Search Documents

By U. C. Tainton

The paper reviews the evolution of electrolytic zinc, describing some of the major obstacles that have been encountered and overcome. The chief remaining limitations of present-day standard practice are: Difficulty of obtaining high extractions, especially from low-grade and complex ores; filtration troubles, if ores contain much soluble silica; and irregularities in electrolysis, if more than certain limited amounts of antimony, arsenic, cobalt; etc., are present in the ores. These factors dictate a relatively high-grade plant feed, thus subjecting the electroytic-zinc process to the same kind of limitations as the zinc smelter. The "high acid process," in which both acid strength and current density are raised to about four times the usual figures, is described. This system concentrates both plant and operations, allows a simple type of flow sheet, and. minimizes some of the difficulties above mentioned. The large-scale operation of the process is described. SOME time ago, at a meeting of the Institute Prof. J. W. Richards1 said, "I take exception to the statement that all the factors in the production of electrolytic zinc were known long ago. There is possible in my opinion an improvement of perhaps 50 per cent. in the new electrolytic processes." An examination of the literature of the subject will show the justness of Professor Richards' contention. The work of the earlier investigators was carried on in the shadow of difficulties that, today, are hardly more than memories. We seldom hear now of zinc sponge, that bête noir of the electrolytic-zinc pioneers. Zinc sponge, was the name given to a peculiar, soft, black, non-adherent form of zinc deposit that was utterly useless for melting into ingot form. At the big plant at Cockle Creek, New South Wales, in 1897, Edgar Ashcroft2 said that the solution in the plant would suddenly, and without ascertainable cause, go "bad" and begin to deposit spongy zinc. The most delicate, even spectroscopic, methods of analysis were unable to detect any difference between "good solution and "bad" solution; either form might change to the other on being kept for a while. When it is added that an outbreak of this insidious disease (which. appeared to

Jan 2, 1924

By J. L. Byers

THE metals of the platinum group, ruthenium, rhodium, palladium, osmium, iridium and platinum, have certain characteristic effects on the structure and appearance of the cupellation bead. These effects are dependent not only on the particular member of the group present, but also on its concentration. The extent and quality of these effects, however, have not received investigation worthy of their importance. Bannister and Patching investigated the qualitative effects of some of these metals in 1913. This paper deals with a possible rapid qualitative and quantitative delineation of the effects exerted on the cupellation bead by the various metals in the platinum group. More particularly, the paper correlates quantitatively the presence of such metals with their characteristic structural and surface effects. Inasmuch as the various members of the group are more frequently associated with gold than with silver, the gold button will receive the greater consideration. The past two decades have seen a much greater interest manifested in this particular group of metals, because of the increased industrial demand for them. There has been a constantly growing need for assays that will not only express the silver and gold values but also indicate the presence of the other metals. If these indications could be used as quantitative criteria, their value would be greatly increased, because the methods of chemical analysis for the elements of the platinum group are lengthy, involved and unreliable except when the analysis is made by an expert and experienced rare-metal analyst. The materials for the X-ray diffraction examination of the different beads were prepared by filing and screening to -200 mesh, and were then rayed in a General Electric X-ray diffraction unit. The filed samples were treated magnetically in order to eliminate any possible steel particles. Sodium chloride and C.P. gold and silver were used as standard comparators.2 (See Table 1.)

Jan 1, 1932

By L. F. Porter, G. J. Dienes

The effect of netltron irracliation on mart ensite transformation in the iron-nickel system was studied by means of electrical resistance and magnetic induction measuremertts. Irradiation lowers the M8 temperature of the allay, causes changes in the electrical and magnetic properties of the high-temperature phase which indicate either an accelemtion of phase separation or a clustering of nickel, and enhances strain recovery on specimens partially transformed to martensite prior to irradiation. Possible mechanisms to account for the effects; are discussed. 1 HERE is a good deal of evidence to indicate that neutron irradiation is capable of affecting phase changes in metals and alloys. While the mechanism of the effect often is not well understood, such studies hold promise of improving our understanding not only of irradiation damage but also of the phase changes involved. To date systematic and detailed investigations have been made of the influence of nuclear radiation on order-disorder transformations in the copper-gold, nickel-manganese, and copper-zinc system, and on precipitation reactions in the copper-iron, nickel-beryllium, and copper-beryllium system.' These solid-state reactions are diffusion controlled and the results can be explained largely on the basis of enhanced diffusion due to the presence of irradiation produced vacancies. Reversion of low-temperature eutectoid phases to the high-temperature phase has been observed in the fissionable uranium-molybdenum and uranium-niobium systems.2 Here it is necessary and logical to postulate a displacement spike mechanism to account for the effect. The only case of irradiation induced diffusionless transformation in metals is the allo-tropic transformation of p to a tin reported by Flee-man and Dienes. 3 They found that neutron irradiation markedly decreases the incubation time for transformation, but the mechanisnl is not well understood. Another diffusionless transformation which one might expect to be influenced by neutron irradiation is the martensite transformation. This transformation is known to be sensitive to strain, structural hardening, and the presence of imperfections.6,7 Since the transformation is diffusionless, one also

