Search Documents

By Eric R. Jette, A. G. H. Anderson

During the period from 1910 to 1920, there was a lively interest in the subject of grain growth and many papers were published, followed by interesting discussions. Questions dealing with the fundamentals of grain growth—why and how it occurs—were treated in papers by Howe,' .Jeffries,² Sauveur,² Chappell,4 Beilby,5 Mathewson and Phillips,6 Mc Adam,7 Carpenter and Elam,8 Stead and Carpenter,9 Ruder,l0 as well as others. Since that decade, much indeed has been done regarding the practical control of grain size, and it is now possible to order material having a size sperified in the American Society for Testing Materials classification. Also, one hears the term "spliced boundaries" used in connection with certain tungsten lamp filamentsl1 wherein it is important to prevent "sagging" between consecutive turns of the helix.12 This practical use of a certain knowledge as to the things that influence or even control grain growth should not be taken as an indication that the subject has been thoroughly explored and that nothing more can be learned from fundamental studies. It is with this thought in mind that a report of some observations and experiments on abnormal grain growth is believed justified. Abnormal grain growth is met with in heat-treating high-carbon steels in the austenitic range,'? in the hardening treatments of high-speed steels,14'15 and in the low-temperature annealing of strained low-carbon steels as in sheet and wire material. Entirely aside from these instances of exaggerated grain growth, which some might consider as freakish, it is thought that abnormal growth is simply a case where the growth forces and the forces tending to prevent growth are almost, but not quite, evenly balanced. Information relative to grain-growth characteristics in general may be obtained from a study of abnormal growth conditions, which may be looked upon as vantage points. Reflection will show at once that if growth forces are great in comparison to inertia or resistance to growth, many crystals start to grow and, for that reason, none becomes very large, whereas if inertia is

Jan 1, 1938

By J. Alfred Berger, O. M. Katz

The Zv-Hf-H ternary system was studied between 500° and 900°C at pressures less than 1 atm of hydrogen gas between 1 and 60 at. pct H. A new and unique microgravimentric apparatus was used. Cizanges of slope on pressure-hydrogen composition isothernis delineated phase boundaries. These boundaries separatecl the three regions, a, 0, and y—so designated to correspond to the regions of the Zr-H binary system—from the multiphased areas between them. A eutectoidal decomposition was found with the ß region phase or phases decornposing into a lamellar product on quenching to rool ter,zperatuve. Reproducible decomposition-pressure hysteresis occilrved lnainly at lower hydrogen cornpositions and at lower temperatures across multiplzase vegions between a and 0 and a and y. Tire effects of hqfniur7z on the hydriding charactevistics of zirconiurrz weYe as follows: 1) stabilization of the a and y vegions while destabilizing the 0 region; 2) a/?preciable elevation of the decomposition pressrkres in the multiphase region between the a and /3 field; 3) ~nouenzent of the eutectoid reaction to high te~nperatures; 4) reduction in the total qiiantity of hydrogen absorbed under one atmospheve of Hz p7-essure; and 5) introduction of a split deconzposilion at the eiitectoiclal poinl in pa?? of the ternavy. Assuru~ptions based on an ir-2terstitial vandonl-solulion rtioclel 0.f hydrogen in metals slzowed that the bindit~g energy between solute sites prednnzinatecl at low /i?!dvogen concentrations. However, at high hydrogen contents the entropy was the predorninatlt factor in determining the stability of the Zr-Hf-H al1o.s. This was interpreted to mean a scarcity of filrtlzer itltevslilinl solute sites caused by hydrogen-hydvogen intet-actions in the metal lattice. INTEREST in the reaction of hydrogen with metals has increased in the past few years for the following reasons: 1) the formation or use of high hydrogen potential environments in nuclear reactors; 2) the reaction of hydrogen with alloys in nuclear reactors with the accompanying deleterious effects on the mechanical and corrosion properties; 3) the theoretical implications of thermodynamic data on the theory and rules of alloy formation in the metal-hydrogen systems; 4) the use of hydrogen-containing fuels in rocket engines; 5) the need for a process of making fine metal powders of high-melting reactive metals; and 6) the beneficial impregnation of superconducting alloys with hydrogen. In nuclear pressurized-water reactors, the problem exists of limiting the hydrogen pickup of zirconium alloys which are utilized as fuel cladding, heat shields, and support members. In general, zirconium alloys have good mechanical and corrosion-resistant properties in high-temperature water. However, hydrogen is absorbed from the corrosion reaction between metal and water, greatly accelerating the formation of the corrosion product ZrOz as well as mechanically embrittling the underlying metal. In addition, recent observations1 at zirconium to hafnium welds showed that secondary elements in zirconium can have an appreciable, and somewhat unexpected, effect on hydrogen absorption. This paper lists the thermochemical data in the range 500" to 900°C for the equilibrium reaction of four high-purity Zr-Hf alloys with hydrogen. Phase boundaries and thermodynamic functions are determined while the structural data will be presented in a future paper. In general, the Zr-Hf-H system approximates the well-known, eutectoidal, Zr-H diagram2,3 with modifications introduced through the behavior of hafnium.4,5 The Hf-H system,' published while this work was in progress, provided a consistent trend with the Zr-Hf-H data. PREPARATION OF Zr-Hf ALLOYS Table I presents a complete flow chart of the preparation procedure. The zirconium and hafnium crystal bars were completely immersed in high-purity kerosene and slowly cut into thin wafers. Wafers were then cold-sheared into approximately 1-g pieces, thoroughly cleaned, weighed, and inserted into the furnace. The alloys, B-2, B-4, B-6, and B-8, were then nonconsumable arc-melted under 500 mm of purified argon. Additional purification of the argon was accomplished by melting a large titanium button each time an alloy was re-melted or a different alloy melted. Each alloy button, which weighed 25 g, was remelted four times in an approach to complete homogeneity. Material losses were less than 0.02 wt pct. Alloy buttons were alternately cold-rolled and vacuum-annealed into 10- and 20-mil sheets. Table I1 gives the composition of the four alloys used. Very little elemental segregation existed be-

