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By Axel Hultgren

In ingots of silicon-killed carbon steel-made without addition of aluminum, transparent spherical or nearly spherical inclusions, up to about 0.15-mm diameter, are generally present. They may be glassy or crystallized, wholly or partially In steels of high carbon contents they consist of silica,4,12 in steels of lower carbon contents and, in consequence, higher oxygen contents, of iron-manganese-silicate with higher iron and manganese contents the lower the carbon content. Small silicate inclusions, about 0.01 mm or smaller, may occur, just as sulphide inclusions do, in portions that have been the last or nearly the last to freeze,l0,12,17 this location being verified by primary etching.10 Hence, it may be concluded that they have separated from liquid steel that has become enriched in oxygen through segregation during freezing. Silicate inclusions with a diameter larger than about 0.02 mm appear, at least in ingots of moderate size, to be distributed at random in relation to the primary structure and therefore may be assumed to have existed as droplets in the liquid steel before freezing started at the point in question. The average size of such large inclusions has been found to increase toward the axis1' and their amount to increase from the top downward in the axial region.2,10 It has been assumed that this accumulation is a result of the droplets being brought down by the slowly settling metal crystals.2 It has been suggested that large silicate inclusions generally form,6 or may form,'' by coagulation of smaller ones, or by small drops being caught by larger ones during the rising of the latter.9 The occasional occurrence of a large inclusion surrounded by a swarm of small ones has been taken as a proof of this mechanism but, as suggested later, may be explained in another way. It has been pointed out that an inclusion droplet in the liquid steel may grow at the expense of an adjacent one, which implies diffusion of soluble inclusion-forming elements in the liquid steel.'= Further, it has been pointed out that, on cooling in ladle or mold, the metal-silicate equilibrium is displaced in the direction of separation of silicate. Thus inclusions may grow in size and amount.2,3,8,12,13,17,19,20 The exposure of the liquid steel to the oxygen of the air during tapping and pouring, and the consequent reaction, has been mentioned as a source of the formation of inclusions19.20 but this possibility has not received particular attention. Further, the admixture of furnace slag during tapping1 or of refractory material from various sourcesl,20 has been considered. The inclusions due to the latter may be recognized, however, through their resistamce to hydrofluoric acid, imparted by the alumina content. Silicate inclusions are often referred to as the natural and inevitable product of reaction on adding silicon or silicon

Jan 1, 1949

By Axel Hultgren

In ingots of silicon-killed carbon steel-made without addition of aluminum, transparent spherical or nearly spherical inclusions, up to about 0.15-mm diameter, are generally present. They may be glassy or crystallized, wholly or partially In steels of high carbon contents they consist of silica,4,12 in steels of lower carbon contents and, in consequence, higher oxygen contents, of iron-manganese-silicate with higher iron and manganese contents the lower the carbon content. Small silicate inclusions, about 0.01 mm or smaller, may occur, just as sulphide inclusions do, in portions that have been the last or nearly the last to freeze,l0,12,17 this location being verified by primary etching.10 Hence, it may be concluded that they have separated from liquid steel that has become enriched in oxygen through segregation during freezing. Silicate inclusions with a diameter larger than about 0.02 mm appear, at least in ingots of moderate size, to be distributed at random in relation to the primary structure and therefore may be assumed to have existed as droplets in the liquid steel before freezing started at the point in question. The average size of such large inclusions has been found to increase toward the axis1' and their amount to increase from the top downward in the axial region.2,10 It has been assumed that this accumulation is a result of the droplets being brought down by the slowly settling metal crystals.2 It has been suggested that large silicate inclusions generally form,6 or may form,'' by coagulation of smaller ones, or by small drops being caught by larger ones during the rising of the latter.9 The occasional occurrence of a large inclusion surrounded by a swarm of small ones has been taken as a proof of this mechanism but, as suggested later, may be explained in another way. It has been pointed out that an inclusion droplet in the liquid steel may grow at the expense of an adjacent one, which implies diffusion of soluble inclusion-forming elements in the liquid steel.'= Further, it has been pointed out that, on cooling in ladle or mold, the metal-silicate equilibrium is displaced in the direction of separation of silicate. Thus inclusions may grow in size and amount.2,3,8,12,13,17,19,20 The exposure of the liquid steel to the oxygen of the air during tapping and pouring, and the consequent reaction, has been mentioned as a source of the formation of inclusions19.20 but this possibility has not received particular attention. Further, the admixture of furnace slag during tapping1 or of refractory material from various sourcesl,20 has been considered. The inclusions due to the latter may be recognized, however, through their resistamce to hydrofluoric acid, imparted by the alumina content. Silicate inclusions are often referred to as the natural and inevitable product of reaction on adding silicon or silicon

Jan 1, 1949

By R. C. de Lange

An improved replica technique is developed for a nondestructive study of the nucleation and growth of fatigue cracks. Three different growth stages of a fatigue crack were observed. An initial stage of high growth rate is followed by a long period of very slow growth, until finally a third period of high growth rate leads to fracture. THE metal-fatigue process is of a very complex nature. The experimental approach towards fatigue is susceptible to many minor influences which, together with the unavoidable scatter in testing results, makes the interpretation a matter of extreme care. The generally accepted fact that, in a homogeneous material, the fatigue cracks start at a free surface indicates that surface studies may contri-but in an essential way to a better understanding of the fatigue mechanisms. Indeed metallographic observations of the fatigue phenomena have been very fruitful in the last decade. The study of extrusions, intrusions, and persistent slip bands has contributed much to the present knowledge of the early processes in fatigue.' The drawback of the current metallographic techniques lies in the fact, however, that either observations of fatigue phenomena had to be made in a destructive way after an only partly executed test or a sequence of observations on the same test bar was only possible on notched specimens.2 4 In both cases an uncertainty exists as to the nucleation and growth of the fatigue crack in an unnotched test bar, where the crack can start at any place of a large area and develops in a test bar, which is macroscopically

