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By H. Kreye, E. Hornbogen

An increase in electrical resistivity during aging or after quenching from high temperatures is known in a large number of alloys. Plastic deformation of alloys in this condition leads to a decrease in resistivity. This anomaly has been called the K effect. This effect has been reported in Ni-A1 alloys by several authors.'-4 The nickel-rich solid solution y is in equilibrium with Ni,A1 (y') with CkAu structure. With this alloy it is therefore possible to produce a homogeneous solid solution as well as a coherent mixture of y and y. The changes in resistivity should be determined for these two conditions. In the present literature on the K effect no distinction is made between them, because of experimental difficulties to determine small coherent particles. In the present work transmission electron microscopy and image formation by (100) and (110) superlattice reflection of the y' phase was employed for a detailed determination of the temperature dependence of solubility of aluminum in nickel and the resulting microstructures with alloys up to 15 at. pct Al. The solubility line follows the equation indicating a smaller solubility than given in the literature at temperatures below 700°C. X-ray scattering was used to get evidence for short-range ordering in the homogeneous solid solutions.5 Parallel to the determination of the microstructure electrical resistivity was measured as a function of aging tempera- ture, time, and aluminum content, which led to the following results: 1) Resistivity increases slightly during aging of a homogeneous solid solution which was quenched from a high temperature as a consequence of short-range order, Fig. 1, curve c,. 2) Resistivity shows a maximum during aging of all alloys which are heterogeneous according to the phase diagram, but which were obtained as a supersaturated solid solution by quenching, Fig. 1, curve c,. The increase in resistivity is mainly due to coherent precipitation of y'. In Fig. 2 the effects due to short-range order pN and due to precipitation of particles APT are shown for an aging temperature of 450°C and different aluminum contents. 3) The increase in resistivity depends on the size !f the coherent particles. A critical diameter of 15 5A for the maximum was found for all alloys and aging treatments by extrapolating particle sizes determined by electron microscopy. 4) The maximum of resistivity of the heterogeneous conditions does not occur if particles, larger than the critical size, have formed already during quenching. 5) Plastic deformation destroys short-range order as well as small coherent particles. Therefore a decrease in resistivity is observed if the increase in resistivity caused originally by these phenomena is

Jan 1, 1967

By R. G. Davies

Isothermal annealing of quenched Fe3Al reveals that the superlattice forms by the nucleation and growth of ordered domains. The activation energy for isothermal ordering and initial domain growth is -29 keal per mole, which is thought to he the activation energy for vacancy migration. Domain growth, in the absence of excess vacancies, has an activation energy of -100 keal per mole. A maximum in the hardness is observed when the structure consists Of ordered domains, -35-4 in size, growing into a short-range ordered matrix and, also as a function of degree of order, S, with an infinite domain size at S = 0.4 to 0.5. It is proposed that both maxima are due to a transition from deformation by unit dislocations to deformation by superdislocntions. FE-AL alloys containing more than -23 at. pet Al, when quenched from above the critical temperature for the Fe3A1 type of order, always contain the maximum possible amount of FeAl (B2) type of order' (see Fig. 1). Thus when the Fe3A1 superlattice forms, in a 25 at, pet A1 alloy, it does so from an imperfect FeAl structure. The formation of the Fe3Al superlattice has been the subject of several recent X-ray,1, 2 electrical-resistivity,3"6 and dilatometric3 investigations. McQueen and Kuczynski3 concluded that the isotherma1 ordering of Fe3Al, after quenching from 800°C (T,, the critical temperature for order, is 55O°C), takes place by the nucleation and growth of ordered domains. However, Feder and Cahn4 found that the isothermal ordering was complex but more like a homogeneous nucleation process than a nucleation and growth process. Lawley and Cahn2 reported that the degree of order is a continuous function of temperature which is a characteristic of a homogeneous nucleation process. Thus, the evidence is conflicting as to the nature of the Fe3Al ordering reaction. Lipson7 was the first to suggest that ordered domains should exist in the Fe3A1 structure since, as