Jan 1, 1965

By W. R. Shimer, F. O. Kichline

The investigation herein described was carried out for the purpose of studying, both by chemical and metallographical means, the extent of over-oxidation of steel that can be accomplished by excessive overblowing in a Bessemer converter. 111 addition there are appended some results of basic open-hearth and crucible steels. Remarks on Chemical Methods. The test block used was hemispherical in shape, 4.5 in. in diameter and 2 in. thick. Drillings for analysis were obtained by drilling through the entire solid part of the block, carefully avoiding all cracks and blow-holes and mixing these drillings thoroughly. Drillinga were carefully examined for scale and the finest were gone over with a microscope to be sure all scale from whatever source had been removed. Oxygen was determined by the well-known method of Ledebur, modified according to Cushman,1 with, however, the removal of the preheating tube, which was found unnecessary. This method has been found to give concordant results. Blanks run from 0.0003 to 0.0006 g. Duplicate determinations varied not more than 0.001 to 0.003 per cent. It seems to be generally supposed that the hydrogen method gives only that oxygen which is combined with iron. Whether this be true or not, the content of oxygen existing as iron oxide no doubt predominates sufficiently, so that the result obtained is sufficient to indicate whether the steel is over-oxidized or not. We are well aware that the oxides of manganese and chromium, when existing alone, are not reduced by hydrogen at a red heat. To learn of their action in steel the following work was carried out: A sample of ferro-manganese and of ferro-chrome were each exposed

Jan 1, 1914

By T. L. Fritzlen

EXTRUSION of aluminum alloys in this country is performed mainly by direct extrusion, therefore this paper is confined only to factors affecting the rate of extrusion by this method. Many factors affect the rate of extrusion of aluminum alloys. It is the purpose of this paper to present a discussion on the factors that are subject to technical control. Most of them are closely related, so that a clear-cut distinction between them is rather difficult. However, an attempt will be made to evaluate the effect of each factor separately, as far as possible. The most important factors only will be considered, and these will be discussed in detail under three general headings: (I) characteristics of the equipment, (2) size and structure of the ingot, and (3) temperature of ingot, cylinder, and other parts. It is a well established fact' that the pressure and the speed of extrusion are related and are exponential functions of one another. Therefore, whenever a certain condition tends, for instance, to increase the pressure required for extrusion, it will reduce the extrusion speed if the pressure is maintained constant. For this reason, in this paper the effect of various conditions is expressed either in regard to their effect on the pressure required or the extrusion speed. Another established fact is the relationship between the temperature of the material being extruded and the pressure required to extrude. According to Schishokin,2 this relationship is of the exponential type, and as the temperature is increased the pressure required is decreased. It would seem from the preceding that the greatest rates of extrusion can be obtained by the use of higher temperatures. However, as is well known, other factors intervene to limit the rate of extrusion of aluminum alloys when increasing the temperature. These factors will be discussed in detail later. CHARACTERISTICS OF THE EQUIPMENT The capacity of the press determines the maximum unitary pressure that can be applied for any given cylinder diameter. Pressures used for extrusion of aluminum alloys range from approximately 50,000 lb. per sq. in. to 150,000 lb. per sq. in., dependent upon the alloy and the size of the shape extruded. Higher pressures could be used to advantage, but existing steels limit pressures employed in the production of extrusions. Numerous mechanical characteristics of the extrusion press affect extrusion rate. The most important characteristics of the press itself are: (I) sensitivity of the main pressure valve and accuracy of pressure gauges, (2) accuracy of equipment to heat the cylinder uniformly and to control cylinder temperature, and (3) accuracy of means of measuring the speed of extrusion. The variables subject to control during the extrusion operation are pressure, temperature, and speed, and it is evident that the preceding characteristics

Jan 1, 1945

By J. M. Woodall

GuAs crystals which have been grown in quartz boats by the horizontal Bridgman method in the pvesence of Ga20 vapov have beetz found to have carrier and donor distributions which do not correspound to those expected from simple dopant seg-vegation during directional freezing; Instead, the carrier distribution is determined by the heat-trentnzent history of the crystal, while the donor distribution, zohiclz is principally due to silicon, is fixed by the pozuth rate, the geo)tlet.ry of the crystal growth vessel, and tlze initial Ga20 pressure. WHEN semiconducting materials are made into doped crystals by the normal freezing method,&apos; they usually exhibit doping variations along the growth axis. If a) there is no dopant diffusion in the solid, b) the dopant is distributed uniformly in the melt, and c) the distribution coefficient, k, does not vary with composition, then the doping variation along the growth axis is represented by the equation: where C is the doping level in the solid at a point where a fraction g of the original liquid has frozen, and Co is the mean concentration. When the dopant is either a singly ionized shallow donor or shallow acceptor, C also represents the carrier concentration. Even though this equation accurately describes the dcping profiles of a large number of normal freeze systems, there are several special systems for which Eq. [I] does not apply. One such system is the horizontal Bridgman method for preparing oxygen-grown GaAs crystals using quartz vessels. Several workers2"* have shown that GaAs crystals grown by the horizontal Bridglnan method using quartz vessels are generally contaminated with silicon in concentrations in excess of 5 < 10lG atoms per cu cm. This contamination is ascribed to a reaction: occurring at the walls of the crystal-growth vessel which liberates silicon into the melt and Ga2O vapor. It has been shown4 that this reaction, and, hence, the silicon contamination, can be suppressed by the addition of oxygen to the crystal-growth apparatus. It is the purpose of this paper to describe a special apparatus capable of yielding single crystals of GaAs grown in the presence of oxygen and to describe both the kinetics of silicon suppression in this system and the relationship between the carrier concentration profile and the silicon concentration profile. EXPERIMENTAL-CRYSTAL GROWTH A schematic diagram of the apparatus used in the oxygen addition experiments is shown in Fig. 1. The most important features of this apparatus are: a) the use of a sand blasted quartz boat, h) a quartz rod of length 1, with a hole of cross section A, that is placed near the boat to limit the free space volume V, over the melt during the growth, and c) the temperature gradient at and near the solid-liquid interface. Sand blasting the boat is necessary to prevent wetting of the melt. The quartz rod retards the diffusion of the GazO vapor away from the melt to the colder portions of the ampoule. A crystal is prepared by first loading a 5.5-in. boat, which has been cleaned in aqua regia, with 40 g of 99.9999 pct Ga along with 1 to 8 mg of Ga203 powder. GazO3 is a convenient source of oxygen since it reacts with gallium at the melt temperature to form Ga20 vapor, the species which apparently controls the suppression of silicon contamination. The loaded boat, the quartz rod, and the 99.9999 pct As are placed in the ampoule as shown in Fig. 1. Generally, GaAs seeds were not used since most of the unseeded growths resulted in monocrys-tals. The ampoule is evacuated to 10-5 Torr and sealed off. The GaAs melt is synthesized by placing the ampoule into a two-zone horizontal Bridgman furnace. The two zones are separated by several sandwiched layers of -in. Fibre-Fax board drilled with holes slightly larger than the OD of the ampoule. This causes a very large temperature gradient between the two zones, which is necessary for single-crystal growth. The two-zone furnace is mounted on a stand fitted with a roller bearing which allows uniform motion of the furnace in a horizontal direction. A uniform velocity of the stand is achieved by the use of a windlass device which winds up a wire attached to the stand. Movement of the solid-liquid interface is accomplished by fixing the position of the ampoule and moving the stand. Growth rates investigated in this experiment were between 0.4 and 1.2 in. per hr. The composition of the melt is fixed by maintaining an arsenic reservoir at 618°C. The stand is moved away

