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By W. A. Barrows, Carl Zapffe

Jan 1, 1913

By Robert Linton

SPECIFICATIONS for clays serving raw materials in the ceramic industry usually contain the following items: (1) Chemical analysis, sometimes with mineralogical structure determined by microscopic investigation. (2) Maturing temperature and firing range, usually expressed in terms of pyrometric cone equivalent. (3) Shrinkage (drying and firing shrinkage separately). (4) Degree of plasticity. (5) Color after firing. Clays may be classified, according to the purpose of the classification, in several different ways, based on (1) Geological occurrence-transported, altered An place, indurated (shales). (2) Physical properties-plastic (bonding clays, ball clays) and nonplastic (flint clays, kaolins or china clays, bone clays). (3) Color. (4) Use. A common classification of clays for commercial designation, and one fol-

Jan 1, 1936

By W. B. Plank

I HAVE been asked by the Chairman of the Engineering Education Committee to outline what the engineering colleges would like the mining companies to do with the young engineer just, out of college. It has been suggested that, because my own personal experience has been almost entirely within the coal industry, I handle this subject from that view- point. Many thoughtful engineering educators are today wondering what effect the present employment policies of many industrial companies are going to have on the young man just starting out in his profession. Since before the World War large industrial concerns, such as those in the electrical and mechanical industries, have gone to the colleges to interview 'likely graduates until today there has developed in most of our engineering colleges, especially in the eastern part of the country, what appears to be a lively competition on the part of these companies in bidding for the young graduate engineer. The employers of mining engineers, with one or two exceptions, have not yet begun to enter this competition. It remains to be seen how serious this situation will become, how it will affect the training of other engineers, and what effect, if any, it will have on the mind of the young man about to enter college who may be inclined to take up mining engineering, but who gains the impression from this well advertised employment service just mentioned that there is no opportunity for him in the mineral industry as com- pared with the opportunities offered in say civil, mechanical, and electrical engineering.

Jan 1, 1929

By H. W. Weart, S. Sarian

The effect of temperature, impurities, and capillarity on the morphology of solid particles in a liquid matrix is investigated. For the NbC-liquid iron system, at least, it is found that those particles which are dissolving will assume a spherical shape; all others will take the equilibrium shape. An aniso-tropic-isotropic transition in the interfacial tension is observed at about 1725°C and is believed to be due to cooperative roughening of the interface on an atomic scale. The presence of boron as an impurity lowers the transition temperature. It is the purpose of this paper to present some observations on the shape of particles in a system consisting of solid particles of nonuniform size in a liquid matrix and to propose an explanation for these observations. Although qualitative only, these observations, which were made largely in the course of a study of particle-growth kinetics,' are presented separately to draw attention to them. They reveal the effect of particle size, temperature, and surface-active impurities on particle shape in the solid NbC-liquid iron system, and are unique in demonstrating for the first time the reversible effect of temperature and impurity concentration on the anisotropy of the energy of the solid-liquid interfaces involved. Some of the changes are readily demonstrated with the few particles whose shapes can be clearly shown in a single photomicrograph, but others are more subtle, and their recognition depends upon some experience and the study of a much larger sample. For this reason, a certain amount of judgment must be acknowledged to underlie their detection, but every effort was made to maintain proper scientific objectivity in evaluating the observations. The observations will be presented first, and then an explanation rationalizing them will be offered. Brief reminders of the relevant theories will be inserted into the discussion at the appropriate points. EXPERIMENTAL The observations were made of the shape of NbC particles dispersed in liquid iron. All specimens were prepared by pressing NbC powder and car- bony1 iron powder treated according to standard powder-metallurgical procedures. Except as noted, all specimens were heated to 1900°C for 1/2 hr. In investigating the effect of temperature, specimens were quenched from 1900°C to various lower temperatures, and held long enough (48 hr) to allow shape changes to be completed. Boron, in the form of amorphous powder, and oxygen, in the form of niobium oxide powder, were added as impurities to some specimens. Nitrogen was introduced into some specimens as an impurity by heat treating the samples under a partial pressure of nitrogen gas. Effect of Temperature. It was first observed that increasing temperature had a marked effect in rounding the corners and edges of the particles. At 1900°C the particles are all very nearly spherical and remain so even when the holding time is extended to 48 hr. As can be seen from Fig. 1, there is virtually no evidence of the presence of low-index planes. At 1300°C, however, most of the particles are prismatic with very slightly rounded edges and corners, Fig. 2. (In making these shape comparisons, attention should be focused on the larger particles, even though the smaller ones would be expected to be the first to change shape, owing to the shorter diffusion distances involved. The cause of the mixed particle shape will be explored later.) It was established that these structures are re-

Jan 1, 1965

By Harvey S. Mudd

WHEN talking about Europe it is well to endeavor to keep politics and economics apart but they have become so intermingled in recent years that the discussion of one topic inevitably leads to the other. Economic influences formerly prevailed over politics but today national and international economies are governed by political forces. Commerce in raw materials, food, and manufactured products is now so restricted by governmental controls that the capacity of the world to produce and to consume, to exchange goods and to pay for them, cannot be gauged by the past or predicted by a multitude of experts. In prewar days there were some yardsticks but today there are none. Thus it is hazardous to make forecasts, for in all fields of human effort the world has been cast upon an uncharted sea.

