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By B. R. Berger

The Getchell area in the northern Osgood Range, Humboldt County, Nevada, was first prospected in the late 1800s for silver, copper, and gold; the first recorded mention of mining activity concerned a silver discovery. The September 17, 1874 edition of the Winnemucca Silver State notes that the Crystal Mining District was newly organized in the Fremont Range (as the Osgood Range was then called) to include discoveries of silver, lead, and antimony made on the Louisa claim. The name Crystal was derived from the plentiful quartz crystals found in-cavities near the prospects. This same area, occurring on the northwestern margin of the Osgood granodiorite stock, was later mined as the Richmond tungsten property, and crystals may still be found there in cavities in a massive garnet tactite. By 1878, how-ever, the name Crystal had been abandoned in favor of Potosi. The next flurry of activity in the Osgood Range was reported in 1899 when a copper strike on Granite Creek in the southern portion of the range received attention. The copper occukences were in a granite contact zone on the southeastern edge of the Osgood stock and, in 1916, the first tungsten in the district was discovered in the same deposit. During the last part of the nineteenth century, cattle and sheep were brought into the Kelly Creek Valley east of the Osgood Mountains by various ranchers. Alphonso and Adrian Bernard were among these ranchers, and in the first decade of the twentieth century, one of them recognized the altered rock where the Getchell Center it is now located. A 3-m (10-ft) deep prospect pit was dug on the outcrop, but the highest assay ran about $11/t ($10 per st). No further work was done on the gold-bearing rock. In 1934, Emmet Chase of Battle Mountain and Ed Knight of Golconda prospected in the area of the Bernard prospect pit and located several lode claims. At this time, Noble H. Getchell, a state senator, was looking for new mining opportunities to develop. He held interests in the Gold Circle Consolidated at Midas, which was becoming a losing venture. Chase and Knight brought their discovery to his attention, and he grubstaked them to prospect further in the area. In need of financial backing, Getchell solicited the help of George Wingfield, Sr., who held a minor interest in the Gold Circle. Wingfield was a successful banker and mine operator during the heyday of Tonopah and Goldfield, and he had become good friends with Henry C. Frick and Bernard Baruch. Wingfield let Baruch in on several very successful mining ventures in Goldfield, and Baruch returned the favor by underwriting the financing of the Getchell mine. Eleanor Roosevelt and Jack Dempsey were early stockholders in the Getchell operation. George Wingfield brought his brother-in-law, Roy A. Hardy, a consulting engineer who had constructed the American Flat mill near Virginia City, to look at the property. Upon inspecting the Chase-Knight claims, Hardy declared the whole area to be mineralized. A crosscut was driven to check the extent of the ore, and the results were better than anticipated. The locators were reportedly given $50,000 for the claims. The Getchell Mine, Inc., was organized on November 5, 1936. A small railroad that ran from Battle Mountain to Austin was purchased, and the ties and rails were moved to Getchell to be used in underground exploration. Enough reserves were found in the Center and South pit areas to construct a 40 t/d (44 stpd) mill. The financial and engineering help for the mill was provided by the Newmont Mining Co. for a 17 1/2% interest in the corporation. Newmont set about trying to solve the metallurgical problems with the sulfide ore, and also directed the early development work. Newmont wanted to grind the sulfide ore underground and oxidize it in the atmosphere by spreading salt on it. Detrital Flat in Arizona was purchased to provide salt for

Jan 1, 1985

By Nicholas J. Grant, John S. Benjamin

A series of composite alloys containing a high volume of NiAl, Ni3Ah or CoAl, bonded with 0 to 40 vol pct of a ductile metal phase, were prepared by powder blending and hot extrusion. The binder metals were of four types: pure nickel or cobalt, near saturated solid solutions of aluminum in nickel and cobalt, type 316 stainless steel, and niobium. Sound extrusions were obtained in almost all instances. Studied or measured were the following: interaction between the alunzinides and the binders, room-temperature modulus of rupture values, 1500° and 1800°F stress rupture properties, hardness, structure, and oxidation resistance. Stable structures can be produced for 1800°F exposure, with interesting high-temperature strength and good high-temperature ductility. Oxidation resistance was excellent. A large number of experimental investigations have been made of the role of structure on the properties of cermets and composite materials. Gurland,1 Kreimer et al.,2 and Gurland and Bardzil3 have indicated the preferred particle size in carbide base cermets to be about 1 µ, with a hard phase content of 60 to 80 vol pct. The optimum ductile binder thickness was noted to be 0.3 to 0.6 µ.1 Complete separation of the hard phase particles by the binder is important in reducing the severity of brittle fracture.' The purpose of the present study was to produce structures comparable to the conventional cermets, using a series of relatively close-packed intermetal-lic compounds rather than carbides as the refractory hard phase, and to study the effects of binder content and composition on both high- and low-temperature properties. The selected intermetallic compounds were particularly of interest because of the potential they offered in yielding room-temperature ductility. The highly symmetrical structures are known to possess high-temperature ductility and room-temperature toughness. Based on a ductile binder, the alloys were prepared by the powder-metallurgy route to avoid melting and subsequent alloying of the matrix, and were extruded at relatively low temperatures. It was expected that the composite alloy would retain useful ductility. In contrast, infiltration and high-temperature sintering led to alloying of the matrix and to decreased ductility. The systems Ni-A1 and Co-A1 were selected for this study. In the Ni-A1 system the compounds NiA1, having an ordered bcc B2 structure, and Ni3Al(?1), having an ordered fcc L12 structure, were chosen. In the system Co-A1 the intermetallic compound CoAl with an ordered bcc B2 structure was used. ALLOY PREPARATION The intermetallic compounds, see Table I, were prepared by using master alloys of Ni-A1 and CO-A1, with additions of either cobalt or nickel to achieve the desired compositions. The master alloy in crushed, homogenized form, was melted with pure nickel or cobalt in an inert atmosphere, cold copper crucible, nonconsumable tungsten arc furnace. The resultant intermetallic compounds were homogenized at 2192°C in argon, crushed, and dry ball-milled in a stainless mill to -100 and -325 mesh for the Ni-A1 compounds and to -325 mesh for the CoAl compound. Finer fractions were separated for some of the composite alloys. Several ductile binders were utilized. These included Inco B nickel, 5µ ; pure cobalt, 5 µ, from Sher-ritt Gordon Mines, Ltd.; fine (-325 mesh) niobium hydride powder; fine (15 µ) type 316 stainless-steel powder; and near-saturated Ni-A1 and Co-A1 solid-solution alloys, also in fine powder form. The niobium hydride was decomposed above about 700°C in processing of the compacts in vacuum to produce niobium powder. The Ni-7.1 pct A1 and the Co-5.3 pct A1 solid-solution alloys were prepared from pure nickel or cobalt and pure aluminum by nonconsumable tungsten arc melting under an inert atmosphere. The ingots were homogenized, lathe-turned to fine chips, and dry ball-milled in air to -325 mesh powder. These solid-solution alloys are designated NiSS and CoSS; see Table I. Subsequently the hard and ductile phases were dry ball-milled as a blend. Experiments clearly established the need to coat the hard particles with the ductile binder to optimize subsequent hot compaction by extrusion. Ordinary dry mixing usually resulted in nonhomogeneous alloys which were quite brittle. Conventional cermets are consolidated by liquid phase sinteiing or infiltration, which resulis in undesirable and uncontrolled alloying of the binder phase. For this study, a loose (unsintered) powder-extrusion process was emploved, minimizing reactions between binder and hard particle, thereby permitting much greater control of composition and structure. The constituent powders were first mixed in the desired

