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A. 11. P. WYNNE, San Jose de Gracia, Sinaloa, Mex. (communication to the Secretary): In the comparative tests reported by Mr. Tays, the stamp-batteries were provided with various styles and mesh-sizes of screens, while the Bryan mill was run throughout with a slot-punched Russia iron, equivalent to a 30mesh wire screen. This is a somewhat onc-sided method of comparison. Should not the Bryan mill also have had the chance to show what it could do with other screens? I cannot wholly approve Mr. Tays's proposed improvements upon the Bryan mill. In my judgment, the inventor has given us a machine well-adapted to the work for which it was designed, namely, the crushing and amalgamation of soft ores. In my judgment, the principal cause of failure to amalgamate satisfactorily in the Bryan mill is not, as Mr. Tays thinks, the shallowness of the basin, or the lack of copper plate about the inside periphery of the mill, but rather the want of experience and practical knowledge on the part of most amalgamators. Having had experience, not with one Bryan mill only, but with several, and taking into consideration the experience of other capable mill-men, I am positive that the proper way to save gold ill this mill is: first, to use enough quicksilver to keep the amalgam liquid, thus allowing it to accumulate in the annular space between the dies and the rims of the basin; secondly, to replace the old dies, when a little more than half worn-out, with new ones; and, thirdly, to clean up the mill at least every two weeks—or oftener, if rich ore is being treated. It is evident that Mr. Tap endeavored to amalgamate quite " hard," as is done in a stamp-battery, in which case the swash of the pulp would scour off portions of the hard amalgam adhering to the inside plates and dash it through the screens— after which a considerable amount would be lost through floating on, or in, the pulp in its journey over the outside plates, and

Jan 1, 1900

By Hans-Jürgen Schaller, Horst A. Brodowsky

Activity coefficients of boron in palladium were determined at concentrations up to PdB0.23 by reducing B2O3 between 870" and 1050°C in a controlled H2-H2stream and measuring the resulting weight gain. The deviations from ideal behavior closely resemble those of the system Pd-H and are interpreted in terms of three principles: 1) The solute atoms occupy octahedral interstitial positions. 2) They donate their valence electrons to the 4 d and 5s bands of palladium, raising its Fermi energy. 3) The lattice strain energy is lower for two nearesl -neighbor interstitial particles than for two farther separate ones. SOLID solutions of hydrogen in palladium are a useful subject for studying thermodynamic aspects of the formation of alloys and of nonstoichiometric systems.1-3 The activity of hydrogen is readily measurable to a high degree of accuracy,4'5 even at low temperatures where the deviations from ideal behavior are more pronounced, and its simple structure facilitates an interpretation of these deviations in terms of a detailed model. Two effects are discussed to account for the non-ideal properties:3 An "electronic" effect, connected with the rise of the Fermi energy, as electrons of the interstitial hydrogen atoms enter the electron gas of the metal, and an "elastic" effect, due to an interaction of the regions of strain around each interstitial atom. The electronic effect is based on the idea that the lowest energy levels of the dissolved hydrogen atoms are higher than the Fermi energy, so that the electron will not occupy a localized state but enter into the electron band of the metal.6 The elastic effect is based on the observation that dissolved hydrogen distorts and expands the palladium lattice. The hypothesis is put forward that the elastic strain energy is lower for two adjacent dilatational centers than for two separate ones; i.e., they attract each other. The resulting pair interaction can be used to calculate an elastic contribution to the thermodynamic excess functions by means of one of the statistical methods. This model permitted a detailed description of the solution properties of hydrogen in palladium3 and in palladium alloys.798 An extension of the approach to describe the excess functions of substitutional palladium alloys is possible.9 In order to further test and refine the model, an investigation of other interstitial alloys was started. Palladium dissolves considerable amounts of boron in homogeneous solid solution.10 The palladium lattice expands linearly up to nB = 0.23 (nB = B/Pd atomic ratio), the highest concentration studied." The expan- sion, extrapolated for 1 mole of interstitial per mole of palladium, is 17 pct of the lattice constant of pure palladium vs 5.7 pct in the case of hydrogen.12 The fact that the lattice expands rather than contracts is a strong indication that interstitial positions are occupied. According to neutron diffraction experiments, hydrogen occupies the octahedral sites of the fcc lattice.13 Unfortunately, this direct evidence is not available for the Pb-B system, mainly because of the high-reaction cross section of boron with thermal neutrons. However, by way of analogy and on the grounds of the rather close similarities between the two systems to be reported here, it seems safe to attribute octahedral positions to the dissolved boron, too. At higher boron contents, compounds of stoichiomet-ric compositions are reported such as Pd3B, which has the structure of cementite,14 so that a close structural relationship seems to exist with the system r Fe-C. In their study of hydrogen absorption in Pb-B alloys, Sieverts and Briining noted that alloys with an atomic ratio of about nB = 0.16 are no longer homogeneous15 This observation was confirmed in an extensive X-ray investigation.11,16 The phase boundaries of two miscibility gaps were established. One two-phase region was stable below a transition temperature of about 315°C and extended from nB = 0.015 to 0.178. The other one extended from nB = 0.021 to 0.114 slightly above the transition temperature and had an apex at nB = 0.065 and 410°C. All phases involved have the fcc structure of pure palladium with lattice expansions proportional to their boron contents. The occurrence of miscibility gaps, i.e., the coexistence of dilute and concentrated phases, points to an energy of attraction between the dissolved particles, in the Pb-B system as well as in the Pd-H system. The filling up of the electron bands seems to be analogous, too, in the two systems, as indicated by the hydrogen absorption capacit15,17,18 and by the suscepti bility of Pd-B alloys.l8 In both types of experiments, boron acts as an electron donor. A chemical method was used to measure the activity of boron in palladium. Boron trioxide was reduced in a moist hydrogen stream: B2O3 + 3H2 = 2B + 3H3O [l] At known activities or partial pressures of boron trioxide, hydrogen, and water, the activity of boron could be calculated from the law of mass action. The equilibrium concentration of boron corresponding to this activity was determined as the weight gain of the sample. EXPERIMENTAL The samples consisted of small pieces of foil of 0.1 mm thickness and about 100 mg weight. The palladium was supplied by DEGUSSA, Germany, and stated to be