Jan 1, 1960

By Axel Hultgren

In ingots of silicon-killed carbon steel-made without addition of aluminum, transparent spherical or nearly spherical inclusions, up to about 0.15-mm diameter, are generally present. They may be glassy or crystallized, wholly or partially In steels of high carbon contents they consist of silica,4,12 in steels of lower carbon contents and, in consequence, higher oxygen contents, of iron-manganese-silicate with higher iron and manganese contents the lower the carbon content. Small silicate inclusions, about 0.01 mm or smaller, may occur, just as sulphide inclusions do, in portions that have been the last or nearly the last to freeze,l0,12,17 this location being verified by primary etching.10 Hence, it may be concluded that they have separated from liquid steel that has become enriched in oxygen through segregation during freezing. Silicate inclusions with a diameter larger than about 0.02 mm appear, at least in ingots of moderate size, to be distributed at random in relation to the primary structure and therefore may be assumed to have existed as droplets in the liquid steel before freezing started at the point in question. The average size of such large inclusions has been found to increase toward the axis1' and their amount to increase from the top downward in the axial region.2,10 It has been assumed that this accumulation is a result of the droplets being brought down by the slowly settling metal crystals.2 It has been suggested that large silicate inclusions generally form,6 or may form,'' by coagulation of smaller ones, or by small drops being caught by larger ones during the rising of the latter.9 The occasional occurrence of a large inclusion surrounded by a swarm of small ones has been taken as a proof of this mechanism but, as suggested later, may be explained in another way. It has been pointed out that an inclusion droplet in the liquid steel may grow at the expense of an adjacent one, which implies diffusion of soluble inclusion-forming elements in the liquid steel.'= Further, it has been pointed out that, on cooling in ladle or mold, the metal-silicate equilibrium is displaced in the direction of separation of silicate. Thus inclusions may grow in size and amount.2,3,8,12,13,17,19,20 The exposure of the liquid steel to the oxygen of the air during tapping and pouring, and the consequent reaction, has been mentioned as a source of the formation of inclusions19.20 but this possibility has not received particular attention. Further, the admixture of furnace slag during tapping1 or of refractory material from various sourcesl,20 has been considered. The inclusions due to the latter may be recognized, however, through their resistamce to hydrofluoric acid, imparted by the alumina content. Silicate inclusions are often referred to as the natural and inevitable product of reaction on adding silicon or silicon

Jan 1, 1949

By A. H. Underwood, M. –Ls. Trépanier

INTRODUCTION The King and Beaver mines of Asbestos Corporation Limited (ACL) were joined in 1957 to form the King-Beaver mine. In 1964, ACL purchased the Johnson's Asbestos Company. The combined underground workings are a part of the King- Beaver mine. Asbestos ore is mined at the rate of 4080 metric tons per day (mtpd) to provide 15% of the company's daily requirement in the Thetford Mines area. Figure 1 shows the location of Thetford Mines. Two shafts service the underground. No. 3 shaft, 457 m deep, is used for personnel and material transportation, downcast ventilation, and ore hoisting. No. 4 shaft, 820 m to the south, is 389 m deep. It is used for material transportation and upcast ventilation. Figure 2 shows the relative locations of the shafts and the underground workings. The ore is mined by block caving. HISTORY The orebody outcropped at surface and was first mined as an open pit. Underground mining commenced when open pit mining became impractical. Shrinkage stoping, used initially, was unsatisfactory in the highly fractured ore. Sublevel stoping improved ore recovery and safety, but did not overcome the adverse effects of the un- stable rock. Observations of the rock failures indicated that block caving would be satisfactory. This was confirmed by careful investigation and testing. A method of block caving similar to that used at the Miami Copper Company was adopted. The first block commenced production in June 1933. BLOCK CAVING General A large block of ore is undercut above pre- pared drawpoints. The ore above the undercut caves by its own weight. Caving progresses up- ward as the caved ore is drawn. A large proportion of the rock breaks up in the caving action so that it will pass through the drawpoints. Oversize blocks are broken by drilling and blasting or mudcapping. Figure 3 shows the progressive stages of block caving. In the central part of the ore zone, the vertical dimension of a block is determined by the 132 metres between production levels. When isolated ore pockets are developed, the vertical dimensions of a block depend on the extension of the ore zone above the undercut. The minimum horizontal dimension of an under- cut determines the cavability of the overlying block. When necessary, caving is assisted by drilling long horizontal holes into the block above the undercut and blasting them. This procedure speeds up or restarts caving. It will not ensure complete caving of blocks having undercuts less than 30 m wide. Orebodies suitable for block caving are highly fractured and wide enough to cave freely without excessive dilution. Undercutting must

Jan 1, 1981

By G. K. Manning, A. R. Elsea, A. B. Westerman

INTRODUCTION A CONSIDERABLE number of high-temperature alloys, that is, alloys which have load-carrying ability at elevated temperatures, have been developed on an empirical basis. In order to determine why these alloys possess this particular ability, a project was started at Battelle Memorial Institute on the fundamental factors promoting high-temperature strength of alloys. At the time of inception of this study, the cobalt-base alloy, Vitallium, was widely used because of its good high-temperature properties. One phase of this study was to include an investigation of the structures observed in Vitallium after various heat treatments. However, before investigating the structural features of Vitallium, it appeared logical to examine the phase diagrams of the related binary systems in order to develop information as to which phases- might be present in Vitallium. A preliminary survey of the pertinent alloy systems was made, with particular emphasis on the cobalt-chromium system, which is the base for the Vitallium series of alloys. The early work of Lewkonja1 and Guertler2 indicated that cobalt and chromium are completely miscible in the liquid state, and, within a narrow temperature range, in the solid state; that the liquidus shows a minimum; and that the alloys containing 45 to '85 pct chromium undergo a transformation at 1200 to 1250°C. More recently, Wever and Haschimoto,3 Wever and Lange,4 and Matsunaga5 carried out experimental work which led to the diagrams shown in Fig I and 2. Both diagrams are based on thermal analysis, metallographic, magnetic, and dilatation data; in addition, X ray diffraction results were used by Wever and Haschimoto, and electrical resistance measurements by Matsunaga. It is apparent from a comparison of the two diagrams that there was a difference of opinion as to the locations of the boundaries of the various phases, and as to the existence of a phase containing approximately 47 pct chromium. With these exceptions noted, it might be pointed out that the two diagrams are similar, and that, in general, any particular reaction which is shown on both diagrams occurs at a lower temperature on the Matsunaga diagram than on that of Wever-Haschimoto-Lange. Thus, in view of the gaps in the data and the questionable features in the published diagrams, it appeared desirable to start the fundamental study of Vitallium by determining the diagram for the cobalt-chromium system. EXPERIMENTAL WORK For the convenience of the reader, the diagram as determined by the present investigation is presented in Fig g. The liquidus and solidus data for the low-chromium alloys were taken from the work of Wever and Haschimoto,3 while the, melting point for pure chromium is the