Jan 1, 1968

By Albert J. Phillips

LEAD is an inexpensive metal produced and consumed in large quantities with certain chemical and physical properties that permit handling the molten metal in bulk. These properties are its low melting point, resistance to oxidation, and its negligible solvent action on iron and steel. Thus, lead is almost always melted in large iron or steel kettles, heated from beneath, and is seldom melted in either directly or indirectly fired furnaces or electric induction or arc furnaces. Even with a modest operation a kettle holding 2 or 3 tons is employed, and the size will range up to 215 tons in the large lead refineries. The methods of handling tin and its alloys are essentially the same as for lead because of the similarity in melting point and resistance to oxidation, although because of the cost of the metal the scale is more modest. In melting lead and its alloys preliminary to a casting operation it is seldom that the melting unit is completely drained of its contents. The usual procedure is to continuously charge solid metal and continuously remove small amounts of liquid metal, so that the kettle is maintained essentially full at all times. For type metals, linotype machines usually employ automatic means of adding the solid alloy to the melting pot in proportion to the removal of the liquid metal in casting slugs of type. In the casting of grids for storage batteries, frequently automatic means of charging solid metal in direct proportion to the consumption of liquid metal in the casting is used. In the manufacture of bearing alloys, new alloy is charged in direct proportion to the consumption of metal in the bearings being cast, so that at no time does the level of the metal in the melting pot lower to any great extent. Even in the manufacture of bulky commodities such as pipe, cable. sheathing, and slabs for rolled sheet, it is the practice to employ a melting kettle of several times the capacity of the casting being made. A substantial level of liquid metal in the melting kettle is thus maintained even after the casting is poured and the original metal level is quickly restored by promptly charging new metal, even though there is no immediate need to proceed with additional casting. This procedure is permissible because the amount of fuel consumed in keeping the lead molten and immediately available for use is negligibly small and the cost is more than offset by the convenience. However, this convenient procedure necessitates study and control of certain properties of lead and tin alloys that are not of serious consequence in most other alloys and for that reason are not well known. These properties are "drossing rates" and "melting segregation." DROSSING RATE Users of lead long ago observed that the rate of dross formation on the surface of the lead in their kettles or pots varied greatly from time to time. This was particularly common in places where both antimonial lead alloys and commercially pure lead were being handled.

Jan 1, 1946

By Harry R. Hall

THE proposition to make pig-iron from magnetic concentrates and cobbed ore with charcoal-fuel weighing from 12 to 20 lb. per bushel is, on the face of it, not inviting; but the successful work that has been done under these adverse conditions warrants a permanent record in our Transactions, especially as it furnishes a comparison with, coke blast-furnace practice under similar conditions. The record of the furnace-practice is briefly summarized as follows:-The charcoal was made of hemlock, pine, spruce, balsam, maple, beech and birch woods. All the concentrates containing from 62 to 65 per cent. of iron would pass in. screen, of which about 34 per cent.. would pass an 80-mesh sieve. The cobbed ore, varying in composition from 32 to 47 per cent. metallic iron and from 23 to 32 per cent. of. silica, would pass through a I-in. ring, of which about 25 per cent. would pass a 1/8 in. screen. The furnace was blown in October 20,1903, with charcoal which was supposed to weigh when dry 18 lb. to the bushel. Some charcoal similar to this was stored in the stock-house for about a year and was then found to weigh 12 lb. per bushel. During the first 10 days of the blast, the furnace produced 362 tons of pig-iron, using 200 bushels of charcoal per ton of iron. From November 1st to 12th, the volume of the blast was increased gradually from 2,808 to 5,460 cu. ft. per minute, with practically no increase of output. The silicon-content of the resultant pig-iron averaged 0.75 per cent., with sulphur about 0.035 per cent. A supply of freshly-burned dry charcoal was then used which increased the silicon-content and lowered

Sep 1, 1905

The Institute should feel honored that one of its distinguished members, Auguste J. Rossi, has been awarded the Perkin Medal of the Society of Chemical Industry. The medal was presented at a meeting of the New York Section of the Society on Jan. 18, at the Chemists&apos; Club. Mr. F. A. J. Fitzgerald, Past President of the American Electrochemical Society; described the work of Mr. Rossi, while the presentation address was delivered by William H. Nichols, Past President of the Society of Chemical Industry.

