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Report of the Americanization Society, New York Section During the past two months the Americanization Committee has held social meetings of various kinds for foreign-born women, as a preliminary to future effort along the lines of Americanization work. At the same time, a thorough investigation was made as to the work being done in New York by organizations devoted to this interest. The committee will be very glad to help any other committees being formed, in mining centers, with advice and information that has been gathered together as to classes and meetings of all kinds best suited to the needs of a foreign-born population. The work has been so extensively taken up in New York that for the present it would seem to be, for this committee, the most profitable and interesting work to continue the social meetings and to undertake the care of one or more immigrant families, to be friend and adviser in joy and sorrow, work and play, and in this way do some fundamental Americanization work. Respectfully submitted, LAURA G. BURGER, Chairman.

Jan 6, 1918

By W. E. STOKES

COMPANIES operating airplanes have had a relatively prosperous year, permitting them gradually to re- place old types of equipment. The pre-eminence of American-made planes, engines, and accessories has been challenged, either by foreign designs or by domestic patterns that have, for one reason or another. been licensed for manufacture abroad.

Jan 1, 1938

By Rossiter W. Raymond

(Read at the Wilkes-Barre Meeting, May, 1877.) THE 12th of April, 1877, witnessed the celebration, at Freiberg, Saxony, of an event profoundly important for the ancient mining industry of that district, and full of interest to mining engineers throughout the world-namely, the connection of the famous Roth-

Jan 1, 1878

By J. H. Brophy

Experimental evidence in recent years shows that nickel coated hydrogen reduced tungsten powder can be sintered to 98 pct of theoretical density at 1100°C. New data indicate that the sintering rate is the same for nickel contents ranging from 0.125 wt pct (mon-atomic layer) to 5.0 wt pct Ni coated on 0.561. tungsten powder. Nickel tungsten made by coreduction from oxides sinters in a kinetically similar way, but the rate tends to increase with higher nickel contents. The activation of sintering can be accomplished with a minimum amount of nickel if coated powder is used. Transverse rmpture strength was found to increase in specimens containing 0.25 pct Ni as density increased during the initial stage of the carrier-phase sintering process. Upon the onset of grain growth in the final stage, strength was found to decrease. With increasing nickel contents up to 4 pct, strength increased. Maximum strengths in three-point loading were observed at 73,000 psi for 0.25 pct Ni and 93,000 psi for 5 pct Ni. These were comparable to that of massive tungsten recrystal-lized in the presence of nickel and tested in the same way. The results were sensitive to the mode of powder treatment and ductility was negligible in all cases. The study and application of sintered bodies of nickel and tungsten is not new; however, the analysis of the mechanism by which small additions of nickel increase the sintering rate of tungsten promises to Contribute new information regarding sintering theory in general. Initial experimental evidence was presented by the authors showing that nickel-coated hydrogen-reduced tungsten powder may be sintered to 98 pct of theoretical density at 1100°C in 16 hr.' Simultaneously, a kinetic analysis of the process showed that it took place in two distinct stages. In the first stage, nickel evidently served as a carrier phase through which tungsten atoms could move. This stage was kinetically similar to the 'solution-precipitationo step postulated when the carrier phase is liquid3 although in the nickel tungsten case, there is no evidence of the existence of a liquid phase either in equilibrium phase relationships or in mi-crostructures at the temperatures under consideration. In the present work, new data for nickel-coated tungsten powder are offered in support of the proposed mechanism, and a more explicit study of the second stage of densification is presented. Concurrently, the sintering kinetics of these coated powders are compared to Ni-W powder made by coreduction of oxides. In this way the present work can be compared to earlier results in nickel activated sintering of tungsten in which the coreduction process was emplyed. The ultimate utility of such an activated sintering process will be determined to some extent by the mechanical and physical properties of the final prod-