Jan 1, 1964

By R. Wachtell

IN order to study better the phenomena at work in various phases of diffusion of the Fe/Si system when compounded and alloyed by powder metallurgy methods, several attacks have been planned. Electrical, hardness and other measurements have been discussed e1sewhere.l,2 The metallographic phenomena will be considered here. Concerning the metallography of the ferrosilicon alloys, some initial discussion is in order. As this laboratory discovered to its chagrin, there are many traps that lie in wait for even the experienced metal-lographer when he deals with alloys of high silicon content. Chief of these is probably the anomalous behavior of these materials under conventional etching techniques. We have before us, for instance, Corson's paper" of 1928, wherein is mentioned and illustrated (probably for the first time) the peculiar "barley shell" structure sometimes found in these alloys. Again in 19412 he same author reports the structure in 5-14 pct silicon irons. Houghton and Becker mention its presence' but make only tentative effort to explain it. In 1943, Hurst and Riley6 discussed the "barley shell" and declared it to be a false structure, resulting from the peculiar and characteristic film-forming propensities of the alloy. At about the same time, Wrazej7 proved correct the contentions of Hurst and Riley, establishing pretty well the mechanisms of formation and identifying the constituents involved. Hurst and Riley6 a1so describe in their paper (1943) a "cracked film" structure, also a pseudo-morph, the precise nature of which they do not suggest. Careful examination of such a film reveals that it not only cracks but begins to curl away from the surface of the metal and peel, much in the manner of flaking paint. The careful microscopist will be able to focus on the topmost edge of the flake, and, by optical sectioning, follow it down to the metal surface. We have been able, by oblique illumination of such specimens, to observe the raised segment of the film, the shadow that it casts, and the mating segment on the metal into which it fits. As late as 1946, Hurst and Riley8 found occasion to correct misinterpretation of this "cracked film" structure in a work by others on a 12 pct Si/Fe alloy. The "barley shell" and "cracked film" pseudo-morphs are not the only ones encountered in studies of these materials. A third pseudomorph, which may, however, have some structural significance is a fine striation of the surface. Closely spaced and parallel striae extend entirely across single grains. Each grain shows a different direction for the striae, and different closeness of the spacing. The regularity of this structure, and the fact that the striae do not cross from one grain to another, suggest that the structure is a film cracking controlled in some way by the crystal plane orientation of the surface being etched. In general it is wise for the metallographer dealing with these alloys to cultivate a feeling of uncertainty not only of the more complex questions of interpretation, but also of the basic question of whether that which he sees is indeed a true representation of the structure. We have taken, as a first test of the validity of a structure, its reproducibility "in situ" after repolishing and re-etching of the samples. Thus we knew the "barley shell" to be false (fig. 1) even before encountering the definitive works on the subject. The Metallography of Silicon Diffusion in Laminate Bars: As a first step in the study of the metallography of the diffusion process, we have pressed a series of laminate bars. These bars, or "sandwiches" as we have termed them, were pressed with covering layers of iron powder, and a "filler" of the Fe/Si master alloy under study. Bar dimensions were approximately %, x 3 x 1/4 in. The total

Jan 1, 1951

By E. T. Peters, L. E. Tanner

In the course of a study concerned with the rnechanical behavior of the intermetallic compound CO3V, both structural observations and anomalous electrical-resistivity behavior indicated the presence of a new high-temperature phase in an alloy with the stoichiometric composition, Co-22.4 wt pet V. The evidence, which is given below, suggests that the new phase is situated between the low-temperature ordered hexagonal phase, k,1, 2and the high-temperature disordered fee solid solution, a,3 within a small range of temperature in the vicinity of 1070°C. Structural studies included X-ray diffraction analysis* and electron-transmission microscopy examination* of stoichiometric composition speci- mens homogenized for 1 hr at 1200oC and either air-cooled or water-quenched. The air-cooled specimens yielded X-ray diffraction evidence of the k phase, while electron microscopy revealed a microstructure of both heavily and moderately faulted regions as shown in Fig. 1. The high density of stacking faults (or fine twins) in region A is reminiscent of hexagonal cobalt following its allotropic transformation.4 Selected-area diffraction from this region revealed distinct superlattice reflections of the ordered hexagonal phase, as well as continuous rel-streaks which were normal to the fault planes in the micrograph. The latter is indicative of extremely small spacings between the faults.' Regions such as ß were found to be only partially transformed to the k phase. Attempts to study the antiphase domain structure by either bright- or dark-field techniques were unsuccessful in these initial studies. This was due to the low intensity of the superlattice reflections arising from the closeness of the cobalt and vanadium scattering factors. Quenching Co3V from above Tc appears to suppress the formation of the k phase. Both X-ray and electron-diffraction results indicated that the disordered phase, a, was partially ordered during severe cooling and that the structure of the ordered phase was of the L12 (Cu3Au) type,6 designated here as a'. As in the case of the hexagonal phase, examination of the domain structure was unsuccessful; however, significant information concerning the morphology of this condition was revealed by the appearance of dislocation arrays. In particular, it