Jan 2, 1918

By L. D. Lamont, W. C. Muehlhof

Throughout the anthracite industry&apos;s history, the problem of handling and disposing of refuse material has been one of major importance. In the early days of the industry&apos;s activities, only the large sizes of anthracite were satisfactory for market requirements and the smaller sizes were deposited as refuse. As time went by, more efficient coal-burning equipment was developed and a market demand was created for the finer sizes of coal. This demand made it possible to recover marketable coal from banks deposited as refuse in the earlier days of mining. The same condition holds good today and we can see a definite trend toward the objective of utilizing the No. 4 and No. 5 buckwheat sizes. Under present conditions, we know that a limited market exists for No. 4 and 5 buckwheat; we know that the finer sizes of anthracite can be burned as pulverized fuel; we know that these fine sizes can be processed and sold as briquettes, and we also know that the heat content of these fine sizes of coal is comparable with that of the sizes now merchandized. Therefore, it should be the objective of the industry to conserve these fine sizes of coal in such manner that if and when processing facilities are developed, and market demands exist, the material can be salvaged and made available to meet these demands. From the standpoint of economical future reclamation, the ideal method of disposing of refuse containing the finer sizes of coal is to deposit it in banks separate and apart from other breaker refuse. By this procedure, the minimum amount of material will be processed in future salvage operations to recover the No. 4 and No. 5 buckwheat sizes. Many collieries in the anthracite region are operated under leases that have a specific term. These lease agreements provide that, upon expiration of the lease by term, exhaustion or otherwise, the ownership of the refuse material reverts to the lessor. Therefore, a question was raised in the minds of these operators: Why should I assume the expense of such procedure when I have no assurance that I will benefit by the future processing and sale of these fine sizes? Under present conditions, the answer to this question is not one of choice but of necessity, as restrictions imposed by legislation make it mandatory on operators of preparation plants to take necessary steps to properly treat silt-laden waters before permitting them to enter the streams. By reason of these restrictions, the disposal of fine sizes becomes a secondary rather than a primary operating consideration, and the procedure for disposing of this refuse becomes a matter of choice governed principally by topographical and geographical conditions. Methods of Handling Refuse There are three principal methods of handling and disposing of fine-size refuse; namely:

Jan 1, 1947

By W. C. Muehlhof, L. D. Lamont

Throughout the anthracite industry&apos;s history, the problem of handling and disposing of refuse material has been one of major importance. In the early days of the industry&apos;s activities, only the large sizes of anthracite were satisfactory for market requirements and the smaller sizes were deposited as refuse. As time went by, more efficient coal-burning equipment was developed and a market demand was created for the finer sizes of coal. This demand made it possible to recover marketable coal from banks deposited as refuse in the earlier days of mining. The same condition holds good today and we can see a definite trend toward the objective of utilizing the No. 4 and No. 5 buckwheat sizes. Under present conditions, we know that a limited market exists for No. 4 and 5 buckwheat; we know that the finer sizes of anthracite can be burned as pulverized fuel; we know that these fine sizes can be processed and sold as briquettes, and we also know that the heat content of these fine sizes of coal is comparable with that of the sizes now merchandized. Therefore, it should be the objective of the industry to conserve these fine sizes of coal in such manner that if and when processing facilities are developed, and market demands exist, the material can be salvaged and made available to meet these demands. From the standpoint of economical future reclamation, the ideal method of disposing of refuse containing the finer sizes of coal is to deposit it in banks separate and apart from other breaker refuse. By this procedure, the minimum amount of material will be processed in future salvage operations to recover the No. 4 and No. 5 buckwheat sizes. Many collieries in the anthracite region are operated under leases that have a specific term. These lease agreements provide that, upon expiration of the lease by term, exhaustion or otherwise, the ownership of the refuse material reverts to the lessor. Therefore, a question was raised in the minds of these operators: Why should I assume the expense of such procedure when I have no assurance that I will benefit by the future processing and sale of these fine sizes? Under present conditions, the answer to this question is not one of choice but of necessity, as restrictions imposed by legislation make it mandatory on operators of preparation plants to take necessary steps to properly treat silt-laden waters before permitting them to enter the streams. By reason of these restrictions, the disposal of fine sizes becomes a secondary rather than a primary operating consideration, and the procedure for disposing of this refuse becomes a matter of choice governed principally by topographical and geographical conditions. Methods of Handling Refuse There are three principal methods of handling and disposing of fine-size refuse; namely:

Jan 1, 1947

By W. Rostoker

INTRODUCTION THE process of diffusion alloying of mixed powders offers attractive possibilities both commercially and experimentally. It avoids the fabrication difficulties arising from high alloy contents. Indeed, the process is facilitated by larger alloy additions. The direct formation of the part or test piece and the opportunity to perform machining operations under more favorable conditions are of no small importance. However, it is necessary to make only brief mention of the interesting features. As yet there have been few attempts to study this process in detail and to carry it to completion, if such is possible. It was intended not only to measure the intermediate and optimum magnetic properties of sintered alloys but to use magnetic testing itself as a means of following the progress of alloying under various conditions. Having established that alloys could be successfully synthesized by this technique, the opportunity was taken to study the magnetic properties of several alloys of iron and silicon which had hitherto been too difficult to prepare in suitable test piece form, While avoiding some of the difficulties inherent in the fabrication of alloys by orthodox methods, the diffusion alloying process introduces a few of its own. The first of these is the evaluation of the approach to homogenization. In recent years some notable contributions have been made to this question, but considerable uncertainty remains in its definition and evaluation. This work has concerned itself with further study of this problem and it is hoped that more light has been shed on it, The second of these difficulties arises from the presence of porosity. It is inevitable that in the consolidation of a loose bulk of powder, there should remain some voids unless extremely high pressures are repeatedly applied. It was therefore of importance to investigate their effect on magnetic properties. Recently this effect has been studied and certain predictions made of its behavior. It is believed that the experimental results presented here provide a sound substantiation of the theoretical treatment. The attempt has been made to synthesize five alloys of the iron-silicon system and three each of the iron-nickel and iron-cobalt systems. A series of experiments was also conducted on some of these alloys to test their susceptibility to heat treatment. Some unfortunate experiences have provided opportunity to report data on the deleterious effects of nitrogen on magnetic properties when present as a contaminant in protective atmospheres. MATERIALS, PREPARATION OF SAMPLES AND TESTING The purest materials available were employed for this work. A specially purified,