Jan 1, 1967

By D. W. Fuerstenau

Recently R. J. Charles1 showed that comminution of brittle or semi-brittle materials in batch operations is described more appropriately by a variable energy relationship than by the specific relationships proposed by Rittinger,2 Kick,3 and Bond.' That is, the value of n in the empirical differential operation which applies to crushing and grinding (Eq. 1) is variable. Independently, Holmes' developed an identical generalized equaton for size reduction. The experimental data of Charles and Holmes show that the value of n is determined by the material and the comminuting machine and is constant for a given combination of these two conditions. The general objective of the present article is to extend the Charles-Holmes concepts to continuous ball milling. The specific objective is to analyze the operation of a vibratory ball mill dry-grinding chromite in a continuous open circuit and to develop an equation for the comminution of chromite in this mill. Basic Principles: In a batch mill the length of time during which the material is ground is a direct measure of the energy expended on the material, except possibly for exceedingly fine grinding. However, in a continuous operation, the fraction of the total expended energy which is useful for grinding will vary according to feed rate and amount of material in the mill. If the amount of material within a continuously operating vibratory mill is low enough so that only part of the expended energy is used for grinding, the extent of size reduction at constant power input must depend upon the length of time the material is retained within the mill, namely, upon retention time. The following analysis pertains to a continuous mill operating under these conditions. The empirical differential equation that applies to the comminution of brittle materials is:" dE = -Cdx/x" [1] where dE is an infinitesimal energy expenditure, C a constant, dx an infinitesimal size change, x the particle size, and n a constant. For continuous grinding in an open circuit at constant power input, P, Eq. 1 must be written as (dtr)P = -C dx/x" [la] where dt, is an infinitesimal time during which a particle of size x is in the mill, and C' is a constant. A ground product never consists of particles of single size but always comprises a distribution of sizes. In comminution, the expenditure of energy causes a lateral shift in the size distribution.', ' " The general objective of any engineering study of comminution is to correlate the magnitude of this size distribution shift with the energy expended, the ultimate aim being to cause the greatest shift with the least amount of expended energy. Although a number of criteria can be used to measure this shift, Schuhmann's size modulus has been found convenient.' Mathematically, the Schuhmann size distribution can be expressed as follows:" y = l00(x/k) [2] or in differential form: dy = [F- xa-' dx ] K, where y is percentage of material finer than size x, a is a constant that determines weight distribution among particles of various sizes, and k the size modulus that denotes the theoretically maximum-sized particle in any size distribution. By considering the energy required to reduce an element of weight of material, dy, from size x, to size x, the element of time during which the material must remain in the mill at at power will be (dt..).= J'-Cax!(dy) [3] Summing up all the elements of weight from zero to 100 for a product that conforms to the Schuhmann size distribution yields C C f Cdx,\( 100«x°-'dx\

Jan 1, 1960

By A. Lawley, H. W. Schadler

TWINNING has long been recognized as a possible mode of deformation in crystalline solids and has been studied in a wide variety of crystals.&apos; Recently, deformation markings which have the topographical characteristics of twins have been seen in tungsten" and the body-centered cubic alloys of tungsten and molybdenum with rhenium.3-7 These markings have been identified as twins because: 1) they are visible after polishing and etching;2,3 2) their formation is accompanied by an audible click and a drop in load;4 and 3) the habit plane is {112}, the plane which has been identified as the twinning plane for several body-centered cubic metals.&apos; However, there has been no published evidence for the specific determination of the twinning direction in rhenium alloys of tungsten and molybdenum, because usually the deformation twins were not broad enough to be detected with a conventional X-ray beam. Consequently, [111] is inferred as the twinning direction since the latter must lie in a plane of symmetry perpendicular to the twinning plane (112). As Sims and Jaffee3 and Lawley and Maddin4 have shown, the addition of rhenium to molybdenum and tungsten increases the ductility of the alloy over that of the pure metal, but twinning is much more profuse in these alloys than in the pure elements, particularly at Mo-35 pct Re*and W-30 pct Re. Since twinning is often associated with brittle fracture in iron8 and the silicon-irons,9 direct experimental proof that the deformation markings in tungsten and molybdenum alloys of rhenium are twins seemed to be needed. Using the Laue back-reflection technique, the deformation markings in a Mo-35 pct Re alloy are shown to be deformation twins. The twins have a (112) habit plane and the twinning direction is <111>. The Mo-35 pct Re specimen used was taken from a 0.060-in. zone-refined single crystal deformed 6 pct in tension at 78°K. The Laue back-reflection technique, with a film to specimen distance of 2 cm, and white radiation from a copper target were used.