Jan 1, 1962

By F. W. C. Boswell, R. L. Cunningham, H. M. Weld

Chemical analysis shows the partition of the magnesium addition between the metal and the graphite is independent of whether the graphite is in the form of nodules or flakes. The bright central spot observed in nodular graphite polished sections is shown to be due to the structure of the nodules. DURING the past few years numerous papers have been published on the production1-7 and heat treatment" of nodular cast irons. However, very little information is available on the structure of the nodules and the reason for their formation. It would appear that an explanation of the phenomenon of nodulation must await a detailed knowledge of the nature of the nodule. Disagreement exists among various workers on such a fundamental point as to whether or not a nodule actually contains a microscopically visible, nongraphitic nucleus.9-11 Work was initiated with a view to obtaining definite information on this and other aspects of the problem. Chemical Separation and Analysis Microscopic appearance had indicated9-11 that a nongraphitic core might exist in nodules (Fig. 1). It appeared that a chemical comparison of separated nodules and flake graphite would be profitable, for if the core is nongraphitic a relatively large amount of certain elements would be found in the nodules as compared to the flake. This comparison could best be made by obtaining nodules and flake from the same melt of cast iron, part of which had been in-noculated with a suitable addition reagent for the formation of nodules. A 300 lb melt was produced. About half of this was innoculated in the ladle with an Fe-Si-Mg alloy to produce nodules and cast into keel blocks; the remaining half of the melt had silicon added in an effort to compensate for the silicon in the innoculant used in the other half of the melt and this was cast into similar keel blocks. The analyses of the resultant irons are given in Table I. Nodules and flake were chemically extracted from 2 kg samples of metal as described in the appendix. While this method dissolved the iron and chromium carbides, it did not dissolve the titanium, vanadium, and molybdenum carbides. These latter carbides could have been removed by nitric acid, but in other experiments it was found that it attacked pools of metallic iron entrapped in the nodules which had, until the nitric acid treatment, been protected from the attack of the sulphuric, hydrochloric, and hydrofluoric acids. It will be shown in the next section that entrapped iron exists in nodules but not in flake. On microscopic examination many tiny irregular particles were observed to accompany both the nodules and the flake. These were identified by tabling a sample of nodules from which a sufficient fraction was obtained for X-ray analysis. A good pattern of titanium carbide was obtained indicating that this was the predominant compound present. Since it was not possible to determine how completely the carbides could be removed mechanically and since their removal from flake is much more difficult, it was decided not to resort to any separation as this would render further comparisons more difficult. The samples were ashed in a combustion

Jan 1, 1953

University of Alabama-University, Alabama, Mining and Metallurgical Society, William L Mason University of Alaska-College, Alaska, Mining Society of the University of Alaska, Doug¬las W Huber University of Arizona-Tucson, Arizona, Miners Society, S L Smith University of California-Berkeley, California, Mineral Technology Association, Jack Washburn Carnegie Institute of Technology-Pittsburgh, Pennsylvania, The Metals Club, Gerhard Derge Case Institute of Technology-Cleveland, Ohio, Case Metallurgical Engineering Society, Robert F Hehemann

Jan 1, 1958

By L. H. Sjodahl, J. B. Conway

of each particle on grain boundary movement is proportional to the area it occupies on the grain boundary and this increases with the square of the particle radius. In contrast to the impression gained by Tiwari the effectiveness of a particle, as a grain boundary obstacle, increases rapidly as the particle size increases. The counterbalancing factor in a dispersion of constant volume (e.g., those considered by Zener) is of course the rapidly diminishing number of particles, a situation which does not exist in thoria dispersed materials where the volume fraction of thoria remains constant. Although it has been assumed that the Zener approach is valid for the sake of the above argument, it is open to the following objections: 1) It assumes that the effect of many small particles along a grain boundary is cumulative and does not take into account the possibility of each small particle being overcome separately. In this case the boundary will not be able to move forward until the largest particle on the boundary has been overcome. This means that the rate controlling factor will be the rate at which a few relatively large particles are overcome with the remaining small particles playing a negligible role. 2) When the dispersion consists of very small particles the interparticle spacing is usually small compared to the grain size and most particles are located away from grain boundaries where they can play no part. The most efficient dispersion, from a purely geometrical viewpoint, is one where the particle spacing is equal to the grain size so that all the particles are on grain boundaries. Tiwari's second point that voids cannot prevent re-crystallization seems to be unfounded as several workes20-23 have shown that voids alone can retard grain boundary movement and recrystallization. To further convince Tiwari, Fig. 17 shows a void in deformed TD Nichrome holding back a recrystallization interface. Tiwari's explanation for the effect of voids on recrystallization, i.e., they reduce local strain energy and hence lower the driving force for recrystallization may be valid in itself but is irrelevant to the present case where even complete recrystallization after ONE of the more recently proposed expressions for describing stress-strain behavior is that due to Voce.1,2 This equation is usually written: s =s8 ?(s8?so)e??/?c [l] where s is the true stress, ? is true strain, cC is a constant called the characteristic strain, s8 is the asymptotic stress approached at high strain values, and so is the threshold stress at which the plastic deformation begins. While this expression has not been universally accepted as yet it does contain certain features which are consistent with experimental stress-strain behavior. In a more recent publication Hsu, Davies, and Royles3 employed the concept of generalized strain proposed by seth4 and identified a method for linearizing the stress-strain plot. In this approach the generalized strain was given by: where 7 is the generalized strain, 1 is the instantaneous length, 1, the datum length, and n is a constant called the "coefficient of strain measure". A value of n was said to exist which would yield the linearity: s —s+ S[.-(tf] [3] In recent studies of both these approaches, it has been found that while the Voce equation and the gen-

