Search Documents

By M. Kitada, U. Kawabe, F. Ishida, T. Doi

The relations between micros tructures and critical current density in transverse magnetic field were experimentally investigated due to each transformation of the 0 to 0' + P" phases at 700' C for superconductmainly examined using replication electron microscopy. The ß' or a precipitates were found to pin down magnetic flux lines in these alloys. The effects of precipitation upon the critical current density were discussed in relation with the size, spacing, and characler of these precipitates. HIGH magnetic field superconductors, such as Nb-Zr, Nb-Ti, and Nb-Zr-Ti alloys, have been recently put to extensive practical use as winding materials for superconducting magnets.13 The critical current density of these hard superconductors under an applied magnetic field is an important characteristic for magnet materials and is very sensitive to metallurgical structure. It is generally known that the critical current density is increased by introducing dislocations and precipitates into a superconductor; that is, dislocations and precipitates are presumed to be barriers that hinder quantized flux lines from moving.4'5 Theoretical6'7 and experimental analyses of the motion of flux lines and the interaction between flux lines and various defects have already been reported by many authors. Metallographic analysis of high magnetic field superconductors such as Nb-Zr and Nb-Ti is difficult, so that no quantitative relationship between microstruc-ture and critical current density has been established yet. In this paper, the effect of precipitation on the critical current density in magnetic field was investigated for two superconducting alloys, Nb-40Zr-10Ti and Nb-5Zr-60Ti. In these alloys the resistive critical field H, at 4.2oK was about 100 kG and the critical current density Jc at 80 kG was of the order of 104 amp per sq cm.13-l5 The superconducting properties were examined in relation to the microstructural changes due to transformation of i) the ß to ß' + ß" phases at 700°C for Nb-40Zr-10Ti alloy and ii) the ß to a + ß phases at 500°C for Nb-5Zr-6OTi alloy. The effect of size, spacing, and character of precipitates on flux line pinning was in particular examined. The microstructures were studied by means of residual resistivity, microhardness, and tensile strength measurements as well as by X-ray diffraction, optical, and replication electron microscopies. I) EXPERIMENTAL PROCEDURE Pure niobium, zirconium, and titanium, in the form of rods 0.8 cm in diam, served as raw materials. Results of chemical analyses of these rods are given in Table I. Ingots of the alloys, 0.4 cm in diam and 3 cm in length, were prepared by means of levitation melting, utilizing a copper mold in an argon-gas atmosphere. Samples from the ingot then were cold-worked by grooved mill to 0.2 cm in diam, heat-treated homogeneously (in the ß phase region) for 5 hr at 1100° in a vacuum of 1 x 106 Torr, and finally cold-drawn to 0.025 cm in diam. For heat treatments, samples were wrapped in niobium foil and sealed in an argon-gas atmosphere in fused quartz capsules. Water quenching was done after each heat treatment. Subsequently H-J, were performed at 4.2° by slowly transporting the current through the samples 4 cm long, under transverse magnetic field, until the least detectable resistive terminal voltage was observed. The resistive critical field H, was taken as the field at which 100 pv appeared at 4.2°K across a sample 3 cm in length, with a current of 5 ma. The critical temperature T, was measured by means of a conventional four-probe resistivity technique and taken as the temperature at which the sample resistance reached one-half of full restoration of the normal-state resistance with a current l ma flowing through a sample 2 cm in length. Precipitates were observed by means of optical microscopy and carbon replication electron microscopy. The etching solution consisted of 5 ml HF, 10 ml H2SO4 10 ml H2O2, and 50 ml H2O, and shadowed carbon replicas were examined in a itachi HU-11 electron microscope operated at 50 kv. X-ray diffraction photographs were taken by a 11.46-cm-diam Debye-Scher-rer camera using copper Ka radiation. The micro-hardness was measured under a load of 200 g using

Jan 1, 1969

By W. S. Reid

In March, 1938, E. P. Fleming, metallurgist for the American Smelting and Refining Co. inaugurated an investigation into the possibilities of recirculating the gases from Dwight-Lloyd sintering machines operating on lead charge, with the twofold object of concentration of the SO2 content and reduction in volume of total gas produced. The possibility of recovering a commercial grade of SO2 gas from D & L machines operating on lead charge had previously been considered by several investigators. Early History of Recirculation The Selby Smelter Commission Report, published by the Bureau of Mines in 1914, contains a chapter by A. E. Wells regarding results obtained at Selby, wherein some of the richer gas was recirculated through a hood over the pallets. Oldright and Miller of the U. S. Bureau of Mines had also made tests at Trail, B.C., and at Kellogg, Idaho. R. C. Rutherford, while at the Chihuahua, Mexico, Smelter of the American Smelting and Refining Co., in May, 1937, proposed recirculation of D & L gases to decrease the volume of gas handled by the baghouse. At none of these plants, however, was the operation commercialized. In July, 1938, Mr. Fleming, in correspondence with the Selby Plant, inquired regarding the possibility of obtaining 6 pct SO2 gas from the Selby D & L machines. At that time, the writer advised that there was slight possibility of obtaining 6 pct SO2 gas without re- circulation, but believed that it was possible with recirculation, and that experimental work toward that end should be tried at some plant where spare D & L machines were available. The foregoing statement was based on the following information then available— 1. Tests on Selby first-over machines showed 2.28 pct SO2 from first windbox and 1.03 pct SO2 from second windbox, and corresponding figures for second-over charge of 0.81 pct SO2 for first windbox and 0.08 pct SO2 for second windbox. 2. Oldright and Miller (US. Bureau of Mines) in 1932 at Bunker Hill, on 42 in. X 22 ft machines found: a. First-over charge—Maximum SO2 concentration (leaving cake) of 9.5 pct. b. First-over charge—SO2 concentration of over 8 pct from the 4 ft to the 12 ft points beyond the front dead-plate and that the concentration then dropped rapidly. c. That approximately 80 pct of the total sulphur eliminated on the second-over machines occurred during the travel of the pallets from the 1 ft to the 6 ft distances from the dead-plate. d. That approximately 94 pct of the total sulphur eliminated on the second-over machines occurred over the first windox. 3. Oldright and Miller (US. Bureau of Mines) in 1932, at Trail, on 42 in. X 50 ft machines found: a. First-over charge—SO2 varied from 2.0 pct to 5.5 pct (leaving cake) from the 12 ft to 28 ft points from the deadplate. b. That the average SO2 increased from 1.0 pct at 7 ft from dead-plate to maximum of 3.3 pct at 20 ft, then dropped to 1.5 pct at 40 ft. c. That on a 22 ft, second-over machine with an 11 in. bed the SO2 varied from 4.5 pct to 6.5 pct from the 2 ft to the 8 ft points from the dead-plate. d. That on the final roast, the SO2 concentration also varied across the pallets; that is, 2 1/2 pct at the side, increasing to 6.0 pct, 5 in. in, and to 7.0 pct at 10 in. to 20 in. in, then decreased vice versa at the opposite side. e. Concluded that most of the sulphur on a 22 ft machine was removed over the first 7 ft of the first windbox; therefore, they partitioned the first windbox so that the exit gas from the second 4 ft section was returned to the surface of the pallets over the fist 7 ft section, and during a seven-day trial the gas from the 4 ft section averaged 2.4 pct SO2, while the recirculated gas from the first 7 ft section only increased to 3.8 pct SO2. (Excess suction over the 7 ft section to prevent escape of SO2 laden gas from the 4 ft section caused dilution by air.)