Jan 1, 1947

By H. R. Spedden

As part of its broad program of educational assistance-including grants, fellowships, and scholarships -Kennecott Copper Corp. offers summer employment opportunities for college students at each of its operations. Originally, summer job programs were intended primarily to help students defray their educational expenses. But as the nation's economy expanded with increased production and a growing demand for well-trained engineers, summer employment became an important means to help qualify future professionals and fill industry's need for them. Both the employer and the student thus derive benefits from such training.

Jan 6, 1967

By John R. Freeman

THIS paper relates entirely to the casting of brass under fluid pressure in steel dies Die castings of metals and alloys of low melting point have been available for many years but the development of brass die castings has been relatively recent, occurring within the past five years The successful development of brass die castings has resulted from the availability of alloy steels that would give satisfactory life as dies and the development of suitable die-casting machines The melting point of copper is 1084° C (1983° F ) and of a high brass about 900° C (1652° F ), as compared to 327 5° C (6215 °F) for lead, 419 4° C (786 7° F) for zinc and 658 7° C (1217 7° F) for aluminum, temperatures that ale indicative of the casting temperatures required for the various types of alloys, of which the melting points are somewhat lower than those of the pure metals TYPE OF DIE-CASTING MACHINE In the well-known plunger and gooseneck types of die-casting machines used for casting the alloys of relatively lower melting point, much of the working mechanism is immersed in the molten alloy For alloys having the relatively high melting point of brass, this obviously is impossible Machines with a separate melting chamber are essential Two rather distinct types are in extensive use at the present time, the Pack and Polak machines, the former designed by Charles Pack and developed at the American Brass Co , the latter developed by Zoseph Polak, of Prague, Czechoslovakia The principles of operation of these machines have been described in articles by Pack1,2,3 and by Sieg 4

Jan 1, 1935

By Carl F. Lutz, Robert F. Mehl

A layer of brass was formed on 0 brass using a vapor-solid reaction technique. The variation in composition with distance within the phase layer and across the a -ß interface was determined by an electron probe. The diffusivities were calculated as a function of concentration; the diffusion coefficient in brass was found to change by a factor of twenty-five over a compositional range of 8 at. pct. An explanation is proposed to account for the large change in the diffusivity, based on possible stresses in the diffusion zone; it is suggested that the concentration dependence of the thermodynamic factor might decrease the activation energy with increasing zinc content. ALTHOUGH the body of data on diffusion coefficients (D- values) in terminal solid solutions and in isomorphous systems is quite large, there is but little quantitative information for intermediate alloy phases. This paper recounts measurements of D-values for the ? phase in the Cu-Zn system. The experimental method employed consists in the forming of alloy layers and determining and analyzing the concentration-distance (c-x) curve. The experimental method employed could have been applied to several systems; the Cu-Zn system was chosen because the D-values in the a and ß phases are well known, and can thus be used for purposes of comparison. The rates of growth, of intermediate alloy layers are known for a number of systems. In nearly all cases the thickness varies parabolically with time, as to be expected if the rate of growth is controlled by the D-values. In a few cases, nonparabolic behavior has been noted.1 Nonparabolic growth may be the result of a) a variation of the composition at the phase interface with time, b) interface-controlled kinetics, c) short diffusion times coupled with long vacancy lifetimes (in vacancy diffusion),2 and d) crystal reorientation during diffusion in anisotropic systems.3 In the Al-Ni system4 and in the Al-u5 system parabolic growth kinetics are slowly approached after an initial transient period. In general those phases stable at a given temperature in a system A-B appear as separate phase layers when A and B are put in contact and given a diffusion he at-treatment. In most cases the compositions at the interface of adjacent phase layers are those read from the phase diagrams for the termini of the respective phases. This discontinuity in concentration is taken as independent of time, and the growth of one phase at the expense of another is assumed to be independent of interface kinetics, i.e., the rate of interface movement is controlled by volume diffusion in the phases concerned. wagner6 has given many solutions for multi-phase diffusion processes, assuming that the chemical diffusion coefficient, D, is independent of concentration, as have others. Jost7 has pointed out that the familiar Boltzmann-Matano solution is as valid for a (c-x) curve exhibiting concentration discontinuities at phase boundaries as for the usual single-phase case. Heumann and associates8,9 have applied this solution to the multi-phase case, but lacking concentration data within the several layers were forced to assume linear concentration gradients, thus calculating only average D-values. The purposes of the present study were: 1) to determine the dependence of D on concentration, 2) to calculate the intrinsic diffusion coefficients D?Cu and D?Zn, 3) to establish the time-law for the movement of the interface, 4) to determine the concentration limits at the ß-? interface, and, 5) to attempt to clarify the mechanism of diffusion. EXPERIMENTAL PROCEDURE The ? phase is exceedingly brittle; conventional solid couples and conventional sectioning methods were thus inapplicable. Vapor-solid couples were accordingly employed. Unfortunately the equilibrium vapor pressures of Zn for the ? phase are unknown; to escape this awkwardness, samples were enclosed in a capsule containing powder of an alloy of known composition, sufficiently large in quantity as to be effectively an infinite source, maintaining the concentration of Zn at the sample surface constant with time. The sources of Zn-vapor which can be employed in reaction with Cu, or a brass, or ß brass, to form a layer of ? brass satisfactorily wide in concentration range, are alloys of the phases ?, or ? + E or ? + d (depending on the temperature).

