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By D. T. Peters, J. C. Bisseliches, J. W. Spretnak

The pain boundary relaxation phenomenon in high-purity copper, 0FHC copper, and a precipitation-hardenable alloy o-fCu-2 uit pct Co has been studied by internal ,friction and elastic aftereffect techniques. The data were analyzed by assuming the broad distribution of relaxation times observed obeys a lognormal distribution law which enables calculation of the relaxation strength. Grain size had no effect on relaxation strength, although the peak width increased with increasing grain size and decreasing temperature. A double grain boundary peak was produced in high-purity copper by annealing to produce secondary recrystallization. The activation energy measured for 99.999 pct Cu is 37.5 * 3 kcal per mol and 46.5 kcal per mole for OFHC copper. A small peak in the Cu-Co alloy at about 220°C at 1 cps was deduced to be the grain boundary peak. Its height increased and its activation energy decreased with increasing aging time. Except for erratic peak positions the observations were consistent with the theory of pain boundary sliding. BEFORE the mechanism of the grain boundary internal friction peak in metals can be satisfactorily explained, the detailed characteristics of the relaxation process must be documented. Of interest are the effect of grain size on relaxation strength, the high-temperature background, and peak width and position. The effect on the grain boundary peak of a precipitate at grain boundaries and of the change in distribution of the particles with aging time must also be explained by any existing or forthcoming theory. The model which has received the most attention was proposed by KG' and zener2 and proposes that the anelastic strain results from sliding of adjacent crystals at grain boundaries. The sliding is limited by the interlocking of grain corners. It follows, according to this model, that the relaxation strength is independent of and the relaxation time is proportional to the grain size. The proportionality between grain size, D, and relaxation time, t, has been confirmed by K&' and Bisseliches . In this case, the internal-friction peaks of various specimens can be superimposed by plotting Q-' as a function of the parameter vD exp H/RT where v is the frequency and H the activation energy. However, the data of Starr et 01.' and Leak6 fit better to a parameter containing D to an exponent near two. A meaningful relationship is not always deducible from the shift of peak temperature with grain size. KG7 demonstrated that a twofold variation in grain size obtained by varying the amount of cold work followed by identical recrystallization anneals will cause a peak shift that is too large to be accounted for even by an exponent of two. Since the decreased grain boundary area existing with a larger grain size is exactly compensated for by an increased sliding distance, the relaxation strength should not decrease with increasing grain size. The height of the internal-friction peak has generally been observed to be rather insensitive to grain size as long as the average grain diameter does not approach the specimen diameter. Exceptions have been reported by Koster, Bangert, and Lang8 for copper, and by Leak6 for iron. The observations for copper could well have been confused by the presence df 0.3 pct0. The possibility that the peaks may simply be increasing in width with increasing grain size and that the relaxation strength, which is proportional to the area under the peak, may remain constant was not examined. his point has been checked in this investigation using high-purity copper. Atomic mechanisms to account for damping by sliding at grain boundaries have been proposed by KG7 and Mott. KG predicts the activation energy to be identical to that for self-diffusion. Mott's "island" model suggests the activation energy is nL where n is the number of atoms in an island and L is the latent heat of fusion per atom. Other authors have felt that a value equal to the activa-

Jan 1, 1964

By D. Turnbull, J. H. Hollomon

THERE is a large body of evidence'" indicating that solidification during the liquid-solid transition is usually induced by heterogeneities present in the liquid. By dispersing liquid metals into small droplets, the impurities responsible for catalyzing solidification are isolated within a small number of these droplets. The effect of the foreign body therefore is restricted to a single drop by this technique. Thus upon cooling below the melting temperature, solidification is initiated by homogeneous nucleation in the majority of the droplets that do not contain impurities. In the case of solidification of liquid metals, the activation energy for nucleation is so great that its rate changes by orders of magnitude for a change in temperature of only several degrees centigrade.' Effectively homogeneous nucleation occurs at a critical temperature upon continuous cooling. Thus by microscopic observation of single particles during cooling, a temperature at which the rate of homogeneous nucleation becomes sensible can be determined.3 since at the temperatures at which nucleation occurs in the absence of impurities the rate of crystal growth is extremely rapid, the temperature at which the entire particle solidifies is very nearly the temperature at which the nucleation of the solidification occurs. Thus for liquids that freeze at high temperatures the onset of nucleation can be established by simply observing the temperature at which the marked heat evolution and increase in brightness of the particle occur. For liquids that freeze at lower temperatures the onset of nucleation can be determined by a rumpling and change in shape of the particle resulting from its solidification. The microscopic technique for observing the solidification of small particles has already been described." In earlier papers the nucleation of solidification of pure metals 5,6 and of alloy systems7 showing complete liquid and solid solubility have been described. In the present paper, the observations are extended to a simple eutectic system (Pb-Sn) where the possibility of the formation of two solid phases exists. Metals for the investigation were obtained from the American Smelting and Refining Co. in the form of pure lead and pure tin, 99.8 and 99.9 pct purity, respectively. An ingot of each of the pure metals was made into shot by heating the metals at a temperature about 50 °C in excess of the melting point and pouring the liquid slowly into a container of water at 15°C. Samples of the shotted pure metals were weighed out to make alloys containing 5, 10, 15, 20, 25, 30, 40, 50, 60, 70, 80, and 90 atomic pct Pb. Samples of each alloy were then melted in separate beakers. Each melt was poured through a pyrex funnel into a cylindrical mold (% in. ID). The casting solidified in 10 to 20 sec. The inside of the mold as well as the funnel through which the metal was poured were coated with graphite to eliminate adherence of the metal. Analyses were performed on some of the compositions and are given in Table I. The compositions also were checked for these samples and for those that were not analyzed by determining the spread between the liquidus and the solidus upon melting the small metal particles. These measurements agreed as well with the nominal compositions as the analyses listed above. Results The results of the supercooling experiments for the several alloys are summarized in Table II and plotted on the constitution diagram in Fig. 1. Data for the pure lead and pure tin were taken from earlier investigations. The values for the maximum supercooling of the several alloys are the average of several determinations on a number of drops of each alloy. The maximum value in any determination was within about 2 pct of the average. For the alloys containing from 20 to 60 atomic pct Sn, inclusive, two marked changes of the surface structure were observed upon cooling. At the higher temperature, after the first appearance of the solid phase it continued to grow slowly at a constant temperature and then stopped. At the lower temperature the alteration of surface structure was abrupt. For the alloys containing from 70 to 95 atomic pct Sn, inclusive, an abrupt change in surface structure was observed at a single critical temperature.