Jan 1, 1970

By H. Doi, Y. Fujiwara, K. Miyake

In order to throw light on the mechanism of plastic deformation of WC-Co alloys, compressive tests of WC-(7 to 43) vol pct Co alloys have been carried out at room temperature, and stress-micro strain relation has been investigated in detail. The analysis of the factors affecting the yield stresses reveals that the yield stresses can be predicted by modified Oro-wan's theory if one properly estimates the planar in-terfiarticle spacings. Conzpressive straining of some of the alloys by 0.066 to 0.17pct increases the decrements by a factor of as much as 3.4 to 14, whereas the corresponding increase in the electrical resistivities is less than 10 pct. The analysis of the decrement data in terms of -Gramto and Lücke theory shows that the marked increase is attributed to increased dislocation darnping itt the binder (cobalt) phase. By cornbilling the decrement data and the conzjwession duta, one obtains the relation between flow stress in shear (?t) and increase in dislocation density (p): At = const . v6 . This is interHeted to mean that the mechanism of strain hardening of CirC-Co alloys is essentially sarne as the one for dispersion strengthened alloys. The possible effect of bridge formations between the carbide particles has also been examined. OWING to the combination of hardness, strength, and other physical and chemical properties, WC-Co alloys have opened the way for unique fields of applications, the recent ones being, for instance, anvils for super-high-pressure generation apparatuses. In such applications, the alloys are frequently subjected to very high compressive stresses: these stresses may cause the alloys to deform plastically and eventually to fail. However, much remains obscure regarding the nature of the plasticity of the alloys. Evidently, the alloys owe their high strength to the hard carbide particles which frequently occupy as much as 80 to 90 pct in volume fraction, whereas the ductility required for practical applications is provided by the small amount of the binder phase between the carbide particles. When the volume fraction of the carbide phase is not very large, deformation behavior of the alloys may be described by some of the current dispersion strengthening theories. However, greatly increasing the carbide phase is thought to lead to some carbide skeleton structure or bridge formations owing to the increased chances for direct contacts between the carbide particles;1,2 this may appreciably affect the plasticity of the alloys. Regarding the effect of formation of the carbide skeleton structure, it is interesting to note the work by Ivensen et al.3 on compression tests of the alloys containing somewhat large carbide particles; they observe extensive generation of slip bands in the carbide particles after application of some preliminary compressive stresses. They interpret the results in terms of plastic deformatiot: of the carbide particles which are supposed to have formed a skeleton structure; the binder phase plays only a passive role, at least in the early stages of the deformation. That carbide crystals exhibit microplasticity at room temperature is apparent from the work of Takahashi and Freise4 and French and Thomas5 on indentation of WC single crystals. On the other hand, Dawihl and coworkers6-10 maintain that even when volume fraction of the carbide phase is very large (for instance, more than 90 pet), a very thin binder layer generally exists between the carbide particles. They interpret the results of the extensive mechanical tests in terms of the plasticity of such a layer. Gurland and Bardzil11 point out that decrease in ductility of the alloys with increase in the carbide phase is caused by the effect of plastic constraint exerted by the dispersed carbide particles. Drucker12 further develops this concept from a continuum-mechanics approach on an assumption that a continuous thin binder layer separates the carbide particles. A common feature of the studies reported so far on the plasticity of the alloys is that the information deduced is invariably qualitative in nature. Thus, very few systematic experiments for obtaining reliable and sufficiently detailed stress-strain curves of the alloys varying widely in the microstructural features have been carried out. In particular, it may be of special interest to investigate in detail the early stages of the plastic deformation of the alloys in order to shed light on the strengthening mechanism. However, such work appears to be extremely rare. Doi et al.13 recently reported a first brief account of the results of some quantitative analysis of the plasticity of the alloys in terms of dislocation theory. Their experiment was rather limited in the composition range covered (volume fraction occupied by the carbide phase: 79 to 83 pct), and thus they could not necessarily elucidate the controlling mechanism of plastic deformation of the alloys of a more general composition range. Consequently, in the present investigation, deformation behavior and some other physical properties of the alloys were investigated and discussed in more detail over a much wider composition range. SPECIMEN PREPARATION WC-Co alloys used in this experiment were prepared in cylindrical or rectangular form by sintering in vacuo compressed mixtures of tungsten carbide and cobalt

Jan 1, 1970

By K. R. Carson, J. Weertman

An attempt is made to ascertain the mechanism of tangle and cell formation and its dependence upon dislocation-interstitial carbon interactions. The strain-hardening behavior of single crystals of 3.25 pct Si-Fe was determined at 300° and 425°K and under conditions of both continuous and interrupted tensile strain. Significantly enhanced hardening was observed in crystals deformed at the elevated temperature, and it was further accentuated by interrupted straining. Transmission electron microscopy was used to study the resultant dislocation structures. Strain aging was found to aid tangle and cell formation at 425°K, but at both temperatures embryo tangles formed solely from primary glide dislocations, presumably by a process involving cross slip and "mushrooming". IN the course of plastic deformation all bcc metals and alloys develop a dislocation structure characterized by loose-knit groups of tangled dislocations. With increasing strain the tangles become more tightly knit and grow larger; finally a three-dimensional cellular substructure is formed:1 This process has been observed with the transmission electron microscope.'-l7 However, most investigations were confined to the study of nearly pure polycrystalline metals at relatively low temperatures. At intermediate temperatures, 0.17 to 0.14 Tm where T, is the melting temperature in degrees absolute, the mobility of interstitial impurities such as carbon is high enough to permit migration to nearby glide dislocations but is still low enough so that a significant drag force is exerted.18,19 it is also in this temperature range that a hump occurs in the curve of work-hardening rate vs temperature for iron. Analogous plots for tantalum" and columbiumzo show a definite upward trend in the work-hardening rate. Keh and Weissman1 have pointed out that this behavior may be explained solely on the basis of changes in the dislocation configuration: at low temperatures the dislocations tend to be relatively straight and uniformly distributed, but at intermediate temperatures tightly knit tangles and cellular substructure appear. The interference of these tangles with glide dislocations causes the observed increase in the work-hardening rate. This explanation appears reasonable, yet one might ask what factors cause tangle formation to be so favorable at intermediate temperatures. It seens likely that the strong dislocation-interstitial interactions which are known to occur in this temperature range are at least partly responsible," with the magnitude of the effect being proportional to the interstitial concentration. The purpose of the present work is to study the relationship between tangle formation and strain hardening in a bcc metal in the temperature range 0.17 to 0.4 Tm. Particular emphasis was placed upon a study of the effects of interstitial-dislocation interactions. Single crystals of 3.25 pct Si-Fe containing about 200 ppm of C in solid solution were used in the investigation for the following reasons: 1) The mobility of interstitial carbon in 3.25 pct Si-Fe is negligible at 300°K but increases rapidly at slightly elevated temperature22. Hence, differences between the flow curves and dislocation structures of crystals deformed at 300°K, 0.17 T,, and crystals deformed, say, at 425°K, 0.24 Tm, should be appreciable because of the enhanced dislocation-carbon interactions at the elevated temperature. This effect was accentuated in some samples by interrupted straining, thereby introducing a certain amount of aging. 2) Near room temperature, slip in suitably oriented 3.25 pct Si-Fe single crystals is largely confined to the (110) planes.23'24 Dislocation structures formed under conditions of single glide are the least complicated and their method of formation is the most easily discernable. 3) Dislocations in Si-Fe can be tightly locked with carbon atmospheres by a low-temperature aging treatment. The subsequent thinning of samples to foil thicbess causes little or no rearrangement in the dislocation structure.25 EXPERIMENTAL PROCEDURE Large single-crystal sheets of 3.25 pct Si-Fe were donated by Dr. C. G. Dunn of the General Electric Research Laboratory, Schenectady, N. Y. The orientations of the sheets were determined and slabs 1.0 by 0.25 by 0.05 in. were cut such that the desired tensile axis corresponded to the long dimension. The slabs were mechanically polished and subsequently decar-burized by heating at 1000°C for 3 days in a flowing wet-hydrogen atmosphere. A carbon content of about 200 ppm was introduced by heating at 805°C for 25 min in a flowing atmosphere of dry hydrogen containing heptane vapor. Shaped copper tools were then used to spark-machine at 0.125 by 0.50 in. gage length onto each slab. Vacuum annealing at 1225°C for 2 days followed by a quench into the cold end of the furnace to retain carbon in solid solution concluded the soecimen preparation. Continuous tensile flow curves for crystals of severa1 orientations Were obtained both at 300' and 425°K. A strain rate of 6.67 x 10-4 Per set was used in these and all other tests. Crystals oriented for single glide, B and D in Fig. 1, were subjected to a 3.5 pct plastic elongation to insure uniform slip along the gage length; they were then immediately subjected to interrupted strain cycling as indicated in Fig. 2(a). Each cycle consisted of unloading to 1.5 kg per sq mm, holding