Jan 1, 1962

By S. Shapiro, J. A. Ford

The crystallo graphy of several grains of the lamellar Ni-Ni3 B eutectic solidified under conditions appvoachirt equilibriun has been examined.. A preferred interlace relationship has been observed in this system which may be described as: A detailed study of this interface by means of an atomistic model reveals the interface Plane to be puckered and more accurately describable as: Analysis of the atomic configuration on these planes reveals that this interface conforms to the orientation criteria proposed for directionally solidified lanzellar eutectics. In addition, a model is proposed to help explain the growth of a "puckered" interface. DURING the past several years an increasing number of unidirectional solidification experiments on eutectic alloys have been reported in the literature. The primary emphasis, in most cases, has been the effect of growth variables on microstructure. Some authors have, however, studied the crystallographic orientations and interfacial relationships resulting from these experiments. Four crystallographic studies have been published;'-4 in each of these there is a factor which tends to prohibit the extension of the observations beyond the particular system. In the Al-CuA12 system, the approximately equal volume fractions of the phases have resulted in sorrie controversy as to which phase constitutes the matrix. In the Mg-Mg,Sn eutectic the presence of several interconnecting lamellar systems2 complicates the analysis. The two remaining eutectics. 1nSb-b and Al-Ali, are both rodlike (the latter exhibits a plate morphology at low solidification rates): the matrix in each of these alloys is the simpler (fcc) phase. Studies of nucleation in eutectic alloys indicate that the more complex constituent is the nucleating -.phase.' For the purpose of an analysis of interface crystallography it was deemed desirable to study a lamellar system in which the matrix was the more complex and, presumably. the nucleating phase. Recently it has been shown that there is a eutectic in the Ni-B system between Ni3B and i; preliminary investigation by the authors indicated that the matrix in this eutectic is the boride phase. The structure of Ni,B is reported to be orthorhombic, the compound being isostructural with cementite. This eutectic is therefore analogous to that between austenite and cementite. Furthermore, the Ni-B system is not susceptible to a competing eutectic reaction as occurs in the Fe-C alloys; nor does it exhibit a eutectoid decomposition. For these reasons the study of the interfacial relationship in the Ni-Ni3B eutectic was undertaken. PROCEDURE Ingots of nominal composition 3.5 and 3.6 wt pct B were induction-melted in boron nitride crucibles within a graphite susceptor under a boric oxide slag; a dynamic argon atmosphere was maintained over the melt to prevent rapid oxidation of the susceptor. The raw materials used were 99.97 pct Ni from Huntington Metals Division of International Nickel Co. and 99.8+ pct B chips supplied by the United Mineral and Chemical Co. The analyses of the starting materials are listed in Table I. The melt was held at 1500cC for 2 hr to insure complete reaction and allowed to cool to room temperature. The castings were chemically analyzed for boron and were metallographically examined for uniformity. The metallographic samples were mechanically polished through 0.1 p A1,0, and etched with Carapella's reagent. Each master casting was cut up to make several specimen blanks which were remelted, in boron nitride crucibles, to form ingots of 1/2 to 3/4 in. in diarn by 7 in. long. These ingots were unidirectionally solidified using a modified vertical Bridgman technique. The rate of solidification was assumed to be equal to the uniform rate at which the crucible-susceptor assembly was withdrawn from the induction coil. This rate was varied from 1.18 to 2.21 cm per hr. In order to perform these experiments it was found necessary to redetermine the eutectic composition. This was accomplished by the zone-melting technique described by Yue and lark.' An ingot of 4.5 wt pct B was zone-melted in a horizontal boron nitride boat 7-1/2 in. long with a 112-in.-square cross section. A