Jan 1, 1948

By L. S. Darken

Measurements of the electrical conductivity, the thermal electromotive force, and the deviation from stoichiometry by thermogravimetry were made on ferrous oxide (wüstite) single crystals as well as on poly-crystalline samples in the temperature range of -900" to 1330°C. In the region defined by T 2 1000°C and 1.05 5 O/Fe 5 1.10 the electrical conductivity was shown to be proportional to the sixth root of- the oxygen partial pressure and to be linearly dependent on the deviation from stoichiometry. The absolute Seebeck coefficient of single and polycrystalline wüstite was shown to be independent of temperature in the above region and a p to n transition of the Seebeck coefficient was observed at O/Fe = 1.09 in polycrystalline wüstite. Analysis of the results showed that the predominant defects in wüstite in the above region were doubly ionized cation vacancies. The analysis also ruled out the existence of complexes of the type postulated by Roth, except that below 1000°C there is some evidence of 'association of defects. Of the three oxides of iron (FeO, Fe3O4, Fe2O3) only FeO (wüstite) can exist within a wide range of the ratio O/Fe. This has been shown by several investigations of the phase diagram using for instance chemical analysis of quenched samples'-4 or thermogravimetric measurements. 5,8 Darken and Gurry's1 data for example in the temper- ature range 1100o to 1400°C show that the lowest O/Fe ratio obtainable for FeOl +. is about 1.045 and that the existence range extends to O/Fe = 1.20 at 1400oC. On the basis of detailed analysis of thermogravimetric measurements made by Vallet and Raccah6 the existence of three varieties of wüstite at elevated temperatures was suggested by these authors and by Kleman. It is well-established that wüstite has a NaCl structure with vacant cation sites.3, 8 From a neutron-diffraction study of wüstite samples quenched to room temperature Roth9 concluded that defects in wüstite are mostly complexes composed of two cation vacancies and one interstitial cation in a tetrahedral site. The atomic arrangement in the vicinity of such a defect is similar to that in magnetite, where each occupied tetrahedral site is surrounded by four vacant octahedral positions disposed at the corners of a tetrahedron, so arranged that each vacancy is shared by two tetrahedral cations. Thus there is an average of two vacancies per interstitial. Recent X-ray diffraction measurements made by Smuts10 on quenched wüstite samples confirmed the results of Roth's neutron-diffraction experiments. Roth's model of defects in FeO can be regarded as a solid solution of Fe3O4 in FeO; salmon1' used this defect model to calculate the activity of Fe3O1 in FeO at high temperature and claimed good agreement with the results obtained experimentally by Darken and Gurry.' Each missing cation in wüstite is accompanied by two trivalent cations, thus preserving charge neutrality. The "holes" which differentiate between the Fe 2+ and Fe3' states of the ions in octahedral positions can overcome the attractive field of the cation vacancies, and are then free to move through the crystal as charge carriers. A similar picture holds for COO, NiO, andMnO.l2, 13 Localization of the charge carriers in transition

Jan 1, 1968

By C. Laird, C. E. Feltner

An investigation into the effect of a material's slip character on its high strain fatigue properties has been carried out using copper and Cu-7.5 wt pct A1 as representative wavy and planar slip mode materials, respectively. In copper, the constant term and exponent in the Coffin-Manson equation, ?EpNf1/2 = const, are unaffected by the temperature of testing or the history of the metal. On the other hand, in Cu-7.5 wl pcl Al, while the exponent is invariant, the constant is increased with respect to copper and increased further with decrease in temperature. The life in Cu-7.5 wt pct A1 is double that of copper at room temperature and about four times at 78°K. Fractoflaphic studies indicate that these diffevences in life are associated with the early stages of cracking. In copper, severe surface rumpling provides crack nuclei at pain boundaries independent of' testing temperature or the history of the metal. In Cu-7.5 wt pct Al, the planar slip mode of deformation prevents grain boundary folding; crack nucleation is therefore mainly transgranular and the increase in life with respect to copper is associated with a prolonged linking-up stage of such transgranulay cracks. At the lower temperature the increased difficulty of cross slip inhibits multiplication of trans-granular cracks and retards the linking-up process, giving a further enhaancement in life. It is expected that such behavior will be found general in materials of wavy and planar slip mode, respectively. STAGE II crack propagation in high strain fatigue takes place by a repetitive process of crack tip blunting in the tension part of a fatigue cycle followed by resharpening of the crack in the compression part.1-3 This mechanism has come to be called the "plastic relaxation process".3 At very high strains (lives to failure less than 103 cycles), most of the life of a specimen is spent in crack propagation by the plastic relaxation process.4,5 Since the differences in the cyclic strain hardening behavior of metals and alloys are comparatively small in this strain range,6,7 the amount of blunting at a crack tip depends primarily on the applied strain amplitude. The rate of crack propagation in this life range is therefore nearly constant with respect to material and temperature of testing. In the longer life portion of the high strain fatigue range (l03 to 103 cycles to failure), a considerable fraction of the lives is spent in crack initiation and Stage I growth. Grain boundaries have frequently4,5,8 been found to be the sites for this growth which is also associated with a surface folding process.4'5 Such folding behavior may also be general in any material having a grain structure and capable of plastic deformation. It has therefore been suggested475 that the universality of the Coff in-Manson law,9'10 constant,* where ?Ep is the plastic strain range and Nf the number of cycles to failure, is reflected in the approximate invariance of both the crack initiation and propagation behavior. It happens that, in most investigations of the phenomena associated with failure,4,5,8 materials with a wavy slip mode have been used as vehicles for the studies, or else "normally" planar slip mode materials have been tested under conditions giving rise to wavy glide deformation, e.g., stainless steels cycled at high temperature.'' The purpose of the present investigation was to determine whether or not different cracking phenomena would occur when the slip mode of a material is planar and if the associated life to failure would be different than that of a wavy slip