Jan 1, 1964

By C. Ernest Birchenall, Robert F. Mehl

Application of diffusion laws in the customary form to experimental studies in binary metallic solid solutions has shown the diffusion cocfficient to vary with concentration for all systems investigated. The factors determining this variation have not been ascertained. However, the question of whether the diffusion coefficient would become a constant, independent of concentration throughout a single phase if the concentration gradient in the diffusion equation were replaced by an activity gradient has been an attractive subject for speculation. Wagner¹ derived an expression for the force acting on a diffusing particle in terms of the chemical potential. This was also the approach chosen by Jost2 in his monograph. From a somewhat different point of view, Eyring3 and his co-workers have used activity gradients when dealing with diffusion in concentrated nonideal solutions. The only actual test with experimental data is that given by Darken,4 though this is more qualitative than quantitative. Many others opposed the application in any way of thermodynamic concepts to kinetic phenomena, including diffusion processes. A rigorous examination of this principle requires good data expressing the relationship between activity and concentration for solid solutions for which reliable diffusion work has been done. The literature provides such information only for the interstitial solution of carbon in gamma iron (austenite) and for the substitutional copper-zinc solutions (alpha brass). Activity data for carbon in austenite at 800°C and 1000°C are available from the recent work of Smith5 on the equilibrium of hydrogen-methane and carbon monoxide-carbon dioxide mixtures with gamma iron. Combining these with the diffusion studies of Wells and Mehl6 makes possible a quantitative test for this interstitial solid solution. The vapor-pressure determinations for zinc over brass by Hargreaves7 allowed the calculation of activities in this system over a wide range of temperatures and concentrations. Extensive investigation of diffusion in this substitutional solid solution has been carried out by Rhines and Mehl.8 Activities in Brass9 If the vapor phase of zinc is assumed to behave as an ideal gas, any solid or liquid solutions that are in equilibrium with the same partial pressure of zinc at the same temperature have the same thermodynamic activity of zinc. In order to assign a numerical value to the activity azn, it is necessary to choose some standard reference state to which an activity value of unity is arbitrarily assigned. It is customary to choose the liquid or solid state of the pure material. Then azn is defined numerically by

Jan 1, 1947

By C. Ernest Birchenall, Robert F. Mehl

Application of diffusion laws in the customary form to experimental studies in binary metallic solid solutions has shown the diffusion cocfficient to vary with concentration for all systems investigated. The factors determining this variation have not been ascertained. However, the question of whether the diffusion coefficient would become a constant, independent of concentration throughout a single phase if the concentration gradient in the diffusion equation were replaced by an activity gradient has been an attractive subject for speculation. Wagner¹ derived an expression for the force acting on a diffusing particle in terms of the chemical potential. This was also the approach chosen by Jost2 in his monograph. From a somewhat different point of view, Eyring3 and his co-workers have used activity gradients when dealing with diffusion in concentrated nonideal solutions. The only actual test with experimental data is that given by Darken,4 though this is more qualitative than quantitative. Many others opposed the application in any way of thermodynamic concepts to kinetic phenomena, including diffusion processes. A rigorous examination of this principle requires good data expressing the relationship between activity and concentration for solid solutions for which reliable diffusion work has been done. The literature provides such information only for the interstitial solution of carbon in gamma iron (austenite) and for the substitutional copper-zinc solutions (alpha brass). Activity data for carbon in austenite at 800°C and 1000°C are available from the recent work of Smith5 on the equilibrium of hydrogen-methane and carbon monoxide-carbon dioxide mixtures with gamma iron. Combining these with the diffusion studies of Wells and Mehl6 makes possible a quantitative test for this interstitial solid solution. The vapor-pressure determinations for zinc over brass by Hargreaves7 allowed the calculation of activities in this system over a wide range of temperatures and concentrations. Extensive investigation of diffusion in this substitutional solid solution has been carried out by Rhines and Mehl.8 Activities in Brass9 If the vapor phase of zinc is assumed to behave as an ideal gas, any solid or liquid solutions that are in equilibrium with the same partial pressure of zinc at the same temperature have the same thermodynamic activity of zinc. In order to assign a numerical value to the activity azn, it is necessary to choose some standard reference state to which an activity value of unity is arbitrarily assigned. It is customary to choose the liquid or solid state of the pure material. Then azn is defined numerically by

Jan 1, 1947

By John Wulff, Richard P. Seelig

The importance of the pressing operation in the forming of articles by powder metallurgy depends to a great extent on the type of product to be made. While in some few cases the pressing is merely a means of compacting loose powder particles into a solid mass preliminary to a sintering treatment, in most applications it constitutes a molding operation with the purpose of shaping the raw material in powder form into a certain predetermined, accurately defined part not altered in shape by subsequent sintering. In still other cases, no pressing at all but sintering in molds gives useful articles. A limited number of products are made by the simultaneous application of pressure and heat. The factors that control the product of the pressing operation are raw material, lubricant, type, size and shape of die and compact, magnitude and time of application of pressure, as well as temperature of pressing. All such factors need to be controlled; yet, depending on the function of the pressing operation, greater emphasis may be placed on processing factors other than the pressing operation. In any complete discussion of pressing, subsequent sintering cannot be disregarded. For the sake of brevity, the present paper has been limited to the mechanics of pressing and only occasionally refers to sintering. For the same reason, hot-pressing has been omitted as beyond the scope of this paper In any survey of the literature on the pressing of metal powders, the 1938 paper of Balshinl on the theory of the process of pressing deserves mention. The first attempts at an interpretation of pressing phenomena were made by Balshin' and RakovskL2 The experimental results of both are of great interest. Since the experimental results are not described, the work was repeated, using our own interpretation of the possible techniques, and is reported in the following pages. The basic theory of Balshin,' based on a hydro-dynamic analogy, however, leads to hut another way of plotting density and pres- , sure against height of a pressed compact. Since it does not give any new insight into the phenomena, nor account for the influence of side-wall friction, a discussion of the theory has been omitted from this paper. Density Gradients in Compacts The process of pressing metal powders in a die (disregarding for this consideration the powder filling and ejection operations) may be divided into three phases: (I) packing, (2) elastic and plastic deformation, (3) cold-working with or without fragmentation. It must be emphasized that these three steps do not follow each other in sequence; on the contrary, in practice