Jan 1, 1962

WHEN Magma Copper Company in 1952 set about finding $100,000,000 to finance the project of developing and equipping the property of its lusty progeny, the San Manuel Copper Corporation, A. J. McNab, president and guiding genius for both companies (and incidentally a mining engineer by profession), determined that he would avoid the mistake frequently made in such circumstances, of underestimating the cost, and thereby running short of funds before the job was completed. The fear that this might happen was one of the reasons why insurance companies, one potent source of "big" money, were reluctant to finance the enterprise. How, they asked themselves, would Magma, a relatively small company, get additional money in such a contingency? Insurance officials&apos; thumbs were turned down. Conditions Imposed by R.F.C. The upshot was the negotiation by Mr. McNab for San Manuel of a contract with the Reconstruction Finance Corporation, dated July 10, 1952, whereby needed monies in amount up to $94,000,000 would be advanced as a loan, bearing 5 per cent interest. Arrangements were made for the separate financing of the new $10,000,000 town, needed for the project, through the Del E. Webb Construction Company of Phoenix, Arizona, and M.O.W. Homes, Incorporated, of Los Angeles. Thus the $94,000,000 drawing account was believed ample to provide for building the railroad, the mine plant, the concentrator, and the smelter, as well as for the underground development of the mine needed to start production. As a matter of fact, only

Jan 1, 1957

By F. B. Foley

With the possible exception of high-speed tool steel, the service demanded of rock-drill steel is probably more precarious than that of any other tool steel. Unaided by the helpful influence of alloys and dependent solely on the efficiency of the heat treatment applied for its ability to stand up under the abuse to which it is necessarily subjected, rock-drill steel is worthy of intensive study. For its heat treatment the smith has his fire, an oil fire perhaps, and a quenching bath; he judges the temperature by his eye. There are shops where excellent drills are produced by this method, but there are more where the product varies between very good and very bad. Pyrometers with specially designed equipment are used in but few cases. But the responsibility for rock-drill failures does not rest entirely on the smith. Little has been published concerning the efficiency of drill steel, though the Canadian Mining Journal1 published the breakage record of a large number of hollow drill steels. This record indicates that most of the breaks occur during the early life of the steel and that as a given batch of steel is used, its percentage of breakage decreases. This fact, the article points out, is at variance with the much-talked-of theory concerning failure from crystallization and fatigue. It seems to point to the fact that the greater part of drill-steel failure from breakage is traceable to faulty manufacture, resulting in defects in the bore of the hollow steel from inclusions of slag or oxides, etc. As this defective material meets with early failure, it is weeded out. In addition, some breakage may be caused by a lack of proper condition with respect to heat treatment of the steel as furnished by the manufacturer. In the article mentioned, the breakage of 1/14-in (3.17 cm.) hollow steel is given as 4.62 per cent. for the early period of use, with a decrease to 3.35 per cent. and 2.67 per cent. with time in use. For 1-in. (2.5 cm.) hexagonal hollow steel, the figures for breakage are 0.18 per cent. and 0.06 per cent.; attention is called to the great decrease in breakage by the use of smaller diameter drills. The figures are given as so many drills sharpened, but there is no way of determining the percentage of

Jan 1, 1922

By H. Stuart, L. Zwell, N. Ridley

specimen that was subjected to 700°C for 250 hr is shown in Fig. 2(c). The precipitates in this foil appear round and exhibit only very faint strain contrast. From this evidence, along with that described above for the symmetry of the strain field, it may be concluded that this coherent precipitate is spherical, which is generally the favored shape for coherent precipitates when the lattice parameter mismatch is small.12 These observations are consistent with the assumption that this precipitate is the y&apos; (Ni3Al, Ti) phase. y&apos; is known to be spherical in the early stages of precipitation. It is fcc and is similar in composition to the matrix and thus would have contrast similar to that of the matrix as do these precipitates. Direct identification of the spherical precipitate did not prove possible because of its low volume fraction and small size. The results of this investigation suggest that the increases in yield strength and the decreases in elongation observed in 45-Fe 30-Ni 20-Cr alloys containing about 0.4 pct A1 and 0.4 pct Ti on prolonged aging at 550" to 700°C are very likely due to the presence of a coherent, spherical precipitate, thought to be y&apos;. Mi-crohardness measurements indicated that the precipitation of semicoherent M23C6 resulted in some matrix hardening. However, this precipitation process occurs much more rapidly than the changes in mechanical properties at these temperatures. 1C. L. Clark, J. J. B. Rutherford, A. B. Wilder, and M. A. Cordovi: J. Eng. Power, 1962, vol. 84, p. 258. 2 F. A. Comprelli and U. E. Wolff: USAEC Report GEAP-4745, General Electric Company November, 1964. 3C. N. Spalaris: A Monograph, USAEC Report GEAP-4633, General Electric Company, July, 1964. THERE have been many X-ray investigations of the room temperature lattice parameters of austenite which has been retained in quenched steels of various carbon contents. The results of these investigations have been collected by Wrazej,1 Wever et al.,2 and by Roberts3 and are included in Fig. 1 due to Pearson.4 The dashed line in Fig. 1 was considered by Roberts3 to most reliably describe the variation with carbon content of the lattice parameter of retained austenite. This line is given by the equation: a (Å) = 3.555 + 0.044 wt % C and recent unpublished measurements by one of the authors (LZ) support this relationship. The object of this communication is to draw attention to a relationship which describes the variation of the room temperature lattice parameter of &apos;equilibrium&apos; austenite with carbon content, and which is different from that for retained austenite. The relationship a(Å) = 3.573 + 0.033 wt % C has been obtained from lattice parameter measurements made on equilibrium Fe-C austenites at elevated temperatures, which have been extrapolated to room temperature. The room temperature lattice parameter of 3.555Å which the Roberts&apos; relationship gives for pure y-Fe is no: in agreement with the value of 3.573Å obtained when the high temperature lattice parameter data of Basinski et aL.5 and of Goldschmidt6 is extrapolated to room temperature. A linear extrapolation was