Jan 1, 1967

By W. F. Domis, F. von Gemmingen, R. W. Whitmore, F. Garofalo

Constant-load creep-rupture tests at 1100°, 1300° and 1500°F were made on a Type-316, 18 Cr-8 Ni-ZMo, austenitic stainless steel to determine the relationship between ruptzire life and other aspects of creep behavior, The test program was designed on a statistical basis to check the variability of various creep properties and to determine emprically the dependence of minimutn creep rate and ruptzire life on initial stress, The dependence of the rupture life, tn, on the initial stress, a,, is found to be related to the stress dependence of the wrinimum creep rate, i,, and secondary creep strain, The instabilities or breaks found in the conventionual log - logt, plots can be traced to either a change in the linear dependence of log i, on log a, or to a change in secondary creep strain. In the alloy tested, rupture is found to be predowmanantly of the inter crystalline type, For this material, the secondary creep strain is -found to depend on the nature of the grain boundary precipitate, which affects grain boundary migration. It has been shown for a variety of metals and al tested at elevated temperatures under creep conditions that the empirical relationship exists between rupture life, tr, and minimum creep rate, i,. This relationship, in which C and a are in some cases independent of stress or temperature, shows in a general way that creep rupture depends on creep behavior prior to tertiary creep stage. To determine more fully the factors which control creep rupture requires, therefore, more knowledge of the relationship between rupture and creep behavior. To obtain such information, constant-load creep-rupture tests have been made on a Type-316, 18 Cr-8 Ni-2 Mo, austenitic stainless steel at 11000, 1300°, and 1500°F. These tests were designed on a statistical basis to determine the variability of the various creep properties measured and to determine empirically the stress dependence of minimum creep rate and rupture life. As a result of this work, various empirical relations have been established defining more closely the factor, C, in the relation between rupture life and minimum creep rate. This factor is found to be proportional to the amount of strain during secondary creep. The variation of secondary creep strain with rupture life and temperature is determined and its effect on rupture behavior discussed. The variability in minimum creep rate and rupture life is also dis- cussed and the dependence of these quantities on initial stress is interpreted. TEST MATERIAL AND TEST PROCEDURES The material tested is an austenitic stainless steel, AISI Type 316, of the following composition: C, 0.07 pct; Mn, 1.94 pct; P, 0.01 pct; S, 0.021 pct; Si, 0.38 pct; Cr, 18 pct; Ni, 11.4 pct; Mo, 2.15 pct; All 0.003 pct and N, 0.043 pct. The material was received as a hot-rolled 0.5-in. diam bar which was sectioned into 3.25 in. long blanks. Each blank was heat treated by holding 1/2 hr at 2000° F followed by water quenching. A 3/8-in. coupon was removed from each heat-treated blank and the microstructure was examined, the grain size determined, and the hardness measured. The microstructure was found to be uniform from blank to blank and the grain size was found to range from 4 to 6 ASTM numbers. Hardness, measured using a 20-kg load, varied between 127 and 148 Vpn. Tensile creep-rupture specimens having a 0.25 in. diam within a 1.5-in. reduced section were machined from the heat-treated blanks. Constant-load tests were made on these specimens at 1100°, 1300°, and 1500 F. Six specimens were tested at each stress level. Three creep-rupture machines were employed for tests at each temperature, two specimens being tested in each machine for each stress level. To minimize any effect due to local inhomogeneities along the length of the as-received bar the specimens were randomized before test. Each specimen was tested to rupture and in most cases an autographic extension-time curve was obtained. For very low stress levels at 1500°F no auto-

Jan 1, 1962

By R. C. Himes, K. Zanio, J. B. Wagner, J. W. Moody, J. F. Miller, W. Johnson

The recent research with which this paper deals has been concerned with the preparation of pure PbTe crystals suitable for control of radiative processes and othe.v electronic applications. The research has been directed toward the developrnent of methods for the preparation of puve, homogeneous, loul-ca,wier-density PbTe crystals suitable for use in the fabrication of optically pumped and electrically pumped diode lasers. Steps in the preparation of the crystals include 1) purification of the starting materials, 2) synthesis and purification of the PbTe, 3) growoth of PbTe crystals, and 4) adjustt~zertt of composition of the crystals This Paper describes the general preparation procedure employed and discusses some investigntiotzs of purification processes. THE as-received lead and tellurium are nominally 99.999 to 99.9999 pct pure. Assays furnished by the suppliers, obtained by emission spectrographic analysis, show only the presence of about 2 ppm Mg in the tellurium. Such analyses are not very informative, however, since they give no information on the concentrations of normally gaseous impurity elements such as oxygen. Both the lead and the tellurium in the as-received state usually contain appreciable concentrations of oxygen—primarily in the form of oxides on the surfaces of the material. To reduce the oxygen content, the lead is held in the molten state in hydrogen at -700°C for several hours. For this treatment, the lead is contained in a graphite boat which is rf-heated, and the hydrogen is passed through a liquid-nitrogen trap prior to use. Tellurium is subjected to multiple distillation in hydrogen to remove oxygen and other foreign impurities. If, in conducting the distillation, approximately the last 10 pct of the charge at each "stage" is left as the discard cut, purification of the tellurium proceeds as indicated in Fig. 1.l Tellurium used in this work has been distilled two or three times to reduce carrier concentration to about 10" per cu cm. The PbTe is prepared from the purified elements by reacting them in the molten state in sealed, evacuated, or hydrogen-filled silica ampoules. The charge is usually held above the melting point of PbTe (923°C) at 1000 C for 24 hr to ensure that it has been homogenized. Two methods have been employed for the growth of the required single crystals of PbTe: growth from the melt, which will be discussed first, and growth from the vapor phase. Large single crystals of p-type PbTe can be grown readily from the melt by the Bridgman method. For this method of crystal growth, the PbTe is synthesized in situ from proper proportions of the purified elements. Growth (dropping) rates in the range 1.3 to 6.5 mm per hr are employed. Material near the maximum melting composition (about 50.012 at. pct Te) is utilized to avoid composition change on freezing, and thus to obtain crystals of nearly uniform composition. Crystals prepared by use of procedures described up to this point were sufficiently pure and perfect that they could be used for the preparation of laser crystals in studies at Lincoln Laboratory.2-5 ADJUSTMENT OF COMPOSITION Good n-type crystals of PbTe cannot readily be grown from the melt. Further, growth from the melt of crystals of any compositions (n or p type) other than those near the maximum melting composition is difficult because of the large composition changes that must occur on freezing.'" However, melt-grown crystals can rapidly be brought to other desired compositions by isothermal equilibration through the vapor phase with a two-phase Pb-Te stock material, essentially as described by Brebrick and Allgaier.1 The work to be discussed has been concerned with the preparation of n-type crystals; thus, the two-phase stock which was employed contained excess lead, and all ensuing discussion is of the preparation and properties of n-type crystals. The equilibration was designed to the shift composition to the lead-rich limit of stability at the selected temperature. Initial treatment of the crystals at high temperature was employed to minimize the total time required for the equilibration. Final adjustments of the composition and state of the crystals were made by use of low-temperature heat treatment at temperatures in the range 200" to