Jan 1, 1950

By B. B. Rath, Hsun Hu

In wedge-shaped bicrystals of zone-refined aluminum it is observed that (111) pure tilt boundaries migrate under the driving force of their own inter-facial free energy. The boundary velocity is a power function of the driving force. The driving force exponent decreases with decreasing angle of misorien-tation. For example, at 64O°C, the exponent decreased from 4.0 for a 40 deg to 3.2 for a 16 deg tilt boundary. An evaluation of the driving force acting on the boundaries during their motion indicates that for low driv-forces, up to about 2 x l03 ergs per cu cm, the velocity is relatively independent of misorientation, whereas at higher driving forces a 40 deg tilt boundary exhibits the highest velocity. The measured activation energy for boundary migration approaches that for bulk self-diffusion at low driving forces, decreasing from 33 to 27 kcal per mole as the driving force is increased from 1 x l0 to 5 x l03 ergs per cu cm. These results are compared with current theories of grain-boundary migration. In previous experimental studies of grain boundary migration the driving force has been limited to a difference in stored energy across the boundary. This stored energy has been introduced into the crystal either by prior deformation1-3 or by grown-in lineage structure. A part of the energy stored in the deformed crystal is released by recovery either prior to or concurrently with grain boundary migration, thus introducing an uncertainty as to the magnitude of the driving force responsible for grain boundary migration. The grown-in lineage structure, though thermally stable during annealing, neither provides conditions under which different levels of energy may be stored in the imperfect crystal nor provides a control of orientation difference across the migrating boundary of a growing grain. Furthermore, because of variation in the lineage structure, it is difficult to determine accurately the energy stored in the imperfect crystal. Several investigations of grain boundary migration during normal grain growth have also suffered from difficulties in estimating the driving force because of uncertainties in the principal radii of curvature.~ In the present investigation the velocity of pure tilt boundaries in zone-refined aluminum bicrystals of selected orientation (40, 30, and 16 deg around the [Ill] tilt axis) has been measured in the absence of a dislocation density difference across the moving boundary, thus eliminating the previous experimental difficulties. The driving force for boundary migration is derived from a gradient of the total interfacial free energy of the migrating boundary in wedge-shaped bicrystals. A similar method was attempted by Bron and Machlin in a study of grain boundary migration in silver. However, they found that one of the crystals was deformed and consequently the motion of the boundary was partly due to a difference of stored energy across the boundary. The observed behavior of boundary velocities as affected by the driving force is examined in the light of the predictions of the current theories of grain boundary migration.7"10 The effect of boundary misorientation on velocity is compared with the theory of " which is based on a dislocation core model for high-angle boundaries. EXPERIMENTAL METHOD Seed-oriented bicrystals of zone-refined aluminum, 2.5 cm wide, 0.5 cm thick, and 12 cm long, containing tilt boundaries with a common (111) axis, were grown from the melt in the direction of this axis. Spectro-graphic analysis, reported earlier,'' indicated the purity of the crystals to be 99.999+pct. Three such bicrystals containing 16, 30, and 40 deg tilt boundaries were used. Wedge-shaped specimens were prepared from these bicrystals by spark cutting followed by electrolytic polishing. The angle of the wedge was usually 40 deg and the specimens were usually 0.25 cm thick. The intercrystalline boundary was located within 0.2 to 0.5 cm from the tip of the wedge. Fig. 1 shows a section of an oriented bicrystal containing an outline of a wedge-shaped specimen. The crystallographic directions shown in Fig. 1 represent the orientation of one of the crystals (the larger section of the bicrys-tal); the orientation of the other crystal differs only by rotation around the common [lil] axis. The parallel faces of the wedge always corresponded to the common (171) planes in both crystals, whereas the orientation of the side faces varied, depending on the misorientation angle. The bicrystal orientations were determined