Jan 1, 1970

By R. J. Charles, G. Agar

Experiments on single particles show that the amount of material created during impact that is finer than any chosen size is proportional to the energy of the impact. As the underlying principle of comminution, it might be stated that each unit of energy input to a given comminution system tends to add to the system an identical assembly of new particles and subtract an equivalent volume of larger particles. Thus, two of the consequences of the fact that each unit of energy tends to add identical assemblies of new particles to the system are: 1) on continued application of energy the total charge in any batch comminution system tends to assume the characteristics of this assembly, and 2) the process of comminution may be described by a relationship such as E = Ak-a. Recently Schuhmannl presented a theory of comminution which puts rational perspective on the extensive controversy surrounding most of the existent theories of comminution. With hindsight the origin of controversy can be ascribed to a relatively simple circumstance. In general, a progressive shift of a size distribution curve towards finer sizes is observed as energy is added to a comminution system. It is also observed that, with continued grinding, the size distribution curves usually change and appear to approach some constant shape. Inves-tigators generally felt that an appropriate method of relating and generalizing such data would be to consider the hypothetical case where particles in the system were uniformly subjected to comminution such that each of these particles, receiving its proportion of input energy, was transformed into a number of smaller particles of uniform size. Since the changes in shape of the size distribution curves during size reduction were clearly associated with the coarse particles in the system, it was also assumed that these changes accounted for little of the input energy and one could deal solely with the equilibrium shape towards which the size distribution curves tended. Although the unrealistic nature of the foregoing points of view was realized some justification was felt for such treatments since, in most cases, the resulting mathematical formulae could, by the adjustment of at most two parameters, satisfactorily describe experimental observation. Difficulties arose, however, in that the formulae required a specific influence of the feed size on the comminution process. In actual fact, the application of a feed size term in fitting experimental data was often found to be unnecessary or contrary to the above formulae. Secondarily, much discussion arose concerning the appropriate energy number that would relate the size change of one particle to the amount of comminution energy it received. As has been previously shown3 all the above theories that treat comminution as a continuous process of size reduction may be related to a proposition of the following form: dE = -Cfn [1] where dE is the increment of energy, dx is the size change, x is the particle size, and C and n are constants. It is now clear, as pointed out by Schuhmann, that Eq. 1 is inapplicable to comminution processes, since comminution cannot be considered other than as a discontinuous mixing process in which the overall size distribution of the product results from mixing, in various proportions, finished material with unfinished, and sometimes untouched, material. A striking illustration of this mixing phenomena may be obtained by comparing the size distributions of grinds in a typical batch ball mill with size distributions calculated on the basis of mixing various proportions of feed material with a ground material which obeys a power law size distribution relationship. Figs. la and lb illustrate such a comparison and the similarity of these figures indicates that mixing plays an important role in determining the overall shape of the product size distributions. Figs. la and lb illustrate, additionally, the underlying basis of the comminution theory presented by Schuhmann.&apos; In the preliminary grinding of relatively close sized feed particles in a batch ball mill, one may visualize that, when a successful encounter of a ball with a feed particle occurs, then the particle undergoes severe size reduction and the size range in which the resultant particles reside is far removed from the size of the feed particles. After a number of feed particles are reduced by independent but similar impacts and one examines the size dis-

Jan 1, 1961

By AIME AIME

Report of Committee 24, of the International Association for Testing Materials, presented at the Brussels Congress, 1906. Republished for use at the 94th Meeting of the American Institute of Mining Engineers, New York, N. Y., February, 1908. Discussion by letter is requested by the Secretary of the Institute. The Committee appointed for the solution of Problem 24 consisted of the following members Prof. H. M. Howe, New York, N. Y. (Chairman); L. Levy, Paris, France (Vice-Chairman) ; Prof. D. Tschernoff, St. Petersburg, Russia (Vice-Chairman) ; Prof. Albert Sauveur, Cambridge, Mass. (Secretary) ; Van Drunen, Ixelles, Belgium; Lieut.-Col. H. Tusen, Copenhagen, Denmark ; E. P. Martin, Abergavenny, England ; A. Pourcel, Paris, France ; H. Brauer, Dortmund, Germany ; Dr. Prof. H. Wedding, Berlin, Germany; P. Verole, Milan, Italy ; A. Baalsrud, Christiania, Norway ; A. R. v. Dormus, Vienna, Austria ; A. Sailler, Vienna, Austria; Bureau des Forges, St. Petersburg, Russia; Prof. N. Belelub-sky, St. Petersburg, Russia ; N. Jossa, St. Petersburg, Russia ; Maj.-Gen. M. Korobkoff, St. Petersburg, Russia ; S. Smirnoff, St. Petersburg, Russia ;.Ladislaus v. Tetmajer, Budapest, Hungary ; H. H. Campbell, Steelton, Pa. Early in February, 1905, a list of species and varieties of iron and steel, with definitions, drawn by the Chairman and Secretary of Committee 24, was sent to every member of this Committee with a request for comments and suggestions. A special request was sent to the following gentlemen . for the equivalents in their several languages of these species and varieties : A. , Pourcel, Paris, France; Prof. H. Wedding, Berlin, Germany ; J. A. Brinell, Stockholm, Sweden ; Prof. H. I. Hannover, Copenhagen, Denmark; Prof. P. Kley, Delft, Holland; P. Verole, Milan, Italy; Col. Cubillo, Trubia, Spain. This present report, which is based upon our original list and definitions, and upon the replies to these requests, consists of four parts I. A polyglot, which gives the names of the chief classes of iron and steel in English, French, German, Swedish, Danish and Dutch. We .regret that we have failed to get the Italian and Spanish names. II. English definitions of principal classes of iron and steel. III. A glossary of special sizes and shapes of iron and steel. IV. A note on the boundary between steel and cast-iron. I. POLYGLOT, IN ENGLISH, FRENCH, GERMAN, SWEDISH, DANISH AND DUTCH, OF THE CHIEF CLASSES OF IRON AND STEEL, AND OF THE APPARATUS IN WHICH THEY ARE MADE. This list is not intended to be complete, but rather to include only those varieties which are of interest to the Association.