Jan 1, 1969

By Charles E. Lundin

The character of the Er-D system was established by determining pressure-temperature-composition relationships. A Sieuerts&apos; apparatus was employed to make measurements in the temperature range, 473" to 1223"K, the composition range of erbium to ErD3, and the pressure range of 10~s to 760 Torr. The system is characterized by three homogeneous phase regions: the nzetal-rich, the dideuteride, and the trideuteride phases. These phases and their solubility boundaries were deduced from the family of isotherms of the system zchich relate the pressure-temperature-composition variables. The equilibrium plateau decomposition relationships in the two-phase regions were determined from can&apos;t Hoff plots to be: The differential heats of reaction in these two regions are AH = - 53.0 * 0.2 and -20.0 *0.1 kcal per mole of D2, respecticely. The differential entropies of reaction are AS = - 36.3 * 0.2 and - 31.0 * 0.2 cal per mole D2. deg, respectively. Relative partial molal and intepal thermodynamic quantities were calculated from the pure metal to the dideuteride phase. The study of the Er-D system was undertaken as a logical complement to an earlier study of the Er-H system.&apos; The primary interest was to compare the characteristics of the two systems and relate the difference to the isotopic effect. Studies of rare earth-deuterium systems by other investigators have been very limited in number and scope. Furthermore, there is even less information available wherein an investigator has systematically compared a binary rare earth-hydrogen system with the corresponding rare earth-deuterium system. The available information consists primarily of dissociation pressure measurements in the plateau pressure region of a few rare earths. Warf and Korst&apos; determined dissociation pressure relationships for the La- and Ce-D systems in the plateau region and several isotherms for each system in the dideuteride region. They compared these data with those of the corresponding hydrided systems. The study of these systems as a whole was very cursory and did not give sufficient data for a thorough comparison of the effect of the hydrogen vs the deuterium in the respective rare earths. The heat capacities and related thermodynamic functions of the intermediate phases, YH, and YD2, were determined by Flotow, Osborne, and Otto,~ and the investigation was again repeated for YH3 and YD3 by Flotow, Osborne, Otto, and Abraham.4 This investigation studied only these specific phases. Jones, Southall, and Goodhead5 surveyed the hydrides and deu-terides of a series of rare earths for thermal stability including erbium. They experimentally determined isotherms of selected hydrides and plateau dissociation pressures for deuterides. These data allowed comparison of the enthalpy and entropies of formation of the dihydrides and dideuterides. To date, no one rare earth has been selected to thoroughly establish the complete pressure-temperature-composition (PTC) relationships of binary solute additions of hydrogen and deuterium, respectively. The objective in this investigation was to provide the first comparison of a complete family of isotherms of a rare earth-deuterium system with those of a rare earth-hydrogen system. This would allow one to determine what differences exist, if any, in the various phase boundaries and the thermodynamic relationships in various regions of the systems. I) EXPERIMENTAL PROCEDURE A Sieverts&apos; apparatus was employed to conduct the experimental measurements. Briefly, it consisted of a source of pure deuterium, a precision gas-measuring buret, a heated reaction chamber, a mercury manometer, and two McLeod gages (a CVC, GMl00A and a CVC, GM110). Pure deuterium was obtained by passing deuterium through a heated Pd-Ag thimble. A 100-ml precision gas buret graduated to 0.1-ml divisions was used to measure and admit deuterium to the reaction chamber. The reaction unit consisted of a quartz tube surrounded by a nichrome-wound furnace. The furnace temperature was controlled by a recorder-controller to . An independent measurement of the sample temperature in the quartz tube was made by means of a chromel-alumel thermocouple situated outside, but adjacent to, the quartz tube near the specimen. Pressure in the manometer range was measured to k0.5 Torr and in the McLeod range (10~4 to 10 Torr) to *3 pct. The deuterium compositions in erbium were calculated in terms of deuterium-to-erbium atomic ratio. These compositions were estimated to be *0.01 D/Er ratio. The erbium metal was obtained from the Lunex Co. in the form of sponge. The metal was nuclear grade with a purity of 99.9+ pct. The oxygen content was reported to be 340 ppm and the nitrogen not detectable. Metallographically the structure was almost free of second phase (<i vol pct). A quantity of sponge was arc-melted for use as charge material. The solid material was compared with the sponge in the PTC relationships. They were found to be identical. Therefore, sponge material was used henceforth, so that equilibrium could be attained more rapidly. The specimen size was about 0.2 gr for each loading of the reaction chamber. The procedure employed to obtain the PTC data was to develop experimentally a family of isothermal curves of composition vs pressure. First, a specimen

Jan 1, 1969

By L. W. Blau

RESEARCH work in exploration and production was further reduced during 1943 owing, partly, to difficulties in the acquisition of apparatus and, principally, to the exodus of research men to governmental agencies conducting research and to the armed forces.