Jan 1, 1970

By K. L. Meissner

IT is well known that the so-called pure aluminum contains noticeable amounts of impurities, chiefly iron and silicon, and many investigators have studied the forms in which these impurities exist. Hanson and Archbutt1 state that iron exists in aluminum as FeAl3 and silicon as elementary Si. The solubility of FeAl3 in solid aluminum is practically zero, or at, least so small that the smallest trace of this compound is always visible as a separate constituent. Under the microscope an untrained eye can hardly distinguish, without etching, the two constituents FeAl3 and Si, but they can be differentiated by suitable etching reagents. Anderson2 also found the same results. At the same time, Merica, Waltenberg, and Freeman3 found a constituent of unknown composition. They found a small thermal effect, at 610° C., on the cooling curves of aluminum with small amounts of iron and silicon. The intensity of this effect increased with increasing silicon content, until an effect occured at 576° C. due to the eutectic between aluminum and silicon. The authors ascribed the effect at 610° C. to a compound, which they call "X constituent." It was thought to be a ternary compound containing iron and silicon and could be observed under the microscope as dark particles. Dix4 has published unetched photomicrographs of an aluminum with 0.51 per cent. iron, 0.99 per cent. silicon, and 0.32 per cent. copper that shows three constituents: the dark Si (black in the figure), the lighter FeAl3, which appears in half tone and crystallizes in needles, and a constituent of the same color as FeAl3 but of different crystal form, often resembling Chinese script. This last constituent, Dix thought, is a compound of Fe and Si; not the compound FeSi but a compound containing a relatively small amount of silicon. Therefore an alloy of aluminum, iron, and silicon (the copper should not be regarded in the following considerations) must contain the following four constituents: Al, FeAl3,

Jan 1, 1926

By Lyall Zickrick

NUMEROUS brands of virgin pig lead are used in the manufacture of lead and lead-alloy cable sheath for commercial cables. In changing from one brand of pig lead to another a readjustment of press operating conditions is nearly always necessary. This is due largely to different flow characteristics, and some brands extrude at a much lower pressure than others; for example, copper-bearing lead requires much more pressure than a low-bismuth lead under the same operating conditions. With the modern type of lead press the extrusion practice in general differs very little. Molten lead at a temperature of 375° to 400° C. (710° to 750° F.) is cast in a water-cooled cylinder and allowed to cool or solidify into a slug for six to nine minutes. . This slug is forced into the die block and around the core tube through which the cable runs. The die block is held at an optimum temperature by means of gas flames, usually 175° to 225° C. (350° to 4.35° F.). Measurements of the tem¬perature of the cable sheath taken at the time of extrusion have covered a range of 190° to 225° C. (375° to 435° F.). According to Tammann, the rate of extrusion for a metal under constant pressure through a simple orifice is almost doubled for each 10° C. rise in temperature. Because of the irregularity of die-block construction, such a great increase is not obtained in lead-cable extrusion practice. To obtain more" definite information about the effect of small per¬centages of metals usually present as impurities in virgin pig lead, a series of compression tests were made covering several different brands of lead and various alloys. Since the usual temperature of the sheath is approxi¬mately 200.° C. as it emerges from the die block, these compression tests were all performed at 200° C. Five different brands of virgin pig lead were selected, as follows: Brand A: A desilverized western lead of very high purity. This lead contains neither copper nor bismuth. Brand B: A bismuth-free southeastern Missouri lead, usually con¬taining 0.05 to 0.08 per cent copper. Brand C: A copper-free mid-western desilverized lead of very good grade, usually containing around 0.06 to 0.10 per cent bismuth.