Mar 1, 1908

By C. H. Samans, G. F. Tisinai, J. K. Stanley

The addition of nitrogen (0.10 to 0.20 pct) to Fe-Cr-Ni alloys of simulated commercial purity results in a real displacement of the u phase boundaries to higher chromium contents. The effect is small for the (Y + s)? boundary, but is pronounced for the (y + a +s)/(y + a) boundary. Although there is an indication of an exceptionally large shift of the n boundaries to higher chromium contents, especially in steels with nitrogen over 0.2 pct, the major portion of this apparent shift results from the fact that carbide and nitride precipitation cause "chromium impoverishment" of the matrices. The effect of combined additions of nitrogen and silicon to the Fe-Cr-Ni phase diagram is demonstrated also. Nitrogen can nullify the effect of about 1 pct Si in shifting the (y + o)/? phase boundary to lower values of chromium at all nickel levels from 8 to 20 pct. NItrogen can nullify this U-forming effect of about 2 pct Si at the 8 pct Ni level, but not at the 20 pct Ni level. The alloys studied were in both the cast and the wrought conditions. There are indications that the u phase forms more slowly in the cast alloys than in the wrought alloys if both are in the completely austenitic state. The presence of 6 ferrite in the cast alloys accelerates the formation of U. Cold working increases the rate of o formation in both cast and wrought alloys. THE major improvement in Fe-Cr-Ni austenitic alloys in recent years has been in the addition or removal of minor alloying elements to facilitate better control of corrosion resistance, sensitization, and heat resistance. One shortcoming of the austenitic Fe-Cr-Ni alloys, which never has been completely circumvented, is their propensity toward u formation. In the AISI-type 310 (25 pct Cr-20 pct Ni) and type 309 (25 pct Cr-12 pct Ni) steels, sufficient amounts of u phase can form, if service or treatment is in a suitable temperature range, to cause severe embrittlement. Also, there is a growing conviction that this phase may be contributory to some unexpected decreases in the corrosion resistance of certain of the 18 pct Cr-8 pct Ni-type steels. The present paper discusses the effect of nitrogen additions on the location of the (r+u)/d and the (y+a+u)/(y+a) phase boundaries in the ternary Fe-Cr-Ni system, for cast and wrought alloys of simulated commercial purity, and in similar alloys containing up to about 2.5 pct Si. The objective is to define compositional limits for alloys which will not be susceptible to u formation when used near 1200°F (650°C). An excellent review of the early studies of the u phase in the Fe-Cr-Ni system has been compiled by Foley.1 Rees, Burns, and Cook2 have determined a high purity phase diagram for the ternary system, whereas Nicholson, Samans, and Shortsleeve3 are- stricted themselves to a portion of the simulated commercial-purity phase diagram. Both groups of investigators show almost an identical position for the commercially significant (y+u)/y phase boundary. Further comparison of the work of the two groups indicates that, below the 8 pct Ni level, the commercial alloys have a decidedly greater propensity toward u formation than the high purity alloys. The two groups of workers agreed that both the AISI-type 310 (25 pct Cr-20 pct Ni) and the type 309 (25 pct Cr-12 pct Ni) steels are well within the (y+~) region and that the 18 pct Cr-8 pct Ni-type alloys straddle the U-forming phase boundaries. Nicholson et al.3 showed, in addition, that these boundaries shift toward lower chromium contents if greater than nominal amounts of silicon or molybdenum are added. The effect of nitrogen on the location of the s phase boundaries in the Fe-Cr-Ni system has not been known with any certainty. In 1942, an approach to this problem was made by Krainer and Leoville-Nowak,&apos; but at that time they apparently were unaware of the slow rate of s formation in strain-free samples and aged their samples for insufficient times, e.g., 100 hr at 650°C (1200°F) and 800°C (1470°F). For this reason, it would be expected that their (y+ u) /y boundary would be shifted toward lower chromium contents (restricted ?-field) when equilibrium conditions were approximated more closely. Procedure for Studying the Alloys The alloys used were prepared in the following way: Heats of 200 lb each were melted in an induction furnace. A 5 lb portion of each heat was poured into a ladle containing an aluminum slug for de-

Jan 1, 1955

By E. Motica

Ores containing uranium and vanadium minerals have been mined from the Salt Wash Member of the Morrison Formation from many localities in the Colorado Plateau region since about 1900. The most productive deposits are in a relatively small area in Southwestern Colorado referred to as the Uravan mineral belt. The mineral belt is a narrow elongate area in which uranium-vanadium deposits in the Salt Wash have a closer spacing, larger size, and higher grade than those in adjoining areas. Pre- 1948 ore production from the mineral belt was 655,000 tons averaging 1.91 per cent V2O5 and 0.28 per cent U3O8. Production during the 1948 to 1964 period was 7,900,000 tons of ore averaging 1.46 per cent V2O5 and 0.27 per cent U3O8. The ore deposits occur principally in the uppermost sandstone unit of the Salt Wash. This unit consists of sandstone lenses formed by a system of aggrading braided streams flowing in an easterly direction generally normal to the mineral belt. These lenses, referred to as channels, are as much as one mile in width and average about 50 feet in thickness and can be traced for several miles along their courses. Areas favorable for one deposits are recognized principally by the following criteria: ( 1) a host sandstone thickness of over 30 feet, (2) the presence of carbonaceous material in the host sandstone, and ( 3) gray mudstones and clays associated with the ore-bearing sandstone. The ore minerals are believed to have been precipitated from laterally migrating solutions; there is no apparent genetic relationship between ore deposits and tectonic structural features. Exploratory and development drilling from the surface is conducted in three phases: first, wide-spaced drilling to find and outline favorable channels; second, moderate-spaced drilling in favorable ground to locate ore deposits; and third, close-spaced drilling to block-out ore. Long-hole drilling from underground stations is used to find and delineate extensions of ore bodies being mined. Due to intense production since 1948 from shallow deposits, best potential for new deposits exists in areas where the ore-bearing unit is 600 to 800 feet deep.