Jan 1, 1968

By E. J. Dulis, R. M. Fisher, K. G. Carroll

IT is well known that ferritic steels containing more than 15 pct Cr when subjected to temperatures in the range of 700" to 1000°F exhibit increasing hardness and decreasing ductility. The phenomenon has been widely termed the "885 °F embrittlement," after the temperature of most marked effect.'.' In view of the excellent review articles available in the literature3-6 only a brief account of experimentally established facts need be given here. The extent of changes in physical characteristics during embrittlement depends on chromium concentration and time at temperature, higher alloy content and longer time both promoting more rapid and extensive changes. In a 27 pct Cr steel, changes in impact strength and in angle of fracture in bending can be detected after only a 1 hr exposure at 885°F; after 50 hr this steel becomes quite brittle. Hardness increases slowly with time during thousands of hours exposure and may attain a maximum hardness number twice as large as that of the unexposed steel. Microstructural changes accompanying embrittlement have been described as an initial widening of grain boundaries followed by eventual darkening of ferrite grains. Embrittled steels etch more readily, e.g., the weight loss of a 27 pct Cr steel in acid solution may occur at a rate one hundred fold greater following exposure at 885 OF. Marked changes which accompany embrittlement have been observed in electrical resistivity, specific gravity, and magnetic coercive force. Changes in physical properties may be readily removed by heating at temperatures above the embrittling range, such as a treatment at 1100°F for 1 hr. It has frequently been noted that the 885 °F embrittlement suggests precipitation on a submicro-scopic scale of a chromium-rich constituent, the nature of which has not been revealed by X-ray diffraction. Progressive broadening of the body-centered cubic diffraction lines during embrittlement has been observed," and recent observations by Lena and Hawkes' upon single crystals have shown early asterism in X-ray photographs, disappearing within an hour at 900 °F. Many workers have ascribed8-13 the phenomenon to a precipitation of a phase (FeCr), which is known to cause embrittling effects at temperatures much higher than 885°F. Two general observations, however. suggest that a precipitation cannot be responsible for the 885°F phenomenon: 1—prior cold work greatly enhances a formation, whereas it scarcely affects the 885 °F embrittlement, and 2—the presence of an alloying element such as nickel or manganese may have an effect on the 885 °F embrittlement which is opposed to its effect upon a formation. The slight enhancement of a formation and 885°F embrittlement observed in the presence of elements with strong carbide and nitride forming tendencies' is probably a consequence of lessened chromium depletion of the matrix. The bar graph in Fig. 1 shows a typical example, taken from two 27 pct Cr steels used in this work, of the hardness after exposure for 10,000 hr at 900°, 1050°, and 1200°F. Steel A (0.03 pct C, 3.13 pct Mn) showed marked hardening at 900" and 1200°F, whereas steel B (0.12 pct C, 0.63 pct Mn) exhibited only the 900°F hardening. The cr phase was found in steel A at the higher temperatures but not in steel B. Presumably a formation is enhanced by the low-carbon and high-manganese concentrations in A," Thus there are two distinctly different hardening phenomena present which cannot both be ascribed to v precipitation without invoking a transition phase possessing remarkable properties. Materials A number of chromium steels exposed for long periods (5000 to 34,000 hr) at 900°F, as well as unexposed samples of one of the steels, were available for this investigation. Table I gives the chemical compositions and aging treatments of these steels. In addition to these steels exposed in the elevated temperature test furnaces of the National Tube Division, a number of high-chromium steels were heated for short periods in small laboratory air furnaces and lead baths. Supplementing these commercial steels, a sample of high-purity (0.018 pct C, 0.002 pct N) 28 pct Cr iron, exposed 1000 hr at 887°F. was furnished by the Union Carbide and Carbon Corp. In addition, an alloy of iron and chromium of high purity containing 46 pct Cr was used. This