Jan 1, 1968

By D. J. McAdam, R. W. Mebs

The technical cohesive strength of a metal means, not the interatomic forces, but the technically estimated resistance to fracture. An example of such resistance to fracture is the so-called "true" breaking stress of a tension-test specimen, the breaking load divided by the sectional area at fracture. According to the prevalent view,8 the technical cohesive strength of a metal. in any particular state as regards prior mecharlical treatment and heat-treatment, is determined by a limiting value of the algebraically greatest principal stress. In two previous papers by one of the authors, evidence was presented to indicate that the technical cohesive strength of a metal cannot be represented by a single stress value, but that it comprises an infinite number of values corresponding to the infinite number of possible combinations of the principal stresses. The technical cohesive strength thus comprises an infinite number of technical cohesion Limits, each representing a specific stress combination at fracture. In the same papers, evidence was presented that plastic extension causes a continuous increase in the technical cohesive strength. Such a variation is not in accordance with the prevalent view that plastic extension first increases, then decreases, the technical cohesive strength.' In this paper, as in the previous papers,14.15 the algebraically greatest principal stress will be designated S1, the least principal stress will be designated 33, and the intermediate principal stress will be designated S². The technical cohesion limit under polarsymmetric tension (S1 = S2 = S3) will be termed the "disruptive stress," and the cohesion limit under unidirectional tension (S2 and S3 = o) will be termed the "severing stress."l4, .l5 The yield strength of a metal, according to the generally accepted theory, may be represented by a constant value of the shearing-strain energy, and hence by a constant value of the sum of the squares of the three principal stress differences. In the second of the previous papers,15 however, evidence was presented that the limiting yield stress decreases with increase in the volume stress, the algebraic average Of the three principal stresses.* This in fluence of the volume stress becomes prominent only in the field of triaxial tension. In the two previous papers,".'5 diagrams were presented to show the variation of the technical cohesion limit, yield stress, and ultimate stress with the combination of principal stresses and with plastic extension. Those diagrams and the resultant • Any stress system ma? be resolved into a polarsymmetric stress causing pure volume strain and a combination of pure shearing stresses. which cause no change of volum

Jan 1, 1945

By F. P. Bullen

Empirical relationships between fracture stress, orientation angle, and diameter of crystal have been determined at 77°K. Orientation ranges of markedly different behavior were found—a law of constant normal stress&apos; of value a (diameter)-1/2 for the fracture of ductile crystals, and a condition of shear stress (or strain) for more brittle crystals. The observations are not consistent with current theories. An interpretation is advanced which is also applicable to observations on the effect of prestrain at room temperature on the subsequent fracture stress at 77°K and to the effect of cyclic stressing on the cleavage strength.&apos;&apos; The law of constunt normal stress&apos; and the brittle ductile transition are also explained. The interpretation is more consistent with the initiation of cracks by intersecting dislocations than with theories based on stress -concentration by dislocation arrays. ZINC single crystals are particularly suited to the study of cleavage because fracture occurs on the basal plane over a wide range of crystal orientations. Analysis of the conditions of stress and strain at fracture in crystals of different orientations should indicate which parameters control the cleavage process. Unfortunately, controversy has arisen over the correct empirical relationship between tensile fracture stress and orientation. Schmid&apos;s observations1,2 favored a &apos;law of constant normal stress&apos;, as observed in other materials.2 For zinc, however, the observed values are far below the theoretical strength and cannot represent the true limit of cohesion between neighboring atomic planes. Hence, the interpretation of such a &apos;law&apos; is not straightforward. Deruyttere and Greenough3,4 found a complex variation between tensile fracture stress and orientation; this variation did not agree with a &apos;law of constant normal stress&apos;. Two theories have been advanced to account for their observations: a) the propagation of cracks from low-angle boundaries,5 and b) the release of energy from piled-up dislocations during crack-propagation.4 The present work resolves the apparent discrepancy between the observations and shows that neither of the above theories are applicable to the tensile fracture of zinc single crystals. A phenomenological explanation, along the lines suggested by Gilman,&apos; is advanced and successfully applied to previously unexplained effects. EXPERIMENTAL DETAILS &apos;Crown Special&apos; redistilled zinc was used, except for one comparison series op tests using &apos;Tadanac&apos; electrolytic zinc. Crystals of 1 mm diam, subsequently called &apos;1 mm crystals&apos;, were grown from the melt in vacuo in precision-bore Pyrex tubes internally coated with graphite. Several specimens 1 in. long were cut from each crystal and chemically polished. Jigs were used to minimize handling strains, and crystals were mounted in the Polanyi machine the day prior to testing to allow recovery from any such strains. One-mm crystals were chemically polished for long periods to obtain 0.1 mm (approx) crystals. One-mm crystals were cemented into miniature gimbals by &apos;Araldite&apos; casting resin. The Appendix gives the reasons for using gimbals and the results obtained by other methods. More complete details of all techniques are given elsewhere.7 The symbols and terminology used are as follows: X = orientation angle (angle between tensile axis and line of greatest slope in the basal plane). T = tensile stress (on true cross-section) S,N = shear and normal stress (components of T with respect to the basal plane) ? = shear strain D = crystal diameter. The subscript &apos;f&apos; will be used to denote values at fracture. PART I-ANNEALED CRYSTALS EXPERIMENTAL OBSERVATIONS One-mm crystals were used to establish the variation of fracture stress at 77°K with orientation at fracture (Xf), Fig. 1. For 18 deg = Xf = 55 deg, a &apos;law of constant normal stress&apos; was observed. For Xf > 55 deg, the fracture condition approximated to a constant shear stress. At Xf< 18 deg, twinning occurred before fracture so that the results were not typical of homogeneous single crystals,4,8— such specimens will not be considered herein. The dependences of fracture stress upon Xf were of similar type for 6 mm,* 1 mm, and 0.1 mm crystals,

Jan 1, 1963

By F. Keller, A. H. Geisler

Some of the important changes that take lace in the structure of aluminum alloys are largely submicroscopic in character. This is especially true of the changes that accompany age-hardening and recrys-tallization. Although improved metal-lographic practices have been helpful in indicating some of these changes indirectly, results are limited by the resolving power of optical microscope lenses. The electron microscope, however, provides a new and important means for investigating the fine structure of metals in a range not possible heretofore. It is anticipated that this microscope will yield new and useful information to metallurgists when suitable techniques are developed to utilize its very high resolving power. Several unique methods have been devised to permit the examination of the structure of opaque samples by the electron microscope. These are all based 011 the principle of producing a very thin film that represents the prepared surface of the opaque metal sample, which can be examined in the transmission microscope; however, the available methods fall into three categories, depending on the manner in which the film is obtained. With the oxide-film method developed by Mah13-l7 a very thin surface layer of the metallographic sample is converted into oxide by a thermal, chcmical or anodic treatment. .After removal from the metal sample by one of the special techniques described by Evans,l9 this oxide film is examined in the electron microscope, and is found to portray the structure of the original metal surface. Films formed on aluminum,8,l7 nicke1,7 iron3,8,9 and a nickel-beryllium alloyl1 by heating; films formed on iron8,9 by chemical action; and films formed on aluminum3-&apos;" by anodic oxidation, have been examined. The other two are more recent methods developed principally in this country; they consist in making the thin film in the form of a mold or replica of the surface contour of an etched sample. With the negative replica method the thin film is made directly on the sample; with the positive replica method it is made on a negative reproduction of the sample. With both methods, variations in the contour of the metal sample, caused mainly by etching, produce variations in the thickness of the surface replica. Because of the direct dependence of the amount of scattering of the illuminating electrons upon the thickness of the film penetrated, only the variations in thickness of the replica are responsible for tones in the electron image. Negative replicas are made by forming a thin film of some suitable material directly on the polished and etched metal sample, and derive their name from the fact that after removal from the sample they have the high and low points in the sample reversed, as shown in Fig. I. Nega-

