusion. Thus the velocity varies as r-1 jor surface-dijf1~sion controlled migratzon. is independent of r for volume diffusion in either phase, and varies as if the rate is determined by an interfacial reaction (n is the order of the interfacia1 reaction). These same results hold tor the migration of a cvlinder of length much greater than its radius. and .for other types of potential gradients, e.g., an electrical field. These equations are combined with recent electron-microscopic ohservations to show that the rate of migratzolz of small helium-silled bubbles through copper is determined by surface diffusion of the metal atoms. With these equations and the temperature gradients attainahle in an electron-microscope joil, the dominant transpar/ mechanism (or any migrating pnrtzclos can he determined. AS a result of fission, rare-gas atoms are created in the hot fuel element of a nuclear reactor. These insoluble atoms precipitate to form bubbles of the gas in the fuel. The subsequent migration and coalescence of these bubbles is thought to play a dominant role in the swelling of fuel elements which in turn can limit the fuel-element life. As a result of this practical problem and because of the relative simplicity of the system, workers in several laboratories have imbedded rare-gas atoms in metals with the aid of an accelerator and studied the formation and behavior of the bubbles that form on annealing. Recently Barnes and Mazey have studied the migration of such helium-filled bubbles in copper foils using the electron microscope and the temperature gradient that can be induced in the foil by the electron beam. They found that the smaller bubbles moved faster than the larger ones. It is shown below that, if this is true, the rate-controlling step in their migration is surface diffusion of metal atoms instead of volume diffusion, vapor transport, or an interfacial reaction. The analysis given is in no way limited to gaseous inclusions so the variation of velocity with particle size for any relatively insoluble precipitate particles could be used to obtain information about the relative importance of surface diffusion, volume diffusion, and interfacial reaction in such two-phase systems. VOID MIGRATION IN A TEMPERATURE-GRADIENT ANALYSIS Barnes&apos; studies indicate that helium does not dissolve or diffuse in metals to a measurable extent, so that when the voids move they must do so through the movement of metal atoms from the leading to the trailing sides of the void.&apos; Thus voids might move by surface diffusion, diffusion of vacancies in the surrounding lattice, or vapor transport. We consider first the case of flow by surface diffusion alone. We assume that there is a force, in the sense of irreversible thermodynamics, tending to move the atoms around the bubble in a given direction. We designate this direction as the x axis and the force per atom as Fa. The net rate of flow of atoms from one side of the bubble to the other, under the influence of this force, is the surface flux, Js, times the cross-sectional area available for flow, A,. A, is taken as the circumference times the thickness of the high-diffusivity surface layer 6. Thus where 51 is the atomic volume. For bubbles to advance a distance dx, a volume equal to nr : dx must flow around the void so that The minus sign enters because the bubble moves in the direction opposite to that of the force on the atoms. Equations for volume diffusion and diffusion through the vapor can be obtained in the same manner. For vapor transport we have Ag = and the ratio of the densities of metal atoms in the gas and the solid (pg/pl) enters so that