Jan 1, 1970

By R. P. Agarwala, M. S. Anand

Using residual activity technique, the diffusion of palladium, silver, cadmium, indium, and tin in alunzinum has been studied in the temperature range of 400" to 630°C. The diffusivities (in units of square centimeters per second) have been expressed as: IMPURITY diffusion in aluminum,1-9 silverand lead5 for cases of low solid solubility of the impurity in the host metal has yielded frequency factors in the range of l0-6 to l0-9 sq cm per sec whereas the activation energy is practically half the self-diffusion activation energy value. From the observed values of frequency factor, activation energy, and entropy of activation, it has been suggested' that these solutes are not diffusing by vacancy or interstitial mechanisms but by a mechanism which should be consistent with such low values of the diffusion parameters (Do and Q). However in spite of extensive work on these types of systems, the mechanism of diffusion is still not well understood. The present investigation on the diffusion of palladium, silver, cadmium, indium, and tin in aluminum has been carried out to throw further light on the diffusion mechanism in systems, where the solid solubility is very low (except for the case of silver). The results are discussed on the basis of solid solubility and the structural changes involved owing to the presence of the solutes in aluminum solid solution. An attempt has also been made to apply the existing theories of charge5-8 and size8 difference between the solute and the solvent. EXPERIMENTAL PROCEDURE Specimens (1/2 in. diam by 3/8 in. high) were machined out of pure aluminum (99.995 wt pct) rod obtained from Johnson Mattheys. They were sealed under vacuum in quartz tubes and annealed at 620° C for several hours; the grains thus developed were sufficiently large to eliminate the possibility of diffusion along the grain boundaries. The flat ends were prepared carefully after polishing as described previously,10 Radioactive nitrates of cadmium, indium, and tin and chloride of palladium containing, respectively, cd115, 1n114, sn113, and pd103 were dissolved in distilled water and mixed with 30 pct acetone. By means of a micropipet a drop of this solution was placed on a smoothly polished and lightly etched surface of the specimen. Due care was taken to see that the solution spread uniformly on the surface of specimen without trickling down its sides. Radioactive silver was elec-trodeposited using a AgCN-KCN bath. The amount of sample deposited in all the cases was not more than 0.1 µ thick. The samples were then sealed in quartz tubes in vacuum. The cadmium samples were sealed in a purified argon atmosphere to avoid evaporation. The samples were then diffusion-annealed. The temperature of annealing varied between 400° and 630°C and was controlled to ±5°C. On heating to -400°C,the deposits of cadmium, indium, and tin, which were of the order of 0.1 p in thickness, were converted to their respective oxides. The contribution of oxygen present in the lattice of aluminum due to these oxides has been calculated and found to be less than 10 ppm in all cases. Oxide method has already been used by other workers11'12 in diffusion studies without any controversy on the issue. However, in some of these investigations, metallic deposition was also tried. The diffusivities calculated from these measurements were found to agree very well with the diffusivities obtained by using the oxide method. Thus it is assumed that the measured diffusivities represent true diffusion coefficients. Since palladous chloride decomposes at about 500°C, the deposited samples which were to be diffusion-annealed below 500°C were heated in vacuum for a very short time at 500°C to allow the decomposition of palladous chloride to palladium metal. Time taken in decomposition of nitrates to oxides and chloride to metal was negligibly small as compared to the period of the diffusion anneals. The residual activity technique13 was used to study the diffusion profiles where thin layers from the specimen surface were removed by grinding it parallel to a flat surface on a 600-grade carborundum paper. The specimen was washed, dried, and weighed, the differ -ence of the weight being the measure of the thickness of the layer removed. After each such abrasion and weighing, the total residual activity on the surface of the specimens was measured by counting 0.656, 0.94,

Jan 1, 1968

By Lawrence H. Van Vlack, Otta K. Riegger

Periclase-type oxides were examined microscopically after being exposed to siliceous liquids. The rate of grain growth was found to be inversely proportional to the grain diameter. Grain growth proceeds more rapidly at higher temperatures, but is retarded by increasing liquid contents. aMag-nesiowiistites with higher MgO contents grow less rapidly than those with higher FeO contents. The growth rate is reduced by the presence of a second solid phase. The silica-containing liquid penetrates as a film between the individual magnesiowus tite grains. This is independent of time, temperature, amount of liquid, or the MgO/ Fe0 ratio. When present, olivine and spinel-type phases can provide a solid-to-solid &apos;&apos;bridge" between magnesioustite grains. THIS paper presents the results of a study of the microstructures of periclase type oxides in the presence of a silicate liquid. The purpose was to learn more about the effect of service factors such as 1) time, 2) temperature, and 3) liquid content upon A) grain growth, and B) liquid location among the solid grains. This study was prompted by the fact that periclase refractories are known to have very little solid-to-solid contact when the phases which are present are limited to periclase and liquid. Such a micro-structure gains industrial significance because it permits fracture during service when stresses are applied at high temperatures. The details of ceramic microstructures have not received extensive attention. This is in contrast to the extensive attention given to a) the phase relationships pertaining to refractory compositions, and b) the details of the microstructures of comparable metallic materials. A brief review will be made of the pertinent phase relationships and microstructural considerations in general, as well as of refractory compositions. a) Phase Relationships. This investigation was limited to those compositions in which (Mg, Fe)O was the solid phase. MgO and FeO form a complete series of solid solutions. Pure MgO has the name of periclase. The related FeO structure is called wustite. Both have the NaC1-type structure: however, wustite possesses a cation deficiency so that the true composition is Fe<10 even in the presence of metallic iron. The phase relationships involving solid (Mg, Fe)O and a silicate liquid are shown in Fig. 1. In this case. the liquid is saturated with (Mg, Fe)o. There-fore its SiOz content is below that encountered in orthosilicate liquids. As a consequence the liquid phase specie:; are primarily the following ions: and 0-&apos; plus occasional Fe+ ions. Two features are of importance: a) the liquid contains relatively small species and b) the liquid contains large quantities of the same species as the solid. viz., Fig. 2 shows the system, FeO-SiOz, which will be used in some of the discussions that follow. This diagram is the right side, vertical section of Fig. 1. Here, as pre\iously, the composition at the FeO end of the diagram is nonstoichiometric, varying from Feo.950 when the liquid oxide is in contact with the solid iron, to about Fe 0, when the solid oxide is in equilibrium with an atmosphere of equal proportions of CO and C02 at the solidus temperature. The Fe/O ratio will be maintained in wustite in the presence of SiO,. However, the FeM/Fe++ ratio in the liquid will be lower because of the effect OIF the SiO, on the activity of the FeO. With the addition of MgO to wustite, the over-all composition (IvZg, Fe)@, has a value of x lying between 0.9 and 1.0 when the COz/CO ratio is 1.0&apos;. b) Microstructures. In general, published attention to refractory microstructures has been directed toward the phase analyses that accompany compositional variations. This is illustrated by Harvey6 in his work on silica brick and by wells7 in his work on periclase brick. In each case, a series of altered zones is encountered which provides a sequence of phase associations. If due consideration is given to reaction kinetics, such an examination reveals phases that are compatible with equilibrium studies. Admittedly, however, it is often necessary to determine more complicated polycomponent systems to account for all the phases present.8 Relatively little attention has been given to microstructural geometry in ceramic materials. Certainly less attention has been given to this aspect of ceramic microstructures than to the size, shape, and distribution of the constituent phases in metals. Burke has pointed out that the grain size of oxides follows the same growth rules as for metals, viz.,