Jan 1, 1948

By F. V. Lenel, M. V. Rose, A. B. Backensto

WHEN aluminum-flake powders are compacted and hot pressed and the resulting compacts are extruded or hot forged, a group of materials with unique properties is obtained. Not only do they exhibit high room-temperature strength of the order of some of the aluminum alloys, but they maintain this strength even after prolonged heat treatment at temperatures up to 900°F. Moreover, their strength decreases much less with increasing temperature than that of aluminum alloys of comparable room-temperature properties. Irmann,&apos;" who discovered and first described this group of materials, which was called SAP, showed that the room-temperature mechanical properties of the extrusions are a function of the oxide content of the powders and therefore the oxide content of the extrusions. The higher the oxide content, the higher are their tensile strength, yield strength, and hardness, and the lower their ductility. Lyle" investigated the relationship between mechanical properties at elevated temperature and oxide content in the course of the development work of the Aluminum Co. of America on materials, called APMP, similar to those developed in Switzerland. He found that at 600 °F, just as at room temperatures, tensile strength and yield strength increase and elongation decreases with oxide content. A careful study of his data (Fig. 3, loc. cit.) shows, however, that the scatter is more pronounced than at room temperature. Gregory and Grant&apos;s data on the creep5 and stress-ruptureb roperties of the one commercial grade of the Swiss SAP and the various grades of APMP indicate that creep strength and stress to rupture also generally increase with increasing oxide .content. Ever since this new group of materials was discovered, attempts have been made to explain the mechanism by which they obtain their high room-temperature and particularly their elevated-temperature strength. In most of the attempts the new material was considered as a dispersion of aluminum oxide particles in a matrix of aluminum. Ir-mann, Von Zeerleder, and Rohner&apos; applied a theory on dispersion strengthening developed by Rohners to the flake-powder extrusions. According to this theory the elastic limit S, of a dispersion would depend upon the average distance between the dispersed particles L according to the equation: 2 E ¦ a where E is the modulus of elasticity and a the distance between nearest neighbors in the lattice of the matrix material. Irmann and co-workers assumed

Jan 1, 1958

By William L. Fink, H. R. Freche

The constitution of the aluminum-cobalt alloys has been the subject of relatively few investigations. Brunckl analyzed the residue which remained after treating the 17 per cent cobalt alloy with dilute hydrochloric acid. The results of the analyses indicated the existence of a compound Co3Al13. Guillet2 published a liquidus for the aluminum-cobalt system which contained two maxima corresponding to CoA16 and CoAl2 and two eutectic points at 30 and 74 per cent Co. The most extensive investigation of the system was carried out by Gwyer13 whose results are summarized in Fig. 1. It will be noted that at the aluminum end there is a suppressed eutectic at 644&apos; C. By extrapolation of the duration of the eutectic arrest to zero, the composition of the compound was placed at Co3Al13. Another thermal arrest was noted at 550" C., which Gwyer attributed to a peritectic reaction. The present investigation of the aluminum end of the aluminum-cobalt system is the eleventh of a series dealing with the constitution of aluminum-rich alloys prepared from electrolytically refined a1uminum.l Preparation of Alloys.—The alloys used in this work were all prepared from electrolytically refined aluminum, the analysis of which is shown in Table 1. In preparing some of the alloys, metallic cobalt (98 to 99 per cent) was used either directly or in the form of an intermediate alloy, the composition of which is given in Table 1. In order to obtain alloys of higher purity, cobalt was introduced into some of the alloys by means of chemically pure cobaltous chloride which was readily reduced by the aluminum. A small electric furnace with nickel-chromium heating elements was used in the preparation of the alloys. Chemical Analysis.—An analysis sample from each melt was chill-cast in the form of a plate 3/16 in. thick. In all cases the casting temperature was well above the liquidus. Cobalt was determined electrolytically in all alloys containing more than 0.05 per cent cobalt and by the nitroso-

Jan 1, 1932

By N. Tiner

In an article on magnesium and its alloys, Gann and Winston! stated that manganese has a limited solubility in the liquid state. W. Schmidt2 showed a diagram according to Joseph Ruhrmann indicating that the solubility of manganese in liquid magnesium falls from about 2.65 per cent at 790°C. to about 0.5 per cent at 675°C. The curve was extrapolated to show zero solubility at the melting point of pure magnesium (Fig. I). This work suggested that at the lowest manganese contents it was not possible to detect either a eutectic or the formation of a solid solution, so that manganese was stated to be present in the structure as large individual crystals. Two years after the presentation of, these results, Gann3 brought evidence to show that some manganese is capable of dissolving to form a solid solution, any excess of manganese present then assuming chiefly the form of small blue-gray particles scattered throughout the primary magnesium crystals, occasionally in clusters but seldom at the boundaries. He presented micrographs of cast alloys containing 0.1, 0.4 and 1.0 per cent manganese that show a small amount of coring around the individual particles due to solid solution, and indicate that the location of the manganese particles is closely associated with the original dendritic structure of the primary crystallization. He also pointed out that the amount of manganese that can be dissolved in magnesium depends on the amount of other alloying elements present, the effect of the second alloying element being to decrease the amount of manganese retained in the alloy. E. Schmid and G. Siebe14 have determined the solid solubility of manganese and its variation with temperature by means of precision X-ray diffraction measurements. They found that the solubility is 3.4 per cent at 645°C. and decreases to practically zero at 200°C. It is apparent that the X-ray diffraction measurements of Schmid and Siebel are not in agreement with the results of Ruhrmann as presented by Schmidt. This disagreement, and also the lack of data on magnesium-aluminum-manganese and magnesium-zinc-manganese systems, led the author to the present investigation, to determine accurately the liquid solubility of manganese and its variation with temperature. Experimental Method The experimental method for making solubility determinations consists of saturating the molten magnesium or the alloys with manganese at various temperatures and analyzing the dip samples taken from these melts. Melting is carried out in a cylindrical steel pot heated by an electric resistance furnace. The pot had a metallic outside coating to protect the steel against oxidation, and a graphite inside lining to prevent the contamination of the melt by prolonged contact with iron. The inside diameter of the pot was 4 in. and the depth 12 inches.