Jan 1, 1968

By J. T. Dewan

An experimental setup is described whereby the responses of Neutron logging instruments have been determined opposite formations of different porosities and under various borehole conditions. The results of these tests are presented in several charts which show how the Neutron log is affected by the following factors: hole diameter, casing, cement, salinity and weight of drilling mud, and position of logging tool with respect to the wall of the hole. The responses in the case of empty holes are also shown. A chart is furthermore presented which shows the shapes of the Neutron log, as a function of logging speed and time constant, at the boundaries between two formations with different porosities. These were derived mathematically using as a basis laboratory determination of the shape of the curve under static conditions, i.e. at very low logging speed. Corrections for the readings made opposite thin beds are given by these computations. The correction is found negligible under usual logging conditions for beds thicker than about 4 ft. INTRODUCTION The Neutron log is used in borehole surveys both for the delineation of porous formations and for the quantitative estimation of their porosities.1,2,3,4,5 It is common practice to calibrate the response of the Neutron log in terms of porosity values for each given geological unit by means of empirical comparisons between the deflections of the neutron curve and the porosity values given by the core analysis in a sufficient number of wells. Usually, the plots of the porosity values vs Neutron log deflections show a definite trend. It is therefore possible to trace an average line across the plotted points. This average line constitutes a calibration curve which can be used to evaluate the porosities from the Neutron logs in the other wells drilled in the same region. Futhermore, when the porosity values are plotted on a logarithmic scale and the Neutron log deflections on a linear scale, the calibration curve is practically a straight line. Accordingly the calibration curves are often determined by simply tracing a straight line between two points corresponding to two widely different porosities. Different choices are made for these two points, depending on local conditions. For example, one point will be defined by plotting the deflection of the Neutron curve opposite a formation known to be very tight, for which an average porosity value of 1 or 2 per cent is assumed. The other point may correspond to shale sections, to which a porosity of around 40 per cent is ascribed. The information provided by the eletrical logs — in particular the Laterolog, Limestone log or Micro-Latero-log — is also very often a help in the calibration of the Neutron logs. Empirical approaches such as these make possible the derivation of fairly reliable porosity values from the Neutron logs in many cases, at least in open holes and in essentially clean formations. There are, however, several causes of uncertainty which may sometimes entail appreciable errors. On the one hand, the Neutron log readings are influenced by the borehole conditions: hole diameter, presence or absence of casing and cement, presence or absence of drilling mud, characteristics of the mud, and position of the logging instrument with respect to the wall of the hole. The amplitudes of the curves opposite thin beds are furthermore atfected by the bed thickness,

Jan 1, 1957

By R. T. Mishler

Silver-gold. mine, in northern Mexico, with arrow veins in rhyolite. One fourth area developed has been ore. Deposits average 20 in. wide and 73 oz. silver per ton; diluted in mining to 40 oz. per ton. Developed by drifts and raises at 100-ft. intervals. Samples cut every 5 ft. Ore and waste sampled separately to make combined stoping width 3 or 3 1/2 ft. Ore reserves figured to sloping width; high assays not modified. Tonnage and assay of ore extracted higher than estimates. Ore stoped principally by cut and fill method. Cars from each stope counted and sampled. Cars check tonnage to concen¬trator within 1 per cent.; silver assay 9 per cent. high. Tonnage and assay from scopes corrected to agree with feed to concentrator. Profit or loss for each slope estimated monthly. THE Lucky-Tiger mine is a silver-gold property, situated at El Tigre, in the northeastern part of Sonora, Mexico, at, an elevation of 6000 ft. in the Sierra Madre Mountains. It is 30 miles by wagon road from the Esqueda station on the Nacozari railroad. Since its discovery in 1903, the mine has produced forty million ounces of silver. The ore has averaged 40 oz. silver and 0.25 oz. gold per ton. Copper, lead and zinc have been of minor importance, having averaged 0.4 per cent., 1.1 per cent., and 1.5 per cent., respectively. The richness of the ore,, combined with its occurrence in narrow veins, required the use of the more expensive methods of mining.

Jan 2, 1925

By J. A. Solomon, H. E. Shafer

As a result of the great demand for metallic aluminum and the scarcity of domestic bauxite, numerous processes have been proposed and developed to extract the oxide of this metal (Al2o3) from clays, shales, and other low-grade raw materials.1-4 An exploratory study was made to extract alumina with sulfuric acid from relatively impure coal-associated draw-slate and to establish an optimum calcination temperature for maximum recovery. Draw-slate is the term used to describe the hard, slaty material directly overlying many coal seams and may range from a few inches to a few feet in thickness. Alumina content of these draw-slates may range from as low as 15% to over 35%. Although the following results are limited in scope, some of the data should prove beneficial for further work. Coal-associated clay, shale, and gob heretofore have been considered nuisance products; any successful alumina extraction process would greatly enhance the worth of these almost limitless resources. This research demonstrated that an optimum calcination temperature for alumina extraction could be established for a coal-associated draw-slate and that this temperature was lower than those suggested in previous work with unspecified clay types. It should be noted that slightly higher acid concentrations and somewhat finer particle sizes were used in this series of tests as compared to previous work. EXPERIMENTAL WORK Essentially, alumina extraction proceeded in four steps: 1) grinding, 2) calcining, 3) acid leaching, and 4) precipitating of aluminum sulfate. An initial series of reference tests to determine the optimum temperature of solubility for alumina was run using relatively pure kaolin. Kaolin forming clays are hydrous aluminum silicates of the approximate composition 2H2O A12O3 2SiO and kaolinite is the mineral that predominates in most kaolins. 5 The kaolin was ground to -30 mesh and subsequently calcined at temperatures ranging from 100°C to 700°C. Each sample was calcined for 1 hr in a muffle furnace and in turn subjected to sulfuric-acid leaching (H2SO4, 9 molar) for a 2-hr period. The filtrates were analyzed for the amount of dissolved alumina recovered.637 Also, each sample of leached or barren residue was