Jan 1, 1952

By W. R. Hibbard Jr., R. W. Guard, N. G. Ainslie

Evidence is presented that copper-base solid solutions of different solutes having equal grain sizes, no preferred crystal-lographic orientation, equal electron-atom ratios, and, within experimental scatter, identical initial yield strengths, need not have identical stress-strain curves at strains larger than about 0.04. The stress-strain behavior is rationalized in terms of the proposed Suzuki chemical interaction between solute atoms and extended dislocations using what is thought to be a somewhat different means of representing stress-strain data. ALTHOUGH the effect: of alloying element upon the strength characteristics of sold solutions is a subject which has received considerable attention in the past, the exact relationships between the common deformation parameters and certain common variables are not really known in some cases. As a result some of the experiments reported in the literature in which these variables are inadequately controlled lose some of their persuasion regarding underlying principles. Nonetheless, facts are known which bear pointing up: When the true stress, a, and true plastic strain, E, of tensile deformatic~n are plotted on a double logarithmic coordinate system, one may observe a straight-line relationship at strains greater than 0.02. The form of the curve in the linear region is given by a = Ken! where a represents true stress, E, true strain, and K and tn, constants. If the relationship holds, K and m define the flow characteristics of the material being tested. m and K, however, may vary with other parameters. Hollomon found that in a-brass, m is influenced by grain size. French and HibbardZ found in alloys of copper that inverse relationships existed between m and 1) the solute concentration for a given solute, 2) the 0.01 yield strength, and 3) the constantK. Lacy and Gensamer3 observed (du/d~) (= U/Em) to increase with increasing values of K in systems of alloyed ferrites (although with considerable scatter of data which may be attributed to uncontrolled grain size). Brick, Martin, and Angier* deduced in copper-base alloys a straight-line relationship (with some scatter) between the change in the Dph number due to solid-solution strengthening and the change in the Dph number due to work hardening which suggested that copper-base alloys having equal yield strengths might have identical stress-strain curves in the plastic flow regions. French and HibbardZ concluded that the yield strength of copper-base solid solutions is the proper basis for comparing the effects of solute elements. Also, Allen, Schofield, and ate' showed that, within their experimental variation, copper-base alloys of zinc, gallium, germanium, and arsenic having the same electron-atom ratios have the same true-stress true-plastic strain curves. Dorn, Pietrokowsky, and ~ietz' also found that with aluminum-base alloys the stress-strain curves in the flow regions are approximately the same if "equivalent" concentrations of alloying elements are used. Solute valence and lattice parameter distortion were the parameters used to determine equivalency. The present report describes an investigation in which an attempt was made to obtain copper-base solid-solution alloys of four solute elements having within close tolerances equal grain sizes and yield strengths, and to see if the level of yield strength does indeed define the flow curve regardless of solute type. During analysis of the data certain unexpected features of the stress-strain curves became apparent which gave rise to some speculation and are discussed at length in the paragraphs that follow. EXPERIMENTAL PROCEDURES Alloy Preparation—Using the data of French and HibbardZ as a first approximation, four different binary copper-base alioys were designed so as to have the same yield strength. In addition, other alloys were prepared in which the solute aoncentrations varied slightly from those calculated above so as to span a range of yield strengths, see Table L The yield strengths of all alloys prepared except the copper-tin alloys were subsequently found to Lie fairly close to one another. The copper used in the alloys was produced by the American Smelting and Refining Co. and was of very high purity (99.999 pct). The alloy additions and their initial purities are as follows:

Jan 1, 1960

By D. V. Edmonds, C. J. Beevers

Tensile tests were carried out on a! titanium containing 850, 1250, and 2700 ppm 0, and up to -500 ppm H. The tests were performed at -196", -78", 20°, 150°, and 300°C at a strain rate of -1.0 x 10??3 sec-1. Increasing oxygen content, increasing grain size, and decreasing test temperature resulted in enhanced embrittlement of the a titanium by the hydrogen additions. Metallographic observations showed that this can be correlated with the influence of these parameters on the introduction of cracks into the a! titanium by fracture of titanium hydride precipitates. CRAIGHEAD et al.1 reported that the hydrogen content normally found in commercial-purity a! titanium (60 to 100 ppm) was sufficient to cause a substantial lowering of the impact strength, and they attributed this embrittling effect of hydrogen to the precipitation of titanium hydride. Lenning et al.&apos; found that in commercial-purity a titanium there is an almost complete loss of impact strength at about 200 pprn H, which is approximately half the value needed to eliminate the impact strength of high-purity a titanium. They also showed that the presence of 3000 ppm hydrogen reduces the room-temperature tensile ductility of commercial-purity material to a value of the order of 10 pct; the corresponding hydrogen concentration for high-purity titanium is over 9000 ppm. It thus appears that the detrimental effect of hydrogen on the mechanical properties of commercial-purity titanium becomes evident at much lower hydrogen contents than for high-purity titanium. The main difference between the two types of a titanium might be expected to be the higher level of interstitial impurity in the commercial-purity grade. Jaffee et a1.3 studied the influence of temperature and strain rate on the hydrogen embrittlement of high-purity and commercial-purity ! titanium. In general, the behavior was the same for both materials; embrittlement was enhanced by decreasing temperature and increasing strain rate. Recent results from tests on commercial-purity a titanium containing 850 ppm O and varying amounts of hydrogen up to -500 ppm showed that the degree of embrittlement by hydrogen is intimately related to the fracture characteristics of titanium hydride precipitates.4 The present paper considers the interrelationship between the mechanical properties and micro-structural features of commercial-purity a! titanium containing 850, 1250, and 2700 ppm 0 and varying amounts of hydrogen up to -500 ppm. 1. EXPERIMENTAL PROCEDURE Three types of commercial-purity titanium supplied by IMI* were used in the investigation, and for the *Address: Witton, Birmingham 6, United Kingdom. purpose of this paper are designated Ti 115, Ti 130, and Ti 160. The principal impurity elements are given in Table I. The material was received in the form of 12.7 mm diam bars having a fully recrystallized structure. Tensile specimens with a round cross-section of 4.5 mm diam and a gage length of 15.2 mm were machined from the bars. In order to develop the same grain size (mean linear intercept of grain boundaries) in each of the three types the specimens were annealed under a dynamic vacuum of <10?5 mm Hg, Table 11. Specimen hydriding was carried out in a modified Sieverts apparatus;&apos; hydrogen was taken into solution at 450°C and after holding the specimens at this temperature for 24 hr they were furnace-cooled to room temperature at an average rate of -100 C deg per hr. By this method nominal hydrogen contents of 0, 50, 100, 250, and 500 ppm were introduced into specimens of Ti 115, Ti 130, and Ti 160 (100 ppm (wt) -0.5 at. pct). The actual hydrogen contents were calculated from the weight differences obtained by weighing the specimens before and after the hydriding treatment. Tensile tests were carried out at temperatures of -196", -78", 20°, 150°, and 300°C on a 10,000 kg In-stron machine at a nominal strain rate of -1.0 x 10-3 sec-1. Fractured specimens were sectioned in planes parallel to the tensile axis, mechanically polished to 0.25 µm grade of diamond paste, and then attack polished using a solution containing by volume 99 parts H2O, 1 part HF, and 1 part HNO3. Although the latter treatment unavoidably opened out cracks and voids visible after mechanical polishing, it did reveal the grain structure, titanium hydride morphology, and deformation twinning structure.

Jan 1, 1970

By E. J. Rapperport, E. D. Levine

The boundaries of the (a +ß) field in the Al-Li system were determined between 150°and 550°C utilizing quantitative metallography and lattice-parameter measurements. The solubility of lithium in aluminum decreases from 12.0at. pct Li at 550°C to 5.5 at. pct Li at 150°C. P Al-Li is saturated with aluminum at 45.8 at. pct Li and has this boundary value constant over the temperature range 150°to 550°C. THE solid solubility of lithium in aluminum has been determined by several investigators, 1-6 but, as shown in Fig. 1, there is little agreement among the various determinations. The earliest investiga-tions&apos;-&apos; are suspect because of the use of impure materials. Although high-purity materials were employed in more recent work,4&apos;5 the experimental techniques may have led to contamination of the specimens. Probably the best work has been that of Costas and Marshall,6 who obtained close agreement between results obtained by two independent phase-boundary techniques: electrical resistivity and mi-crohardness. No detailed studies of the solubility of aluminum in the bcc ß phase, Al-Li, have been reported. Cursory investigations1,2,6 have indicated only that the (a+ß) -p boundary lies between 40 and 50 at. pct Li and is relatively independent of temperature. The present work was undertaken in order to provide an independent check on Costas and Marshall&apos;s determination of the solubility of lithium in aluminum, to extend knowledge of this solubility limit to temperatures below 225°C, and to make an accurate determination of the solubility of aluminum in Al-Li. EXPEFUMENTAL Alloy Preparation. In view of the difficulties encountered in previous investigations of the A1-Li system, close attention was paid to the use of methods of alloy preparation and treatment that would minimize contamination. Aluminum sheet (99.99 + pct Al) was vacuum-induction melted in a beryllia crucible to remove hydrogen. Lithium (99.9 pct Li) was charged with pre-melted aluminum into a beryllia crucible, in a helium-filled drybox. The crucible was sealed in a Vycor tube and transferred from the drybox to an induction furnace. Melting of alloys was performed by induction heating in a helium atmosphere. Solidification was accomplished by means of a suction apparatus, shown in Fig. 2, in which the alloy was forced by changes of pressure into a 3/16-in. inside diam closed-end beryllia tube. This technique produced rapid solidification of a small portion of the melt, resulting in alloys with a high degree of homogeneity. Typical lithium distributions are presented in Table I. Transverse sections 1/8 in. long were cut from the alloy rods, and each section was split in half longitudinally. One half of each section was analyzed for lithium, and the opposing halves were employed for phase-boundary determinations. Lithium contents were determined by flame photometry with an accuracy of 1 pct of the amount of lithium present. Thermal Treatments. Homogenization and equilibration heat treatments were performed in electrical-resistance furnaces with temperatures controlled to ± 2OC. Calibrated chromel-alumel thermocouples were employed to measure temperature. Homogenization was performed in helium-filled l?yrex tubes for 1 hr at 565°C. The encapsulated specimens were then transferred directly to furnaces maintained at lower temperatures for equilibration. Equilibration times were 2 hr at 550°C, 8 hr at 450°C, 27 hr at 350°c, 90 hr at 250°c, and 285 hr at 150"~. These times were chosen on the basis of conditions employed by previous investigators. Alloys were quenched from the equilibration temperatures by breaking the capsules into a silicone oil bath. By performing all possible operations either in sealed capsules or in a helium-filled drybox, the specimens were given minimum exposure to the atmosphere. Quantitative Metallography. Metallography of Al-Li alloys is difficult because of the atmospheric reactivity of the ß phase. It was found possible, however, to prepare surfaces of good metallographic quality by preventing contact with moisture during preparation. Grinding through 4/0 paper was performed in the drybox. The specimens were then transferred under kerosene to the polishing wheel. Three polishing stages were employed: 25-p alundum with kerosene lubricant on billiard cloth, 1-µ diamond paste on Microcloth, and 1/4-p diamond paste on Microcloth. Between stages the samples were cleaned by rinsing in trichloroethylene and buffing