Jan 1, 1945

By N. S. Stoloff, M. Gensamer

The effects of temperature, grain size, and magnesium content on the strength and ductility of cadmium were studied in the range -269° to 23 °C. A sharp drop in ductility between -140° and -190°C marked a transition in fracture mode from ductile shear to inter gvanular fracture. The addition of up to 15.35 wt pct Mg in solid solution raised the transition temperature, and for two compositions produced basal cleavages. The relatively high ductility of unalloyed cadmium at cryogenic temperatures is attributed to temperature independence of the yield stress and the absence of dislocation locking. ALTHOUGH numerous studies have been made of the mechanical properties of cadmium single crystals,1-5 no comprehensive investigation of deformation and fracture in polycrystalline cadmium has been carried out. This investigation was initiated to determine the mechanisms of deformation and fracture in polycrystalline cadmium, and to attempt to correlate these mechanisms with the behavior of other metals of hexagonal close-packed structure. A second objective of this work was to determine whether the slip systems and mode of fracture of polycrystalline cadmium could be altered by reduction of the axial ratio, which is 1.886, through alloying. Dorn and co-workers&apos; have been able to induce prismatic slip in magnesium through alloying to reduce the axial ratio. Single crystals of cadmium are so highly ductile that glide strains of 400 pct are observed at room temperature, and 80 pct at -269°c.l Failure in every reported case is by necking down and shearing off of the crystal. Cleavage in cadmium single crystals has never been observed. Magnusson and aldwin reported a ductile to brittle transition for polycrystalline cadmium at about —160 C for specimens tested at a strain rate of 0.05 in. per in. per min in tension. However, even at -196° substantial ductility was retained. The transition temperature was strain rate dependent, impact loading at 19,000 in. per in. per min raising the transition temperature to about -85°C while leaving the low temperature ductility unchanged. No information was given on microstructures or fracture mode, nor were stress-strain characteristics reported. In the present investigation, tensile tests were carried out on polycrystalline cadmium and solid solution Cd-Mg alloys to determine flow stresses, work-hardening behavior, fracture strength, and fracture ductility. Variables employed included temperature, alloy content (and c/a ratio), and grain size.Some compression data also have been obtained for comparison purposes. I) EXPERIMENTAL PROCEDURES A) Polycrystalline Cadmium. Two purities of cadmium were employed in this investigation. Cadmium of 99.95 pct purity was supplied by Belmont Smelting and Refining Co. Cadmium of 99.99+ pct purity was supplied by American Smelting and Refining Co. Threaded cylindrical tensile specimens were machined to give approximately a 1/2 in. long straight gage section with a reduced diameter of about 0.110 in. Specimens were heat treated in evacuated pyrex capsules in the range 140" to 290°C to produce desired grain sizes. All specimens were polished in a solution containing 320 g Cr03 and 20 g NaS04 in 1000 cc of H20. Occasionally this solution caused staining of the specimen, in which case a further polish with a solution of one part HN03 and two parts H202 in two parts ethyl alcohol was employed.&apos; Tensile testing was carried out on a screw-driven Instron machine, at a constant rate of crosshead movement, which for most tests was 0.01 in. per in. per min. Load vs time, and therefore load vs total strain, were recorded. Temperatures from -196" to -40°C were obtained by passing nitrogen gas through a copper coil immersed in liquid nitrogen, and then feeding the gas into a jacket around the specimen. Tests at -269°C were carried out in a helium cryostat, at a crosshead rate of 0.1 in. per in. per min, with specimens of 1 in. gage length. Compression specimens were tested in a jig similar to one designed by ilman. The jig was mounted on the under side of the Instron crosshead. Load was transmitted by means of a rod connected to a movable part of the jig at one end, and the load cell at the other end. A downward movement of the crosshead applied a compressive load to the specimen.

Jan 1, 1963

Marius R. Campbell, Washington, D. C. (Communication to the Secretary): Having spent considerable time in a systematic study of this field,&apos;both as regards the details of its structure and stratigraphy and the broader question of its relation to the Appalachian basin, I wish to add a word concerning the correlations made by Mr. Hodge in his paper. From A commercial standpoint, the question of correlation, even along the strike, is of little or no importance, for the variability is such that each opening should be tested individually, regardless of the horizon at which it appears; but these correlations assume an entirely different aspect when their scientific bearing is considered; they then become of vital importance, and, before being accepted, should be subjected to the most rigid scrutiny and severe tests. Mr. Hedge states positively that the Imboden seam of coal is equivalent to the Elkhorn coal of Pike and Letcher counties, Kentucky. In the writer&apos;s opinion, such equivalency is, at the utmost, only a possibility. This correlation has been generally adopted by the geologists of the Kentucky Geological Survey, but the writer has been unable to find on what grounds they base the statement. The evidence of identity seems to be in the character of the coal, as both seams furnish fine coking coal; also, in the similarity of the interval between them aid the " Bee Rock," or top of the great conglomerate. It is probably an easy matter to determine the thickness of this interval on Elkhorn creek, and we will assume that the measure of 300 to 350 feet is correct. In the Big Stone Gap field, the question of measuring the interval between the Imboden and the " Bee Rock " is exceedingly difficult. The line of outcrop of this mass of strata is along the northern base of Stone mountain, and, in the uprising of the great arch, the strata on its flanks have been thrown into many minor folds, and, in places, have been broken by local faults. This complicated structure, together with the absence of ally well-marked stratum, renders it exceedingly difficult to obtain a correct measure of the intervening sandstones and shales. The interval of 300 to 350 feet, assigned by Mr. Hodge, is mani-