Jan 1, 1968

By L. J. Boucher

The beneficiation process at Northampton is described and reasons are given for installing a flotation plant. The economics of running the plant, the difficulties of operation, and subsequent remedial action are discussed. CEMENT rock from a quarry operated by the Northampton plant1, 2 of the Universal Atlas Cement Co. has a relatively low average CaCO³ content. To meet specifications required for cement manufacture it was formerly necessary, over succeeding years, to add increasingly larger amounts of high-lime, low-alumina rock, which was obtained from a quarry operated by the company at a point along the New Jersey-New York state border, 79 miles from Northampton. Previous to 1940, 8 pct of total kiln feed consisted of high-lime rock. From 1940 to 1943, this figure jumped to 13 pct. Today it would be necessary to use 18 pct of this high-lime to meet current specifications. Increasing transportation costs made this an expensive operation. In addition, this high-lime rock quarry was contaminated with granite and high magnesia pockets. As there was no available high-lime rock in the Lehigh Valley, it was decided in 1936 to install a pilot plant where experimental work was done on processing local cement rock with flotation and burning the resulting product in a small kiln. This plant was in operation for six months and at the end of that period a suitable method of flotation had been established. From these studies it was also found that if the proportion of shipped-in rock rose slightly above 10 pct it became economically advisable to use the flotation process. It is the object of this paper to present briefly the operation of the flotation system at the Northampton plant. It seems advisable to comment on the raw material used in this process. The Northampton quarry lies geologically in what is known as the Jacksonburg formation of Trenton limestone.&apos; In Pennsylvania this extends in a southwesterly direction from Portland, through Nazareth, Bath, and Northampton and across the Lehigh River to Ormrod. This argillaceous limestone was formed from marine deposits laid down during the early Ordivician period when a shallow sea covered most of Pennsylvania. Carbonaceous matter in the original deposit formed graphitic carbon under later metamorphic forces. From this graphite the cement rock of the Lehigh Valley derives its dark color. Most of the graphite must be removed before lime flotation can proceed. The quartz in this rock is desirable from the chemist&apos;s point of view, as it furnishes him with a form of silica necessary in the calcining process, but its abrasive properties damage pumps and piping. Aluminum silicates of potassium, sodium, iron, and magnesium are found in the form of sericite. Cal-cite averages from 69 to 76 pct. Iron in the form of iron pyrites is also present and has no deleterious effect. Talc is present in minute quantities. The

Jan 1, 1954

By Harry M. Bendler, W. A. Wood

Electron metallography has been applied to the study of the torsional fatigue of polycrystalline OFHC* copper. The development of slip bands, fis- sures, and microcracks has been followed as a function of strain amplitude and number of cycles. Those factors that have been found to be relevant to a geometrical theory of the fatigue process are described. From the point of view of this theory it is important to distinguish between the effects of cycling at large amplitudes and at small amplitudes (true fatigue deformation). EARLY investigations showed metal fatigue cracking is caused by continued reversal of a small plastic strain. Fundamental study of fatigue, therefore, has been largely concerned with the cumulative effects of these plastic reversals on the metal structure. In particular, recent work has shown that the plastic reversals produce slip movements that are finer than the slip movements produced by unidirectional increments of plastic strain; individual slip movements in fatigue, for instance, are too fine to see by optical microscopy.2 It has also been shown that the fine slip movements of fatigue tend to concentrate in bands and that in this way they can build up, after enough cycles, a progressive surface disturbance. It is this disturbance which, in effect, makes the slip movement visible. It is also in this disturbance that the first cracks appear to form. From this point of view, therefore, the problem of fatigue becomes one of finding why the fine slip movements should concentrate in bands and why the. bands should finally open into cracks. Thus far, study of the slip bands has been limited to the changes they produce at the surface only. These studies have given interesting results. Ewing and Humfrey3 many years ago showed that the bands in iron could extrude debris. Forsyth found that bands in aluminum alloys could extrude thin tongues of material. Thompson showed that, after lightly dissolving away the surface of copper specimens, some bands persisted as superficial cracks might, and that the life of a specimen under fatigue deformation might be prolonged by dissolving away the surface at intervals. Thus these and related results have suggested that the metal structure in fatigue bands may disintegrate and that this may happen long before obvious cracking. These surface studies, however, have not shown definitely how cracking may begin, for they do not show sufficiently what is happening in depth. However, some advance in this direction has been made recently by a new approach, summarized later, which reveals slip bands and their surface disturbances in section. In the work done thus far, the sections have been studied by optical microscopy only. In the present work they have also been studied by electron microscopy with extended results. This paper also describes some observations on the surface disturbances prior to sectioning. These observations are relevant to a suggested theory of fatigue cracking. REVERSED SLIP MOVEMENTS Slip movements in fatigue deformation must reverse direction in each cycle. The question arises: DO the "backward" and "forward" movements take

Jan 1, 1962

By Frank L. Nason

RECENT explorations with a diamond-drill, originated and conducted by Mr. J. A. Van Mater, Superintendent of the Lehigh Zinc and Iron Company, have revealed some very important geological facts, which, it is believed, will be of special practical interest to mining engineers. During these explorations the writer was frequently called in consultation, and hence the facts here recorded are almost entirely the results of personal observation. The franklinite-deposit upon which the explorations were conducted has a long continuous outcrop. Figs. 1, 2 and 3 give a map and sections. Beginning at the extreme southern point,