Jan 1, 1963

By R. W. Mancantelli, J. R. Woodward

IN 1947 a large low grade deposit of uranium was located in the northwest corner of Saskatchewan, in the Beaverlodge property of Eldorado Mining & Refining Ltd. Most of the values occur as thin seams and as coatings on other minerals, and the ore, which is light and friable, is not amenable to the usual gravity methods of concentration. The acid leaching process developed to retreat the company&apos;s Port Radium tailings was also impractical, as the percentage of carbonates is high and the percentage of sulphides relatively low. Construction of the mill building and installation of equipment, both scheduled for 1952, awaited selection of a satisfactory ore dressing process on or before Oct. 1, 1951. An extensive research program on dressing of ores from the Ace mine therefore was undertaken by the Mines Branch in Ottawa, a research team at the University of British Columbia, and an ore dressing group at the company&apos;s Port Hope refinery. Pilot plant operations were conducted at the Ottawa plant of Sherritt Gordon Mines Ltd. under the general direction of C. S. Parsons, consultant on metallurgy and ore dressing for Eldorado. In mid-September, having compared results of the several research programs, Dr. Parsons recommended a process employing a carbonate or basic leach. He reported that, while the chemistry of the process had been proved, engineering of the flowsheet was incomplete, and he suggested that under normal conditions further pilot plant work at the Ace mine might be profitable. However, since this would involve a delay of 12 to 18 months in bringing the property into operation, he recommended that design of the concentrator proceed immediately. Concentrator capacity was determined by two considerations: 1) the amount of ore then available and 2) the expectation that continued improvements would be made in the ore dressing process as a result of the intensive research being carried out by the Mines Branch of the Department of Mines and Technical Surveys. As these improvements could affect both the chemistry and the en-

Jan 1, 1956

By BLAllEY STEVENS

(Canal Zone Meeting, November, 1910.) I. INTRODUCTION. TEH object of this paper is to show how igneous intrusion is governed by definite mechanical laws. A distinction is made between dikes and fissures, and the various characteristics of intrusions are explained on an exact scientific basis. The modification of rock-stresses by intrusion is then considered and extensive isostress regions are recognized. Finally, a complete classification of intrusions is constructed covering all pos-sible cases. The generalizations of observed data, taken from Geikie&apos;s Textbook of Geology, conform to the laws of intrusion herein stated, and the deductions based upon these laws. No authority, however, is quoted for the type which I have named a corrugated intrusion. Its existence is a necessary deduction from the laws of intrusion, and even the entire absence of examples in technical publications cannot disprove its reality. In all probability, corrugated fissures would have remained undiscovered, but for the working of mineral veins. As intrusions of tabular form are seldom worked for commercial purposes, it has been left for theory to anticipate the discovery of corrugated dikes. In the light of these laws of intrusion and the corresponding laws of fissures,1 there seems to be a remarkable uniformity of stress-distribution in rocks at moderate depths. This renders .possible the mathematical calculation which should help us to approximate exact science in this and kindred branches of applied geology.

Jan 1, 1911

By B. C. Allen

Surface free energy and surface self-diffusion of solid molybdenum were studied in the temperature range 1600" to 2400°C using pressure-sintered bi-crystals. Comparison of groove angles formed in various surfaces perpendicular to the grain boundary indicate a maximum of 1 pct variation in surface free energy with crystallographic mientation. This anisotropy tends to decrease with increasing temperature. The surface diffusion of the bicrystals is equivalent to that of sheet with a mild (100) Preferred orientation. Anomalously low values found for bi-crystals with surface orientations of (OOl), (012), and (011) are rationalized in terms of anisotropy in surface free energy. THE effect of crystallographic orientation on surface free energy1,&apos; and surface self-diffusion3,4 has been primarily studied in fcc metals. The object of this work was to study the effect of orientation on surface diffusion and surface free energy of bcc molybdenum using pressure-sintered bicrystals. EXPERIMENTAL WORK Materials and Crystal Preparation. Arc-melted molybdenum rod was obtained commercially and electron beam zone refined at 50 cm per hr at 10- 5 torr to form single crystals about 8 cm long and 0.65 cm diam. Three crystals were prepared with axial orientations about 1 deg from [001.], [011], and. [111]. To reduce the carbon content, the crystals were annealed 2 hr in 1.4 atm flowing wet hydrogen at 2050°C. Then the oxygen content was reduced by annealing for 2 hr in -30°C dewpoint hydrogen at 2020°C. The resulting impurity analysis is given in Table I. Bicrystal Preparation. The single crystal rods were cut into transverse slices with a thin silicon carbide abrasive wheel to produce specimens about 0.6 cm long. They were mounted in epoxy and surrounded by stainless steel washers. Cutting in half was done longitudinally at various angles to known crystallographic planes containing the cylinder axis according to Fig. 1. To reduce surface deformation resulting from the cutoff wheel and thus reduce parasitic grain boundary formation on subsequent annealing, about 0.003 cm was manually ground off each cut surface with 600 grit paper. Care was taken to keep the surface flat. After removal from the mounts, one half was generally ro-tated 180 deg with respect to the other to give a po- tential symmetrical tilt grain boundary between the two halves. In the other cases when low misorienta-tion angles were desired, the crystals were not rotated. On the basis of symmetry, sufficient bicrystals were prepared to cover the entire range of misorientations for symmetrical tilt boundaries. The misorientations, +, ranged from 0 to 45, 0 to 90, and 0 to 60 deg for [001], [011], and [111] bicrystals, respectively. One [Ill] twist bicrystal was prepared from 2 single crystal discs rotated 17 deg relative to each other. Each specimen consisted of two pieces which were placed in a cylindrical tantalum can. Sharp edges were rounded and the fit was made as snug as possible to reduce subsequent deformation during bonding. The assembly and crystals generally were vicuum outgassed at 900" or 1700°C and then electron beam welded in the can at l x 10-4 torr. After being leak checked, the samples were placed in an autoclave and hydrostatically gas-pressure bonded5 in four batches under helium at 10,000 to 18,000 psi at 1650°C for 3 hr. Satisfactory bonds were obtained in many cases, and most of the crystals bonded after two exposures. The results did not appear to be affected by the various pressures used, preannealing conditions, crystal orientation, or time-pressure-temperature route taken to the final bonding condition. After bonding, the tantalum cans were selectively removed in cold concentrated HF. Measurements indicated overall deformation was under 1 pct. The bicrystals were metallographically ground and polished flat and perpendicular to the axis. Examination showed the boundaries were straight and almost free of parasitic grains caused by extraneous local deformation. Annealing. In preparation for thermal grooving, the bicrystals were cleaned and annealed by outgassing at 10-5 torr at 1900°C and heating at 2300°C under 1 atm 99.996 Ar for 0.5 hr. The crystals were held in a closed 4-deck box made of molybdenum sheet, and were heated in a Ta-1OW resistance furnace. The ar-