Jan 1, 1893

By J. D. Price, G. V. Woody

MANY producers of by-product coke have spent considerable time and given considerable thought to the use of a substitute for low-volatile coal as an admixture with high-volatile coking coal for charging high-temperature by-product coke ovens. Generally speaking, there are two principal reasons for this, one being occasioned by the lack, in certain localities, of low-volatile coal delivered at a cost sufficiently low to permit its use, and another being the desirability of finding a substitute product with a delivered cost lower than the cost of the low-volatile coal. It was for the first of these reasons that the Colorado Fuel and Iron Corporation was particularly interested in this subject. This paper deals with the experimental work and the results obtained by that company in producing and using low-temperature char. OVERCOMING DEFECTS IN COKE FROM COLORADO COALS Colorado coking coals are all of the high-volatile type and were all laid down during he Cretaceous period, which makes them something like 80 million or so years younger, geologically, than eastern coals. They behave, when coked alone, something like the Pittsburgh-seam coals. They make a very brittle, highly cross-fractured, fingery coke of fair shatter value but quite low in resistance to abrasion. Mixed with a low-volatile coking coal such as Pocahontas or certain Oklahoma coals, they make a coke that is equal to virtually any of the eastern cokes. The inferior quality of Colorado coke has long been recognized, and a considerable amount of time, effort and money has been spent upon investigations of means for improving its physical properties. While it has been found that some improvement in quality of coke can be secured by the adjustment and control of such variables as pulverization, moisture content, oven temperature and bulk density of charge, and through the addition of certain inert materials such as pulverized coke breeze, pitch, high-volatile noncoking or semicoking coal, the benefits so gained have been comparatively small. There are no low or medium-volatile coking coals available within a reasonable freight-cost distance. Therefore a substitute for low-volatile coal was developed through low-temperature carbonization-or, more correctly speaking, through the partial devolatilization-of high-volatile coals. The blending of this lower volatile char with the high-volatile coking coal distinctly improves the quality of coke made from Colorado coal. There is nothing especially new about this idea. In 1908 a Japanese patent was taken out by Kotaro Shimomura, whose claim was: The method of making a nonfingery coke out of bituminous coals, without using natural coals of low-volatile matter; according to which a certain coal is heated at a temperature between about 300°C. and about 600°C. so as to leave about 15 to 25 per cent volatile matter in the coal; this coal is mixed with the

Jan 1, 1944

By T. L. Berger, M. C. lnman, M. F. Rose

AFTER shock loading, mechanical twinning has been observed by several authors1-4 for a variety of fcc metals and alloys. It has been shown2,3 that low stacking fault energy materials deform primarily by the formation of mechanical twins and stacking faults during dynamic loading. Further, Inman, Murr, and Rose2 have shown that a correlation exists between the average hardness and the volume percent of mechanically twinned material in stainless steel. It has been observed in nickel&apos; that mechanical twins are produced during shock deformation if the peak pressure associated with the shock wave is greater than 350kbar. It is the purpose of this work to discuss shock-hardening mechanisms in nickel and to determine to what extent the initial condition influences the final response. The "flying plate technique" was used to generate planar shock waves with peak pressures as high as 710 kbar. The details of the technique5" have been discussed elsewhere and will not be repeated here. The material used in this study was polycrystalline nickel of 99.98 pct purity and had an initial grain size of 0.18 mm. Prior to any deformation, all samples were given a stress-relief anneal of 6 hr at 600°C in a vacuum furnace. After annealing, the initial condition of the samples was changed by cold rolling individual samples from 0 pct to a maximum of 80 pct reduction in thickness. This provided samples with controlled initial histories for shock deformation. Immediately after shock loading, the specimens were slowed and quenched by suitable arrangement of buffer material and water bath. This technique minimized specimen damage and the effects of shock heating on the samples. Since deformation twins in high stacking fault energy fcc metals are only produced at high stress levels, the initial defect configuration prior to mechanical twinning should greatly influence the number of twins formed. In fact, Zukas and Fowler7 have found that prestraining in iron and steel greatly suppresses the formation of mechanical twins during the shock deformation process. Following the methods of Warren and warekois8 and Cohen and Wagner,9 X-ray diffraction profile analysis was performed to determine the incidence of stacking faults and mechanical twins in the deformed nickel. Analysis was made on samples in the annealed, cold-rolled, shock-loaded, and cold-rolled—shock-loaded states. The estimated accuracy of the peak position measurements is *0.01 deg in 2?. Therefore the accuracy in the measurement of the relative peak shifts between the (400) and (222) diffraction line profiles is on the order of *0.04 deg in 2?. None of our measurements indicated peak shifts greater than 0.04 in 28. This would correspond to a fault density of -1 pct and sets the limit on the sensitivity of our instrument. Similarly, analysis of the shift in the center of gravity of the diffraction line profiles yields the minimum measurable value of the twin fault probabilitya to be -0.002. Fig. 1 shows mechanical twinning in nickel shock-loaded to 400 kbar. In this photograph, fault intersections are clearly visible. The thickness of the individual fault lamellae was determined from the image widths of their lines of intersections by the