Jan 1, 1944

By D. R. de Fontaine

By computing the values of the resolved shear stress for a great many orientations of the tensile axis on all 12 (111) <110> slip systems, Taylor and Elam&apos; were able to map out a stereogram of slip systems, given here, Fig. 1, in Diehl&apos;s notation (see, for example Hauser and Jackson&apos;). The purpose of the present note is to derive such a stereogram by mathematical reasoning and to give a simple method of defining the system with highest resolved shear stress for any direction of the tensile axis, thereby eliminating tedious computation of the Schmid factor. Let the tensile axis, of direction cosines [xyz], make the angles 0 and relative to the slip plane normal [HKL] and the slip direction [hkl] respectively. Fig. 2 shows a simple cubic cell whose edges define a Cartesian coordinate system. Let A be a unit vector of origin 0 supported by the tensile axis. Its end point [xyz] then lies on a sphere of unit radius, on the positive hemisphere of which the planes z = 0, x = y, y = z limit the so called primitive triangle of a standard [loo] stereographic projection. For any point of this triangle, the following hold: 0«2 «y *sx « 1 [l] Now, for given [xpz], the most highly stressed system will be the one for which the following expression of the Schmid factor has the largest possible value: (HKL) belonging to the form (111) and [hkl] to the form <110>, the denominator on the right hand side of Eq. [2] is equal to 1.fi.a. The unknowns H, ... h, . . . can now be determined by correctly distributing +1 and -1 to H, K, L, and +1, -1,O to h,k, 1 giving the numerator of Eq. [2] its greatest positive value (the angles 0 and cpare chosen acute) subject to the condition: Hh+Kk+Ll = 0 [3] since the slip direction must lie in the slip plane. Condition [3] can be satisfied only by having one of its terms vanish and the other two equal to +1 and -1. Taking the case of the tensile axis lying in the primitive triangle, x2y, z and therefore Hh will be chosen equal to +1 with H = h = +1 in order to retain a maximum of + signs. There remain two alternatives: either 1 = 0 or k = 0. 1) take 2 = 0 Then Kk = -1 and the numerator of [2] becomes: Retaining a maximum of + signs gives L = k = +1 and becomes: 2) take k = 0 Then L1 = -1and, likewise, the numerator of [2] becomes: xz - z2 + xy + yz [5] But the difference [5] - [4] can be written: (y - z)(x+y +z) [6] positive or zero by virtue of [I]. If [6] is positive (y > z), then the largest possible value for the numerator of [2] for any xyz inside the primitive triangle is given by [5], i.e., for: That is to say, the active slip system corresponding to an orientat:~on of the tensile axis inside the primitive triangle is:

Jan 1, 1962

By R. C. Ku, A. J. McEvily, T. L. Johnston

The microplastic response of a series ofas-quenched Fe-Ni-C martensites has been measured at 77°K. At strains less than JO&apos;3 the flow stress is governed primarily by the transformation-induced dislocation structure of the martensite. Only at strains in excess of 10-3 is the influence of carbon manifested in the flow stress. At these macroscopic strains, typically 10-2, the solid-solution hardening is proportional to (wt pct C)1/3, and, in an alloy containing 0.39 wt pct C, amounts to 50 pct of the flow stress. THE technological significance of high-strength ferrous martensite has stimulated many investigations of its structure and properties. Although our knowledge of the characteristics of martensite has increased immensely, especially with the advent of high-resolution techniques, an understanding of the basic strengthening mechanism still remains elusive. The purpose of the present paper is to consider certain aspects of micro-plastic behavior of Fe-Ni-C martensite which we feel can help to resolve this important problem. Such alloys are particularly suitable for experimental investigation because their compositions can be adjusted to reduce the M, to a temperature low enough essentially to eliminate the diffusion of carbon in the freshly formed martensite.1 The mechanical properties in this condition are of interest inasmuch as they reflect a state that is free of the important but complicating influence of precipitation processes. In this virgin martensite the carbon is distributed as it was inherited from the parent austenite; i.e., it is present interstitially, and gives rise to tetragonality through strain-induced ordering.&apos; In order to determine the source of strength of such alloys, Winchell and Cohen1 investigated the low-temperature macroscopic stress-strain behavior of a series of virgin martensites of increasing carbon content but of common M, temperature (-35°C). They found that the flow stress increased rapidly with carbon content up to 0.4 wt pct; beyond this point the flow stress increased at a much slower rate. It was concluded that martensite is inherently strong. To account quantitatively for the strength of virgin or as- quenched martensite in terms of the role of carbon, Winchell and cohen3 suggested that the carbon atoms, trapped in their original positions by the diffusionless martensite transformation, interfere with dislocation motion according to a model akin to that of Mott and Nabarro. 4 In this treatment, individual carbon atoms are considered to constitute centers of elastic strain and thereby generate an average stress resisting the motion of dislocations throughout the lattice. The additional stress necessary to move dislocations, over and above that necessary for motion in a carbon-free martensite, is given by where L is an effective length of dislocation capable of motion. L was assumed to be limited to the spacing between the twins that are an essential structural element of Fe-Ni-C martensites. They assumtd the spacing to be invariant and of the order of 100A. However, recent work5 has shown that L is variable and can be in excess of 1000Å, so that the assignment of an appropriate value of L is not straightforward. In contrast to the above conclusion that there is an intrinsically high resistance to plastic flow, it has been suggested by Polakowski6 that freshly quenched martensite is in fact "soft" in the sense that dislocations are initially free to move upon application of stress. The high indentation hardness and macroscopic yield stress of ferrous martensites are then a consequence of rapid strain hardening that depends upon carbon in solution. Consistent with this point of view are the results of Beau lieu and Dubé who measured the rate of recovery of internal friction as a function of aging (tempering) temperature in a freshly quenched steel containing 0.90 wt pct C, 0.37 wt pct Mn, 0.1 wt pct Cr, and 0.07 wt pct Ni. The kinetics were clearly consistent with the idea that many dislocations are unpinned in the as-quenched state and that during aging they become progressively pinned by carbon at a rate controlled by carbon diffusion in the body-centered martensite lattice. In order to provide a basis upon which to distinguish between the "hard" and "soft" interpretations indicated above, we have made studies of the initial stages of plastic deformation in Fe-Ni-C martensites similar to those&apos;used by Winchell and Cohen. It will be shown that the results support the contention that dislocation segments in as-quenched material are indeed

Jan 1, 1967

By Charles H. Grant

Some of the limiting factors relative to explosive crushing of rock and ways to overcome a few of these problems are presented. Relationships between borehole diameters, bench heights, and spacings, along with a review of the influence geometry has on energy as these are changed, are discussed. Efficiency in use of explosives and the decay of energy as it moves through rock and is absorbed and dissipated, is described, along with fragmentation as a function of spacings and energy zoning, etc. Communications are one of the major problems encountered. In an effort to provide a better understanding of the use of explosives, it is necessary to take a little different view of what explosives are, how to look at them as tools to fragment rock, and some of the problems encountered in doing so. First, take the explosive: although there are many factors involved, consider these as being reduced to only two — shock-strain imparted to the rock by the high early development of energy, and the gas effect which is a combination of heat, moles of gas formed, rate of formation of these gases which develop pressures, etc. First, consider shock energy by itself and assume there is no gas effect in the reaction. Fig. 1 illustrates a block or cube of rock, in the center of which is detonated an explosive charge which is 100% shock energy. Tensile slabbing would be seen on the surface and probably the cube of rock would generally hang together even though microcracks were formed. If the situation is reversed and an explosive whch has no shock energy and only gas effect (Fig. 2) is considered, the cube of rock would act as a pressure vessel and contain the pressure from the gas effect until it exceeded the rock-vessel strength; then the rock would break in a few large pieces. If these two kinds of energy are put together and the area of shock-strain around the explosive (Fig. 3) is considered, the two energies will be seen working together to furnish broken rock. The gas effect applies pressure to the microcracks formed from the shock energy to weaken the rock-pressure vessel and propagate these cracks to break the rock apart. It not only will be broken more finely, but will break apart at a lower pressure than the gaseffect case, since the shock energy has first weakened the rock vessel. Although tensile spalling from the shock-strain imparts momentum to the rock, the main source of displacement comes from the gas effect. The term "rock" is being used to mean any material to be blasted. These energies are absorbed by the rock in different ways. First, classify rock into two main categories: "elastic" and "plastic-acting." Elastic rock should be thought of as rock which can transmit a shock wave and is high in compressive strength, such as granite or quartzite. Since this elastic rock transmits a shock wave well, it makes good use of the shock energy from the explosive-forming cracks, etc., for the gas effect to work on. Plastic-acting rocks are rock masses which are relatively low in compressive strength and absorb shock energy at a much faster rate, thereby making poor use of the shock energy by not developing as extensive a cracked zone for the gas effect to work on. Rocks of this type are generally softer materials such as some limestones, sandstones, and porphyries. For the most part, the shockenergy part of the explosive reaction is wasted in plastic-acting rock, leaving most of the work to the gas effect. Since the ratio of gas effect to shock energy is different in different explosives, it is easy to understand why some explosives perform well in elastic rock and poorly in plastic-acting rock, and vice versa. Some of the most difficult blasting situations arise when mixtures of plastic-acting and elastic rock are encountered (Fig. 4). Fig. 4 shows an example of granite boulders cemented together with something like a decomposed quartz monzonite which is plastic-acting. The elastic granite boulders will transmit the shock-strain within itself, but when this shock tries to move through the monzonite to the next boulder, its intensity is absorbed by the monzonite and little shock-strain is placed on the adjoining boulder. In addition to this loss by absorbtion, shock reflection at the surface of the boulder will effect tensile spalling. The net effect is poor breakage of the boulders which do not have drillholes in them as they simply will be popped out with the muck. The same is true (Fig. 5) when layers and joints make