Jan 1, 1970

By Pradeep K. Rohatgi, David N. French, Surendar M. Jain, Clyde M. Adams

The effect of solute concentration and diffussivity on dendritic solidification of dilute binary aqueous solutions has been investigated; chlorides of sodium, potassium, lithiurn, and hydrogen were used as solutes. Dendrite spacings were measured in spherical droplets and in unidirectionally solidified specimens. In the droplets, the freezing rate is uniform over the entire specimen and remains constant from the beginning to the end of solidification. In the unidirec-tionally frozen specimens the freezing rate varies within a given volume of liquid during the solidification cycle; in addition, it varies with location uithin the specimen. Mass transport analysis has been used to calculate the supersaturation, AC, supercooling, AT, and transcerse growth velocity, V , of ice dendrites during solidijication from the dendrite spacings measured in the droplels. The ice dendrite spacing in a given solute increases lineary with the solule concentralion at a fixed freezing rate. The ualues of AC, AT, and V increase with an increase in solule concentration. The value of AT in the aqueous solution is of the same order of magnitude as in Al-Cu alloys of the same concentration; the dendrite spacings in Al-Cu alloys also increase with an increase in solute concentralion. For the jamily of aqueozrs solulions incestigaled, the dendrite spacing increases linearly with solute diffusivity at fixed values of concertralion and freezing ,rate. The values of ?C and ?T decrease, whereas the value of V increases with an increase in solule diffusicity. The measured dendrite spacings , L, can be expressed by the relatiun: L = K;(dfa/db) 4 K2C +K3D - K4 where dfs/d? is-change in fraction solid with time (freezing rate), C is solute concentration, D is solute diffusivit K1 = (14 * 4)pmin1&apos;2, K2 = (70 ± 10) pl (moles)- K3 = (24 r 3) . l05 psec (cm)-~. K. = -132 ± 10) 1. FREEZING of solutions is almost always an ehunmix-ing" process because the composition of crystallizing solid generally differs from that of the parent liquid. " Solute concentration and transport are therefore in]portant considerations in solidification of solutions. The influence of these factors on the dendritic solutions produced from metallic solutions is of practical significance since the properties of the solid aggre-gate are often strongly influenced by dendrite spacing.&apos;&apos;&apos; Alexander and Fthines3 found that dendrite spac- ings generally increase with solute concentration in several metallic systems. Michael and Bever1 have observed an increase in a aluminum dendrite spacings in dilute A1-Cu alloys. Rohatgi and AdamsJ have reported a linear increase in a aluminum dendrite spacings with increasing copper contents up to 28 pct in A1-Cu alloys. Bolling and Tiller"&apos; have theoretically examined the effects of solute concentration and diffusivity on the growth of cells and dendrites. Their analysis of dendritic growth assumes a steady-state growth and the maximum possible velocity for dendrite propagation. In addition, the quantitative analysis is restricted to dendrites of rotational symmetry and is therefore not directly applicable to the plate-shaped dendrites of ice obtained in the present study. There are practically no experimental observations of the effect of solute diffusivity on dendrite spacings in metallic systems. The reason appears to be the lack of experimental data on solute diffusion coefficients in liquid metallic solutions. The present work is concerned with the effect of solute concentration and diffusivity on dendritic solidification of dilute aqueous solutions: the solutes examined are potassium chloride. sodium chloride, lithium chloride. and hydrogen chloride. There were several reasons for selecting these solutes. The variation of dendrite spacing with freezing rate in these solutiolls is similar to that observed in A1-Cu alloys.8 These aqueous solutions have almost identical phase diagrams at dilute concentrations. Fig. 1. and the diffusion coefficients show very little variation with solute Concentration.9,10 Dendrite spacings were measured as a function of freezing rate. solute concentration. and solute diffusivity. Dendrite spacings were measured in two different types of freezing systems. droplet and unidirectional, respectively. In the droplet freezing system