Jan 1, 1944

By H. Ries, R. E. Somers

Bodies of pyritic ore in schistose rocks have long been known in different parts of the world. The several occurrences resemble each other in being usually of more or less lenticular shape, inclosed in walls of schist or gneiss, and carrying pyrite, chalcopyrite, and pyrrhotite in varying proportions, as the chief ore minerals. The deposits differ, however, in that sometimes one, sometimes another, of the three sulphides mentioned may predominate. The orebodies in general are more or less closely conformable with the schistosity of the wall rocks, although in cases this may not hold true. The boundaries may or may not be sharp, but in general are fairly distinct. The lenses vary in size, may occur singly, or in groups; in the latter case they may be in line, overlapping, or occasionally parallel. Pinching and swelling of individual lenses is not uncommon. The genesis of these pyritic bodies has provoked much discussion, and they have been variously classed as sedimentary, igneous intrusions, hydrothermal replacements or impregnations, etc. Indeed, in some cases different geologists have most positively assigned widely diverse origins to the same deposit, as in the case of Rio Tintol and Rammels-berg.2 It is true that in the light of modern criteria, we have been able in a few cases to decipher clearly the process of origin and so there seems no. doubt now regarding the genesis of the deposits at Rio Tinto, Spain;3 Rammelsberg, Germany;4 Ducktown, Tenn.;5 southwestern Virginia;=

Jan 1, 1918

By Thomas O. Ziebold, Robert E. Ogilvie

Experimental analysis of diffusion samples covering the entire Cu-Ag-Au system at 725°C has been carried out. Experimental coefficients are computed at twenty-eight points in the lermary system. It is found that the direct coefficienl for copper is nearly independent of the silver, content for 1ow-silver alloys and that the cross coefficient for silver correlates with the ther-nzodynatnic properlies. This lalter observalion leads to the conclusion that the mobililies of copper and gold are nearly equal. Thermodynarmic actiuilies for Cu-Ag-Au were computed from binary data and adjusted to be consistent with neasured tie lines across the two-phase region which extends into the ternary diagram at 725°C. Application of these therknodynarnic calculaliorls to the dijftcsion data terifies the Onsager reciprocal relations to the extent that experimental uncertainlies will allow. It is jound that diffusion neav the critical point of the 1200-phase field in the isothermal section at 725°C causes the composition gradients for copper and silver to approach infinity while the direct and cross coeficients become equal. From the standpoint of experimental Procedures, it has been shown that microanalysis of ternary samples may be carried out using a simple expression for the conversion of X-ray zntensily to composilion. The accuracy of&apos; this method has been demonstrated by the analysis of standard alloys and by the reproducibility of composition determinations using differed X-ray lines. DIFFUSION in a ternary alloy system is basically and significantly different from diffusion in binary alloys. When only two elements are present there is just one independent composition variable, and the diffusion profile of composition vs distance across a sample must change monotonically through all solid solutions bracketed by the terminal alloys. If two-phase regions are stable for compositions between the terminal values, these must appear as planar interfaces in the diffused sample, with the composition profile exhibiting sharp discontinuities. Adding a third element may alter this picture completely. Because of the additional degree of freedom in the composition, we may see composition profiles which do not change monotonically through solid-solution ranges, and stable, nonplanar phase interfaces. Even though two-phase regions may lie between the terminal compositions of a diffusion sample on the constitution diagram, they will not necessarily appear in the diffusion profiles. In binary diffusion only one coefficient is needed to describe the interdiffusion process. From Fick&apos;s law this coefficient relates the diffusive flux of one component to its own composition gradient. In ternary diffusion we must allow for cross interaction between the two independent species. By a linear extension of Fick&apos;s law we introduce "cross" or "off-diagonal" diffusion coefficients which relate the flux of one component to the composition gradient of the other independent constituent. Thus, we require four interdiffusion coefficients to describe the transport process in a ternary system. In its most general form,1,2 the rate equation for mass transport by isothermal diffusion in a system of s components states that the diffusive flux of component i, Ji, is a linear summation of all driving forces, Xk, each force being multiplied by an appropriate compliance coefficient:

Jan 1, 1968

By V. L. Elliott, C. L. Meyette

The conditioning of semifinished steel products such as billets, blooms, and slabs to remove surface defects before further processing to finished products is a necessary accompaniment to steel mill rolling operations. A knowledge of the kinds of defects which occur and of the conditions in melting, teeming, and rolling that may lead to their occurrence is favorable to the production of quality products. Surface defects may be classed as either of mechanical origin such as those that occur in the heating and rolling operations and include breaks, tears, guide marks, mechanical laps and the like, or they may have been in the ingot before heating and rolling. Examples of the latter are ingot cracks, blow-holes, scabs, and inclusions which lead to the numerous known types of seams and surface flaws. Most of the common types of surface defects and methods for their identification by visual examination or by the aid of macro-etch tests are well known. The conditions which caused their formation are also appreciably understood and this knowledge has been instrumental in the establishment of methods for metallurgical control. There are, however, defects of doubtful origin. In any endeavor to trace the causes for their occurrence, the question of first concern is: I. Whether they were formed in the rolling procedure. 2. Whether they existed in the ingot prior to rolling. In the case of semifinished products which are rerolled to finished products there is also the question of whether they were formed in the primary or in the secondary rolling. In metallurgical laboratories where numerous or routine examinations of the quality of rolled products are made over the course of time, there will doubtless be many puzzling cases of defect origin. Microscopic examination of the defects may, however, reveal a characterizing condition that can be used to classify them broadly with respect to their origin. Such a condition has been observed and, after detailed investigation, is described here with the thought that it may offer a usable and relatively simple solution to a bothersome problem. Basis For Classifying In this discussion the term "mechanical type defects" will be used to designate those that have occurred as a result of the rolling operations such as in rolling from the ingot to billets, blooms, or slabs or in rerolling to finished products. The term "steel type defects" will be used to designate those that were present in the ingot prior to primary rolling or that were in the semifinished product prior to secondary rolling to finished product. The basis for distinguishing between the two types of defects lies in the relative degree of subscale formation surrounding them. Subscale is defined1 as a zone of