Jan 1, 1932

By James F. Watson, Jack L. Christian

The austenite to martensite solid state phase transformation in the 300 series austenitic stainless steels was studied over the temperature range of 788° to -423°F. The primary variables controlling this transformation are temperature, plastic strain, and chernical composition. Consequently, temperature effects were studied by measurements at 78", -100°, -320°, ad -423°F; plastic strain was controlled by cold-rolling the alloys varying amounts between 40 and 80 pet; and chemical compositions were varied by using 301, 302, 304 ELC, and 310 stainless steel alloys which exhibit varying degrees of austenite stability. The martensite contents of various alloys were correlated with several mechanical properties including yield and ultimate tensile strength, elongation, and notched tensile strength (Kt = 6.3). The yield and tensite strengths generally increased with increasing martensite content, but ) no correlation was observed between elongation and martensite content. The notched/un-notched ratio (Kt = 6.3) at -423°F exhibited a maximun at about 65 pct martensite. Attempts to correlate austenite stability (as measured by the Ms temperature) with both martensite content, and incremental increases in martensite content resulting from tensile testing were generally unsuccessful because the variable of cold Work could not be separated. As a result, only a general correlation of these variables was obtained. A (111) pole figure of austenite in a cold-rolled 301 stainless steel mas prepared an attempt to further investigate directionality as exhibited by mechan- ical properties. Yield strength was correlated with this pole figure, but ultimate and notched tensile strengths were not. 1 HE literature contains a wide variety of references pertaining to the austenite to martensite solid state phase transformation in austenitic stainless steel, see references 1 through 18. Discussions of the crystallography and kinetics of this transformation are given in references 2, 7, 8, 'and 11, while further discussions of austenite stability and hardening mechanisms are given in references 1, 3, and 10. Two important processing variables involved in this reaction are the degree of cold work and the temperature at which cold work is performed, and the nature of these effects are outlined in references 4 and 12, and 15 and 16 respectively. The latter two references deal primarily with cold rolling conducted at subzero temperatures. The effect of chemical composition on the austenite to martensite reaction was treated in detail by Eichelman and Hull, and they determined an equation for the M, temperature of this type of steel which included the effects of nickel, chromium, carbon, nitrogen, silicon, and manganese.5 Previous studies have shown that austenitic stainless steels of AISI types 301, 302, 304ELC, 310, and 301-N exhibit mechanical properties between 78" and -423"~ that are microstructure dependent, and the dependence of this microstructure on chemistry, temperature, tensile stress, and prior deformation is in general accord with previous investigations which have measured the effects of each of these variables on structure, or more specifically, the austenite-to-martensite transformation.lg The behavior of fully stable alloys, such as ATST type 310, at cryogenic temperatures is

Jan 1, 1962

By M. G. Benz

The superconducting performmce of diffusion-processed Nb3Sn is influenced by its micro structure. High isotropic transverse current density may be achieved in this material by a process which forms a precipitate of ZrO, within the Nb3Sn. FOR an ideal type-I1 superconductor, little or no transport current can be carried in the mixed state; i.e., little or no transport current can be carried above the lower critical field H,,, where the field penetrates abruptly in the form of current vortices or fluxoids, even though full transition to the normal state does not occur until the upper critical field H,,.' Fortunately, nonideal type-I1 superconductors can be readily obtained and these carry large transport currents up to the upper critical field H. Both theoretical and experimental investigations have attributed this current-carrying capability for nonideal type-I1 superconductors to pinning of the fluxoid lattice by heterogeneities in the microstructure of the superconducting material. These heterogeneities may take the form of dislocations or dislocation clusters,2"5 grain boundaries: structural imperfections introduced by phase transformations; radiation damage,8"10 or precipitates.11"15 Nb3Sn formed by diffusion processing is a type-I1 superconductor. Heterogeneities are needed for high superconducting critical currents above H,,. This paper will cover: a) what the microstructure of diffusion-processed NbSn looks like; b) what changes in the microstructure take place when the system is doped with precipitates, and c) how these changes in microstructure influence the superconducting critical currents. EXPERIMENTAL Preparation of Samples. Diffusion processing was used to form the Nb3Sn. The procedure used was as follows: a) coat the surface of a niobium tape with tin; b) heat-treat this tape at a temperature above 930°C to form a layer of Nb3Sn at the Sn-Nb interface. Such a layer of NbsSn is shown in Fig. 1 The thickness of the NbsSn layer formed was controlled by the time and temperature of the heat treatment. The same general procedure was used for preparation of both undoped samples and samples doped with a precipitate. An additional step was included in the preparation of the doped samples which consisted of internal oxidation of zirconium to form ZrOn. The details of the doping process will be reported in a later paper. Sample Testing. The Nb3Sn tape samples were soldered to a copper or brass shunt. Current and voltage leads were then attached to the sample in the usual four-probe resistance measurement configuration. The sample was cooled to 42°K. In some cases it was cooled in the presence of a high magnetic field and in other cases with the field turned off. The results were the same for both cases. The samples were oriented in a configuration with field transverse to current but could be rotated such that the angle between the field vector and the wide side of the tape sample could be changed. Measurements up to 100 kG were done in a superconducting solenoid and measurements above 100 kG in a water-cooled copper magnet at the MIT National Magnet Laboratory. Once the test field was reached, the current in the sample was increased until voltage was detected across the sample. The critical current was taken as the current at which voltage was first detected in excess of background noise. In most cases this was 1 to 2 x 10~6 v for a— in.-wide sample carrying several hundred amperes with a in. separation between voltage leads and with a 10 "-ohm shunt resistance. RESULTS AND DISCUSSION Microstructure. Examination of the microstructure of the undoped Nb3Sn shows rather large-diameter (1 to 2 columnar grains growing outward from the niobium surface toward the tin surface. As the layer is made thicker by longer diffusion times, these grains grow longer. Few new grains are started. Transmission electron microscopy shows little or no second-phase material within the bulk of the Nb3Sn layer. The microstructure of a diffusion-processed NbsSn layer changes quite drastically when the system is doped so as to form a precipitate within the NbsSn layer. Instead of large-diameter columnar grains of NbaSn forming, smaller-diameter (0.5 to 1 ) equiaxed grains of Nb3Sn decorated with the precipitate form. Fig. 2 shows a transmission electron micrograph of a Nb3Sn layer doped with zirconium oxide. This layer has been etched so that one may look between the grains