Jan 1, 1970

By Cyril Stanley Smith

Jan 1, 1949

By Cyril Stanley Smith

Jan 1, 1949

By F. H. Froes, D. H. Warrington

Electron microscope evidence which indicates that TaC may precipitate at random sites in the matrix is presented. Initially the particles are almost spherical and coherent with the matrix. However, as they grow in conditions in which there are insufficient vacancies to relieve lattice strain, the particles rapidly lose coherency in two directions and continue to grow as plates with approximately the full lattice mismatch strain present perpendicular to the plane of the plate. The necessary relief of strain comes from dislocations loops which do not become visible until the later stages of aging. The rapid decrease of apparent strain to low values of appoximately 1 pct at small particle sizes arises not from a complete incoherency but from applying a model wrong for the particle shape and strain distribution. PREVIOUS work has shown that MC-type carbides may precipitate intragranularly in austenitic stainless steel on dislocations,1&apos;2 in association with stacking faults,3&apos;4 and randomly through the matrix,5-7 In investigations of the matrix precipitate by thin-foil electron microscopy, considerable lattice strain has been found to occur around the precipitating phase.7&apos;8 Attempts have been made to evaluate the amount of lattice strain by using the methods developed by Ashby and brown.9,10 Values of the linear strain, much less than the 17 pct theoretical mismatch (for TaC), have been reported; it has been suggested that this is due to either a loss of coherency1&apos; or vacancy absorption which occurs during either the initial nucleation or growth of the precipitate." This report is an extension of earlier work7 that dealt with the precipitation of TaC from an 18Cr/12Ni/ 2Ta/O.lC alloy after it had been quenched from 1300°C and aged between 600" and 840°C. In particular, the shape of the precipitate particles and the amount of strain in the matrix, due to the precipitate, have been studied. The work described here is part of a wider investigation of factors that affect carbide precipitation in austenitic stainless steel," details of which are to appear elsewhere. RESULTS The present investigation can be conveniently split into two aspects of the strain-fields surrounding the matrix particles: 1) information derived from the strain-field which indicates the shape and habit plane of the precipitate particles and 2) the magnitude and sign of the strain-field. The Shape and Habit Plane of the TaC Precipitate. In the early stages of aging twin lobes (normally black F. H. FROES, formerly at the University of Sheffield, Sheffield, England, is Staff Scientist, Colt Industries, Crucible Materials Research Center, Pittsburgh, Pa. D. H. WARRINGTON is Lecturer, Department of Metallurgy, University of Sheffield. Manuscript submitted November 1, 1968. IMD on white background, i.e., for the deviation parameter, S > 0) that indicate the strained region of the matrix define the position of the particles by bright field transmission electron microscopy. The actual particles were not detected until they were approximately 120Å diam; below this size they were too small to be imaged in the electron microscope. This meant that particle growth that had occurred before this stage had to be inferred from the matrix strain-field contrast. In all cases when diffraction effects were observed from the precipitate particles, a cube-cube orientation relationship (i.e., (llO)ppt Il<llO>matrix and {1ll }ppt {III} matrix) existed between the precipitate and the matrix. From the matrix precipitate particles lying along edge-on {111} planes (e.g., at A, Fig. I), the precipitates are seen to be plate-like with their diameter being roughly 18 times their thickness after 5000 hr at 650°C. However, the exact shape of the particles cannot be determined because of the masking effect of the strain-field contrast. If a dark-field micrograph, using a precipitate reflection, is studied, Fig. 2, a number of the projected images of the TaC particles [on the (110) foil surface] apear to have straight edges parallel to projected f111) planes. Thus, it appears that in the later stages of aging the TaC particles are plate-like with some tendency for the edges of the plate to be bounded by the matrix close-packed {ill} planes (though the general shape of the particles in the plane of the plate is circular and thus the "diameter" of the particles has a real physical significance). It should be noted that the bands of fine discrete particles observed in Figs. 1 and 2 are not the matrix precipitate discussed in this paper but are precipitates associated with extrinsic stacking faults3j4 occurring on (111) matrix planes. **£** ****** \ *x 23 Fig. 1—18/12/2~a/0.1~ alloy. Solution treated at 1300°C for 1 hr, water quenched, and aged 5000 hr at 650°C. The (112) directions shown are the traces of the e&e-on (111) planes. Foil normal [110]; operating reflection (331); bright field micrograph.

Jan 1, 1970

By Harold M. Mooney

WHEN apparent resistivity data are taken with the symmetrical Wenner 4-electrode spread, a fixed center position is used and readings are taken for values of electrode separation. Basic data consist of apparent resistivity plotted against separation of adjacent electrodes . The interpreter attempts to infer geologic structure, such as the depth to discontinuities and the nature of subsurface earth materials. An earlier paper&apos; described methods for interpreting resistivity data. All of these involve a severe assumption, namely that the earth in the region of interest consists of horizontal layers, electrically homogeneous and isotropic. The actual earth never satisfies this assumption exactly and may deviate from it so much that none of the above methods can be applied. Attempts in three directions have been made to modify the assumption so that it approaches known geologic complexity more closely. First, curves have been calculated for dipping discontinuities. Stern&apos; and Aldredge hose a few widely separated dip angles. Trudu confined his attention to small dips. Berel&apos;kovskiy and Zubanov computed gradient curves for widely separated dip angles. Unz as given the most complete solution, with a brief attempt to treat the three-layer case. Second, anisotropy can be taken into account. It seems geologically probable that layered materials have different vertical and horizontal conductivities. Cagniard, Maillet, and Pirson et up methods for finding an equivalent hypothetical isotropic medium. Standard interpretation methods can be applied to this, and the actual medium can then be deduced. Belluigi&apos;" discounts the practical importance of anisotropy; Geneslay and Rouget do not agree with his conclusion. Third, the effect of variable resistivity in a layer can be considered. Keck and Colby" examine the mathematics of an exponential increase in a surface layer. Several authors, for example, Stevenson,&apos;&apos; consider a continuous variation of resistivity with depth. The present paper deals with a linear variation of resistivity in a surface layer. Geologically, surface variations should be expected. Unconsolidated materials such as glacial drift show marked irregularities over short distances. The effects of weathering change with depth. The moisture content of material above the water table may vary from a dry sand to a saturated clay, and both of these will be changed by rainfall. Figs. 1 to 6 present apparent resistivity curves to show the effect of a variable surface layer. In all cases resistivity varies linearly with depth down to a depth of one unit. Material beneath this depth has constant resistivity. Electrode separation is plotted in depth units. Insets on each figure show the corresponding cross-sections, plotting true resistivity against depth. To illustrate, consider curve A of Fig. 1. The true resistivity of material at the surface of the earth is taken as 0.4 units. Resistivity increases with depth, reaching a value of 1.0 at 0.5 depth units and 1.6 at 1.0 depth units. Below a depth of 1.0, all the material has very low resistivity (zero, for purposes of calculation). Curve A in the main part of Fig. 1 shows how apparent resistivity varies for this case as the electrode spacing is increased. To illustrate further, curve E of Fig. 4 corresponds to true resistivity of 1.2 units at the surface, 1.0 at a depth of 0.5 units, 0.8 at a depth of 1.0, and 1.5 for all depths greater than 1.0. Apparent resistivity curves have been plotted logarithmically so that the shape of the curves becomes independent of the units, giving the curves wide validity. A certain drift-covered area, for example, shows a gradual decrease of resistivity from 230 ohm-meters at the surface to 150 ohm-meters at the bedrock surface, 275 ft down; bedrock resistivity is 800 ohm-meters. Curve E of Fig. 5 indicates that true resistivity decreases from 1.2 units at the surface to 0.8 at a depth of 1.0, then increases abruptly to a constant value of 4.0 units. Since resistivity and depth ratios are the same, this can be used to predict the field curve. For Fig. 5, multiply true resistivity and apparent resistivity by