Jan 1, 1962

By E. Miller, J. M. Eldridge, K. L. Komarek

The liquidus curve of the Mg-Pb system was accurately redetermined. The compound Mg2Pb decomposes peritectically at 538.2° ± 0.3°C to liquid and to a compound p&apos; which melts congruently at 35.0 at. pct Pb and 549.0° ± 0.3°C. The solidus curve of ß&apos; was determined. X-ray diffraction studies indicate that 4&apos; has an orthorhombic structure. Activity values of magnesium calculated from the phase diagram agree with those published in the literature. EXPERIMENTAL thermodynamic properties of binary metallic systems have to be consistent with values calculated from the phase diagram. In systems forming intermetallic compounds the shape of the liquidus curve near a compound is determined by the thermodynamic properties of the coexisting solid and liquid phases. Hauffe and Wagner&apos; neglected the temperature dependence of the chemical potentials and obtained the potential differences of the components of the liquid alloys, relative to stoichiometric liquid. Their calculations were based on the liquidus curve and on the heat of fusion of the compound, and were only valid near the congruent melting point. Steiner, Miller, and Komarek2 developed equations which account for the temperature dependence and obtained the chemical potentials of liquid Mg-Sn alloys over the entire phase diagram from the liquidus and solidus curves and from enthalpy values with the pure components as the standard states. The Mg-Pb phase diagram has been studied by several investigators whose results have been compiled and critically evaluated by Hansen.3 Although the liquidus curve was poorly defined, the general features of the diagram, i.e., one congruent melting compound, Mg2Pb, of essentially stoichiometric composition, two eutectics, and limited terminal solid solubilities, seemed to be suitable for a similar thermodynamic analysis. A careful redeter-mination of the liquidus by thermal analysis revealed, however, the existence of another compound. The liquidus curve between the two eutectics was precisely delineated and the structure and solidus curve of the new compound were investigated. The revised phase diagram was thermodynamic ally analyzed to evaluate the activity of magnesium in the liquid alloys. EXPERIMENTAL PROCEDURE The magnesium metal (Dominion Magnesium Ltd., Toronto, Canada) had a purity of 99.99+ pct; lead (American Smelting and Refining Co.) contained 99.999 pct Pb. Most experiments were carried out in graphite crucibles. Several experiments were made in high-purity alumina (Triangle R.R., Mor-ganite, Inc.) and in Armco iron crucibles to test the inertness of the graphite crucibles. Chemical analysis of magnesium and detailed description of the procedure for thermal analysis have been given previously. For the determination of the solidus curve of the compounds, specimens of initial composition Mg2Pb were equilibrated in a closed isothermal system with magnesium vapor. The source of the magnesium vapor was an alloy which had a gross composition lying in the 0&apos; + L field at the temperature of equilibration. As equilibrium was approached, the specimens lost magnesium to the two-phase reservoir thereby lowering the activity of magnesium in the specimens until activity and composition equaled that of the ß&apos;/ß&apos; + L boundary. Crucibles (1.9 cm ID by 2.2 cm OD by 4.1 cm high) and tightly fitting lids were machined from a molybdenum rod; small, shallow trays were fashioned from thin (0.005 in.) molybdenum sheet, and all the molybdenum components were degreased in hot carbon tetrachloride and then dried. The pieces were then degassed in vacuum at 950°C for about 6 hr. The two-phase alloy was placed at the bottom of the crucible and small specimens of the Mg2Pb compound, weighed on an analytical balance, were placed in two molybdenum trays above the two-phase alloy. The crucible was closed by forcing its lid on and then inserted in a titanium crucible. This crucible was evacuated, flushed twice with argon, and welded under argon. The specimens were equilibrated for about 1 week in a resistance furnace regulated by a Celectray controller, and the runs were terminated by water quenching. The specimens were again weighed and the equilibrium compositions were calculated on the basis that the weight losses were solely due to a loss of magnesium to the two-phase alloy. The structure of the B&apos; phase was investigated by the Debye-Scherrer X-ray diffraction technique. Selected ingots from thermal-analysis experiments containing about 35 at. pct Pb were re-melted, slowly cooled, and crushed in an argon-filled glovebox until the entire ingot passed through a 50-mesh sieve. The powder was thoroughly

Jan 1, 1965

By C. D. Williams, C. E. Ells

The susceptibility of quenched and aged Zr-2.5 wt pct Nb alloy to embritt2ement during irradiation has been examined for a number of solution temperatures and aging times. Material quenched from temperatures approximately 40°C below the transus has high tensile ductility, and this ductility is insensitive to aging at 500°C or to irradiation. If, however, the material is quenched from temperatures above the transus it becomes highly susceptible to loss of ductility either from aging at 500 or from irradiation. Inter granular failure is characteristic of the materials having low ductility. The distribution of the equilibrium phase is found to control the susceptibility to embrittlement by restricting 6 grain growth during heat treatment and thus influencing crack propagation. IN zirconium, as in titanium, -stabilizing alloy additions are used to obtain high strengths via quench and age heat treatments, and the Zr-2.5 pct Nb alloy has been developed1 because of its strength advantage over the Zircaloys. Early in the development of the Zr-2.5 pct Nb alloy the problem of 13 embrittlement was appreciated, and for this reason the solution temperature was chosen below the p transus.&apos; In the course of irradiation studies on quenched and aged Zr-2.5 wt pct Nb alloy it was found&apos; that irradiation introduced an important aspect of p embrittlement, riz., material quenched from the phase and aged 24 hr at 500°C was severely embrittled by moderate doses of neutron irradiation. This effect had not been studied in titanium alloys. In titanium the metallurgical features leading to 0 ernbrittlement were found to be structures with: a) coarse a platelets at the grain bondaries, b) finely dispersed a uniformly distributed throughout the (0) matrix,6 c) Widmanstatten a-13 with more than 50 pct P, d) the presence of some metastable p transformation products,3 and e) large prior -phase grain size.5 Alternatively, the presence of a uniform distribution of coarse a was conducive to high ductility and a structure largely of equiaxed a was very dctile. The detailed mechanisms of the embrittlement have not been worked out for all of these conditions, although weakness at either a-matrix boundaries or prior p grain boundaries have been prominent in the eculation. It was proposed that acicular a might act as a mild notch, and low ductility has been associated with easy fracture along its boundary.&apos; There have been two opposing suggestions for the source of the high ductility associated with equiaxed a phase. JaffeeB proposed that this a would accept a large por- tion of the oxygen, thus increasing the ductility of the matrix, whereas after study of a Zr-Nb-Cu alloy Weinstein and oltz proposed that the a phase, softer than the martensitic matrix, acted to blunt cracks formed in the matrix. In the present work we have studied the effect of neutron irradiation on the ductility, particularly the P embrittlement, of the Zr-2.5 wt pct Nb alloy. By a variation of solution temperature and aging time a variety of metallurgical conditions have been examined, and a range of resultant ductilities obtained. The ductility has been related to the material microstructure and mode of fracture. EXPERIMENTAL The alloy used in the present work came from two separate ingots fabricated into rod of 3/8 or i in. diam, Table I. For both batches the P transus temperature was approximately 890° C. Most of the heat treatments were done directly on lengths of the j} in. diam rod, after which the tensile test specimens were machined. Quenching was achieved by dropping rods from a dynamic vacuum into water, the cooling rate estimated to be 2 100°C per sec. For aging the rods were encapsulated in evacuated silica tubes. Round tensile test specimens, with gage diam and length 0.160 and 1.0 in., respectively, were used throughout and pulled at room temperature or 300°C on Instron tensile machines, at a crosshead speed of 0.05 ipm. Specimens were irradiated in the NRX and NRU reactors, in facilities described in previous publications.&apos;0 The metallurgical conditions examined have been: All tensile test specimens were machined with axes in the axial direction of the swaged rod. Although the specimen had a degree of preferred crystallo-graphic orientation with basal plane normals both parallel with and perpendicular to the tensile axis, the material was comparatively isotropic." The techniques of thin foil examination in the electron micro-