Jan 1, 1949

By N. D. Shutler

This paper describes a numerical mathematical model of the steamflood process that depends on fewer restrictive assumptions than models previously reported. The solution, however, is obtained economically. Example calculations are presented that, on comparison with experimental results, tend to validate the model. Results that expose certain process mechanics are discussed. The model describes the simultaneous flow of three phases — oil, water and gas — in one dimension. It includes the effects of three-phase relative permeabilities, capillary pressure, and temperature- and pressure - dependent fluid properties. Interphase mass transfer of water-steam is allowed, but the oil is assumed nonvolatile and the hydrocarbon gas insoluble in the liquid phases. The model allows heat convection in one dimension and two-dimensional heat conduction in a vertical cross-section spanning the oil sand and adjacent strata. The hydrocarbon-steam gas composition is tracked, but the effect of gas composition on water-steam phase behavior is neglected. The model is solved numerically in three separate stages. The three-phase mass balances are solved simultaneously using Newtonian iteration on nonlinearities occurring in the accumulation terms. The energy balance is solved separately by noniterative application of the alternating-direction implicit procedure. Separate solution of the composition balance is accomplished by straightforward solution of the finite difference equations. The method of effecting nonsimultaneous, stable solution of the mass and energy balances is the key to the success of the model. INTRODUCTION Mathematical tools as well as laboratory and field experiments are necessary to help us understand the complex steamflood process. A mathematical model can expose process mechanics and show the relative importance of process variables, but this ability is often limited by restrictive assumptions. Most known models of steam processes,l-5 with the exception of the model of Gottfried,6 are "simplified" in that they involve analytic approximations and require many restrictive assumptions. The primary utility of these methods lies in the routine use as an aid in engineering design. By contrast, the comprehensive model presented here is numerical and requires far fewer restrictive assumptions. It finds its primary utility as a research tool. It serves as an aid in understanding the nature of the process, in interpreting laboratory experiments and in evaluating and developing simpler mathematical models for engineering design. The major reason why previously presented models have been confined to the "simplified" class is evidenced by the one published exception. In 1965, Gottfried6 presented a numerical model for the combustion process of which the steam process is a subset. The result is a comprehensive tool (though it neglects capillarity and two-dimensional heat conduction) that is troubled with convergence problems and that requires 2 to 3 hours of IBM 7094 time to complete a calculation. Though our present model does not simulate combustion, it does consider capillarity and two-dimensional heat conduction and it overcomes the convergence and computer-time problems. MATHEMATICAL DESCRIPTION OF STEAMFLOODING FLUID FLOW The equations employed to describe three-phase fluid flow are of a familiar form. Darcy&apos;s. law provides expressions for the velocities of the three phases (oil, water and gas), which, when combined with oil, water and gas mass balances, give the partial differential equations governing flow of the

Jan 1, 1970

By Dr. O’Neil Thomas J., Donald W. Gentry

"Who is the man who views the mines and promptly turns them down? Who is the one that thinks this is the short cut to renown? Who is it gives the bum advice to the innocent financier? The knowledge-feigning, theory- straining mining engineer." -Anonymous INTRODUCTION Taxes levied against mining properties and operations are a critical cost in the economic evaluation of mining investments. Indeed, taxes represent a substantial cost of doing business in the minerals industry and often have a significant impact on corporate investment decisions. A good example was the postponement of mineral development in the state of Wisconsin, primarily because of what was perceived to be excessive taxes imposed by that state. In many respects mining investments are no different from other industrial investments. Astute taxing authorities should recognize that geologic endowment is a necessary but clearly not a sufficient condition for mineral investments. The fact that a mineral deposit exists does not necessarily mean that it will ultimately be developed-a point that many taxing authorities fail to recognize. While it is true that ore deposits are not mobile in the sense that they cannot be physically moved to a district having more favorable taxes, corporate investment capital certainly is mobile and flows to ventures which maximize wealth to the firm. In short, higher taxes reduce project yields and tend to drive investment capital elsewhere. The appropriate type and level of taxation imposed upon the mining industry continues to be a very controversial and emotion-charged topic. Mining activities have been taxed at various levels over the years due to widely differing taxation philosophies. Location has also influenced taxation policies, as evidenced by the diversity of tax laws and assessment procedures applied to mineral deposits by the various states. Indeed, mineral taxation varies from state to state, from county to county, and from one mineral commodity to another. Whether or not a given tax is appropriate for a specific mineral deposit, or even an industry, is a difficult problem to assess. First, the type or kind of tax which should be imposed must be considered. Second, it is most important to define a fair and unambiguous base against which to levy the tax. Finally, one must consider the tax rate to be applied to this base. It is the combination of these two components