Jan 1, 1946

By John Wulff, Richard P. Seelig

The importance of the pressing operation in the forming of articles by powder metallurgy depends to a great extent on the type of product to be made. While in some few cases the pressing is merely a means of compacting loose powder particles into a solid mass preliminary to a sintering treatment, in most applications it constitutes a molding operation with the purpose of shaping the raw material in powder form into a certain predetermined, accurately defined part not altered in shape by subsequent sintering. In still other cases, no pressing at all but sintering in molds gives useful articles. A limited number of products are made by the simultaneous application of pressure and heat. The factors that control the product of the pressing operation are raw material, lubricant, type, size and shape of die and compact, magnitude and time of application of pressure, as well as temperature of pressing. All such factors need to be controlled; yet, depending on the function of the pressing operation, greater emphasis may be placed on processing factors other than the pressing operation. In any complete discussion of pressing, subsequent sintering cannot be disregarded. For the sake of brevity, the present paper has been limited to the mechanics of pressing and only occasionally refers to sintering. For the same reason, hot-pressing has been omitted as beyond the scope of this paper In any survey of the literature on the pressing of metal powders, the 1938 paper of Balshinl on the theory of the process of pressing deserves mention. The first attempts at an interpretation of pressing phenomena were made by Balshin' and RakovskL2 The experimental results of both are of great interest. Since the experimental results are not described, the work was repeated, using our own interpretation of the possible techniques, and is reported in the following pages. The basic theory of Balshin,' based on a hydro-dynamic analogy, however, leads to hut another way of plotting density and pres- , sure against height of a pressed compact. Since it does not give any new insight into the phenomena, nor account for the influence of side-wall friction, a discussion of the theory has been omitted from this paper. Density Gradients in Compacts The process of pressing metal powders in a die (disregarding for this consideration the powder filling and ejection operations) may be divided into three phases: (I) packing, (2) elastic and plastic deformation, (3) cold-working with or without fragmentation. It must be emphasized that these three steps do not follow each other in sequence; on the contrary, in practice

Jan 1, 1946

By R. E. Rick-Secker, W. M. Baldwin, R. McC. Baker

The development of residual stress in strip during rolling has not been systematically studied. A few scattered papers1 -3 mention the existence of residual stresses in rolled strip or touch upon some of their consequences,4 but there is no report of experimental studies in the literature describing the effect of roll diameter, strip thickness, reduction per pass, metal or metal temper on the magnitude of residual stresses developed in strip. A knowledge of the effect of these factors is of utmost importance since the present trend in some mills to roll thicker castings brings the rolling operation directly into that set of conditions which, as will be brought out in the present paper, can develop residual stresses of a high order of magnitude. Fundamentally, residual stresses are developed when a metal is inhomogeneously unloaded from a plastic stress state. A quantitative treatment of the problem requires a reasonably complete understanding of both the plastic stress state existing in the metal during deformation and the elastic stress state that would exist in the metal if the metal deformed as a perfectly elastic body. Our knowledge of either state during rolling is so meager that we can make no more than certain limited predictions. The external stress state* operating between the roll and the metal consists of Fig i—External stress state operating on metal being rolled (after siebel and LUEG5) compressive stresses acting almost perpendicularly to the strip surface but varying in intensity both across the width of the strip and along the length of contact between the roll and the strip. An example of such a stress distribution is shown in Fig I. The nature of the internal stress state, produced by this distribution, is not known. Various investigators,5-7 however, have felt that a qualitative similarity must exist between the internal stress state so produced and that produced in a wide strip that is compressed between two flat-nosed punches. In the latter case the external stress state is represented by a uniformly distributed compressive stress acting per-pendicularly over the face of the tool. The