Jan 1, 1965

By C. Ernest Birchenall, Robert F. Mehl

APPLICATION of diffusion laws in the customary form to experimental studies in binary metallic solid solutions has shown the diffusion coefficient to vary with concentration for all systems investigated. The factors determining this variation have not been ascertained. However, the question of whether the diffusion coefficient would become a constant, independent of concentration throughout a single phase if the concentration gradient in the diffusion equation were replaced by an activity gradient has been an attractive subject for speculation. Wagner1 derived an expression for the force acting on a diffusing particle in terms of the chemical potential. This was also the approach chosen by Jost2 in his monograph. From a somewhat different point of view, Eyring3 and his co-workers have used activity gradients when dealing with diffusion in concentrated nonideal solutions. The only actual test with experimental data is that given by Darken,4 though this is more qualitative than quantitative. Many others opposed the application in any way of thermodynamic concepts to kinetic phenomena, including diffusion processes. A rigorous examination of this principle requires good data expressing the relationship between activity and concentration for solid solutions for which reliable diffusion work has been done. The literature provides such information only for the interstitial solution of carbon in gamma iron (austenite) and for the substitutional copper-zinc solutions (alpha brass). Activity data for carbon in austenite at 800°C and 1000°C are available from the recent work of Smiths on the equilibrium of hydrogen-methane and carbon monoxide-carbon dioxide mixtures with gamma iron. Combining these with the diffusion studies of Wells and Mehl6 makes possible a quantitative test for this interstitial solid solution. The vapor-pressure determinations for zinc over brass by Hargreaves7 allowed the calculation of activities in this system over a wide range of temperatures and concentrations. Extensive investigation of diffusion in this substitutional solid solution has been carried out by Rhines and Mehl.8 ACTIVITIES IN BRASS9 If the vapor phase of zinc is assumed to behave as an ideal gas, any solid or liquid solutions that are in equilibrium with the same partial pressure of zinc at the same temperature have the same thermodynamic activity of zinc. In order to assign a numerical value to the activity azn, it is necessary to choose some standard reference state to which an activity value of unity is arbitrarily assigned. It is customary to choose the liquid or solid state of the pure material. Then azn is defined numerically by azn = Pza/Pzn° [I]

Jan 1, 1947

By R. Floyd Farris

A method is presented for determining minimum waiting-on-cement time, which takes into account the differences that exist between types and brands of cements and such individual well conditions as depth, temperature, and pressure. The basis for the method was determined by laboratory tests. Being a laboratory development, several steps were required to prove its merit. The first step consisted of laboratory tests designed to determine the minimum cement strength required in wells. Basis was found for setting a minimum value of 8 Ib. per sq. in. tensile strength. Next, it was shown by laboratory tests that the time to 8 Ib. per sq. in. tensile strength may be expressed as a function of consistometer stirring time to loo "poises," the approximate relation being "the time to 8 Ib. per sq. in. tensile strength equals the time to I00 'poises' times three." Next, it was shown that the time of maximum temperature development in cement slurries, due to heat of hydration, is also related to consistometer stirring time to loo poises, but only by a factor of approximately two. It was shown also that the shut-in casing pressure will build up after cement is placed and register a maximum pressure at approximately the same time the slurry down the hole attains maximum temperature. From this and the relationships listed above, the general rule was established that minimum waiting-on-cement time (time to 8 Ib. per sq. in.) after casing cement jobs in any well is equal to the time when the shut-in casing pressure reaches a maximum, as measured from the initial mixing of cement, times a factor of 1.5. Cement plugs drilled in the field at the time prescribed by this formula were found to drill "firm to hard," thus confirming the laboratory tests. These tests prove that many of the present regulations for waiting on cement require a longer time than is absolutely necessary. Use of the method herein proposed offers the possibility of a saving of $1200 per well. Introduction The length of time allowed for cement to set after casing is determined either by state-wide rules, field rules, or self-imposed rules written into drilling contracts. In general, the time is dictated by experience and common practice. However, owing to differences in opinion and in experience of the various groups involved, waiting-on-cement time often varies from one area to the next. For example, an operator in an area where no rules exist may drill out of surface pipe at 24 to 36 hr., while another operator in another area may wait 48 hr. or more to comply with state or field rules, although the depth of the well, hole size, type of cement, and other data are identical. An even greater difference in practices will be found by making similar comparisons with respect to oil-string cement jobs. Differences in waiting-on-cement times of 36 to 48 hr. are common. Further complicating the picture is the rather common practice of allowing more waiting time for cement to set at the greater depths than is allowed at the shallow depths. This practice has existed for years in spite of the common knowl-edge1,2,3 that the temperature of the earth at the usual setting depths of surface