Jan 1, 1969

By R. L. Fleischer

IN 1925 a criterion for localized deformation in a single crystal was derived by Taylor and Elam.' They noted that for a single active slip plane the slip plane area is constant, but the direction of the stress with respect to the slip direction changes during deformation so that, when the increase in stress due to this rotation is large enough relative to the work hardening, instability occurs. (By instability is meant continued slip on the active slip plane.) The instability criterion may be written t>dt/d?/m tan2? [1] where t and ? are resolved shear stress and strain, m the Schmid factor for resolving shear stress, and ? is the angle between the slip direction and the axis of tension. The formula points out that regardless of the atomic mechanism involved, for large flow stresses and low work hardening rates, slip on a single plane—once started—should continue. The case of alloying provides a direct possibility of adjusting the parameters in Eq. [I]: As an alloying element is added to a pure metal, the flow stress is raised and the work hardening in easy glide may be either decreased2 or unaltered.3 Since flow stresses in easy glide may rise to the kg/mm2 range for alloy crystals while easy glide slopes may be less than % kg/mm2, 2 it is possible for the inequality to be satisfied. Indeed Garstone and Honeycombe3 report that slip becomes coarser with increasing stress in the easy glide range, where the flow stress would be increasing and the slope would be essentially constant. Another case where high flow stresses and low hardening rates accompany coarse slip is that of quench hardening.4 Maddin and Cottrell report that quenched aluminum crystals display slip markings resembling those of a brass. A similar effect has been observed after irradiation.'

Jan 1, 1960

By G. V. Smith, R. F. Mehl

The decomposition of austenite in steels, because of its immerse practical importance, has been subjected to extensive study in recent years from the point of view of the mechanism of the process.1"3 Although much information has been acquired, a number of questions, principally of theoretical interest, remain. One is that of the atomic-crystallographic mechanism by which each of the various reaction products forms. The isothermal reaction products of the eutectoid decomposition in plain carbon steels (to which this study is restricted) have conveniently been classified in three general groupings—pearlite, bainite, and martensite. Of these several products, pearlite and martensite have been studied extensively, but bainite has received only scant attention. -As a result, the mechanism of the formation of pearlite and that of martensite are now fairly well understood while a great deal of uncertainty exists regarding bainite. Pearlite Pearlite forms from austenite by the precipitation of alternate platelets of cementite and ferrite, which grow into the austenite matrix.2--5 The rate of reaction may be described in terms of a rate of nucleation and a rate of growth.2.3 Mehl and Smith6 have attempted to determine the lattice relationships existing between the ferrite in pearlite and the' parent austenite. Using a pole-figure method of analysis, it was found that a limited number of ferrite orientations resulted from the decomposition of a single grain of austenite; it was not possible, however, to decide between two assumed ideal relationships, both of which appeared to account for the experimentally determined pole figure. It was definitely shown, however, that the relationship obtaining was quite different from that determined for the proeutectoid separation of ferrite.7 No attempt was made to determine the habit plane (matrix plane to which the precipitate lies parallel) of the pearlite lamellae. Belaiew8 earlier had stated that the number of directions assumed by the pearlite lamellae in a single grain of austenite was 12 (no details were given) and concluded wrong!:!, as Mehl and Smith6 have pointed out, that the lamellae delineate the {2I0} planes of ferrite.* Martensite It is known that martensite is a transiti0nal constituent of body-centered tetragonal lattice,' and that it forms suddenly