Jan 1, 1949

By M. E. Nicholson

IT has been suggested by a number of investigators, I including Hume-Rothery and Raynor,' that certain intermediate phases in metal systems take on interstitial crystal structures because of an appropriate ratio of radii of the two atoms. If the size factor is an important parameter in determining whether or not an intermediate phase will have an interstitial structure, then it would appear reasonable that the relative sizes of interstitial atoms should have a direct relation to their solubilities in interstitial compounds. It was on such a basis that Petch,' in his discussion of the cementite structure, concluded that nitrogen is the only possible element which could replace carbon. Recent X-ray studies by Jacka indicate that the solid solubility of nitrogen in cementite is negligible. If this is the case, then on the basis of size factor it would not be expected that any element can replace carbon in FeaC, since according to Schwarzkopf and Kieffer4 the covalent atomic radii of nitrogen, carbon, and boron are 0.71, 0.76, and 0.87A, respectively. Nevertheless, Vogel and Tammann in 1922 indicated that a quasi-binary relation exists between the interstitial compounds FeaC and Fe2B and that the solid solubility of boron in FeaC at approximately 1150°C is 1.4 pct B by weight. Because of the apparent disagreement between theory and experiment, the solid solubility of boron in FeaC has been reinvestigated. In addition, the influence of boron on the saturation magnetic moment (per gram) and Curie temperature of cementite was determined. Experimental Procedure A series of Fe-B-C alloys which varied in boron content from 0 to 5.4 wt pct B were prepared by arc melting a 50-g ingot using carbonyl iron, Ache-son AUC graphite, and Cooper grade A boron for starting materials. The composition of these alloys is listed in Table I and shown graphically in Fig. 1. The samples represented by the solid circles were used for X-ray analysis and Curie temperature measurements; by the squares, samples used for quantitative metallography; and by the triangles, samples used for saturation magnetic moment measurements and Curie temperature measurements. After arc melting and subsequent quenching from 1000°C, so that all phases in the resulting microstructure were brittle, one half of the 50-g ingots were pulverized until a powder was obtained which was less than 100 mesh. These powders were then homogenized in evacuated Vycor capsules at 1000°C for 24 hr and water quenched. Although the homogenized powder had sintered somewhat, it was suitable for saturation magnetic moment and Curie temperature determinations, for which the 60 to 100-mesh fraction of the powder was used. For X-ray analysis the homogenized powder was crushed again until a 300-mesh powder was obtained. After crushing, the powder was mixed with Alundum of approximately the same grit size to prevent sinter-

Jan 1, 1958

By G. D. Cremer, J. J. Cordiano

Aluminum powder is a well-known article of commerce and in various forms has been marketed widely for use in paint, for pyrotechnic purposes and for exothermic mixtures. For a number of reasons, however, aluminum powder has not been employed extensively in powder metallurgy, although a great many proposals along this line have been offered. Aluminum and alumiuum-alloy powders oxidize readily upon exposure to air and the thin film of oxide (or hydroxide) thus formed is highly refractory, in that ordinary gaseous reducing agents such as hydrogen or carbon monoxide do not reduce the oxide to metal at temperatures below the melting point of aluminum. It has been thought that the presence of these films of oxide on aluminum particles would interfere markedly in powder metallurgical work and a number of suggestions have been made, either for preventing the formation of the film through the production of the powders in inert or reducing atmospheres, or for the removal of the oxide by means of powerful reducing agents, including "nascent" or atomic hydrogen released from metal hydrides. A second and more serious difficulty in the use of aluminum and aluminum-alloy powders in powder metallurgy resides in the fact that during compression such powders tend to adhere or cold-weld to the die walls, thus interfering with the formation and ejection of a compact. A third difficulty, which has been encountered only with aluminum-alloy powders, is their hardness when compared with elemental aluminum. Certain of these powders are susceptible to age-hardening, which further hinders the formation of strong green compacts. These conditions are particularly true of atomized aluminum-alloy powders. Even in the annealed condition, plastic deformation of alloy powders under pressure is low when compared with unalloyed powders. This fact, coupled with the lower corrosion resistance of the alloy powders. is probably the significant factor in our failure to obtain strong sintered pieces. Methods of Puoducing the Powder Many methods have been proposed for manufacture of aluminum powder, only a few of which are used commercially. Several of the more important processes include: (I) atomization of the molten metal; (2) mechanical pulverization and other types of comminution; and (3) graining, in which molten aluminum is stirred during freezing. For pigment and certain pyrotechnic uses aluminum flake powder is desirable, because of its greater proportion of surface to volume; and large quantities of such material are made by a procedure that involves stamping or ball milling. In making flake intended for use as pigment,

Jan 1, 1943

By G. D. Cremer, J. J. Cordiano

Aluminum powder is a well-known article of commerce and in various forms has been marketed widely for use in paint, for pyrotechnic purposes and for exothermic mixtures. For a number of reasons, however, aluminum powder has not been employed extensively in powder metallurgy, although a great many proposals along this line have been offered. Aluminum and alumiuum-alloy powders oxidize readily upon exposure to air and the thin film of oxide (or hydroxide) thus formed is highly refractory, in that ordinary gaseous reducing agents such as hydrogen or carbon monoxide do not reduce the oxide to metal at temperatures below the melting point of aluminum. It has been thought that the presence of these films of oxide on aluminum particles would interfere markedly in powder metallurgical work and a number of suggestions have been made, either for preventing the formation of the film through the production of the powders in inert or reducing atmospheres, or for the removal of the oxide by means of powerful reducing agents, including "nascent" or atomic hydrogen released from metal hydrides. A second and more serious difficulty in the use of aluminum and aluminum-alloy powders in powder metallurgy resides in the fact that during compression such powders tend to adhere or cold-weld to the die walls, thus interfering with the formation and ejection of a compact. A third difficulty, which has been encountered only with aluminum-alloy powders, is their hardness when compared with elemental aluminum. Certain of these powders are susceptible to age-hardening, which further hinders the formation of strong green compacts. These conditions are particularly true of atomized aluminum-alloy powders. Even in the annealed condition, plastic deformation of alloy powders under pressure is low when compared with unalloyed powders. This fact, coupled with the lower corrosion resistance of the alloy powders. is probably the significant factor in our failure to obtain strong sintered pieces. Methods of Puoducing the Powder Many methods have been proposed for manufacture of aluminum powder, only a few of which are used commercially. Several of the more important processes include: (I) atomization of the molten metal; (2) mechanical pulverization and other types of comminution; and (3) graining, in which molten aluminum is stirred during freezing. For pigment and certain pyrotechnic uses aluminum flake powder is desirable, because of its greater proportion of surface to volume; and large quantities of such material are made by a procedure that involves stamping or ball milling. In making flake intended for use as pigment,