Jan 1, 1895

By C. W. Macgregor

The observation of the flow layers, or Luders&apos; lines, produced in mild steel when it is stressed into the plastic range often provides considerable useful information for the study of the fundamental mechanism of the plastic flow of metals. It is a commonly observed fact that as soon as the upper yield point in a mild-st&apos;eel tensile-test bar is passed, an unstable condition develops in the equilibrium, the load on the testing machine drops and if the surface of the bar contains a thin mill scale, the latter begins to flake off rapidly along certain "lines." These "lines" are ill reality traces of wedge-shaped flow layers and are formed by the inter-section of the surface of slip and the surface of the bar. The average plane of the flow layer itself has been found by various experimenters to be inclined at an angle of approximately 45" to the axis of the test bar. In a body of more complicated shape stressed in possibly a different man-ncr, the tangent planes to the flow layers at each point seem in many cases to be inclined at an angle of about 45" to the principal stress direct-tions at the point. As examples of the latter may be mentionecl the logarithmic spirals obtained when a rivet hole is punched in a thin plate, or when a tube is subject,ed to internal pressure only, etc. IIencc if the flow layers are observed from an experiment, the shear-st&apos;ress trajectories —the curves indicating the direct,ions of the maximum shearing stresses in the body—may then be approximately determined. Since the latter curves form an orthogonal family inclined at 45" to the principal stress directions, the family of principal stress trajectories, or the curves irldicating the directions of the principal stresses, may be thus obtained Although the discontinuous type of stress-strain curve at the yield point is commonly observed with test bars of mild steel, more recent experiments have shown that this is by no means the only material exhibiting this interesting phenomenon. A similar effect has been observed in zinc and cadmium crystals, by Becker and Orowai~; in duralumin and other aluminum alloys, by Dawidenkow,~Elam,~ Fell4 an(l others; in copper and copper-alloy crystals, by Elam;~ in brass by Elam;3 in brass wires by deForest;" in higher carbon steels, tungsten steel, and

Jan 1, 1939

By C. W. MacGregor

The observation of the flow layers, or Luders&apos; lines, produced in mild steel when it is stressed into the plastic range often provides considerable useful information for the study of the fundamental mechanism of the plastic flow of metals. It is a commonly observed fact that as soon as the upper yield point in a mild-st&apos;eel tensile-test bar is passed, an unstable condition develops in the equilibrium, the load on the testing machine drops and if the surface of the bar contains a thin mill scale, the latter begins to flake off rapidly along certain "lines." These "lines" are ill reality traces of wedge-shaped flow layers and are formed by the inter-section of the surface of slip and the surface of the bar. The average plane of the flow layer itself has been found by various experimenters to be inclined at an angle of approximately 45" to the axis of the test bar. In a body of more complicated shape stressed in possibly a different man-ncr, the tangent planes to the flow layers at each point seem in many cases to be inclined at an angle of about 45" to the principal stress direct-tions at the point. As examples of the latter may be mentionecl the logarithmic spirals obtained when a rivet hole is punched in a thin plate, or when a tube is subject,ed to internal pressure only, etc. IIencc if the flow layers are observed from an experiment, the shear-st&apos;ress trajectories —the curves indicating the direct,ions of the maximum shearing stresses in the body—may then be approximately determined. Since the latter curves form an orthogonal family inclined at 45" to the principal stress directions, the family of principal stress trajectories, or the curves irldicating the directions of the principal stresses, may be thus obtained Although the discontinuous type of stress-strain curve at the yield point is commonly observed with test bars of mild steel, more recent experiments have shown that this is by no means the only material exhibiting this interesting phenomenon. A similar effect has been observed in zinc and cadmium crystals, by Becker and Orowai~; in duralumin and other aluminum alloys, by Dawidenkow,~Elam,~ Fell4 an(l others; in copper and copper-alloy crystals, by Elam;~ in brass by Elam;3 in brass wires by deForest;" in higher carbon steels, tungsten steel, and

Jan 1, 1939

By R. A. Grange, H. M. Stewart

Man.; steel parts may crack if quenched directly into a bath near room temperature, but not if quenched at a temperature just above the range where martensite forms and then allowed to cool slowly to room temperature. This latter procedure, which is the basis for such modern hardening techniques as "martempering" and "isothermal quenching," may entail some sacrifice in depth of hardening but very little, if any, in intensity of hardening. In planning such treatments, it is necessary to know the upper limit of the martensite range, and often desirable to know the proportion of martensite that would form on cooling to any lower temperature; furthermore, the tendency of a particular steel to crack when quenched is unquestionably associated with the temperature range of martensite formation, and consequently knowledge of this range aids in selecting the optimum composition for a given application. Certain puzzling phenomena associated with the properties of quenched and tempered steel may be at least partly clarified as the manner in which martensite forms is more fully understood. Information on the temperature range of martensite formation, particularly with respect to how much martensite results from quenching to a given temperature, is lacking for most commercial types of carbon and low-alloy steel. These considerations led us to study martensite formation in 14 carbon and low-alloy steels. The resultant data may be directly used in the following ways: (I) for selecting the lowest quenching temperature at which no martensite will form; (2) for selecting the highest quenching temperature at which virtually all marten . site will form, thereby avoiding quenching to an unnecessarily lower temperature with attendant danger of cracking; and (3) in producing a mixture of tempered martensite and bainite, which in high-carbon steels has been found to possess somewhat better ductility than tempered martensite, yet does not require the full transformation time necessary for a completely bainitic (austempered) structure. PART I. OBSERVATIONS OF MARTENSITE FORMATION IN FOURTEEN CARBON AND LOW-ALLOY STEELS Transformation of austenite to pearlite or bainite can be conveniently studied by observing the change that occurs over a period of time at constant temperature; such studies are the basis of the isothermal transformation diagram (S-curve). This isothermal technique is not applicable to martensite formation, which occurs within a time interval too small to measure, at least by any ordinary means. For all practical purposes, therefore, martensite may be regarded as forming only during cooling and not during a period of time at constant temperature. As austenite is cooled below

Jan 1, 1947