Jan 1, 1970

By A. Ghosh, T. B. King

The kinetics of the discharge reaction: 20&apos;- (in silicate melt) = O,(g) + 4e- at a platinum anode in lithium silicate melts have been studied al 1350°C by galvanostatic methods. Plots of the sleady-state overpotential, q, as a function of the logarithm of the current density, i, showed injlections and were linear only at high current densities. The value of the overpotential was influenced by bubbling gas through the electrolyte. The ocer potential was also studied as a function of time. The rise and decay of overpotential were very slow processes. At low current densities transport is the likely rate-controlling process but at high current densities passivation of the electrode, Presumably by an oxide film on the surface, seems to be a contributory functor. IT is well-established that molten silicates behave as electrolytes&apos;5 and, except in a few cases,6 conduction is entirely ionic. Moreover, it is supposed that a possible, and perhaps predominant, mechanism for phase boundary reactions between metals and slags is similar to that in corrosion whereby anodic and cathodic processes occur at unrelated sites, the metal serving to conduct electrons.1&apos;8 Thus electrochemical studies of some slag-metal reactions would seem to be a useful way to diagnose the rate-controlling steps in the overall reaction. The electrochemical method is, in principle, a better diagnostic tool than the direct chemical method for the following reasons: 1) The partial electrochemical reactions, which are simpler than the overall reaction, may be studied individually. 2) The rate of reaction can be controlled at will and independently of the concentrations of reactants. 3) Fast reactions can be studied by relaxation methods.&apos; Esin and his coworkers5&apos;10"12 have pioneered such studies in silicates and have deviloped some ingenious techniques. Not all of their findings, however. can be accepted without a good deal of further work. In this investigation, the kinetics of the oxygen discharge reaction: 202- (in silicate melt) = Oz(g) + 4e- [I] at a platinum electrode were studied by both steady-state and transient galvanostatic techniques. Interest in this reaction was first developed as a result of the findings of Fulton and chipman13 that the reduction of silica, in a silicate slag, by carbon, dissolved in liquid iron, is a very slow reaction. Subsequent work, for example, by Rawling and ~lliott,&apos;~ has demonstrated that the reaction under these conditions must be slow, because the rate is limited by diffusion of oxygen in the iron to the metal-crucible phase boundary at which a CO bubble may be nucleated. Further work by Tarassof,&apos;~ in which the reduction of silica by aluminum dissolved in copper was studied, has shown that under these conditions, where carbon monoxide evolution is not involved, control of the reaction rate resides in diffusion of silica in the slag phase. However, there is no practical way of inducing sufficient convection in the system to make it clear that the phase boundary reaction is indeed fast. The overall reaction of silica reduction involves the discharge of silicon ions at cathodic sites and oxygen ions at anodic sites. In the examples cited, the discharged ions are dissolved in a liquid metal. In the present study of oxygen ion discharge, gaseous oxygen may be evolved at high current densities or oxygen may simply dissolve, possibly as oxygen molecules, in the silicate at very low current densities. The discharge of an oxygen ion at an anode must, in silicates less basic than the orthosilicate composition, be preceded by a reaction in the vicinity of the electrode, such as: which makes oxygen ions available. Platinum was chosen as the working electrode since it is comparatively inert to oxygen and is, therefore, expected to come rapidly into equilibrium with the electrolyte and with gaseous oxygen. Minenko, Petrov, and Ivanova16 have measured the electromotive force at a platinum electrode in molten silicates as a function of the partial pressure of oxygen in the atmosphere, the concentration of oxide ions in the melt, and the temperature. They found platinum to behave as a reversible oxygen electrode. At two different oxygen pressures, Po2 (I) and Pq (11). the electromotive force is given by: where F is the Faraday constant, equal to 23,060 cal per v equivalent, indicating that the electrode reaction is as written in Eq. [I.]. Platinum has been similarly used in molten silicates by other inve~ti~ators. "&apos;~~ In this investigation platinum was used only as an anode, since a current deposits other elements on its surface and changes its characteristics.