Jan 1, 1946

By R. Floyd Farris

A method is presented for determining minimum waiting-on-cement time, which takes into account the differences that exist between types and brands of cements and such individual well conditions as depth, temperature, and pressure. The basis for the method was determined by laboratory tests. Being a laboratory development, several steps were required to prove its merit. The first step consisted of laboratory tests designed to determine the minimum cement strength required in wells. Basis was found for setting a minimum value of 8 Ib. per sq. in. tensile strength. Next, it was shown by laboratory tests that the time to 8 Ib. per sq. in. tensile strength may be expressed as a function of consistometer stirring time to loo "poises," the approximate relation being "the time to 8 Ib. per sq. in. tensile strength equals the time to I00 'poises' times three." Next, it was shown that the time of maximum temperature development in cement slurries, due to heat of hydration, is also related to consistometer stirring time to loo poises, but only by a factor of approximately two. It was shown also that the shut-in casing pressure will build up after cement is placed and register a maximum pressure at approximately the same time the slurry down the hole attains maximum temperature. From this and the relationships listed above, the general rule was established that minimum waiting-on-cement time (time to 8 Ib. per sq. in.) after casing cement jobs in any well is equal to the time when the shut-in casing pressure reaches a maximum, as measured from the initial mixing of cement, times a factor of 1.5. Cement plugs drilled in the field at the time prescribed by this formula were found to drill "firm to hard," thus confirming the laboratory tests. These tests prove that many of the present regulations for waiting on cement require a longer time than is absolutely necessary. Use of the method herein proposed offers the possibility of a saving of $1200 per well. Introduction The length of time allowed for cement to set after casing is determined either by state-wide rules, field rules, or self-imposed rules written into drilling contracts. In general, the time is dictated by experience and common practice. However, owing to differences in opinion and in experience of the various groups involved, waiting-on-cement time often varies from one area to the next. For example, an operator in an area where no rules exist may drill out of surface pipe at 24 to 36 hr., while another operator in another area may wait 48 hr. or more to comply with state or field rules, although the depth of the well, hole size, type of cement, and other data are identical. An even greater difference in practices will be found by making similar comparisons with respect to oil-string cement jobs. Differences in waiting-on-cement times of 36 to 48 hr. are common. Further complicating the picture is the rather common practice of allowing more waiting time for cement to set at the greater depths than is allowed at the shallow depths. This practice has existed for years in spite of the common knowl-edge1,2,3 that the temperature of the earth at the usual setting depths of surface

Jan 1, 1946

By E. B. Swanson

There is only one way to transport natural gas and that is by pipe lines. In the past few years, these lines have been extended rapidly into areas which previously had been served mainly by solid and liquid fuels. In the subject given me, this extension is referred to as an invasion, which implies a rather forceful entry. Whether or not such an implication may be applied to recent natural-gas developments is unimportant. It is the fact of the extension, not its definition, which is important. For centuries, natural gas has been utilized for various purposes in many parts of the world. It was the foundation of fire worship at Baku. The Chinese used it for the evaporation of brine in salt manufacture. In our country, its history seemingly dates back 110 years to the vicinity of Fredonia, N. Y., where seepage gas was assembled and used for lighting purposes. A few years later, a %inch pipe was laid to a lighthouse on Lake Erie, and the early records of the Lighthouse Service tell of the difficulties surrounding this use. With the discovery of oil in Pennsylvania, there followed further development; the town of Titusville was piped with natural gas and, some few years later, gas was made available for manufacturing plants adjacent to the oil fields. For more than a half century, the industry was concentrated in the