Jan 1, 1970

By W. G. SWART

THE best estimate on reserves of iron ore in the Lake Superior district is that made, in 1920, by Mr. R. C. Allen, amounting to 2,947,225,000 tons of assured and probable ore. This includes direct- shipping or so-called merchantable ores, and those capable of beneficiation by demonstrated processes. De- ducting from this the 287,000,000 tons shipped since 1920, leaves a present reserve of approximately 2,500,- 000,000 tons, or; about 50 years&apos; supply. The assured ore in this estimate is given as 1,560,500,000 tons, which checks Mr. Lake&apos;s estimate of 1,528,948,000 tons, or about 30 years&apos; supply. This is not a long time in the life of the steel industry, and it can be readily appreciated that the matter is important enough to war- rant the attention it is beginning to receive. The first question that engages our attention is, What is merchantable ore? In 1902 the average of all Mesabi shipments was 56.07 per cent iron and 4.35 per cent silica. Twenty-five years later it had dropped to 51.99 per cent iron, with silica up to 7.92 per cent. It has remained around these figures for&apos; ten or twelve years and, except for a probable rise in silica, can stay there for an indefinite period. Most of these figures I am giving cover the Mesabi Range. Including the other ranges does not change the figures greatly, and the Mesabi Range commands the situation, since it produces about 65 per cent of the total annual tonnage and, according to Mr. Allen, contains 81 per cent of the assured reserves, 50 per cent of the probable reserves, and 66 per cent of both the assured and the probable reserves.

Jan 1, 1926

By O. G. Paasche, Yuan-Shou Shen

THERE is a disagreement among the various authors about the exact manner of transformation of ZrCr2. Rostokerl and others2 stated that ZrCr2 had a C-14 (MgZn2) type of structure below 1000°C and a C-15 (MgCu2) type of structure at temperatures above 1000°C. Alisova3 and others4 reached the opposite conclusion and stated that the transformation temperature is close to the melting point of ZrCr2. A literature survey shows that various investigators3&apos;= who homogenized the specimens at a temperature higher than 1000°C have concluded that ZrCr2 had the C-15 structure at room temperature. Meanwhile, Jordan et al.4 reached similar conclusions without annealing the specimen. Other investigators1,2,6,7 who X-rayed the specimens in the as-cast condition without annealing reached different conclusions. The investigation reported herein was conducted with the aim of exploring the exact manner of transformation of ZrCr2 by various heat treatment tests. The alloys for this examination were prepared from iodide-reduced zirconium crystal bars, 99.9 pct purity, and electrolytic chromium, 99.9 pct purity. They were melted in a nonconsumable electrode arc furnace with water-cooled copper crucible in a helium atmosphere. The melting loss of each alloy was less than 1.5 pct by weight. Chemical analysis of a randomly selected specimen indicated that there was a very close agreement between calculated and analyzed compositions. Before being heat-treated each specimen was encapsulated in a vycor or quartz tube inside which an argon atmosphere was maintained at a pressure of lower than 1 atm. In determining the crystal structure of each specimen with a Debye-Scherrer camera, the standard procedure8 for X-ray quality analysis (Hanawalt method) was followed. The different series of heat treatment tests in this investigation are tabulated in Tables I and 11. The tests in Series I, specimens from 1-1 to 1-9, which were similar to Rostoker&apos;s experiment1 indicated that the transformation temperature seemed to fall between 870° and 900°C and that the crystal structure of ZrCr2 at lower temperature seemed to be of the C-14 type. However, once the compound is transformed to C-15 type, it is impossible to reverse the transformation back to the C-14 type by first heating the specimen above 900°C and then annealing it slowly below 900°C as shown in Experiments II-1 to II-3. Thus, it appears that the specimen of ZrCr2 will transform from C-14 to C-15 structure when heated above 900°C but will not transform from C-15 to C-14 when annealed slowly passing 900° C even after the extremely slow cooling process such as indicated in the experiment of Specimen II-3. As a valid transformation temperature is a temperature at which the transformation is reversible, therefore the temperature 900°C (or other temperature close to 900°C) is not the transformation temperature for ZrCr2 and the C-14 structure is not the stable structure of ZrCr2 at lower temperatures. The C-14 structure is retained at room temperature because the transformation to C-15 structure is very sluggish and the fast cooling after melting does not allow enough time for the transformation to take place. Additional energy is required to alter the metastable condition of the C-14 structure. The sluggishness of this transformation was again demonstrated through another series of experiments. Four specimens with C-14 structure were taken. Then they were annealed at 900°C but each specimen was soaked for a different period of time, Table 11. X-ray diffraction patterns of this group indicated that the C-14 structure gradually disappeared as the soaking period was lengthened. The figures listed under the column "C-14 Structure, pct" were estimated from the intensity of the d = 2.330 line of the diffraction pattern corresponding to the structure. Notice that the intensity of this line became weaker for longer soaking periods. To determine the transformation temperature of ZrCr2, specimens with C-14 structure (as-cast condition) were annealed at 1300°, 1400°, 1500°, 1550°, and 1600°C, respectively. A final specimen was first heat-treated to 1500°C in order to transform it to C-15 structure, then heat-treated at 1600°C again. From the X-ray analyses of this series of tests, Specimen Nos. III-1 to III-6, it is evident that a transition from C-15 structure at lower temperatures to the C-14 structure occurs at some temperature between 1550° and 1600°C.