Jan 1, 1942

By R. F. Mehl, G. V. Smith

The decomposition of austenite in steels, because of its immerse practical importance, has been subjected to extensive study in recent years from the point of view of the mechanism of the process.1"3 Although much information has been acquired, a number of questions, principally of theoretical interest, remain. One is that of the atomic-crystallographic mechanism by which each of the various reaction products forms. The isothermal reaction products of the eutectoid decomposition in plain carbon steels (to which this study is restricted) have conveniently been classified in three general groupings—pearlite, bainite, and martensite. Of these several products, pearlite and martensite have been studied extensively, but bainite has received only scant attention. -As a result, the mechanism of the formation of pearlite and that of martensite are now fairly well understood while a great deal of uncertainty exists regarding bainite. Pearlite Pearlite forms from austenite by the precipitation of alternate platelets of cementite and ferrite, which grow into the austenite matrix.2--5 The rate of reaction may be described in terms of a rate of nucleation and a rate of growth.2.3 Mehl and Smith6 have attempted to determine the lattice relationships existing between the ferrite in pearlite and the' parent austenite. Using a pole-figure method of analysis, it was found that a limited number of ferrite orientations resulted from the decomposition of a single grain of austenite; it was not possible, however, to decide between two assumed ideal relationships, both of which appeared to account for the experimentally determined pole figure. It was definitely shown, however, that the relationship obtaining was quite different from that determined for the proeutectoid separation of ferrite.7 No attempt was made to determine the habit plane (matrix plane to which the precipitate lies parallel) of the pearlite lamellae. Belaiew8 earlier had stated that the number of directions assumed by the pearlite lamellae in a single grain of austenite was 12 (no details were given) and concluded wrong!:!, as Mehl and Smith6 have pointed out, that the lamellae delineate the {2I0} planes of ferrite.* Martensite It is known that martensite is a transiti0nal constituent of body-centered tetragonal lattice,' and that it forms suddenly

Jan 1, 1942

By R. D. Reiswig, J. M. Dickinson

The 0s-Ir equilibrium diagram was determined. The diagram is of the simple peritectic type, with a peritectic temperature of about 2660°C. The solid miscibility gap is narrower than previously reported, the nearly vertical solvus lines occurring at about 42 and 63 wt pct 0s. Widmanstaetten figures resembling martensite in some of the two-phase alloys account for a previously suspected transformation. In view of the long history of use of such alloys, it is surprising that the 0s-Ir equilibrium diagram has not yet been completely determined. The boundaries of the miscibility gap in this system have been estimated to be between 24 and 39 wt pct 0s and 79 and 100 wt pct OS. It has been predicted3 that alloys of ruthenium, osmium, and rhenium with the fcc platinum metals should solidify peritectically in their two-phase regions. The metallographic examination of natural osmiridium alloys has been said4 to indicate, by the presence of a martensitic structure, the existence of a transformation or a phase separation. On the basis of superconducting transition temperature measurements, a 70 at. pct 0s alloy was found5 to be two-phase prior to annealing. EXPERIMENTAL The elemental components used in this work were obtained from several suppliers as sponge. The purities, stated by each supplier to be in excess of 99.8 pct, were confirmed in each case by a semiquantitative spectrographic analysis. The pure components were premelted in an inert-gas tungsten arc furnace with the evolution of some fume, crushed, and pickled for use as alloy melting stock. Upon arc melting to the final compositions, no fume evolution was observed. Because weight losses of the order of 15 mg or less were observed in preparing alloy buttons weighing 15 g, "synthetic analyses" were used in determining the compositions of the alloys. Alloys used in this investigation contained 15, 20, 33, 40, 42, 43, 45, 50, 55, 60, 63, 64, 66, 70, 79, and 80 wt pct 0s. Heat treatments of the alloy buttons and pieces thereof were performed in a previously described6 tungsten tube furnace under pressures of 105 torr or less. It was found that homogenization treatments on the specimens were essential to remove the severe coring which was present, a 3-hr treatment above 2300°C removing all the metallographi-cally visible evidence of coring. Following the various heat treatments, the specimens were "vacuum-quenched" by melting a tungsten fuse wire and allowing them to drop onto a water-cooled copper plate. In a few instances, specimens were quenched into molten tin for comparison, but no difference was apparent. Vacuum-quenched specimens cooled from 2400°C to below a red heat in less than 60 sec. Typical equilibration times varied from 3 hr above 2300°C to 88 hr near 1100°C. The utilization of successively longer equilibration times and the approach to equilibrium from both directions indicated that the times used were sufficient to give a close approach to true equilibrium. Heat-treatment temperatures were measured with a calibrated Leeds and Northrup optical pyrometer through a 0.120-in. lateral sight hole in the tungsten heater tube. It was known that the assumption of black-body conditions inside the 1-in.-diam by 10-in.-long tube was unwarranted. The empirical results of work with thermocouples and melting