Jan 1, 1932

By John Wulff, J. H. Hollomon

THE opinion has been widely held that the corrosion of alpha brass occurs by the selective solution of zinc. As late as 1939, Fink1 and Evans2 suggested that in the initial stage of the corrosion the selective solution of zinc produces copper-rich areas, which subsequently act as cathodes that facilitate the deposition of copper. Previously Abrams3 had shown that most of the copper found on corroded brass was redeposited and not residual, although he did not consider the nature of the initiation of the corrosion process. Enough data exist to indicate that any metal surface is sufficiently heterogeneous to provide the anodes and cathodes necessary for the initiation of corrosion.4,5 Thus, for alpha brass we need postulate not an initial selective solution of zinc, but only the existence of local differences in electrochemical potential. The heterogeneity responsible for potential differences may arise from oxide and other films as well as from alloy concentration gradients. Even though selective corrosion may be possible in some alloy systems, previous data on alpha brass as well as the experimental work reported in this paper do not require this supposition. Bengough and May,6 in a discussion of Abrams' paper, first suggested the addition of small amounts of arsenic (0.03 per cent) to alpha brass to prevent the redeposition of copper (dezincification). This suggestion was later substantiated by Nixon.7 In their studies of the inhibiting influence of arsenic, Evans and Fink were led to the selective-solution hypothesis in order to explain the initial corrosion of such alloys. They believed that an arsenic film covered anode areas and prevented the formation of copper-rich spots capable of acting as cathodes for the deposition of copper. They did not consider, however, the mechanism by which the arsenical film was formed. Bengough and May,8 in another publication, suggested that the commonly found gray or black arsenical film may act to: (I) carry oxygen to the brass surface, (2) increase the discharge potential of copper on brass, (3) precipitate cuprous chloride from colloidal suspension in chloride solutions. Masing9 assumed that the film found upon arsenical brass was metallic arsenic. His measurements of hydrogen overvoltage of bulk arsenic led him to believe that the inhibition of dezincification occurs because the overvoltage of a brass surface is lowered by the arsenic film. He reasoned that this film would permit the more rapid reduction of hydrogen at the cathode, which would prevent the redeposition of copper. In a recent paper,10 we have pointed out that the film formed during the corrosion of arsenical brass and copper in stagnant solutions is essentially arsenious oxide (As203),

Jan 1, 1942

By H. F. Yancey, O. L. Furse

Although usually coal is screened with punched plate screens having round-hole perforations, there has been an increasing use of wire-cloth screens with square apertures, especially at pneumatic coal-cleaning plants and at other plants that use vibrating screens. It is therefore desirable to establish a relationship between round-hole and square-hole screens for use not only in experimental investigations of coal-preparation problems, but also for use in commercial operations. This need is apparent whenever it is necessary to substitute wire cloth for punched plate and still retain the familiar trade-size product. Although similar work probably has been done by industrial organizations concerned with the screen sizing of coal, the only recorded study of the subject other than the one described herein of which the authors are aware is that of Roesler,1 who screen-sized two samples of limestone and two of zinc concentrates crushed to pass 1/2 in. with a set of square and round-hole screens and found the ratio to be 1:1.23. Procedure In the course of a study of various methods for determining the friability of coal,2 it was necessary to screen-size the products of a series of friability tests of three different coals with both round-hole and square-hole screens. The coals used were from the Roslyn No. 3 and the Bayne mines in Washington and the Rolapp mine in Utah. Twenty 25-lb. samples of each of these three coals at 3 to 2 1/2 in. round-hole size were dropped from a height of 10 ft. to a concrete floor, and then screened with

Jan 1, 1934

By H. Vincent Wallace

ALTHOUGH it is universally accepted that free oxygen is A necessary in a cyanide solution for the dissolution of gold and silver-in accordance with Elsner's equation that 2Au + 4KCN + 0 + H20 = 2KAu (CN), + 2KOH- little has been written and insufficient attention paid to the drop in extraction due apparently to loss of dissolved oxygen in mill solutions. Aeration or oxygenation of cyanide solutions has been more or less successfully accomplished by various types of agitators such as the Dorr (with or without air lifts), Pachuca tanks and more recently "Turbo-mixers" and Denver "Super- agitators"; also by the introduction of chemical oxidizing agents such as Solozone (Na,O,) . At the Wright-Hargreaves and other mines in the Kirkland Lake camp, the introduction of this reagent at the rate of about 1.25 oz. per ton of ore applied to the feed of the ball mills has certainly been beneficial.