Jan 1, 1970

By F. Claisse, R. Angers

Self-diffusion in iron has been measured during rapid a-r transformations using a variant of the Kryukou and Zhukhovitskii diffusion method. The study was performed by thermally cycling iron foils (1 to 6 cpm) through the transformation (=910°C). Some foils have been subjected to over 1000 cycles and some have spent more than 15 pct of their total diffusion time in the process of transformation. The experimental results show that the a-r transformation has no measurable effect on self-diffusion in iron. The study is completed by a quantitative analysis of mechanisms which can affect the diffusion rate during the transformation. The analysis confirms the experimental results. SINCE diffusion is an important factor in many solid-state transformations, it is of interest to study how it is affected by the stresses generated during these transformations. Clinard and Sherby1&apos;2 were the first to make a study along these lines. They measured diffusion coefficients in Fe-FeCoV couples subjected to slow thermal cycling (1.5 cph) through the a-r transformation range. They found an enhancement of diffusion by a factor of about two. The purpose of the present paper is to report measurements of the self-diffusion coefficient of iron during much more rapid thermal cyclings (1 to 6 cpm) through the a-r transformation (-910°C). These more rapid cyclings produce higher strain rates during the transformation and should emphasize any possible influence of transformation upon diffusion. EXPERIMENTAL Iron foils, 25 to 35 µ thick, were cold-rolled from 99.92 pct pure iron and annealed in pure helium for 2 hr at 870°C; the resulting grain diameter was about 150 µ. Specimens 0.5 by 8 cm were cut from the foils and I7e55 was vapor deposited on one of their surfaces. A 38 gage alumel-chrome1 thermocouple was spot welded in the middle of one of the specimen long edges, Fig. 1. Two 38 gage chrome1 wires were also spot welded along the same edge on each side of the thermocouple; they were placed 2.5 cm apart and used for electrical resistance measurements. In order to prevent twisting and crumpling, the specimens were pinched between two quartz plates 0.1 by 1 by 7 cm and the assembly was close fitted into a 1 cm ID quartz tube. Four holes were drilled through the tube to let the 38 gage wires out: these were connected to the recording equipment by means of extension wires. 20-gage nickel wires fixed at both ends of the specimens were used to thermally cycle the foils by Joule heating. The above described device was placed in a 2.7 cm ID quartz tube which in turn was placed in a tubular furnace. Either a pure helium atmosphere or circulating hydrogen was used during the experiments. Specimens were subjected to thermal cycles between a minimum temperature To and a maximum temperature Tm at rates ranging from 1 to 6 cpm. This was obtained by maintaining the furnace at a constant temperature near the minimum temperature To and periodically passing an electric current through the specimen. Cooling was achieved by heat losses to the surroundings. The forms and periods of cycles were varied from one specimen to another; however, each specimen was subjected to one type of cycle only. The temperature and electrical resistance variations of the specimens were recorded as a function of time. The temperature curves were used for diffusion calculations while the electrical resistance curves were used to monitor the transformation and to determine its starting point and its approximate duration. Diffusion was measured by the method developed by Kryukov and zhukhovitskii3 and modified by Angers and Claisse.4,5 In this method a metallic foil is coated on one side with a radioactive isotope and the activity is measured periodically on both sides during the diffusion anneal. The following equation then holds: where: I1 Activity on the surface on which the deposit is made. I, Activity on the opposite surface. t Diffusion time. B Constant. D Diffusion coefficient. d Foil thickness including the deposit. G(t) A function of time; it is a second order correction term which is given graphically in Refs. 4 and 5. The diffusion coefficient D is found by plotting ln[(Il - I2)/(I1 + I,)] -G(t) against t; the resulting slope m leads to an accurate calculation of D through Eq. [2]. The effect of the a-r transformation on diffusion is expressed by the ratio DT/DU where:

Jan 1, 1970

By W. M. McKewan

Magnetite pellets were reduced in flowing hydrogen at pressures up to 40 atm over a temperature range of 350° to 500°C. The rate of weight loss of oxygen per unit area of the reaction surface was found to be constant with time at each temperature and pressure. The reaction rate was found to be directly proportional to hydrogen pressure up to 1 atm and to approach a maximum rate at high pressures. The results can be explained by considering the reaction surface to be sparsely occupied by adsorbed hydrogen at low pressures and saturated at high pressures. PREVIOUS investigation1,2 have shown that the reduction of iron oxides in hydrogen is controlled at the reaction interface. Under fixed conditions of temperature, hydrogen pressure, and gas composition, the reduction rate is constant with time, per unit surface area of residual oxide, and is directly proportional to the hydrogen pressure up to one atmosphere. The reduction rate of a sphere of iron oxide can be described3 by the following equation which takes into account the changing reaction surface area: where ro and do are the initial radius and density of the sphere; t is time; R is the fractional reduction; and R, is the reduction rate constant with units mass per area per time. The quantityis actually the fractional thickness of the reduced layer in terms of fractional reduction R. It was found in a previous investigation2 of the reduction of magnetite pellets in H2-H,O-N, mixtures, that the reaction rate was directly proportional to the hydrogen partial pressure up to 1 atm at a constant ratio of water vapor to hydrogen. Water vapor poisoned the oxide surface by an oxidizing reaction and markedly slowed the reduction. The enthalpy of activation was found to be + 13,600 cal per mole. It was also found that the magnetite reduced to meta-stable wüstite before proceeding to iron metal. The following equation was derived from absolute reaction-rate theory4,8 to expfain the experimental data: where Ro is the reduction rate in mg cm-2 min-&apos;; KO contains the conversion units; Ph2 and PH2O are the hydrogen and water vapor partial pressures in atmospheres; Ke is the equilibrium constant for the Fe,O,/FeO equilibrium; Kp is the equilibrium constant for the poisoning reaction of water vapor; L is the total number of active sites; k and h are Boltzmann&apos;s and Planck&apos;s constants; and AF is the free energy of activation. Tenenbaum zind Joseph5 studied the reduction of iron ore by hydrogen at pressures over 1 atm. They showed that increasing the hydrogen pressure materially increased the rate of reduction. This is in accordance with the work of Diepschlag,6 who found that the rate of reduction of iron ores by either carbon monoxide or hydrogen was much greater at higher pressures. He used pressures as high as 7 atm. In order to further understand the mechanism of the reduction of iron oxide by hydrogen it was decided to study the effect of increasing the hydrogen pressure on rebduction rates of magnetite pellets. EXPERIMENTAL PROCEDURE The dense magnetite pellets used in these experiments were made in the following manner. Reagent-grade ferric oxide was moistened with water and hand-rolled into spherical pellets. The pellets were heated slowly to 550°C in an atmosphere of 10 pct H2-90 pct CO, and held for 1 hr. They were then heated slowly to 1370°C in an atmosphere of 2 pct H2-98 pct CO, then cooled slowly in the same atmosphere. The sintered pellets were crystalline magnetite with an apparent density of about 4.9 gm per cm3. They were about 0.9 cm in diam. The porosity of the pellets, which was discontinuous in nature, was akrout 6 pct. The pellets were suspended from a quartz spring balance in a vertical tube furnace. The equipment is shown in Fig. 1. Essentially the furnace consists of a 12-in. OD stainless steel outer shell and a 3-in. ID inconel inner shell. The kanthal wound 22 in. long, 1 1/2, in. ID alumina reaction tube is inside the inconel inner shell. Prepurified hydrogen sweeps the reaction tube to remove the water vapor formed during the reaction. The hydrogen is static in the rest of the furnace. The sample is placed at the bottom of the furnace in a nickel wire mesh basket suspended by nickel wire from the quartz spring. The furnace is then sealed, evacuated, and refilled with argon several times to remove all traces of oxygen. It is then evacuated, filled with