Jan 1, 1954

By H. C. Chang, N. J. Grant, A. R. Chaudhuri

During the creep of coarse-grained polycrystalline magnesium at elevated temperatures, a nonbasal type of slip was found to play an important role in the deformation processes. The nonbasal slip traces were examined metallographically in eight specimens (13 grains) and the observed glide plane located stereographically for each grain. The tests were run at 500&apos; and 700°F at stresses of 148 to 786 psi. Based on these measurements and theoretical calculations, the crystallographic elements for nonbasal slip were determined. WHEN a metal deforms by slip, the conventional appearance of the slip bands on the surface of a specimen is that of straight lines. Such is the behavior of the face-centered-cubic metals, as for example, aluminum. The particular slip system on which slip occurs in these metals is governed by the criterion of resolved shear stress. Work on the body-centered-cubic metals a-iron,&apos;. &apos; molybdenum, and columbium4 has shown that the process of slip is somewhat more complicated than the simple gliding of one close-packed plane of atoms over another in a close-packed direction and that it results in the formation of wavy and irregular slip traces on the surface. The recent reviews and experiments of Vogel and Brick and Chen and Mad-din3 show that, for the body-centered-cubic metals, the resolved shear stresses along the planes and the degree of close-packing do not necessarily determine the slip system. Theirs and other investigations indicate that a complete understanding of the slip process in the body-centered-cubic metals is not yet possible. However, one behavior displayed by all these metals is that the slip direction is invariably one of the close-packed directions <11l>. The wavy nature of the slip-band traces has been explained on the basis of cooperative slip by planes sharing this same slip direction.1&apos;3l&apos; It is generally acknowledged that, when the hexagonal-close-packed metal magnesium (axial ratio c/a = 1.624) deforms by slip at room temperature, it does so by basal slip in the close-packed direction <1I%>. This type of slip has the conventional appearance of straight lines and is governed by the critical shear-stress law. Schmid and coworkers" showed that basal slip was operative in the temperature range of —185" to 300°C (—300" to 572°F). Work by Servi, Norton, and Grant6 has shown that, during the creep of coarse-grained aluminum at high temperatures, new slip systems come into operation. The existence of a new high temperature slip system at temperatures greater than 225 °C (437°F) for magnesium was suggested also by Schmid. This was later indicated by Bakarian and Mathewson&apos; to be slip along the pyramidal planes (10i1) in the close-packed direction <11%>.* They * Mention will be made in the text of pyramidal planes and prismatic planes. The pyramidal planes referred to are the type I. order 1 of the form {10il); and the prismatic planes of the type 1 of the form (10i0). For the sake of brevity, they are termed in the teat as pyramidal (10il: and prismatic (10i0) planes, respectively. Where reference is made to the pyramidal plane type 1, order 2 of the form {10i2), they are referred to as the pyramidal planes of the type (10i2:. observed that this kind of slip resulted in irregular markings on the specimen surface. According to them, a possible cause for the appearance of the markings was due to "limited accessory slip on the pyramidal {10i2) planes." Burke and Hibbard," on deforming a single crystal of magnesium at room temperature, found evidence of slip on a pyramidal plane {1011). They explained it as being due to the effect of grip restraints. It is pertinent to note at this point that, although basal slip has received a fair amount of attention, only two specimens have been investigated, that of Bakarian and Mathewson and of Burke and Hibbard, to establish the elements of nonbasal slip in magnesium. During the study of the deformation mechanisms

Jan 1, 1956

By R. J. Evans, A. G. Revesz

Silicon films have been grown by chemical reaction on (111) silicon substrates. The surfaces were examined by various microscopic and interfero-metric methods. Surface structures are classified into two groups depending on whether the angles between the bounding planes and the (111) surface are small (<3 deg) or large (> I0 deg). A layer-gvowth mechanism is postulated to explain the topography and the influence of growth parameters. Growth proceeds by lateral expansion of the (111) layers in (211) directions, thereby forming steps. The step edges are perpendicular to the direction of expansion and are thus along (110) directions. Step bunching md surface nucleation both play an important role in modifying the flow pattern and in the step regeneration. These are influenced ver?) much by impurities, especially oxygen. Although the growth is a result of a surface-controlled ga: reaction, it is very similar to growth from an impure solution. OVERGROWTH ("epitaxial growth") of silicon is important and widely used. During the past years several investigators have observed various topographical features or so-called faults on the grown films. A study of the surface morphology gives some insight into the growth mechanism. I) EXPERIMENTAL Silicon films have been grown on silicon substrates by a method described by Theuerer.&apos; Silicon is formed by reduction of SiCl, in hydrogen. The rate of incidence of SiC1, was varied between 5 x 10"6 and 2 x 10&apos;4 moles min-&apos; cm-&apos;. The mole fraction of Sic4 in the Hz-SiCL mixture was about 5 x X Both the rate of incidence and the mole fraction were calculated on the assumption that the Hz bubbling through the SiC1, reached equilibrium saturation. Assuming uniform velocity distribution, the linear velocity of the gas in the reactor tube was estimated to be 2 cm sec-l. The temperature of deposition was varied between 1070" and 1260°C. Chemically polished p-type silicon wafers of 20 to 100 ohm cm resistivity were used as substrates. These were within 1/2 deg of a (111) orientation. The thickness of the grown n-type film was determined by two methods: a) measurement of the depth of the p-n junction taking into account the junction shift due to diffusion, b) measurement of the traces of stacking-fault tetrahedra on the (111) plane.&apos; The agreement between the two methods was better than 20 pct. The rate of growth varied between 0.2 and 4.0 pmin-&apos;. The formed films were examined by reflection and light-section microscopy, and by multiple-beam interferometry. Phase-contrast illumination revealed height differences of about 30A. For measuring heights greater than 1 p a light-section microscope3 was used. The resolution by multipke-beam interferometry4 was between 50 and 100A. Crystallographic directions on the surface were determined by X-ray and from the orientation of etch pits and stacking faults as revealed by Dash etch (I part HI?, 3 parts HNOJ, and 12 parts CH3COOH). Some specimens have been investigated by electron microscopy. Platinum-shadowing and carbon-replica techniques were used to show the fine topography and reflection diffraction to determine the structure of the grown layers. 11) RESULTS The topographical features can be classified into three groups. he first group is comprised of fine ridges (striations). The second group consists of structures, which are bounded by planes making angles of less than 3 deg with the (111) surface. They are, therefore, vicinal planes. Structures which closely resemble growth or etch pyramids from the third group would normally be bounded by perfect low-index crystallographic planes. In this case they are deformed. The angles between the bounding planes and the (111) surface are generally larger than 10 deg. The fine ridges are present, even on the smoothest surfaces, as shown in Fig. 1. They are along a (110) direction. The height differences could not be resolved with interferometric techniques. The basic characteristic structures of the second group are depressions and hillocks. These are bounded by three planes, two of them with larger slopes than the third one. The edges between the planes, as projected to the (111) surface, point in (211) directions. Various habits of these structures are shown