Jan 1, 1964

R. F. McELVENNY.- The next item is: "Center and other methods, vs. side charging." That was all discussed in the Carson case, and I do not know but what many of the plants are still using center charging. Some are using side charging. Did I understand that where you use the wet charge it is all side charging on the wet ore? B.-G. C. HONEYMAN.- We were on center charging for about a year with wet ore. R. F. McELVENNY.- Then you went to side charging. Has anyone any remarks to make on side versus center charging? E. A. BARNARD.- We have one side and one center, both on deep bath. It is an experimental proposition so far, but the side charge looks very favorable.. R. F. McELVENNY.- This is with calcine? E. A. BARNARD.- This is with calcine on a floating charge proposition.. They are both water jacket. The charge acts very similar on both center and side charging. The calcine spreads much the same as it does on the center charge. It will run down 75 ft. on the furnace side as well as the center. The advantage is in taking care of secondaries and things of that nature. C. BARDWELL.- Do you get any side wall protection from it? E. A. BARNARD.- No. R. F. McKENZIE.- Do you not project the side wall up to protect the skewback? E. A. BARNARD.- There is a slot projection down around the top of the bath; but the repairs on the side wall have been practically the same; that is, they, have been practically the same as they have on the center charge. A. H. RICHARDS.- A point I would like to make is that if you have any doubt about the center charging, we feel that we get better tonnages and lower fuel ratios and better slags with center charging than we ever got with side. Different kinds of charges would, of course, vary, and I could readily see that if you had the same advantage with deep bath smelting that you mention you would probably get just as good a slag and as good fuel ratios. E. A. BARNARD.- In fact, I think the calcine charge spreads more on the side than on the center charge. A. H. RICHARDS.- If you maintain a deep bath. E. A. BARNARD.- That is one thing we have to be careful about, to hold the charge back.

Jan 1, 1930

By Robert W. Hunt

PeRhaps of all the scientific economic questions which have been claiming the attention of capitalists, metallurgists, manufacturers, directors of public utilities, and the general public of America, during the last few years, none have been of more importance than the one relating to the wear and safety of steel rails. The discovery of the Bessemer process made possible modern railway development. In that growth, as has generally happened in all progressive movemeuts, all its departments have not remained equally balanced. Iron rails permitted the construction of roadways for the passage of power-driven vehicles. In time the speed and weight of such carriages proved too much for the wearing-qualities of the rails, and the limit of railway capacity seemed to have been reached. Bessemer made his invention, which was soon perfected and developed, especially as related to steel rails, and thereby increased the limit of capacity, and vastly greater railway progress was made pgssible. The Siemens-Martin open-hearth process followed, and was later supplemented by the Thomas-Q-ilchrist basic improvement. For a time the Bessemer steel rails of commerce were more than equal to all that was asked of them, but this did not long continue, as evidenced by the records in our Tansactions. Revolutionary changes occurred in the industrial world, and in no part of it more than in steel-rail manufacture. Great and powerful commercial combinations were made, whose financial resources made possible the supplanting of the then most efficient blast-furnaces, steel-works, and mills, by others of theretofore unimagined capacity. Old traditions and practices of steel-making were wiped away, and, in too many cases, output was apparently the chief aim, with the result that much dissatisfaction with the wear and safety of rails was felt by railway officials, and the general public was concerned as to its safety when traveling. The climatic conditions of the winter

Jan 1, 1913

For the year ending February, 1917 PRESIDENT L. D. RICKETTS,1 NEW YORK, N. Y. PAST PRESIDENTS BENJAMIN B. THAYER,1 NEW YORK, N. Y. WILLIAM L. SAUNDERS,2 NEW YORK, N. Y. FIRST VICE-PRESIDENT SIDNEY J. JENNINGS,1 NEW YORK, N. Y. TREASURER GEORGE C. STONE,1 NEW YORK, N. Y. SECRETARY EMERITUS ROSSITER W. RAYMOND, NEW YORK, N. Y. SECRETARY BRADLEY STOUGHTON, NEW, YORK, N. Y.