Jan 1, 1955

By J. D. Grace

Screening by centrifugal force is a unique subject, for all vibrating screens, with but one exception, depend on gravity alone to pass material through a sizing mesh or screen cloth. Centrifugal screening is one of the greatest advancements made in screening since the development of vibrating screens. It makes possible sizing in a range previously im-practical by mechanical methods. The centrifugal screen is specially designed for fine sizing from No. 4 separations to finer than No. 100 separations. Centrifugal force provides a positive control of factors that make fine particle separation practical and effective as opposed to gravity which is a fixed force acting only in a fixed direction, The centrifugal screen provides a control of the force of material against the screening surface, and also controls the distance of departure of the material from the surface and its rate of return. The centrifugal screen has a rotating drum that becomes a natural fan blowing air outward through the material being screened and carrying float dust and fines through the apertures with it. Screening by centrifugal force is an unique subject, for with the exception of only one vibrating screen presently on the market, all vibrating screens depend upon gravity to pass material through a sizing mesh or screen cloth. This one screen that sizes by centrifugal force and utilizes gravity only to convey material along the screen, is the Symons V-Screen. The requirements of the industrial minerals industry to meet ever-increasing stringent fine sizing specifications by a practical mechanical method accentuates the weakness of gravity, a fixed force acting in a fixed direction, as a prime force for fine sizing. It was apparent a force other than gravity was needed. This led to the development of a centrifugal screen to meet modern demands for practical fine sizing operations. Gyratory-type screens in use today were originally designed in a day when a 1/4-in. separation was normally considered fine sizing and 1/8-in. separations were extremely fine sizing. Although gyratory-type screens are sometimes used for separating below 1/8-in., they are not practical screens for this purpose. They were not designed to do this type of work. A growing demand for a practical means of making positive mechanical separations in a range of from 1/8 in. downward led to the development of the vibrating screen. These screens utilize higher speeds and lower amplitudes of vibration than

Jan 1, 1967

By John A. Mathews

D R. CHARLES W. ELIOT, the great educator and philosopher-he of the five-foot book shelf-recently gave expression to a thought I had long been cherishing as a private opinion, when he said: "It is obvious that standardization has become a dangerous adversary of progress in both education and industry." This idea must have been more widespread than either of us knew, for some five years ago a writer of verse for one of the comic papers dealt with it at great length. I will quote only a few of his stanzas to illustrate his understanding of "A Standardized Existence" : A visitor at large, I stand Beneath a standard sky- Where everything is standardized To please the standard eye. Each tree is of a standard height, Each building is the same; Each man is uniform in height, And has a standard name. Above there is a standard star Beside a standard cloud, And a standard weather here prevails To please a standard crowd. With standard joy, and standard gain, The standard people live; And if there is one bit of pain T?is such as standards give.

Jan 1, 1926

By William B. Plank

FOR the third consecutive year, I have collected the data on enrolment and employment of graduates from the schools in. the United States and Canada that grant degrees in mineral technology. The data were collected in the same way as formerly; namely, by questionnaire, dated Jan. 5, 1935, and are therefore comparable with the figures reported for former years. This year, along with the names of the 48 schools in the United States and the six in Canada, are listed the courses offered and the degrees granted by each, school, together with the names of the professors in charge. The names and addresses of the four schools in Mexico are also given. The list is on pages 450-1, which should he valued as a directory by those interested in mineral technology education. So far as is known, this is the most complete and up-to-date list that has ever been published.

Jan 1, 1935

By Henry D. Hibbard, Edward L. Ford

The process of making washed metal now followed is in principle that described by Holley before this Institute in 1879, as the Krupp washing process. Sir Lowthian Bell also experimented extensively in England during 1877 and 1878 on refining iron and joined his interest with Krupp, the process being often referred to as the Bell-Krupp process. The only plant in this country now operating this process is that of the Brier Hill Steel Co., Youngstown, Ohio. This process may be considered as the sole survivor of the numerous methods conceived for refining crude iron, when the removal of phosphorus from steel-making material was the great desideratum affecting large districts. The east coast of England and all northern Germany, the homes of Bell and Krupp, respectively, could at that time make steel only from imported ores. These processes were in use or in various stages of development 40 years ago and all except the Bell-Krupp received their quietus from the development of the basic steel processes, both pneumatic and open hearth. The aim of the early refining processes was, first, the removal of the silicon, which could be done on a sand bottom, and later, the removal of the phosphorus, requiring a bottom which would give or permit a basic slag. The incidental advantages of increased yield and small fuel consumption were not generally recognized or considered; the aim of the later operations was the saving of the difference in value between low and high phosphorus irons, as well as hastening the removal of the metalloids. The chemistry of the washing process follows in part that of the puddling process, and indeed washing might be considered as the first half of puddling, since all the non-ferrous elements of the crude iron, except the carbon, are eliminated and the iron is maintained molten until cast. In puddling, the iron loses its carbon as well, and because of that, its fusion point rises above the hearth temperature; it freezes into a pasty malleable form, and is worked at once by squeezer, hammer, or rolls. Since Holley&apos;s paper, important improvements have been made from time to time in both plants and methods for washing iron, and the