Jan 1, 1984

By A. Gatti, R. L. Fullman

A very fine dispersion of aluminum oxide is produced by internal oxidation of solid-solution alloy of 0.14 pet A1 in Ag. The particle size of the aluminum oxide is approximntely 50 to 100A in radius. The yield strength of the alloy is increased markedly by internal oxidation. A further increase in strength is produced by cold working the internally oxidized alloy. Recrystallization is retarded by the finely dispessed aluminum oxide particles, so that the strength increase resulting from cold work is retained on annealing at temperatures 14 to about 700°C. MANY workers&apos;-3 in the past have studied various aspects of the internal oxidation of aluminum-silver alloys. This paper is an extension of these studies with emphasis placed on the effect of time and temperature of annealing on the strength of these alloys after oxidation and subsequent cold working. Two general conditions are necessary to internally oxidize an alloy. First, oxygen must diffuse through the base material more rapidly than does the addition; otherwise oxidation will take place as a surface layer. Secondly, the affinity of oxygen for the addition must be greater than for the base material. After internal oxidation of certain alloys takes place, a marked increase in hardness accompanied by higher yield stress and improved creep properties is noted, presumably as a result of the highly dispersed oxide within the base material. Meijering and Druyvesteyn1 also noted that the internally oxidized portion of a partly oxidized alloy failed to recrys-tallize under annealing conditions that led to coinplete recrystallization of the unoxidized part. EXPERIMENTAL-METHODS AND PROCEDURES Few alloys can be made to contain a second phase that is extremely stable at high temperatures. Silver plus aluminum in solid solution was chosen for these internal oxidation studies because of the high rate of oxygen diffusion through silver and the very stable nature of aluminum oxide. Two alloys were vacuum cast. The nominal compositions were: Alloy A—1 pct Al, balance Ag; Alloy B—0.1 pct Al, balance Ag. Chemical analysis, which does not distinguish between aluminum and aluminum oxide, showed the conlposition to be: Alloy A—1.6 pct Al, and Alloy B—0.14 pct Al. The ingots were machined for surface cleaning, swaged and drawn to 0.020-in. diam wire. A sample 20 ft long of the 0.020-in. dianl wire of each composition was annealed 24 hr at 800°C in pure dry hydrogen. Each wire was then cut into two equal pieces. Photomicrographs of the 0.14 pct A1 alloy are shown in Fig. 1, the annealed 0.020-in. wire at the left and the oxidized wire to the right. The oxidation treatment for the first set of data was 1000 hr at 800°C in air. After this treatment the 1 pct A1 proved to be brittle. It is assumed that high alunlinum oxide concentration at the grain boundaries was responsible. The 0.14 pct Al wire remained ductile and all further data were derived using this alloy. One-half of this wire, about 5 ft, plus 5 ft of as-homogenized wire, was then drawn cold to 0.005 in. diam. All tensile tests were conducted with an Instron Engineering Corp. tensile-testing machine, Model TT-B. Unless otherwise indicated, the tests were made at room temperature with a strain rate of 0.1 per min. All metallographic samples were etched with an aqueous solution of 2 pct each of CrO3 and H2SO4 . EXPERIMENTAL RESULTS AND DISCUSSION PARTICLE SIZE DETERMINATION A study was made of the particle size of the aluminum oxide produced in the samples of Ag + 0.14 pct Al, oxidized 1000 hr at 800°C. A cross section of the as-oxidized wire was mounted in bakelite, polished, and etched with an aqueous solution of 2 pct each of CrO3 and H2SO4. The specimen was then thoroughly cleaned by stripping successive coatings made by applying 10 pct nitrocellulose in amyl acetate. The final replica of the cross section was made by applying 2 pct nitrocellulose in amyl acetate. The replica was stripped, transferred to a copper screen, shadow cast with chromium at 10 deg and photographs taken using a Phillips Metallix electron microscope at an accelerating potential of 100 kv. A photograph of an etched sample of the as-oxidized material is shown in Fig. 2. We believe the pits in the photograph are places were A12O3 inclusions were sitting in the matrix. By inspection, it appears that the volume fraction ob-

Jan 1, 1960

By F. S. Pettit

The oxidation of Ni-3 to 25 wt pd Al alloys has been studied in 0. 1 atm of oxygen at temperatures between 900° and 1300°C. These alloys have been found to oxidize by three different mechanisms which depend on the temperature of- oxidation and the alloy composition. Two of the three mechanisms do not permit a continuous layer of Al,0, to be formed on the alloy surface and the oxidation rates are greater than that for pure nickel. The third mechanism results in the formation of an external A1203 scale and lke oxidation rates are about three orders of magnitude smaller than those for pure nickel. The minimum amount of aluminum required for the formation of external scales of A L,0, has been determined. NICKEL-base alloys are currently the main source of materials for use at elevated temperatures in gas turbine engines. These alloys are usually coated to obtain oxidation resistance. Coatings on nickel-base alloys are frequently formed by reaction of the alloy with aluminum whereby alloyed nickel aluminides are formed. The alloyed nickel aluminides provide protection to the nickel base alloy because external scales of A1203 (alumina) are formed during oxidation and mass transport through A120, is slow in comparison to mass transport through most other oxides. In view of the protective properties of A1203, it is important to know how much aluminum is required in these alloys in order to form external scales of AlzO,. The present paper is concerned with the oxidation kinetics and the oxidation mechanisms of Ni-A1 alloys and the minimum amount of aluminum required in these alloys for the formation of external scales of Alz03. THEORETICAL CONSIDERATIONS When a Ni-A1 alloy is heated in oxygen at elevated temperatures, the following reactions can take place on the surface of the alloy where the oxide phases are assumed to be virtually pure: These oxide phases are in the form of nuclei scattered over the surface of the alloy and, in view of their rapid formation, they need not be in equilibrium with the alloy. As the oxidation process continues, equilibrium between the alloy surface and the oxide phases is approached and the stability of the oxide nuclei is determined by the composition of the alloy at the alloy/oxide interface because of the following reactions: 3NiA1204 + 2A1 (alloy) = 4AlzO3 + 3Ni (alloy) [41 4Ni0 + 2Al (alloy) = NiA120, + 3Ni (alloy) [ 5 1 Application of the mass-action law to Eqs. [4J and [5J yields the following equilibrium conditions for these reactions: where aA1 and aNi are the activities of aluminum and nickel, are the standard free energies of formation of NiO, A1203, and NA1204, respectively, R is the gas constant, and T is the absolute temperature. If the composition of the alloy at the alloy/oxide interface is such that (akl/ahi) is greater than the equilibrium values defined by Eqs. [6] and [7], then Reactions (41 and [5] will go to the right as written. Conversely, if the alloy composition is such that the activity ratio (aLl,/aki) at the alloy/oxide interface is less than the equilibrium values, then Reactions [4] and [5] will proceed to the left. The equilibrium activity ratios in Eqs. [6] and 171 can be calculated since values for the standard free energies of formation of the oxide phases are available. Standard free energies of formation for NiO and A1203 have been tabulated by Elliott and ~leiser.&apos; The standard free energy of formation for NiA1204 can be obtained from the data of Tretjakow and Schmalzried.&apos; The results of these calculations are tabulated in Table I. Table I shows that the following inequality is valid over the temperature interval 900" to 1300°C: (Reaction [5]) (Reaction [4]) « 1and therefore the aluminum activity for these compositions can be taken as equal to the square root of the activity ratios (i.e., aNi = 1). If equilibrium is estab-