Jan 1, 1969

By George Newton

When Mr. Bailey asked us to present a paper describing our new BOF shop, he requested that we avoid a presentation heavily laden with detail and statistics. Not only have we attempted to do this, but also we have prepared a number of colour slides which, hopefully, will provide you with a graphic conception of the facility. The installation of the BOF facility at the Steel Co. of Canada&apos;s Hilton Works was delayed many years because of the remarkable production growth of the oxygen open hearth furnaces in No. 3 shop. However, in mid-1960 it became strongly evident that not only would increased ingot capacity be required, but also that N0.2 shop - with vintage non-oxygen open hearth furnaces - would have to clean its stacks in the near future, or be closed down. A decision to embark on the BOF facility was made in 1967 after an extended period of in-depth planning and assessment of industry trends.

Jan 1, 1972

By Elton A., Youngberg

The Holden mine operated by the Chelan Division of the Howe Sound Co. is on the east slope of the Cascade Range in north central Washington on the south slope of Railroad Creek valley at an elevation of 3500 ft. The mine may be reached by a 40 mile boat trip from the town of Chelan which is at the southern tip of Lake Chelan, to Lucerne at the mouth of Railroad Creek and an 11 mile bus ride up Railroad Creek to Holden. All freight and concentrate is moved over this route to Chelan Falls on the Columbia River which is on the railroad four miles below the town of Chelan. The mine is now producing 2000 tons of gold, copper, and zinc ore per day which is treated in the Holden mill. Gold-copper and zinc concentrates are made, the first of which is shipped to Tacoma, Wash., and the latter to Kellogg, Idaho, for smelting.

Jan 1, 1949

By J. R. Campbell

It is the purpose of this paper to review somewhat in detail the literature on the subject that is extant, which ought to provoke considerable beneficial discussion. The value of 1 per cent. ash in blast-furnace practice was brought to the attention of the writer when he was connected with the U. S. Steel Corpn. Various groups of blast furnaces in the Pittsburgh district compiled a statement in which the average saving was about 6 c. per ton on pig iron for every 1 per cent. of ash reduction. Later, the same groups of furnaces, when materials had reached a higher price, compiled a statement in which there was a 30-cent saving on pig iron for every 4 per cent. ash reduction in the Pittsburgh district, which was equivalent to 7½ c. for every 1 per cent. ash reduction. In the Chicago district, on long freight hauls of coal and coke from the Pocahontas district, the set-up was 30 c. per ton of pig iron for every 2 per cent. ash reduction, or 15 c. for every 1 per cent. Some years ago, Ralph H. Sweetser, of the American Rolling Mill Co., presented a paper which provoked considerable comment.&apos; To get the matter again before this body a résumé of his computations is given below: Comparison of Pig Iron Costs Ash in Coal. Ash in CoKe, Fixed Carbon in Coke peR Ton OB PIG Iron peN Day, Per Cent. PeR CeNt. Coke. Per Cent. PIG IRon. Lb. Tons 6.00 8.70 89.80 1800 365.0 7.00 10.15 88.35 1860 353.2 Coke with 8.70 Per Cent. Ash Ore mixture, 4144 It). = 1.85 tons @ $5.50.............................. $10.17 Coke, 1800 lb. = 0.90 ton @ $6.50..................................... 5.85 Limestone, 1000 lb. = 0.50 ton @ $1.50................................. 0.75 Labor, $365 per day on 365 tons....................................... 1.00 Supplies and service, $250 per day...................................... 0.68 Overhead, $200 per day............................................... 0.55 Reserve for relining and depreciation................................... 1.00 Total............................................................. $20.00

Jan 1, 1931

By Elliott J. Roberts

In spite of the fact that the percentage of true colloids in an ore grind is very small,&apos; these pulps exhibit many of the properties of true colloidal suspensions. A quartz particle, appreciably finer than 10 microns in diameter, has a settling rate so low that at a dilution of 5:l or over, the capacity of a settling tank to settle such particles would be uneconomical if the individual particle rate prevailed. However, by grouping the finer particles into aggregates or flocs, the settling rate can be increased to an economical value. The colloidal chemistry of pulp thickening is thus chiefly concerned with particles having diameters of between 1 and 10 microns, although particles as coarse as 40 microns in diameter are more or less involved. These particles must be associated in such a manner that they settle at an economical rate and give a clear, supernatant liquid or overflow. The treatment of the subject may be divided into three phases, cach phase having a more or less different colloidal aspect. The first phase is the problem of getting the pulp or suspension into the proper condition of flocculence; the second is the behavior of the flocculent pulp in its initial and intermediate stages of settling; the third is the behavior of the flocs in the find period of settling or thickening stage. Let us look first into the methods and mechanism of obtaining the flocs and then consider their behavior and properties. Flocculation Colloids are merely ordinary substances reduced to a very small particle size, the upper limit being placed by definition at 1 micron (0.001 mm.) diameter by the soil chemists and at from 0.2 to 0.1 micron by various other colloid chemists. The lower limit is reached only when the particles approach molecular dimensions. The term L&apos;colloid" should not necessarily imply lack of crystallinity. Particles as small as 0.01 to 0.02 micron give perfect X-ray patterns, while with 0.005 micron particles the bands on the powder photograph are merely broaderled and diffused.2