Jan 1, 1964

By L. J. Dijkstra, R. J. Sladek

IN earlier work the effect of manganese on the general behavior of nitrogen in iron was the subject of a careful examination by Fast.&apos; Part of the investigation was made, in collaboration with one of the present authors (L.J.D.), by using the internal friction method. This technique, which has been described elsewhere,&apos; showed some significant changes in the properties of Fe-N alloys by the addition of a small amount of manganese namely: 1—Under standard conditions the maximum of the nitrogen peak in pure iron appears around 20 °C and the half width is about 27 °C. By addition of 0.5 atomic pct Mn a general broadening of the internal friction peak occurs. The half width is increased by about 9°C. Moreover, a shift of the maximum of about 5°C toward a higher temperature takes place, Fig. 2. 2—The precipitation of nitrogen from the supersaturated solution, measured, e.g., at 200 °C, was inhibited by the addition of the amount of manganese given in effect 1. The broadening of the nitrogen peak was confirmed in more recent work3 on samples which were prepared from Westinghouse Puron iron. The effect of manganese on the rate of precipitation, however, was much less striking than in the earlier work&apos; based on Philips high purity iron. Theory To explain the first change in properties just referred to, effect 1, the following two assumptions were made:&apos; First, every manganese atom in sub-stitutional solid solution creates around it six interstices located between an iron and a manganese atom, as is evident from Fig. 1. These interstices, conveniently called Fe-Mn interstices, have for nitrogen a free energy level which is AG* lower than the normal Fe-Fe interstices between two iron atoms. For such low concentrations of manganese as 0.5 atomic pct, the possibility that two manganese atoms will be nearest neighbors can be disregarded. Second, it is assumed that in addition to the normal relaxation time TFe at least one new relaxation time, Mn, Fe, enters in the phenomena which is associated with the elementary diffusion jumps between a Fe-Mn interstice and an adjacent Fe-Fe interstice. Fig. 1 gives schematically the free energy for a nitrogen atom as a function of its position along the line ABC. A tentative analysis of the nitrogen peak in a 0.5 pct Fe-Mn alloy based on these assumptions has been made in Fig. 2. The dotted curve with a maximum at 22°C is caused by jumps between Fe-Fe interstices, the dotted curve with maximum at 32°C by jumps between an Fe-Mn and an Fe-Fe interstice. The two effects added together give the dash-dot internal friction curve for the alloy. In the following these two components will be referred to as the normal peak at 22°C and the abnormal one. (The small secondary peak at

Jan 1, 1954

NEW MEMBERS The following list comprises the names of those persons who became members during the period Nov. 10, 1916 to Dec.10, 1916. BARNSDALL, THEODORE N., Treas. & Mgr., Highland Gas Co., 24 West Corydon. St., Bradford, Pa. BEASLEY, FRED E., Asst. Supt., Copper Refinery, Consolidated Mining & Smelting Co., Trail, B. C., Canada., BREWSTER, THOMAS T., Vice Pres., Gen. Mgr., Mt. Olive & Staunton Coal CO., 1012 Federal Reserve Bldg.,-St. Louis, Mo. BRINDLE ARTHUR ST. C.... Mining Recorder&apos;s Office, Trout Lake, B. C., Canada. BULLOCK, Harry E., Pres Kentucky-Jewel Coal Co., Lothair, Ky. BUTLER, PERCY BENNETT, Mine Operator :..P. 0. Box 335, Joplin, Mo. BUTLER, ROBERT S Supt., Three Johns Mining CO., ,Joplin,.Mo. CABOT, Godfrey, L., Carbon Black Maker, Gas Engr., 940 Old South Building Boston Mass.