Jan 1, 1932

By Harold S. Morgan

INTRODUCTION The Waste Isolation Pilot Plant (WIPP) is being developed in southeastern New Mexico by the United States Department of Energy as a research and development (R&D) facility to demonstrate the safe disposal in salt of radioactive wastes resulting from defense activities. As part of the WIPP R&D program, a series of in situ tests will be performed to determine the behavior of drifts and storage rooms in the creeping salt medium (Matalucci, et al. 1982). Data obtained in these tests will be used to evaluate and improve numerical models used to compute the structural response of these drifts and rooms. Stress has been proposed as one of the parameters to be measured in the tests, and borehole inclusion stressmeters have been included in the instrumentation package. Cages of this type provide an indirect measurement of stress change in that they actually monitor borehole displacements which are then related to changes in the in situ stress away from the borehole. The relationship between stress and displacement is highly dependent upon the constitutive behavior of the rock in which the gage is emplaced. This relationship is well understood for rocks with linear stress-strain behavior (Fossum, et al, 1976), but, for creeping rocks like salt, the relationship has yet to be determined. Thus, inclusion gage response in salt must be investigated fully before stress measurements from the WIPP in situ tests can be used for comparisons with computed stresses The calculations presented in this paper were performed as part of an investigation of inclusion stressmeter performance in salt. They focus on the strain gaged stressmeter (SGS) (Cook and Ames, 1979) and on a hypothetical solid cylindrical inclusion with variable stiffness Issues addressed include the effects of creep and the effects of gage stiffness on stressmeter output for various loading conditions. The

Jan 1, 1984

By Charles E. Glass

A steadily growing body of evidence indicates that earthquake ground motions can cause failure of rock slopes that are otherwise stable under static loading conditions. As a result, the economic optimum working slope angle is reduced when earthquake shaking is considered in surface mine design, especially in regions of high seismicity. When designing surface mine slopes to withstand seismic motion it is as important to consider the probability of experiencing various levels of ground shaking as it is to consider the characteristics of the ground motion. For most slope designs, simple probabilistic techniques such as extreme value or maximum likelihood techniques are appropriate. These techniques, however, should be applied in such a way that a probability value is assigned to characteristic ground motion parameters at the mine rather than to nearby earthquake magnitudes. Several techniques have been developed to analyze rock slopes subjected to dynamic loads. Most of these techniques do not adequately account for the ground motion parameters of concern. The best technique appears to be one that utilizes actual earthquake time histories. A simple technique that calculates block displacement for any input accelerogram, using a linear acceleration approximation between time steps, has been developed that accurately accounts for rock block response. Results using this technique indicate that the following factors are important for assessing rock slope stability in earthquake regions:

Jan 1, 1983

By M. J. Hordon, M. A. Wright

ReCENT work has indicated that the substantial decrease in the rate of fatigue crack propagation for aluminum observed below a critical vacuum level can be attributed to the depletion of oxygen or water vapor from the vicinity of the fatigue crack opening.' At reduced gas pressures, generally below 10-2 Torr with fatigue strain amplitudes in the range *1 to 2 x 10-3 and cyclic loading rates of 25 to 50 cps, estimated rates of oxide or hydroxide monolayer adsorption on newly formed crack surfaces can fall behind the actual rate of crack extension. Changes in the critical pressure required for the transition in fatigue behavior with changes in temperature, strain amplitude, and cyclic frequency can be interpreted from variations in the adsorption and crack growth rates.'" Although the adsorption model generally appears to explain the enhanced fatigue resistance of reactive metals in vacuum, the specific adsorbing gas species and nature of the adsorption mechanism are uncertain. In addition to the direct dissociation and adsorption of O2 on the metal surfaces,3 the dissociation of H20 and subsequent adsorption of H ions have been advanced for aluminum alloys.4 The present note describes some measurements of the fatigue behavior of 1100 aluminum in reduced partial pressures of specific gases including O2, H2O, and H2 in order to provide additional experimental data. Cantilever-type specimens of 1100-H14 aluminum were fatigued in cyclic reverse bending at constant strain amplitude using a multispecimen vacuum apparatus and procedure described previously.' After initial vacuum evacuation to 1 x 10-7 Torr, the appropriate gas species was bled into the system through a calibrated leak valve. The input rate was adjusted to maintain the desired pressure level under continuous evacuation, thus minimizing possible changes in gas composition during testing. High-purity input gases including 99.9 pct O2, 99.999 pct H 2, and water vapor from triply distilled H 2 O were used to provide residual pressures. Trace amounts of contaminant species were limited by passing the Oz and H 2 input gases through liquid-nitrogen-cooled copper coils prior to entry to the test chamber. Care was taken to repeatedly flush the gas lines and valves before testing to minimize the effect of adsorbed contaminants. Residual gas pressures in the range 10-6 to 5 x x3 Torr were monitored by an NRC ionization gage corrected for each gas species. In the higher-pressure range, 5 x 1CT3 to 5 x 10' Torr, an NRC alphatron gage was used. In operation, eight simultaneous fatigue tests were conducted at a constant strain amplitude of *1.2 x 10 at cyclic bending rates of 25 and 50 cps to give an average fatigue life of 6.5x 105 cycles-to-fracture at standard atmosphere and 3.3 x 106 cycles in 1 x 10"7 Torr vacuum. Representative results for the effect of residual O2, H2O, and H2 contents on the fatigue life are shown in Fig. 1 for tests performed at 50 cps. It is apparent that the presence of either O2 or H2O gas in the system at partial pressures greater than 10"3 Torr was noticeably detrimental to the fatigue life. In contrast, the enhanced fatigue resistance typical of high-vacuum environments was retained with pure H 2 through the pressure range up to 1 Torr. For partial pressures of H2 gas above 1 Torr, the fracture life noticeably decreased. The critical Hz