Jan 1, 1916

By Robert Linton

In 1898, Joseph L. Giroux and J. A. Snedaker organized the Pilot Knob Copper Co. and began developing the Pilot Knob mine at Kim-berly, Nev., for high-grade copper ores, carrying good gold and silver values, that the formation had disclosed. At that time the nearest railroad was over 150 mi. distant. The high-grade ores did not prove to be extensive at this point—at least to the depth (240 ft.) to which the work was carried—but development of neighboring claims indicated extensive deposits of sulfide porphyry ore of concentrating grade lying at shallow depth, and deeper deposits of direct-smelting oxidized ores occurring in limestone near the limestone-porphyry contact. The properties passed through the hands of several organizations, with only meager development, until acquired in 1913 by Consolidated Coppermines Co., whose holdings now comprise 3480 acres of mining claims in the Robinson district. Since that time development has been prosecuted more actively; and for the purpose of determining the most efficient and economical method of treating the ores, considerable testing work has been done during the past three years, both in the laboratory and in actual mill operations. At first, the shipments were of ore high enough in grade to be profitable even with the long wagon hauls to the railroad and prevailing freight and smelting charges. In 1905, a blast furnace and a concentrator were erected at Kimberly. On account of insufficient ore supply, the blast furnace was never blown in and later was dismantled. The concentrator as originally equipped was operated for only about two months, both because the ore supply at that time was not sufficient and because gravity concentration did not yield satisfactory extraction. The location and arrangement of the mill further entailed high operating costs. In 1908, the Nevada Consolidated Copper Co. began to mill successfully, on a large scale, its sulfide porphyry ores, which are similar in character to the adjacent Coppermines ores. After that time, all of the ores and concentrates produced by Coppermines Company were smelted and a considerable part of the Coppermines low-grade ores concentrated at the Steptoe works of the Nevada Consolidated Co.

Jan 1, 1921

By G. A. Bole, F. W. Davis

The purpose of this paper is not to report, to the Institute, the results so far obtained in the survey, by the Bureau of Mines, of the metallurgical requirements for open-hearth refractories, but to present an outline that we hope will serve as a basis for discussion. The systematic survey that the Bureau proposes to make has not been outlined, although considerable preliminary work has been done in the way of interviewing open-hearth superintendents in several of the larger steel plants from the Chicago district eastward. At the several conferences so many of the opinions were diametrically oppoeite, when considering the same point, that it was thought that such an opportunity as the Institute offers in the way of a clearing house for moot metallurgical topics should not be overlooked. For this reason, the paper has been prepared as a basis for a discussion which, it is believed, will be of value in outlining the program of the Bureau. D. A. LYoN, Assistant Director. The importance of securing suitable refractories for metallurgical operations is daily becoming more universally recognized; particularly in the manufacture of open-hearth steel, where the life and earning power of the furnace depends largely on the judgment and care exercised in the selection of the correct quality of refractory for each of the parts. Any innovation that would make available to the steel manufacturers refractories with longer life or better suited to their needs than those now employed, and at a cost that would permit of their use, would be of enormous value to the industry at large. It is quite probable that revolutionary changes would not be necessary. Material now used in some parts of the furnace might, by more suitable preparation, be made of much greater service; in other parts, possibly entirely different material would be necessary in order to effect much improvement. Inasmuch as different parts of the furnace structure perform different functions and are subjected to different service conditions, a refractory suitable for one part will be entirely unsuited for another. For this reason, to make a comprehensive study of the refractory requirements of the processes, we should consider the furnace as built up of a number

Jan 1, 1924

By Horace Winchell

IT IS OFTEN interesting to look backward and review the world&apos;s progress in any line of human endeavor. Our pride is flattered by our achieve-ments and our imagination stimulated by the compari-son between our present productivity and prosperity and those of the relatively unenlightened days of our recent ancestors. We love to reflect on the progress of the last quarter of a century in the art of metallurgy, in. the production and consumption of the various metals and minerals which constitute the basis of our industrial and commercial supremacy. We dwell complacently on the figures which prove that we are the world&apos;s leader in nearly all departments of mining and that our production of necessary minerals over-shadows that of other countries. We recall the days when it was considered the proper thing, if not indeed indispensable, to go abroad for the rounding out of a mining engineer&apos;s education. We had no mining schools equal to those of Germany. Our schools of technology were elementary and their equipment limited. RISE OF AMERICAN MINING SCHOOLS Now we point with true American pride to institutions which, are unsurpassed in all the world in their standard of efficiency. We know full well that the American engineer is in demand in many foreign parts and when we read recent scientific treatises published abroad we are amazed to find them unfamiliar with many of the more advanced ideas which have here already come to be an everyday part of our professional knowledge and equipment. Blessed with a country of unrivalled natural resources, dwelling in a latitude whose climate is most conducive to enterprise, peopled with races of inventive and organizing genius with equality of individual opportunity, our industries have grown and expanded until our output has reached staggering proportions and the draft upon our mineral reserves is appalling. Coupled with this development has arisen a demand for engineers which has brought into our schools of technical science larger and larger numbers of young men whose imagination has been stimulated by the success of their predecessors who happened to be contemporary elements in the recent industrial expansion and prosperity.

Jan 1, 1923