Jan 1, 1935

By W. Klement

By quenching liquid alloys, single-phase solid solutions are obtained in the ranges 0 to 42.0 at. pct Co-Au and 0 to 15 and 75 to about 100 at. pet Co-Cu. Metastable solid solutions are also found in multiphase 42.0 to 49.0 and 69. 0 to about 100 at. pet Co-Au alloys. fie inability to achieve complete solid solubiliiies in Au-co and cu-co alloys is rationalized by a conszderation of the rate processes during cooling and solidification. Trends in lattice distortion and shifts in the compositions of the melting-point minima and critical points are discussed for the binary alloys of copper, silver, and gold with some of the transition metals of the first long period. By rapidly cooling alloys from the melt, a continuous series of metastable solid solutions can often be obtained for components of similar crystal structure which normally form eutectic systems.&apos; The eutectic system Au-Co, Fig. 1, and the peritectic system Cu-Co, Fig. 2, were investigated in order to determine the extent of the metastable solid solutions attainable with the present technique.* EXPERIMENTAL PROCEDURES Most of the alloys were prepared by melting weighed quantities of the elements (purity greater than 99.9 pct) under hydrogen in alumina crucibles by means of induction heating. After reweighing, the alloys were cast under argon into 1 mm diam wires. However, it was not possible to prepare some of the Cu-Co alloys with the requisite homogeneity (over about 1 mm3) in this way. For these alloys, reduced powders (purity greater than 99.9 pct) were mixed for 12 to 48 hr, pressed into compacts that were sintered at 1000º to 1050°C for 6 to 48 hr under hydrogen, and then furnace-cooled. Alloy charges of a few mm3, cut from the cast wires or chipped from the sintered compacts, were heated to 1550º to 1600°C in 30 sec in an alumina insert held within a graphite support and then ejected by a pulse of helium onto a lightly polished copper strip on the inner periphery of a rotating wheel. The

Jan 1, 1963

By F. G. BREYER

ZINC metallurgists continue to follow with keen interest reports of successful results from the continuous retort plants at Palmerton, Pa., and Meadowbrook. W. Va. The new process had already demonstrated its ability to reduce the lead content of the metal and to effect savings in labor and fuel. The past year has shown that the retorts themselves can he built to last long enough so that the retort cost per ton of metal is lower than in past practice with small horizontal re- torts. A number of the latest type of retorts are reported to have been in continuous operation for over 18 months without a failure and are said to be still in good condition. This seems a complete answer to the last question which remained before continuous zinc smelting could he pronounced commercially practicable.

Jan 1, 1932

By Harry L. Brown, James M. Dickinson, Philip E. Armstrong

The elastic stiffness coefficients of niobium have been rneaslrred from -150°C to over 650 C and ave in accord with the anomalous temperature dependence previously observed for Young&apos;s modulus. This unusual behavior has been shown to be the result of changes in the relative values of the shear elastic coefficients. The bulk modulus remained essentially constant over the temperature range studied, while changes in the shear coefficients produced a more nearly isotropic behavior as the temperature increased. A number of investigators have reported data on the temperature dependence of Young&apos;s modulus of polycrystalline niobium. The values cited for the room-temperature modulus ranged from 0.71 to 1.17 x 10" dynes-cm-&apos; (10.3 to 17.0 x lo6 psi) and in each case the modulus was observed to decrease with increasing temperature. In our recent study,8 both polycrystalline and single-crystal samples of essentially strain-free high-purity niobium showed a Young&apos;s modulus whose temperature dependence as well as magnitude was strongly dependent on specimen orientation. This was shown to be the result of changes in the relative magnitude of the two shear moduli, C and C&apos;. Young&apos;s modulus of polycrystalline samples with strong [110] textures and Young&apos;s modulus of single crystals whose orientations were away from the [loo] actually increased as much as 12 pct between room temperature and 1000°C. While studies on the residual effects of previous fabrication, thermal treatments, and impurity contents will probably resolve the problem of the widely differing Young&apos;s modulus values observed in polycrystalline niobium, a more complete description of the elastic properties requires an evaluation of the single-crystal elastic constants. A knowledge of these constants can provide a clearer picture of the nature of the temperature dependence, as well as establish a basis for predicting the modulus behavior of polycrystalline niobium, both in the randomly oriented condition and with various fabrication textures. The effects of residual strains on the elastic behavior then can be separated from orientation effects. EXPERIMENTAL METHOD Thin rod-resonance techniques were used to obtain both the Young&apos;s modulus and shear modulus of single-crystal specimens. The determination of these moduli as functions of the crystallographic orientation of the specimen axis provided sufficient information to permit calculation of the single-crystal elastic constants. The necessity for torsional as well as longitudinal measurements led to the design and assembly of a vertical-tube vacuum furnace incorporating probes for the simultaneous determination of both longitudinal

Jan 1, 1967

By C. R. St. John, R. W. Lindsay, Charles R. Austin

MANY of the factors influencing the creep behavior of ferrous alloys have been investigated and reported upon in the literature, including such variables as grain size, steelmaking practice, nature and stability of the microstructure, and the effect of alloying elements in steels. However, with respect to the latter, the microstructures involved are usually heterogeneous and as a result the effect of the addition is related to its effects on the various phases present in the microstructure. In the instances where this structure consists of ferrite and carbide, the behavior of the alloy addition will be related to its action in strengthening ferrite and in opposing temperature softening of ferrite as well as coalescence of carbide. The present paper is the third in a series comprising a comprehensive review of investigations on the properties of iron binary alloys. 1.2 Its purpose is to describe the results of isolating one microstructural phase, namely ferrite, and studying its creep characteristics in both the unalloyed and alloyed conditions. This should aid in establishing to a large degree the importance of ferrite in the creep behavior of steels consisting of the previously mentioned ferrite-carbide aggregate. REVIEW OF PREVIOUS WORK In view of the fact that the present investigation deals only with the effects of silicon, manganese, chromium, molybdenum, nickel, and cobalt, and to a certain extent with the effects of carbon, the review of previous literature will be confined to these elements. Practically no references were encountered dealing with the influence of these elements on the creep characteristics of the isolated phase. However, various contributions have dealt with the creep behavior of low-carbon steels alloyed with the elements under consideration. A particularly concise review in this respect is given by Bullens (Battelle)3 and much of this material has been drawn upon. The definite improvement in creep resistance and high-temperature strength conferred by molybdenum is well known. Published suggestions by the A.S.M.E.4 permit of an increase of some 2000 to 5000 lb. per sq. in. in permissible design loads over a temperature range from 650° to 1000°F. as a consequence of additions of 0.40 to o.60 per cent of this element in carbon steels of the order of o.25 to o.35 per cent carbon content. Babcock and Wilcox Tube Co.5 has published data showing the pronounced increase in stress necessary to cause one per cent creep per 10,000 hr. when additions of 0.5 or 1 per cent molybdenum are made to a plain carbon steel. These data apply to temperatures over a range from 800° to 1000°F.