Jan 1, 1969

By J. D. Prater, H. R. Spedden, E. E. Malouf

With the expansion of copper leaching of the various mine wastes at the Kennecott properties, additional copper recovery facilities have been required. A research and development program has been pursued to determine a suitable economic approach to the processing of the additional leach solutions contemplated. lnvestigations on laboratory and pilot plant scales covered electrowinning, solvent extraction, launder and vat type cells, activated launders, precipitation drums, and precipitation cones. This paper contains information on the development and use of a new precipitation-cone-type copper recovery plant, which provides many advantages over older systems. The application of research findings to the old art of leaching copper from copper-bearing mine waste has resulted in a significant contribution of copper to over-all copper production. For example, at the completion of the current expansion program at the Kennecott properties in the United States, copper produced from waste dump leaching will amount to about 25% of the total production. Formerly only 2% was derived from this source. This planned expansion of copper leaching focused attention on the problem of developing more efficient methods for the recovery of copper from greatly increased volumes of copper-bearing solutions.1 The detailed investigations which have been made in laboratory pilot plant and plant tests for the recovery of copper from copper-bearing solutions have included electrowinning, solvent extraction, ion exchange and cementation with iron in launders, precipitation drums, activated launders, and precipitation cones. The results of these investigations have led to the development of a new precipitation cone-type copper recovery system providing many advantages over older methods. ELECTROWINNING OF COPPER FROM COPPER-BEARING SOLUTIONS The direct electrowinning of copper from relatively dilute solutions of copper-bearing mine water, namely solutions containing 10 to 20 lbs of copper per 1000 gal, has always been an attractive possibility.2,3 By this method, high-purity copper powder can be recovered with a power consumption of 3.5 kwh per Ib of copper. However, once the solution strength de-

Jan 1, 1967

J. B. Cohen (AVCO Mfg. Co.)-The model presented by the authors is only one of several possible mechanisms for low-temperature aging. It is true that the motion of dislocations to the copper atoms, followed by diffusion along the dislocation lines, is one such mechanism. However, the details of such a process which would :result in copper-rich platelets on (100) planes in the f.c.c. matrix are not clear. Furthermore, there is at least one other mechanism that may be used to e.xplain the authors' results, so that the data do not clearly favor the Turnbull model. The zones may form by "normal" diffusion. The diffusion constants would be higher at room temperature than extrapolated data would suggest due to vacancies trapped in the specimens in the quench. The diffusivity would then be multiplied by the ratio of the number of vacancies present at the solution temperature to that present at room temperature. These vacancies could remain in the alloy specimen for long times at room temerature, due to interaction with the copper atoms. ' The activation energy for diffusion is smaller than that for bulk diffusion at elevated temperatures for it is only that for the motion of vacancies existing after the quench; after these are used up, the aging process effectively ceases. It is somewhat questionable if the laws governing transformation kinetics can be directly applied to changes in elastic moduli, especially in the case of aging where the structure and the composition of the zones may be changing during aging, If such laws are assumed, it is possible to explain the exponential factor of % with this vacancy model. It is probable that the structure of the zones and the matrix match only on (100) planes. Certainly the precipitates that form in this system after various aging treatments show orientation relationships similar to those of the zones. Strains on the edges of the zone may set up fields which do not allow the zone to grow in an edgewise manner. The exponent in the growth of a thin plate in a direction perpendicular to its face would be %. It may also be that copper atoms are attracted to dislocations. If this is the rate-controllin factor in the growth process Ifthisthen an exponent of % or % may arise. The writer would like to know if the authors have tried plotting l/a; proportional to as might be expected for "dislocation drainage." R. H. Doremus (General Electric Research Laboratory)—An important new method for following the growth of Guinier-Preston zones is presented in this paper, and I wish to congratulate the authors on their experiments. Their interpretation of the results, however, is doubtful in certain respects. To determine the exponent m in Eq. [4] the authors plotted -log at against t to various powers (their Fig. 6). The initial value of frequency used for these plots was found by extrapolating the data to zero time. In Fig. A the data from their Fig. 2 are plotted against a linear, rather than logarithmic, time scale. To obtain graphs similar to their Fig. 6 for these data one must use an initial frequency of about 52720; such a plot for m = % is shown in Fig. B. The straight line in Fig. B fits the data in an intermediate range only. The curve drawn in Fig. A was mapped from this line; comparison of