Jan 1, 1943

By C. Ernest Birchenall

AS a basis for further progress in several branches of metallurgy, particularly the study of physical properties of solid solutions and the kinetics of solid-solid reactions, a more complete understanding of the structure of solid solutions and the nature of the metallic bond is- necessary. Any investigation contributing to this understanding should be of primary interest to the metallurgist. Since the metallic bond is characteristic of an assemblage of atoms, it can be studied only where this assemblage exists. One of the most promising approaches to this study of the interactions in metals is through the general problem of binary metallic solutions. Although crystal structures of many solid metallic solutions have been determined and solubility limits for many systems are given in phase diagram studies, progress beyond this is irregular. Electromotive force and vapor pressure determinations have provided thermodynamic activity data for small concentration and temperature ranges in the amalgams and a few other systems.' In two systems, however, copper-zinc and iron-carbon, data are available for the calculation of activities over wide ranges of temperature and concentration. This activity data for the brass system is summarized in the preceding paper.= A detailed study of the solubility and distribution of carbon in iron is given by Smith and Darken.3 Their treatment of interaction energies is applicable only to dilute interstitial solutions. The objectives here will be to derive from activity gats given in the previous paper for the brasses and structural considerations as much new and detailed information as possible about the interactions between copper and zinc in metallic solid solution and the effect of these interactions in determining the short range structure of the equilibrium phases; that is, we shall investigate the energy of interaction between pairs of atoms of the same and different kinds, the deviations of the distribution of nearest neighbors from complete randomness, and show that the conclusions found are consistent with the observed behavior of the brass system. The method of treatment is fairly general for substitutional metallic solid solutions and will be applied to other systems. The determination of vapor pressures of the components of alloy systems is being undertaken in this laboratory to supply more data for this purpose. The information obtained from solid solutions may be extended by careful analogy to liquids as long as the metallic type of bonding is preserved. Since it will be necessary to employ the language of thermodynamics, a brief summary defining the terms used extensively in solution theory is given below. Solution Theory4 An ideal solution is one in which the vapor pressure of the constituent molecules

Jan 1, 1947

By C. Ernest Birchenall

AS a basis for further progress in several branches of metallurgy, particularly the study of physical properties of solid solutions and the kinetics of solid-solid reactions, a more complete understanding of the structure of solid solutions and the nature of the metallic bond is- necessary. Any investigation contributing to this understanding should be of primary interest to the metallurgist. Since the metallic bond is characteristic of an assemblage of atoms, it can be studied only where this assemblage exists. One of the most promising approaches to this study of the interactions in metals is through the general problem of binary metallic solutions. Although crystal structures of many solid metallic solutions have been determined and solubility limits for many systems are given in phase diagram studies, progress beyond this is irregular. Electromotive force and vapor pressure determinations have provided thermodynamic activity data for small concentration and temperature ranges in the amalgams and a few other systems.' In two systems, however, copper-zinc and iron-carbon, data are available for the calculation of activities over wide ranges of temperature and concentration. This activity data for the brass system is summarized in the preceding paper.= A detailed study of the solubility and distribution of carbon in iron is given by Smith and Darken.3 Their treatment of interaction energies is applicable only to dilute interstitial solutions. The objectives here will be to derive from activity gats given in the previous paper for the brasses and structural considerations as much new and detailed information as possible about the interactions between copper and zinc in metallic solid solution and the effect of these interactions in determining the short range structure of the equilibrium phases; that is, we shall investigate the energy of interaction between pairs of atoms of the same and different kinds, the deviations of the distribution of nearest neighbors from complete randomness, and show that the conclusions found are consistent with the observed behavior of the brass system. The method of treatment is fairly general for substitutional metallic solid solutions and will be applied to other systems. The determination of vapor pressures of the components of alloy systems is being undertaken in this laboratory to supply more data for this purpose. The information obtained from solid solutions may be extended by careful analogy to liquids as long as the metallic type of bonding is preserved. Since it will be necessary to employ the language of thermodynamics, a brief summary defining the terms used extensively in solution theory is given below. Solution Theory4 An ideal solution is one in which the vapor pressure of the constituent molecules

Jan 1, 1947

IDO not want to omit anything in which fire or its operation has a part, if I know or can tell it, and have already told you so many other useful, powerful, and ingenious effects (finally of powder and other' fire deadly and very injurious to all living things, indeed, composed for the destruction of those who possess it as well as those who do not) but I do not want the end of my writing to be in tragedy. Therefore I wished to tell you how, from these fires composed of forceful and horrible mate¬rials bringing harm and terror to men, a happy and pleasing effect is also produced and, instead of fleeing from it, the people willingly go to see it. Among others I remember that it was once customary to use these in Florence and in Siena, more than elsewhere. The chief feast day celebrated. by one of these cities is the day of the birth of Saint John, and the other is the Assumption of Our Lady, on the fifteenth of the month of August. In the morning they visit the shrines with great devotion and solemnity, displaying their pomp as much as possible, and later they hunt lions, wild bulls, untamed horses, and many other animals. In the same public squares there was also erected as a decoration an edifice fixed in the ground in the middle of this space, or hanging from a heavy rope crossing between the loftiest points of equal height that could be found in that place. This they have done and still do, and perhaps in time to come will continue to do when opportune. This structure was usually of wood; indeed it is necessary that the latticework of its construction be made only of this. It was covered and en¬larged by tying and fastening hay on top to bring it almost to completion and then papier-mâché was applied. [165v] The figures were first appropriately painted and made of a combination of vessels one above the other, developing various decorative figures in relief which should represent some fable or history and should not appear haphazardly made or without understanding. Now the master proceeded to arrange these figures in a composition according to his skill or art of design, and likewise planned