Jan 1, 1945

By Gerald Edmunds

In a paper1 before this Institute in 1940, the writer reported that the surface orientation texture of zinc and cadmium differed from the texture existing within the casting, in that basal planes were parallel to the surface. He was unable to rationalize the , origin of this surface texture. It is conceivable that the surface orientation texture is influenced by the atomic force fields, which are different at the surface and within the body of metal. It was considered that a radical changc in the surface against which the metal is cast might alter the surface texture. In trial of this, zinc has been cast against a surface of molten lead, but the same result has been obtained—i.e., (001) parallel, ± 2oo, to the surface—as was obtained when it was cast against air, Transite or iron.&apos; Into the bottom of an open cast-iron tray, 7 by 14 in. by 2 in. deep, preheated to about 275°C., about 1/4 in. of lead, then 1/2 in. of zinc was poured, and the whole was covered with an electric heater to enforce freezing from the bottom upward. After a lapse of time sufficient for the zinc but not the lead to solidify, the zinc slab was raised free of the molten lead. The bottom surface of the zinc casting was coated with a thin layer of lead. A specimen of the slab 136 in. square was metallographically polished to flatten it and remove the slight film of lead, and the orientation was determined by the method used before,&apos; with the result stated above. Some determinations have also been made on other metals, as follows: Surface orientation determinations were made on the 3/8-in. diameter grip section of two 1/4-in. diameter aluminum-alloy die-cast tensile bars and on a 1/8 by 1/2-in. flat aluminum-alloy die-cast tensile bar. None of these three aluminum die castings exhibited any preferred orientation at the surface. A brass die casting (nominal composition: 60 Cu, 37.75 Zn, 1.00 Sn, 1.00 Pb, 0.10 Al, 0.15 Si) furnished by C. S. Smith was also examined. The surface showed a random orientation of alpha and beta. Summary A zinc casting solidified against a molten lead surface was found to have the same surface orientation texture, (001) parallel to the surface, as other zinc and cadmium castings. Aluminum and alpha-beta brass die castings were found to have random grain orientation textures at the surface. Acknowledgment The X-ray determinations were conducted by M. L. Fuller.

Jan 1, 1945

By J. F. Butler

The degree and type of carbide dispersion resulting from changes in the cooling rate from austenite determine the amount of carbon remaining in solution after slow cooling- from a subsequent subcritical annealing treatment. A difjusion model has been proposecl to explain the effect of carbide dispersion on the amount of carbon retained in solution and the amount of carbon in solution has been related to the degree of carbon strain aging. COTTRELL and Leak&apos; have shown that the rate of strain aging in low-carbon steels is dependent upon the amount of carbon and nitrogen in solution. Carbon can be effectively removed from solution as Fe , C by quench aging at temperatures of 200" C or lower; however, the relatively higher solubility of nitrogen at these temperatures prevents sufficient removal of nitrogen as iron nitrides. For this reason, nitrogen in solution has been considered the major cause of strain aging in rimming grades of commercial low-carbon steels. Nitrogen can be effectively removed from solution and nonaging steel produced through the addition of aluminum to form the stable nitride, AlN; however, the aluminum also causes deoxidation resulting in a killed steel. III an attempt to produce a rimmed steel which was also nonaging, Morgan and Shyne 2 investigated tile strain-aging properties of low-carban steels containing boron and found that strain aging was suplx-essed in steels containing as little as 0.007 pet B. The original purpose of this investigation was to determine the mechanism of strain-aging suppression by boron. Early in the study, the function of boron was discovered to be much like that of aluminum in that the formation of a nitride, BN, removed the nitrogen from solution.&apos; However, during the course of the investigation, an effect of heat treatment upon the strain-aging properties and the amount of carbon in solution was found which could be related to the microstructure of the steel. The purpose of this paper is to examine these relationships among heat treatment, microstructure, carbon in solution, and strain aging in steels where nondetectable amounts of nitrogen are present in solution. EXPERIMENTAL PROCEDURE The investigation was carried out using 0.030 in. by 12. in wire specimens which had been cold swaged and drawn to the final diameter. Table I lists the compositions of the low-carbon steels used. Those with &apos;V&apos; designations were vacuum melted and cast; the alloys designated &apos;AISI&apos; were prepared from high-purity iron obtained from the American

Jan 1, 1962

By H. C. Burghard, F. R. Brotzen

Experiments were performed to determine the effects of mechanical vibrations on ordering in CuAu. The effects of 5, 10, and 15 kc vibrations were investigated over a range of ordering temperatures of 280° to 360°C. A decrease in the rate of ordering was observed at all frequencies and ordering temperatures. THE effects of mechanical vibrations on certain solid-state reactions have been noted by several investigators. In 1930, Mahoux&apos; reported a substantial increase in the case depth of a steel sample nitrided under the influence of mechanical vibrations over that of an unvibrated sample. Similarly, an accelerating effect of vibrations on the diffusion of chromium in steel was noted. Guillet, in the same year, confirmed the results of Mahoux. More recently Schenck and schmidtmann3 studied the effects of low-frequency, alternating loads on the diffusion of carbon in steel. The results of their experiments, conducted at 500 cycles per min and large strain, (0.2 pct), indicated an acceleration of the process. These workers also reported an increase in case depth of a sample vibrated at 430 kc during a carburizing treatment. Effects of high-frequency vibrations on solid-state reactions have also been investigated by Gudzoff and Gavze6 and Fairbanks and Dewez.7 The former reported an acceleration of the age-hardening process in aluminum alloys. The latter noted changes in the structure and mechanical properties of steel annealed under the influence of vibrations. On the other hand, Altenburg4 and Ehringer5 have given some theoretical reasons and experimental evidence as to the improbability of any effect of vibrations on diffusion. Comprehensive surveys of the effects of vibrations on metallurgical reactions are given by Hiedemann&apos; and Bergmann.9 None of the above investigations led to a physical description of the effects of mechanical vibrations on solid-state reactions. In the present study an attempt was made to gain an understanding of the influence of vibrations on nucleation and diffusion. The order-disorder transformation was chosen for investigation since its mechanism has been extensively discussed in the literature. Moreover, the progress of the transformation may be easily followed by observation of electrical resistance changes. EXPERIMENTAL PROCEDURE The specimens used for the experimental observation of ordering rates were prepared from 50 at. pct Cu-Au alloy wire.* Each specimen consisted of a 0.02-in. diam wire approximately 10 in. long to which two copper potential leads were silver soldered, 2 1/2 in. apart, near one end of the wire. A 1/8-in. diam by 1/4-in. long steel stub was silver soldered

Jan 1, 1960