Jan 1, 1962

By C. N. Garman

Homestake-New Mexico Partners operate a 750-tpd carbonate leach uranium concentrate mill near Grants, N.M. The highly mineralized water available as process water leaves much to be desired. The 628 ppm as CaCO 3 makes the use of raw water very troublesome in pipes and on filter cloths. However, the residual sodium carbonate in the final filter cake going to tails makes an ideal softening agent. To take advantage of this fact, all makeup water used in the mill is first used as tailing slurry dilution water and comes to the mill from the tailings pond. The 5-acre tailings pond serves as a thickener and 100 to 150 gpm of nearly clear solution is decanted to a pump to be returned to the mill. Since this tailings water has small quantities of uranium in the solution an ion exchange scavenger unit was installed to remove as much uranium as possible. The ion exchange raffinate is then used as final filter wash ahead of the tailings slurrying step. In spite of the large settling area this return water is not clean enough for ion exchange feed. The solids present are very fine and composed of approximately 15 pct (by weight) burnable carbonaceous material common to the sandstone uranium ores in the area, 40 pct SiOz plus 45 pct CaC03. Laboratory work showed that this material responds very well to flotation. Before deciding to use flotation, various clarifying systems such as pressure leaf filters, sand filters, and continuous vacuum pre-coat filters, were considered. Each of these could have solved the problem but with much more operating labor, more reagents and greater installation costs than the flotation step. About 100 to 150 gpm of fouled water is fed to two 66-in. Fagergren cells, in series. Reagents used at the beginning were Arquad 2HT75 and Arquad C50, at the rate of about 1% lb per 8-hr shift, or about 0.0053 lb each per ton of ore. This did not completely remove the solids but does an acceptable job. Approximately 75 pct of the slimes are a size that can be caught on a 41-Whatman Paper are removed. Removal of these slimes also allows much better settling of the coarse nonfloatable material. Advantage is also taken of this fact in a small settling tank ahead of precipitation. Removal of this amount of the slimes makes the ion t:xchange feasible. PREGNANT SOLUTION CIRCUIT The carbonate? leach-caustic precipitation method of uranium concertration does not provide for any process purification step ahead of precipitation. Therefore, any fine solids getting into the pregnant solution through the filter cloth show up in the final concentrate. This, of course, lowers the grade, and, at times, the slimy nature of these very fine solids rendered final filtration of the concentrate difficult if not impossible At Homestake-New Mexico Partners a 75-ft thickener was available for gravity clarification of 100 to 120 gpm of this pregnant solution. However this did not sufficiently remove the slimes. Laboratory investigation of the whole range of flocculants that were suggested by literature, salesmen, and friends failed to turn up anything of consequence. A continuous vacuum pre-coat filter would do the job and was investigated. The capital cost and the operating labor and materials made this a last chance choice. Following work done in the metallurgical laboratory on the tailings return water, it was found that some changes in the reagent strengths and combinations made a very definite decrease in the solids in the pregnant solution. Concentrate grade improved about 5 pct anti the final product after drying had an appreciably greater bulk density. Compared to a cost of about 2.2e per ton for pre-coat filter opelation for cleaning just one circuit, flotation costs less than 1.0 per ton of ore for cleaning two circuits. While a pre-coat filter would do a more thorough job, the flotation does all that is required for either circuit. Gravity causes the froth produced to run back into the leach circuit. This does not appear to result in a build-up of objectionable slime. No extra manpower is required; the operators in the separate areas can observe the operation of the cells and mix the small quantities of reagents as needed. Normally the 66-in. Fagergren cell requires 15 hp per cell, but this very dilute slurry needs only 10 hp for both cells. Originally, a combination of the two Arquads mentioned previously served as frothers and promoters. As further testing

Jan 1, 1962

By R. M. Waterstrat

The phases occurring in the V-Mn system were studied by means of X-yay diffraction and metallo-paphic techniques, using are-melted alloy specimens annealed in the temperature range 800° to 1150°C and quenched. The bcc solid solution extends at 1250°C all the way from vanadium to 6-manganese. Below 1050°C the a-phase is formed, and the terminal a-manganese phase is stabilized up to about 900°C by vanadium in solid solution. IN the only previous general survey of the V-Mn system Cornelius, Bungardt and Schiedtl reported the existence of three intermediate phases corresponding to the approximate compositions VMn,, VMn, and V5Mn. The phase VMn8 has recently been identified as a o phase2 but the alloy VMn was found to have a bcc structure2 corresponding apparently to the vanadium solid solution rather than to the large cubic unit cell reported by Cornelius et al. 1 Subsequent work by Rostoker and Yamamoto3 has shown that the vanadium-base bcc solid solution extends to at least 15 pct Mn at 900°C. An alloy corresponding to the composition VMn, was examined by Elliott,4 who reported that the as-cast sample as well as samples annealed at 1200o and 1300°C had bcc structures, but that annealing at 1000°, 800") and 600°C produced two phases. One of these phases was apparently the bcc solid solution and the other resembled the o phase structure. Hellawell and Hume-Rothery5 established the phase relationships in manganese-rich alloys above 1000°C, and showed that the o phase in this system is replaced by the 6 Mn (bcc) solid solution at temperatures above 1050°C. These results suggest that a continuous bcc solid solution may exist above 1050°C between vanadium and 6 Mn. The present investigation was undertaken in order to develop more complete information in regard to this system. EXPERIMENTAL METHODS The alloys used in the present work were prepared by arc-melting electrolytic manganese having a minimum purity of 99.9 pct and vanadium lumps with a purity of 99.7 pct. The major impurities present in these metals were carbon, nitrogen, and oxygen and this would account for the small percentage of nonmetallic inclusions observed metal-lographically. The arc-melting was at first performed under a helium atmosphere and it was necessary to keep the melting times as short as possible in order to minimize the loss of manganese by vaporization. It was later found that the evaporation of manganese was considerably reduced when the melting was done under argon atmosphere. The final composition of each alloy was calculated by assuming that the total weight loss during melting was due to evaporation of manganese. Compositions which were calculated in this manner agreed reasonably well with the results of chemical analysis, as shown in Table I. Spectrographic analysis revealed the presence of contamination by tungsten, but in no case was the percentage of tungsten greater then 0.4 at. pct. The specimens were in each case broken in half and the fractured section was examined visually and microscopically for evidence of inhomogeneity. Each specimen was homogenized at temperatures near l100°C, as shown in Table I. After this treatment most specimens consisted of large columnar grains of the bcc vanadium solid solution. The etchant used in most of the metallographic work consisted of 20 pct nitric acid, 20 pct hydro-flouric acid, and 60 pct glycerine. It was found that this etchant would clearly delineate the phases present in these alloys although it does not produce any striking contrast between the phases. For certain manganese-rich alloys, a 1 pct aqueous solution of nitric acid was used. This etchant gave a brown color to the a-manganese phase, whereas the o phase was virtually unattacked and appeared very light as shown in Fig. 1. The etchants used by Cornelius et a1.l were found to produce spurious effects in some of these alloys. In particular, the vanadium-rich alloys etched in hot sulfuric acid often appeared to consist of two phases when both X-ray diffraction and etching with the glycerine-acid mixture indicated the presence of single phase bcc solid solution. A few percent of what appears to be an oxide or nitride phase was found at the grain boundaries and in the interior of the grains, especially in the vanadium-rich alloys. All alloys were annealed in sealed silica tubes containing 1 atm of pure argon and these tubes were then quenched in cold water. Although some manganese loss occurred during annealing, the loss seemed to be confined to the surface of the speci-

Jan 1, 1962

By R. E. HORE, J. B. Tyrrell

IN the vicinity of Kirkland Lake, northern Ontario, several gold mines are producing, and three are at present being operated very profitably. At these three mines and on the adjoining properties east and west of them there have been important developments during the past year, and additional equipment has been installed to permit a large increase in output. At the end of 1924 the total production amounted to $13,688,600. For the year 1926 the production is estimated at about $19,000,000, and during 1926 the output should be much greater than that of 1925. The mines which are being operated most profitably and are producing a large part of the ore now being milled are the Wright-Hargreaves, Lake Shore and Teck-Hughes. Others which have contributed smaller amounts are the Tough-Oakes-Burnside and the Kirk- land Lake Gold mines.