Jan 7, 1916

For the year ending February, 1917 PRESIDENT L. D. RICKETTS,1 NEW YORK, N. Y. PAST PRESIDENTS BENJAMIN B. THAYER,1 NEW YORK, N. Y. WILLIAM L. SAUNDERS,2 NEW YORK, N. Y. FIRST VICE-PRESIDENT SIDNEY J. JENNINGS1 NEW YORK, N. Y. TREASURER GEORGE C. STONE,1 NEW YORK, N. Y. SECRETARY EMERITUS ROSSITER W. RAYMOND, NEW YORK, N. Y. SECRETARY . BRADLEY STOUGHTON NEW YORK, N. Y.

Jan 5, 1916

By Josiah Royce

To evaluate chemically the sample of rock obtained by diamond drilling, it has long been recognized that the analyses of the two components of the sample, core and sludge, must be given appropriate influence in computing the average analysis of any unit of depth. The purpose of this investigation is to set forth what means are available for apportioning the effect of core and sludge on the final analysis, what variables affect the problem, and what combination of applied mathematics will closest approximate the truth under each condition as these variables proceed within their limitations. A drill hole is bored in iron ore exploration principally to test variations in rock composition with depth and is usually directed as nearly normal to the bedding of a horizon to be tested as possible. This practice has a tendency to minimize variation in composition laterally which in any event is not likely to be great. It is obvious that the opportunity for change in analysis of a particular rock is not statistically as great radially in a diamond drill hole where the distance in which such a change may occur is from 0.719 in. (EX bit) to 1.469 in. (NX bit) as there would be longitudinally even in a' run as short as 5 ft. Variations in composition of bedded or layered rocks are usually greater normal to the bedding than parallel thereto. Even in massive rocks, like porphyries, variations are functions of distance. Hence. in either case, variations along the hole are of greater effect than across it. Let us examine Fig 1 briefly to observe a cross section of a unit of depth of a typical diamond drill hole. If we neglect radial change in chemical composition, which we have seen is small compared to lengthwise variation, we can see that if core recovery was 100 pct, the core and its surrounding area, which would be recovered as cuttings (assuming 100 pct sludge recovery), must analyze the same if the sampling is perfect. This is the foundation of this paper, namely, that if core recovery is complete we may assume that the core constitutes as nearly as possible a perfect sample of the ground drilled. Now let us pass to the methods of weighting core and sludge in the final analysis and then develop each in turn. One might employ core analyses only, sludge analyses only, or apportion the influence of each by one of several methods. The most common apportionment is by direct proportion to relative theoretical vol- ' ume occupied by each component in the cylinder hollowed out by the drill, known as the Longyear formula. Another logical treatment would be to consider core and sludge in the final analysis according to the weight of each recovered. This procedure was elaborately set forth and refined by R. S. Moeh1man.l Then to be sure that we have embraced all the possibilities, we must admit that some other method or methods may be devised empirically or by mathematical maneuvers. Should a drill machine be capable of recovering 100 pct of the core in all types of rock, certainly core alone need be analyzed and no complicated mathematics are required; but although some companies have had fair success in obtaining high percentages of core in certain homogeneous ores on the Mar-quette Range, their same methods have proved disappointing on the Menominee, Cuyuna, and Gogebic Ranges where the iron formation often varies widely from flinty chert beds interstratified with soft hematite to wholly leached ore material composed of an unpredictable mixture of hematite and limonite with occasional zones of sugary, recrystallized chert. In spite of numerous mechanical improvements in core barrels, few drilling programs can count on complete core recovery. The sludge alone might be analyzed but for practical reasons this is inadvisable when core is obtained due to the numerous opportunities for contaminating or losing portions of the cuttings. The Longyear formula involves a weighting of the core and sludge in proportion to the theoretical volume of each, that is, where the volume of the core recovered is one-third of the total cubic volume of the cylindrical hole made by the drill, the analysis of the core is given one-third of the weight in the final analysis. This approach to

Jan 1, 1950