Jan 1, 1968

By R. L. Ash, R. R. Rollins, C. J. Konya

An investigation was performed to determine conditions for optimizing the spacing of simultaneously initiated multiple explosive columns. This was done by using models of mortar, dolomite, and Plexiglas with 10-grain mild detonating fuse as the explosive charge. It was desired to simulate blastholes with multiple primers initiated by detonating fuse or when high-velocity explosives are used in low-velocity materials. It was found that optimum spacing between multiple charges was strongly influenced by charge length. At less than optimum charge length, the spacing at which complete shearing was possible between adjacent charges decreased exponentially with a subsequent loss of broken material volume. For charges fired simultaneously, larger burdens and spacings were possible as compared to those necessary for single-crater charges. For each material studied, there was a characteristic optimum charge length and a maximum attainable spacing at any given burden. Proper selection of the spacing distance between charges is fundamental to successful blasting. Its value directly affects the cost of drilling and explosives used per unit of broken material. In addition, the choice of a spacing that is Compatible with a given set of blasting conditions aids in the control of fragmentation sizing, ground vibrations, overbreak, and throw which in turn, influence other production costs. For example, normally loaded blastholes that are spaced too closely invariably promote overbreak and usually give coarse fragmentation. Unless care is taken, airblast and violent flyrock will occur and under certain conditions cutoffs and misfires may result. Too large a spacing, on the other hand, frequently leads to conditions that form bootlegs or toes. The choice of a particular spacicg to use, however, is largely a matter of individual experience and judgment, usually based on trial and error. Very little is known or can be found in the literature with regard to how the spacing between charges is related to field conditions and charge geometry. As a general rule, the firing time sequence of adjacent charges and properties of a material are thought to have the most significant influence on the spacing distance best suited for any given field condition. For example, delayed initiation of adjacent charges usually always requires a closer spacing than when charges are fired at the same time. This should be expected if one considers that the energy normally dissipated and lost in the surrounding ground from charges fired independently would be captured and utilized for breaking material between charges when they are initiated together. Spacing can be extended also when charges are aligned with structural planes of a material, such as jointing, along which shearing is relatively easy. It is customary to relate the spacing (S) between charges to their common burden (B) in the form of a spacing ratio, or SIB. The burden normally is considered as the optimum depth or distance from any single charge perpendicular to the nearest free or open face at which the desired fragmentation and maximum crater yield are obtained. For production blasting, value of the ratio is generally considered to vary from 1 to 2, depending on conditions.1-6 When adjacent charges are fired independent of one another, the value varies from 1 to about 1.4, the closer amount being employed to square corners or produce craters having the ideal 90" apex angle. The larger ratio is the geometric balance value for craters having an apex angle of 135". The basic ideal crater forms in the plane of the charge diameter for charges fired independently are shown in Fig. 1. In the event charges are fired simultaneously, geometric balance in the plane of their charge diameters suggests that a spacing ratio near 2 would be appropriate, as illustrated by Fig. 2. In practice, however, some compromise ratio value must be selected to conform with the specific ground conditions. An example would be where the jointing planes tend to produce 60° or 120° crater angles, the appropriate geometrically balanced charge arrangement being given by Fig. 3. In this condition, the spacing ratio is 1.15, not 1 or 1.4 as suggested for the 90° cratering of independently fired adjacent charges. In view of the foregoing, it would seem logical to assume that whenever charges all having the same burden are fired at the same time, spacing distances always can be greater than those permitted by charges fired independently. In practice this is not the case, however.

Jan 1, 1970

By V. L. Elliott, C. L. Meyette

The conditioning of semifinished steel products such as billets, blooms, and slabs to remove surface defects before further processing to finished products is a necessary accompaniment to steel mill rolling operations. A knowledge of the kinds of defects which occur and of the conditions in melting, teeming, and rolling that may lead to their occurrence is favorable to the production of quality products. Surface defects may be classed as either of mechanical origin such as those that occur in the heating and rolling operations and include breaks, tears, guide marks, mechanical laps and the like, or they may have been in the ingot before heating and rolling. Examples of the latter are ingot cracks, blow-holes, scabs, and inclusions which lead to the numerous known types of seams and surface flaws. Most of the common types of surface defects and methods for their identification by visual examination or by the aid of macro-etch tests are well known. The conditions which caused their formation are also appreciably understood and this knowledge has been instrumental in the establishment of methods for metallurgical control. There are, however, defects of doubtful origin. In any endeavor to trace the causes for their occurrence, the question of first concern is: I. Whether they were formed in the rolling procedure. 2. Whether they existed in the ingot prior to rolling. In the case of semifinished products which are rerolled to finished products there is also the question of whether they were formed in the primary or in the secondary rolling. In metallurgical laboratories where numerous or routine examinations of the quality of rolled products are made over the course of time, there will doubtless be many puzzling cases of defect origin. Microscopic examination of the defects may, however, reveal a characterizing condition that can be used to classify them broadly with respect to their origin. Such a condition has been observed and, after detailed investigation, is described here with the thought that it may offer a usable and relatively simple solution to a bothersome problem. Basis For Classifying In this discussion the term "mechanical type defects" will be used to designate those that have occurred as a result of the rolling operations such as in rolling from the ingot to billets, blooms, or slabs or in rerolling to finished products. The term "steel type defects" will be used to designate those that were present in the ingot prior to primary rolling or that were in the semifinished product prior to secondary rolling to finished product. The basis for distinguishing between the two types of defects lies in the relative degree of subscale formation surrounding them. Subscale is defined1 as a zone of

Jan 1, 1949