Jan 1, 1942

By K. Tangri, M. Chaturvedi

The addition of 0.5 wt pct Cu to Zr-2.5 Cb alloy increases the as -quenched hardness of the hexagonal martensitic a' phase, produced by water-quenching bccß-Zr phase, by about 35 pct. This strengthening has been attributed to the solid -solution hardening of the matrix. On aging ternary martensite, a' phase reverts to equilibrium a and Zr2Cu and ß-Cb precipitate out, mainly at the twin and grain boundaries, causing a secondary hardening of the matrix. COLD-worked Zircaloy-2 pressure tubes have been in use in power reactors for a considerable period of time. The search for a better material led to the development of Zr-2.5 wt pct Cb alloy which in the quench-aged condition develops 50 pct more strength than that of cold-worked Zircaloy-2, however, its corrosion resistance in water and steam in the temperature range of 316" to 400°C, in absence of neutron flux, is inferior to that of zircaloy-2.' Work carried out by Ells et al.1 and Dalgaard2 has shown that the corrosion properties of Zr-2.5 wt pct Cb alloy can be considerably improved by the ternary addition of 0.5 wt pct Cu. This paper is concerned with the effect of 0.5 wt pct Cu on the formation of martensitic a and its aging characteristics in a Zr-2.5 wt pct Cb alloy. MATERIALS AND EXPERIMENTAL TECHNIQUES Zr-2.5 Cb-0.5 Cu (referred to as the ternary alloy) and Zr-2.5 Cb (referred to as the binary alloy) alloys, supplied by the Chalk River Nuclear Laboratories of the AECL were used. The detailed chemical analysis is given in Table I. Cold rolling and swagging with frequent intermediate anneal of 1000°C were used for the initial fabrication of the alloys. All the heat treatments were carried out after the specimens were wrapped in zirconium foils and encapsulated in silica tubes under a vacuum of 5 x 10-6 mm of Hg. For optical metallography and hardness measurements specimens were mechanically and then chemically polished in a 45 pct HNOj, 45 pct HzO, and 10 pct HF solution. Hardness was measured on a Vickers hardness tester using a 10-kg load. For each specimen at least fifteen indentations were made in order to obtain a representative value. The phase identification and structural analysis were carried out using X-rays and electron diffraction techniques. Wires of 1.5 mm diam reduced to 0.12 mm diam by chemical etching were used for making Debye-Scherrer powder patterns using Cu Ka radiation in a 114.6 mm diam camera. Carbon extraction replicas were prepared by etching the specimens, after depositing a layer of carbon on the metallographic specimen, in one part HF and thirty parts ethyl alcohol. Thin films were prepared by electropolishing heat-treated 3/4 by 1/2 by 0.005 in. thick strips using a modified Bollman-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol bath was kept at -50°C and polishing was done at 5 to 10 V. The thinned specimens were washed in ethyl alcohol at -30º to -40°C and dried between filter papers. Replicas and thin films were examined in a Phillips 300 G electron microscope. For resistivity measurements thin strip specimens 0.02 by 0.3 by 10.0 cm long were used. The potential leads were spot welded to the specimens in order to maintain a fixed length for the initial and the final resistivity measurements. The resistivity was measured by a Kelvin bridge in a temperature controlled room. The temperature was maintained at 72º ±1°F and the accuracy of the resistivity measurements was 0.03 µa-cm. RESULTS As-Quenched Structures. In order to produce a homogeneous matrix to study the precipitation reaction the solution-treatments of both the alloys were carried out in the -field region. From the Zr-Cb phase diagram due to Lundin and cox3 ß/a + ß phase boundary for Zr-2.5 wt pct Cb alloy is 820°C. Ells et al.1 have reported this boundary for Zr-2.5 Cb alloy containing 1100 ppm 0 to be at 920°C. Also, the addition of 0.5 wt pct Cu reduces this temperature by 50°C. Therefore, the solution-treatments were carried out at 1000°C to ensure that the alloys were in ß-phase region. The soaking time was 1 hr and the specimens were water-quenched. The as-quenched hardness of the binary alloy was 245 Vpn whereas, that of the ternary alloy was 330 Vpn. The X-ray diffraction studies indicated that the as-quenched structure of both the alloys consists of martensitic hexagonal phase a', with a c/a ratio of 1.591, and some retained ß-Zr. The presence of a' phase was further confirmed by thin film electron microscopy. Electron micrographs of typical ß-quenched structures of the ternary and the binary alloys are shown in Figs. 1 and 2, respectively. Fig. 3 shows the diffraction pattern from an area similar to that shown in Fig. 1. Although, the as-quenched hardness of the ternary alloy is about 35 pct greater than that of the binary alloy, the structure of both the alloys seems to be the same. The matrix of both alloys is heavily twinned and shows very few dislocations. Furthermore, there is no evidence of any precipitation taking place in either of the two specimens during quenching from the solution-treatment temperature. Aging Behavior of Martensitic a'. The aging kinetics of the ternary alloy were followed by resistivity and hardness measurements. The as-quenched values