Jan 1, 1917

By W. Crafts, J. L. Lamont

Electrolytic extraction of carbides from quenched and tempered steel and their examination under the electron microscope were found to be helpful in understanding the mechanism of secondary hardening in alloy steels1 and the same technique has been applied to isothermally transformed steels. A preliminary survey of the utility of the method has indicated that it has promising possibilities for investigating the characteristics of pearlite and bainite. The examination of a few carbon and chromium steels has suggested that pearlite is formed with carbides of two varieties representing the lamellar and nonlamellar forms and that bainite appears to form with a structure analogous to, but significantly different from, martensite. Procedure The method used for examining the carbides was similar to that described previously.1 The chromium steels used in the investigation were made from an Armco iron base in an induction furnace and were not treated with grain-refining deoxidizers. The composition of the steels is given in Table 1. Three-inch square ingots were forged and rolled to 1/8-in. thick sheet for use in the isothermal studies. After rolling, the sheets were sandblasted to remove the rolling scale. Specimens approximately 3/4-in. wide and l1/2-in. long were prepared from the 1/8-in. sheet. They were heated in a salt "bath at 2100°F (1150°C) for 1/2 hour, transferred to other salt baths at 1300°F (704°C), 1000°F (538°C), 800°F (426°C), or 600°F (315°C), held for various periods up to 100 hr and finally quenched in water. After quenching, the specimens were cut in half, one half being submitted to microscopic examination and the other half being used for the electrolytic extraction. The isothermally transformed specimens were submitted to examination under the light microscope to determine the degree of transformation and type of structure and were photographed at 2000 X. Carbide residues were obtained by electrolyzing in 10 pet hydrochloric acid and collecting the residue in glycerine to minimize attack by the acid. As pointed out in the earlier paper, the carbides are attacked to some degree during electrolysis. This attack is relatively minor on the coarser carbides but may be quite severe on the fine bainitic carbides. The residues were washed free of the glycerine with water and alcohol. The alcohol was removed by repeated washings of amyl acetate, and the amyl acetate suspensions were transferred to stoppered vials. Debye-Sherrer X ray diffraction patterns of the residues were prepared using a chromium target. A selection based on the microstructure and X ray data was then made of samples considered most illustrative of the progress of transformation. Electrolytically extracted residues from these samples were submitted to examination by W. D. Forgeng and A. C. Jenkins on the electron microscope at the Research Laboratories of The Linde Air Products Co. and electron micrographs were made at X 5000 and enlarged to 25,000 X. The micrographs obtained with both the light and the electron microscopes are mounted side by side in the accompanying figures. Although no effort was made to establish the T-T-T diagram for the chromium steels, the times for holding in the salt bath were adjusted with the intention of making the samples represent different stages in the transformation at temperatures of 600, 800, 1000, and 1300°F. The degrees of transformation and results of X ray examination are shown in Tables 2—4. The carbide designations refer to orthorhombic or trigonal carbide types and are not meant to imply that the compositions are exactly as indicated. X ray Diffraction Data The occurrence of Fe3C and Cr7C3 in these steels was found to be consistent with published data on comparable chromium steels. Steels A and B with relatively low ratios of carbon to chromium contained Cr7C3 in the pearlitic structures and FeuC in the

Jan 1, 1950

By A. C. Fieldner

IN the Wall Street journal for March 1, 1941, was a tabulation of the construction under way or under negotiation by thirteen iron and steel companies for a predicted increase in annual coke production of 3,165,600 tons. That even this figure was short of the mark is indicated by other plans, then being considered, for a later expansion of 2,610,000 tons additional. Considerable progress has already been made on thin program, for approximately 500 new ovens were put in operation in 1941. and well over 200 should be completed in 1942. These figures include a number of units that replace older and obsolete equipment.

Jan 1, 1942

By F. R. Hensel

THE alloy under discussion is a nickel-silicon bronze and is one of many age-hardening or precipitation-hardening FIG. I.-EFFECT OF QUENCHING AND AGING TEMPERATURES ON HARDNESS OF TWO TYPICAL NICKEL-SILICON BRONZES, NO. 2351 AND No. 2362. copper-base alloys. The hardening con-stituent is a nickel-silicon compound.1 The quenching temperature necessary to pro-duce hardening by a subsequent aging treatment is a function of the composition and extends from about 700&apos;C. to ()50&apos;C.; the aging temperature range extends from 400° to 600°C. The work carried out by the authors of this paper was confined to cast alloys having compositions within the ranges of 1.5 to 3.00 per cent Ni; 0.4 to 0.75 per cent Si; balance, copper.

Jan 1, 1940

By C. H. Herty

The elimination of metalloids in the open-hearth process depends on oxidation of the metal by diffusion of FeO from the slag (or absorption of FeO by the metal), with subsequent reaction between dissolved iron oxide and the metalloids carbon, manganese, phosphoms, and silicon. During an investigation on deoxidation with silicon, data on rate of oxidation of the metal were obtained, and these combined with certain other information, are presented here in an attempt to show the fundamental principles governing the diffusion of FeO from slag to metal. Inasmuch as the aim of the investigation was not a study of diffusion, certain data are lacking and it has been necessary to make assumptions on metal viscosity and effect of temperature on diffusion. These factors are known to exert a considerable effect on diffusional processes. Overoxidation, or supersaturation of the metal with iron oxide, is not taken up in this paper. Overoxidation has little effect on diffusion except under a few special conditions, and the discussion has therefore been limited to diffusional processes. Though the results given in this paper should be regarded as preliminary, certain of the conclusions drawn are definitely applicable to any steelmaking process. Previous Study of Absorption An extensive study of absorption has been made by the Department of Chemical Engineering, Massachusetts Institute of Technology. The results are of such direct interest in the problem of slag-metal reactions that it is feIt worth while to summarize them here. The following quotations are from Lewis and hitmman.1 The reader need simply substitute "slag and metal" for "gas and liquid," with proper modifications to take into account the variations in physical properties of

Jan 1, 1929