Jan 1, 1964

By W. B. Nelligam, J. Tittman

Refinements in radiation logging techniques during recent years have involved increasing usage of scintillation ditectors. These detectors produce voltage pulses whose heights are related to the energies of the gamma rays which initiate them. Analysis of the gamma-ray spectrum, as indicated by the pulse heights, yields information about the chemical elements composing the formations surveyed. Refined scintillation counter techniques can furnish chemical information concerning earth formations in situ, from a study of the gamma-ray spectra emitted by the formation either naturally or as a result of neutron bombardment."&apos; Accompanying the rising interest in gamma-ray scintillation spectroscopy, there has been increased activity in the development of accelerator-type neutron sources (in contrast to encapsulated chemical-mixture sources). Such neutron generators are attractive for several reasons: (1) they greatly reduce radiation hazards to personnel; (2) there is a great reduction in contamination danger if they are lost in the hole; (3) they can produce larger neutron intensities than can conveniently available encapsulated sources; and (4) they are capable of being pulsed, thus permitting new techniques in logging. In the past, both accelerator and encapsulated neutron sources have been used by others in conjunction with scintillation-detector pulse-height analysis. The results have not been too encouraging, due to the interference among different gamma-ray spectral lines and to the fact that the gamma-ray peaks were not too clearly distinguishable above the large and ill-defined background "noise".&apos;.&apos; This paper is a status report on laboratory studies of a technique using a borehole accelerator as a neutron source, which gives an improved scintillation spectrum, thus permitting more accurate chemical analyses of the formations penetrated. The Schlumberger-accelerator neutron source&apos; is presented; the origins of inelastic and of thermal-neutron, capture gamma rays are dis- cussed, and results are given for some laboratory measurements performed in borehole geometry. THE NEUTRON GENERATOR-TUBE ACCELERATOR In the attempt to develop an accelerator neutron source which would have the desirable properties outlined in the "Introduction" and which would operate satisfactorily under borehole logging conditions, a small-diameter, cylindrical, neutron generator tube has been developed. This tube utilizes the principle of accelerating deuterons (heavy hydrogen nuclei) by a high voltage so that they bombard a tritium (heavy, heavy hydrogen) target. The resulting reactions produce large numbers of neutrons of 14 Mev energy. Furthermore, the tube is permanently sealed and, thus, the use of pumping techniques in the sonde is avoided. The tube consists of a pressure control, an ion source in which the deuterons are stripped of their electrons, an accelerating gap down which the deuterons are sped by the high voltage, a secondary electron suppressor and a tritium-loaded target. Obviously, the tube requires auxiliary circuitry in a sonde for the control of the ion source, tube pressure, high voltage and other operating variables. Neutron yields, both continuous and pulsed, have been produced under simulated field conditions in the range between 1 and 10 times those conventionally used in neutron logging applications. For the experiments discussed in this paper, neutron-pulse repetition rates in the range between 500 and 5,000 pulses/sec are adequate. Furthermore, by making suitable adjustments in operating conditions, one can vary the pulse width from a minimum of several microseconds up to dc. As will be seen later, the present application to well logging does not require that one have available pulses of neutrons which are appreciably shorter than the time it takes for fast neutrons to slow down to thermal energy in water. As a consequence, much of our interest has been directed towards the operating characteristics of the tube with pulses longer than about 10 microseconds in duration. Our experience to date shows that tubes can be operated in this manner with reasonably constant, average neutron outputs.

By A. W. Dana, F. J. Shortsleeve, A. R. Troiano

The phenomenon of flake formation which may occur during cooling or room temperature aging of large steel sections is caused by a combination of hydrogen and stress. As such, the transformation characteristics of austenite in a particular section play a major role in the occurrence of flakes. Isothermal and continuous cooling studies demonstrated that flake formation is particularly sensitive to nature, distribution, and relative proportions of the microconstituents in the cooled sections. In the absence of transformation stresses, abnormally high hydrogen contents could be tolerated without flaking. Flakes were not found in the absence of transformation stresses. No correlation existed between average hydrogen content and flake formation where cooling stresses were low. Hydrogen in the steel was a necessary but not sufficient condition for flake formation. DESPITE a long history of laboratory and commercial investigations, the mechanics of hairline crack or flake formation in large steel sections remains unresolved. Recent publications1-5 emphasize that flaking is still a serious problem.* immediate cause of flaking is the pressure exerted by entrapped hydrogen gas that collects in small internal cavities or voids during cooling. When the temperature of steel is low enough to preclude extensive plastic flow, the cavity pressures cause local brittle fractures (flakes). 2—The microsegregation and transformation stress theory. According to this theory, flaking results from stresses arising from the volume changes accompanying the formation of martensite and/or low temperature bainite after a large portion of the section has transformed at elevated temperatures. The proponents of the hydrogen theory suggest that stresses from other sources may be additive to the stresses produced by hydrogen and may, in some cases, provide the increment of stress necessary to produce flakes. However, the critical aspect of this theory is that no stress, except that produced by hydrogen pressure, is believed to be sufficient in itself to produce flakes. For practical considerations, this means that the only safe thermal cycle is one that reduces the hydrogen content to some low critical level below which flaking will not occur regardless of the state of transformation. On the other hand, the transformation stress theory does not rule out the embrittling effect of hydrogen. This theory simply proposes that flaking does not occur in the absence of transformation stresses. With the transformation stress theory as a basis, commercial antiflaking cycles would be designed to complete the transformation of austenite prior to cooling to ambient temperatures. Both of these theories appear compatible with industrial experience. Antiflaking cycles, such as cooling at extremely conservative rates or soaking for periods of time in the upper subcritical temperature ranges prior to cooling to ambient temperatures, may allow the following: 1—diffusion of hydrogen out of the section and, as a consequence, reduction of the hydrogen content to some safe minimum value (hydrogen pressure theory) and 2—complete decomposition of the austenite at rela-

Jan 1, 1956