Jan 1, 1970

By W. J. Locher

One important means of reducing steel consumption in grinding mills is not mentioned in S.G. Malghan's paper. It is simply to reduce the oxygen content in the mill and pulp. While oxygen dissolved in pulp is not readily reduced without additives, reduced aeration and reduced ventilation of grinding mills are often feasible. It is not necessary to completely eliminate the oxygen to achieve substantial benefits, since corrosion rates vary with oxygen content in a manner analogous to combustion rates as oxygen content increases. It is difficult to generate indisputable quantitative data on the relationship of oxygen reduction in a milling system to changes in consumption of grinding steel. Perhaps this accounts for the absence of technical papers on the effects of reducing oxygen content in grinding mills by reducing aeration of pulp and ventilation of grinding mills. One qualitative observation is that mills operating at a lower percent of critical speed will tend to have lower media consumption for similar fineness of grind without necessarily having an equivalent reduction in power consumption per ton of product. Reduced mill speed would also tend to reduce both ventilation of the mill and entrainment of air in the pulp. Larger grinding mills tend to be more media efficient because they operate at a greater ratio of charge volume to pulp surface area, which will tend to reduce the aeration per unit of pulp volume. The mechanical steps required to reduce aeration of grinding mill pulps depend on the configuration of the grinding circuit and related process. Reply by S.G. Malghan Mr. Locher stated that my paper did not mention the reduction of oxygen content of the pulp as a potential method of decreasing the steel consumption. I would like to bring to the attention of Mr. Locher that an indirect reference was made to the fact that pulp oxidation potential controls the corrosion wear of steel. Though dissolved oxygen is one of the factors that controls oxidation potential, additional discussion was deliberately avoided for the reasons mentioned by Mr. Locher. The most important reason being the difficulty of decreasing and controlling the oxygen concentration of the pulp in a commercial mill.

Jan 1, 1984

By A. J. Sedriks

ALTHOUGH the occurrence of a readily visible tarnish on the stress-corrosion fracture surfaces of titanium alloys tested in oxygenated nitrogen tetrox-ide solutions has been reported,' the possibility that the formation of this tarnish is a requirement for cracking to proceed does not appear to have been considered. It has been proposed that failure occurs in this system as a result of some electrochemical-mechanical process.2 However, recent studies3-6 of the stress-corrosion cracking of a brass in certain ammoniacal solutions, which also cause visible tarnishing, have led to the view that failure occurs by the repeated formation and fracture of a brittle surface film of cuprous oxide. The purpose of this note is to point out the similarities which exist between the tarnish which forms on a brass and that found on titanium alloys failed in nitrogen tetroxide, and to suggest, by analogy, that a similar tarnish-rupture mechanism may be operative in the latter case. The titanium alloy selected for the present study contained 5.2 pct Al, 2.5 pct Sn, 0.025 pct C, 0.16 pct Fe, 0.014 pct N, 0.012 pct H, 0.07 pct 0, and 0.002 pct Mn by weight. Vacuum-annealed alloy strips measuring 4 by 0.25 by 0.04 in. were used to prepare self-stressed test specimens. Each specimen was made from two strips by bending the end portions and spot welding them together to produce curved sections of uniform bending stress. (The exact details of preparing such specimens have been described by Lisagor el al.7) The curvature in the strips determines the ma-magnitude of the outer fiber stress, which in the present specimens was 100,000 psi (yield stress = 120,000 psi). The specimen surface in the spot-welded end regions was shot-peened to induce compressive stresses and thereby eliminate the possibility of stress-corrosion cracking in these regions. The test environment was nitrogen tetroxide, initially containing 99.5 pct N2O4, 0.1 pct H2O, and 0.08 pct C1, through which oxygen had been bubbled for 30 min. During testing the tetroxide was subject to an oxygen pressure of 250 psi. The duration of the test was 96 hr, and during this time stress-corrosion cracks propagated through some 50 pct of the thickness of the specimens. After testing, the tarnish was liberated from the fracture surface by immersion in a methanol-5 pct bromine solution, and then examined by electron dif- fraction and transmission electron microscopy using a JEM-7 electron microscope. The possibility that the bromine solution had affected the structure of the tarnish was investigated by an "in silu" examination of the stress-corrosion fracture surface using a reflection electron diffraction attachment designed for handling bulk specimens. Typical areas of the tarnish liberated by the me-thanol-bromine solution are shown in Fig. 1. Fig. l(a) reveals that the tarnish consists of individual platelets, and from Fig. l(b) it can be seen that the platelets are joined together to form a semicontinuous film, the boundaries of the titanium substrate grains being retained in the tarnish. Electron-diffraction patterns taken from these areas exhibited diffraction spots typical of a monocrystal. The diffraction pattern shown in Fig. 2 was taken from the central area in Fig. l(n). Monocrystal patterns obtained from areas containing numerous individual platelets can only be explained by assuming that their growth is governed by the titanium substrate (i.e., their growth is epitaxial). In other areas, the tarnish had folded over or became mechanically damaged, and these areas gave diffraction patterns consisting of "spotted" rings, enabling interplanar spacings of the tarnish lattice to be determined. These spacings, measured using a