Jan 1, 1926

By A. L. Engel

IN the treatment of complex ores by flotation, one of the most important steps is conditioning the feed. Conditioning primarily consists of the addition, in the grinding circuit, of an alkaline reagent, for example sodium carbonate, which neutralizes any acidity of ore and mill water, counteracts the effect of &apos;harmful soluble salts, disperses colloids, and provides the necessary alkalinity for the best metallurgical results. The several minerals of the ore should be put in such condition that they may be floated or rejected as desired, through the action of further reagents thereafter used.

Jan 1, 1932

By Harold V. Pearce

THE plant of the Boston and Colorado Smelting Company, at Argo, Colo., has not received special notice in technical or scientific publications for some time past. Dr. Peters&apos; described the development of the reverberatory furnace used at Argo, but he dealt almost exclusively with the increase in size and the corresponding gain in capacity. Since the time of his description the changed conditions have led to a modification of the former smelting practice, and this modification forms the basis of the present paper. No claim is made for personal originality in connection with these improvements, in fact, I expressly disclaim any such intention, since nearly every successful step in the development of the process now used at Argo has been suggested by Dr. Richard Pearce at various times during the past 12 years, though perhaps in a different manner of exe¬cution. These changes would undoubtedly have been made before, were it not that many problems of greater metallurgical importance and momentarily of more urgent need thus had to be solved. At the time to which I refer, economic conditions demanded mainly an increase of capacity. Moreover, the determination and placing of exact losses in smelting were not so imperative as they now are. About three years ago, the time of my appointment to the - management of. the works of the Boston and Colorado Smelting Co., the great decline in the price of copper occurred; the Argo plant, in common with all others engaged in the copper-industry, lost heavily in the market-value of the stocks on hand of ore and furnace-products. Added to this was the unfortunate circumstance that the copper-supply of the Argo plant from general sources began to dwindle; in fact, it almost entirely ceased. About this time the smelting-plants in the west-

Mar 1, 1905

By Zay, Jeffries

THE effect of thoria (ThO2) on grain growth in tungsten was discussed in some detail in a paper presented before this Institute by one, of the authors in 1918.1 In that paper it " was assumed that the, thoria was present in the tungsten entirely as a mechanical mixture. There was no indication in the experiments described that even a trace of the thoria was in solid solution in the tungsten. Langmuir,2 however, observed that thoriated-tungsten filaments under certain conditions showed electron emissions many thousand times greater than unthoriated tungsten at the same temperature. He concluded from his studies that a slight amount of the thoria had dissociated on heating the filament to a temperature near 220°. C. At this temperature and higher, part of the thorium diffuses to the surface of the filament where it evaporates at. a much faster rate relatively, than the tungsten, and also at a rate, nearly as fast as it is brought to the surface by diffusion: If the temperature is lowered to say 1800°,. C., diffusion of the thorium atoms to the surface is more rapid than evaporation, sown adsorbed film of thorium atoms is produced on the surface of the tungsten filament. The electron emission is governed by the outside layer of atoms so the emission from such a filament might be expected to approxi¬mate that of a thorium filament at the same temperature. PRESENCE OF METALLIC THORIUM IN THORIATED TUNGSTEN If the filament containing the adsorbed film of thorium. is heated to a sufficiently high temperature, the thorium atoms evaporate: rapidly and the electron emission becomes more nearly that of tungsten. On cooling again to 1800° C., or thereabouts, the emission &apos;builds up gradually to a maximum. The filament containing the adsorbed layer of thorium atoms is said to be "activated." By studying the rates of activation under different conditions of prior thermal treatment and at different temperatures, Langmuir was able to arrive at an estimate of the

Jan 1, 1927

By T. A. Rickard

A knowledge of the principles of punctuation is essential to effective and intelligible writing, for the ease and pleasure of the reader, and even his understanding, may depend upon the choice and the placing of punctuation marks. To &apos;punctuate&apos; is to mark with stops or points; the word comes from the medieval Latin punctuare (punctuatus), which in turn is derived from the Latin punctum, a point. The purpose of punctuation is to show the manner in which the writing is to be understood; to indicate the relations between parts of a sentence and between successive sentences. To some extent "it does for the eye what vocal stress does for the ear", but especially it marks logical relations. Some punctuation marks, such as those indicating admiration, exclamation, and interrogation, are rhetorical as well as grammatical, but most of the common ones are logical, or grammatical. My discussion will be confined to grammatical punctuation in ordinary writing. .&apos;The problem of punctuation in text matter", as George Summey * says, "is to employ words, points, and paragraph breaks in such a way as to achieve at the same time clearness; proper distribution of emphasis, and the desired kind of movement." The criticism of the punctuation of a sentence or of a group of sentences is not final until the context, within the paragraph, is read. A paragraph is a sentence or a group of sentences forming an independent unit of the composition; it represents a unit of thought. A multiplicity of paragraphs produces a choppy effect; &apos;on the other hand, paragraphs of. excessive length are tiresome, because the belaboring of one unit of thought wearies the reader. The aim should be to

Jan 1, 1931

By M. J. Sinnott, J. C. Tobin

A technique for collecting and analyzing the small quantities of gases desorbed on heating metal powders has been developed. The gases collected from copper powders of various types of manufacture have been analyzed and found to contain H2O, CO2, and SO2. The quantities of gas evolved vary from 0.05 to 1.0 milliliter, under standard con-ditons of temperature and pressure, per gram of powder. The source of the gas has been traced to surface films on the powder. IN the powder-metal industry it is recognized that the presence of gases in the powders affects the processing and resulting properties of the sintered product. Several sources for these gases are possible; namely, chemical reactions with the sintering atmospheres, inclusion. of gases during compacting, dissolved gases in the metal powder, and sorbed gases on the metal-,powder surfaces.1, 2, 3 Since the quantities of the gases are relatively small, no direct attempt has been made to collect and analyze them nor to isolate their source. The work that has been done has been concerned with the dimensional stability of the product., its porosity, density, or mechanical and physical properties. Conclusions are drawn on the basis of studies of this type as to the probable effect of gases on these properties. 4, 5, 0 In the present investigation an attempt has been made to obtain quantitative data on the analysis and quantities of gases associated with copper powders manufactured in different ways. Apparatus was constructed to collect gases evolved from loosely packed powder samples on heating to various temperatures under a reduced pressure of helium. The powders were outgassed thoroughly in a vacuum prior to starting the collection operation in order to eliminate mechanically entrapped gases and to prevent the possible reaction of the powder with these gases. Experimental Procedures Materials—The major portion of the investigation was carried out on a single heat of water-atomized, oxygen-free, high-conductivity copper. A relatively coarse fraction of this powder, 20 to 100 mesh, was tested as a mixed powder and as various graded size fractions. The microstructure of this powder showed equiaxed grains and was clean and free from oxides and oxide eutectic. The particle size, as determined

Jan 1, 1959