Jan 1, 1968

By Robert M. Asimow, Bisuddhi Riddhagni

THE use of back-reflection Laue photographs for crystal orientation is standard in many research projects. The time required for the orientation of any given crystal depends largely on luck and the skill of the investigator. In any case, where large numbers of crystals are used, this time can be quite considerable when the standard methods involving a Greninger chart and Wulff net are employed. For this reason the technique described below was developed. This new procedure applicable for cubic crystals (without preparation of new charts) largely eliminates the skill and luck factors and makes possible the determination of orientation in minutes. The procedure works in the following way: the Laue photograph is examined and arbitrarily an important hyperbola is picked out. An important hyperbola is defined as one which contains some spots through which several hyperbolas pass. Such a hyperbola inevitably corresponds to a low-index zone axis. The angles between the reflection planes causing all the spots on this hyperbola can be read directly from a Greninger chart in the usual manner. If the hyperbola contains more than five spots, these interplanar angles will positively identify the zone axis and the indices of each of the reflecting planes. However, initially since neither the zone axis nor the Miller indices of a diffraction spot are known, there are a very large number of possibilities that must be considered. Using tables of interplanar angles to check out every possibility would be quite tedious, so that this approach is not generally used. Instead, various poles are plotted on a stereographic projection. In the cubic system, however, it is possible to plot out once and for all the angles between all planes of a particular zone axis up to some upper limit of Miller indices. This has been done for all low-index zone axes in Fig. 1. Given these charts, it is quite a simple matter to identify the spots of the unknown hyperbola. All that is necessary is to plot the angles read off the Greninger chart to the same scale as Fig. l(b) on a strip of paper. Then this strip can be slid along each zone axis of Fig. l(a) until a good fit is obtained. Once this is accomplished, the Miller indices of each spot are given and thus the orientation problem is solved. To check, it is possible to pick a second hyperbola and see whether a compatible solution is obtained. This is most conveniently done by picking for the second hyperbola one that contains a spot common with the first. If the specimen axis is chosen parallel to the incident X-ray beam, then usually the [ loo] hyperbola will appear on the film as one of the more important hyperbolas and consequently will be either chosen initially or as a check. Using the [100] hyperbola, after the (001) or (011) spots have been identified, the latitude and longitude of the specimen axis can be read directly from the Greninger chart so that it is unnecessary to use a Wulff net for the entire orientation procedure. This makes the operation not only quicker but more accurate. Of course, if neither the first nor second hyperbolas correspond to [loo] then a third hyperbola can be chosen.

Jan 1, 1967

By J. Brett, S. Friedman

A study of the diffusion of nickel into both fibrous and recrystallized 0.065-in.-diam silica-alumina doped tungsten wire at 1200°C has been conducted. The diffusion profiles were determined by chemical dissolulion of successive circumferential layers and spectro-photometric determination of nickel in the etchant. It was shown that the diffusion of nickel into tungsten is markedly structure-sensitive. No significant amount of nickel could be introduced at 1200°C into tungsten which had first been recrystallized at temperatures of 2300°C, leading to an upper limit estimate for the bulk diffusivity of nickel in tungsten of 10-11 sq cm per sec at 1200°C. However, nickel did diffuse into initially fibrous wires, caused re crystallization, and was alsays present at the advancing recrystallization front. The accumulation of nickel in an initially fibrous region ceases when the recrystallization front arrives. The characteristics of the resulting diffusion profiles are explained in terms of the relation between diffusion processes and micro structural changes. WHILE it is well-known that the presence of nickel and certain other metals causes low-temperature recrystallization of the cold-worked fibrous structure of doped tungsten,&apos;&apos;&apos; the character of the phenomenon is not well understood. Recently it was shown that the recrystallization reaction can be induced at low temperatures by the presence of solid nickel on the surface of doped wire, but not by exposure to nickel vapors.3 Once recrystallization was initiated, a continued source of nickel was required for propagation of the recrystallization front. The phenomenon appeared to be a complicated solid-state reaction since, in addition to the change in structure of the tungsten from fibrous to equiaxed grains, diffusion of nickel from the surface into the recrystallized and fibrous structures was occurring with possible chemical reaction between nickel and tungsten at the surface and interaction of nickel with the dopants which ordinarily stabilize the fibrous structure within the wire. The objective of the present work was to study the diffusion of nickel into fibrous and recrystallized tungsten structures in order to clarify the relationship of the diffusion process to nickel-induced recrystallization. Studies of interdiffusion in the Ni-W system4"7 have been principally concerned with the mobility of tungsten in nickel. Information on diffusion of nickel in tungsten is meager, and there seem to be no studies of the mobility of nickel in tungsten which elucidate the structural aspects of the phenomenon. The present experiments were designed to determine the diffusivity of nickel at 1200°C in fibrous tungsten, in large-grained tungsten prerecrystallized at high tem- perature, and in tungsten recrystallized at 1200°C due to the presence of nickel. Although the concentration profiles were determined in these experiments, unexpected changes in boundary conditions or complex diffusion paths complicated the analysis of these profiles so that the interpretation could not be readily expressed in terms of simple diffusivities. EXPERIMENTAL PROCEDURE The experiments were conducted on Sylvania 0.065-in.-diam silica-alumina doped commercial tungsten wire, hereafter designated as ordinary doped wire. A special lot of 0.065-in.-diam silica-alumina doped low-nickel wire was also used. Impurity analyses of the wire are given in Table I. The experimental design is shown in Fig. 1. All heat treatments were conducted in vacuum. Two groups of ordinary doped wire were prepared for diffusion with nickel. One group in the as-received (fibrous) condition was elec-tropolished, nickel-plated, and annealed at 1200°C for 20 and 40 hr, or was exposed to nickel vapor at 1200°C for 20 and 40 hr. The second group was recrystallized to an equiaxed grain structure at 2300°C for 3 hr, electropolished, nickel-plated, or exposed to nickel vapor for anneals at 1200°C for 20 and 40 hr.

Jan 1, 1969