Jan 1, 1968

By R. H. Tien

This paper presents an analytical solution of the problem of freezing of a semi-infinite slab with constant surface temperature; in this analysis accoz~nt is taken of the heat evolution during the solid-state transformations in the frozen skin. The regions where transformations occur are treated as plane isothermal heat sources; i.e., the rates of transformations are assumed to be infinite. In the model considered, three plane heat sources are assumed to be present in the frozen solid. The solution yields the temperature distribution in the solid, the rates of movement of the heat sources , and the rate of movement of the solidification front. It is shown that, although the sum of the heats of transformations in the solid is about 20 or 43 pct of the heat of fusion of iron or manganese, respectively, the effect of the heat evolution in solid-state transformations on the rate of solidification is insignzficant. IN several recent mathematical studies of solidification of liquid metals'-3 no account was taken of the heat evolution accompanying the solid-state transformations. In the case of, for example, steel, there are several phase transformations 6 -y, y -a, and the magnetic transformation. Whether or not a significant error is introduced into the mathematical simulation of solidification processes by the omission of the heat effects associated with these phase transformations is not known. The purpose of the present study is to investigate the effect of these phase transformations on the temperature distribution and the rate of solidification based solely on the heat transfer point of view. That is, for the sake of simplicity, freezing of a pure metal with solid-state transformations will be considered. The mathematical model which represents the present problem includes four moving heat sources representing freezing and phase transformations, 6— y — a — ferromagnetic. The thermal properties such as conductivity and specific heat, and so forth, are different from one phase to the other. But they are considered to be uniform over any particular phase. The usual Fourier heat flow law is employed. The concept of conservation of energy is applied at every heat source in order to estimate their rate of movement. The temperature distribution in the solid and the positions where transformations occur are expressed analytically in terms of error function solutions and four coefficients which are determined from four nonlinear algebraic equations. These coefficients can be calculated either by trial and error or by any numerical method. MATHEMATICAL ANALYSIS Fig. 1 is a schematic diagram showing the zones of different phases 1, 2, 3, and 4 in the solidified portion of a semi-infinite pure liquid metal. The heat is subtracted from the system at x = 0 which is maintained

Jan 1, 1968

By John L. Ham, Robert M. Parke

The melting point of molybdenum is 2625° + 50°C. Heretofore the metal has been considered too refractory to be melted in commercial quantities; hence, it has been formed into rod, wire, and sheet by the methods of powder metallurgy, wherein the temperatures do not exceed 90 per cent of the melting temperature. The manipulation of liquid metals at temperatures of about 2000°C. and above presents some difficult problems in heat transfer, in purification, and in container materials. For metals melting at such high temperatures, it is generally regarded as expedient to circumvent the more common melting and casting process by resorting to powder metallurgy. This circumvention is not accomplished without some penalties, the more notable of which are the relatively high cost of the powder metallurgy method and particularly the size limitations of its product. In an effort to overcome the inherent disadvantages of the powder metallurgy method, especially in respect to the limit of size, a project was begun in 1943 at the Climax Molybdenum Co. laboratory to develop a method of melting and casting molybdenum capable of producing larger bodies of molybdenum than were then obtainable. Later, from Feb. I, 1944, to Aug. 31, 1945, the project was under the supervision of the National Research Committee, Division One. In peacetime, too, there are needs for molybdenurn, which make consideration of another process desirable; for example, large parts for heat engines generally, the omnipresent demand for metal of higher quality and lower cost, and the development of molybdenum-base alloys. This paper describes methods for melting and casting molybdenum, with details of the special process needed to make molybdenum castings ductile, and reports some of the properties of what may be regarded as a new, or at least modified, form of the metal. First Process for Melting and Casting Molybdenum The maintenance of a temperature greater than 2625°C. in a 'pace large enough to melt several pounds of molybdenum may involve the loss of heat at a Very considerable rate. For melting molybdenum, therefore, the source of heat must be of high potential and of fairly high power. Evidently, also, the source of heat must not impair the purity of the product. The electric arc operating in vacuum appeared to meet these requirements and was accordingly the first choice as a source of energy. A number of arc devices for melting molybdenum were examined, and several were tried. In the first experiment, a 20-kw., 60-cycle-per-second alternating-current arc was operated between two horizontal and opposed electrodes of molybdenum- In this arrangement the falling droplets of molybdenum solidified below the arc in an irregular mass. This experiment did not produce a useful casting, but the data obtained indicated that

Jan 1, 1947

By John L. Ham, Robert M. Parke

The melting point of molybdenum is 2625° + 50°C. Heretofore the metal has been considered too refractory to be melted in commercial quantities; hence, it has been formed into rod, wire, and sheet by the methods of powder metallurgy, wherein the temperatures do not exceed 90 per cent of the melting temperature. The manipulation of liquid metals at temperatures of about 2000°C. and above presents some difficult problems in heat transfer, in purification, and in container materials. For metals melting at such high temperatures, it is generally regarded as expedient to circumvent the more common melting and casting process by resorting to powder metallurgy. This circumvention is not accomplished without some penalties, the more notable of which are the relatively high cost of the powder metallurgy method and particularly the size limitations of its product. In an effort to overcome the inherent disadvantages of the powder metallurgy method, especially in respect to the limit of size, a project was begun in 1943 at the Climax Molybdenum Co. laboratory to develop a method of melting and casting molybdenum capable of producing larger bodies of molybdenum than were then obtainable. Later, from Feb. I, 1944, to Aug. 31, 1945, the project was under the supervision of the National Research Committee, Division One. In peacetime, too, there are needs for molybdenurn, which make consideration of another process desirable; for example, large parts for heat engines generally, the omnipresent demand for metal of higher quality and lower cost, and the development of molybdenum-base alloys. This paper describes methods for melting and casting molybdenum, with details of the special process needed to make molybdenum castings ductile, and reports some of the properties of what may be regarded as a new, or at least modified, form of the metal. First Process for Melting and Casting Molybdenum The maintenance of a temperature greater than 2625°C. in a 'pace large enough to melt several pounds of molybdenum may involve the loss of heat at a Very considerable rate. For melting molybdenum, therefore, the source of heat must be of high potential and of fairly high power. Evidently, also, the source of heat must not impair the purity of the product. The electric arc operating in vacuum appeared to meet these requirements and was accordingly the first choice as a source of energy. A number of arc devices for melting molybdenum were examined, and several were tried. In the first experiment, a 20-kw., 60-cycle-per-second alternating-current arc was operated between two horizontal and opposed electrodes of molybdenum- In this arrangement the falling droplets of molybdenum solidified below the arc in an irregular mass. This experiment did not produce a useful casting, but the data obtained indicated that

Jan 1, 1947