Jan 1, 1960

By W. M. Duncan, F. S. Kendorski, C. V. Jude

After shotcrete had been in use for a short time al the Climax Mine of the Climax Molybdenum Co., Climax, Colo., it was noted that regular longhole undercutting blasts would severely damage or detach slusher drift shotcrete linings. If this were to be a usual occurrence, the widespread application of shotcrete would have to be reexamined or abandoned. Concrete drift linings are affected less by blasting because of occasional significant gaps between the placed concrete and rock and the greater thickness of poured-in-place linings. Research Project Organized A research project was organized to determine: 1) the seriousness of the problem, 2) more complete in-use shotcrete strength data, and 3) possible solutions.

Jan 12, 1973

By C. Hanford Henderson

The belt of copper-bearing rocks of the South Mountain has attracted Attention principally iii the townships of Hamilton Ban and Liberty in Adams County, Pennsylvania, and in the southeastern portion of Franklin County. As the developments, thus far, have been almost exclusively limited to that area, the examination here described extended only within those boundaries. The region, being one of considerable interest geologically, has been the subject of careful study and description, both on the part

Jan 1, 1884

By G. Das, S. V. Radcliffe

The substructural features developed in tungsten as a function of annealing temperature (up to 2200°C) and type of material [undoped and doped powder metallurgy (PM) tungsten and electron beam melted tungsten] have been investigated by transmission electron microscopy. For doped PM tungsten wires, characteristic "par-ticulate" substructural features developed rapidly with increase in annealing temperature above 700°C. The features consisted of parallel rows of elongated or circular shapes (500 to 1000A diam) lying along the direction of the wire axis and were identified as internal voids by diffraction contrast experiments. In recrystallized doped PM rod, larger voids were observed and were identified by precision dark field analysis to be cubic in shape and bounded by (100) planes. In marked contrast with both the doped PM materials, recrystallized undoped PM rod exhibited only very occasional and randomly arranged voids. Furthermore, no voids were observed in either material after electron beam melting. The high concentration of voids in the doped PM materials is attributed primarily to vaporization of doping additions or their pvoducts situated at the original grain boundaries , whereas the few voids in undoped material are considered to be traces of microporosity which were not eliminated during sintering. A tentative mechanism is suggested for the dezlelopment of the voids in relation to the processing sequences (sintering and working) and to the subsequent annealing. In recent years, a characteristic substructural feature consisting of rows of small elongated or circular regions of light contrast lying along the direction of working has been seen in thin foil electron microscopy studies of annealed sheet, wire, and rod tungsten. These features were present in the published micrographs of sheet by Weissmann et al.1 and of wire by Meieran and Thomas2, although the authors did not draw attention to them. Wronski and Fourdeux3 observed similar features in sintered rod tungsten (it was not specified whether or not the material was doped*) and interpreted them on the basis of their ap- particles, based on extraction replica evidence from the fracture surface of the initial hot-rolled slab material from which the sheet was prepared. No diffraction contrast experiments on the features were reported in any of these studies. The present investigation was undertaken with the primary objectives of: a) identifying the nature of these substructural features in tungsten by electron diffraction contrast experiments, since the contrast for voids can be expected to differ from that for crystalline or glassy particles, and b) elucidating the origin of the features and their development. For the latter purpose, doped and undoped powder metallurgy tungsten was obtained as rod and wire to represent different stages of reduction during final processing. These materials were examined both in the as-processed condition and after annealing to successively higher temperatures. In addition, the same doped and undoped materials were examined after vacuum melting in rod form. I) MATERIALS AND PROCEDURE Doped powder metallurgy (PM) tungsten wire (commercial purity 99.9 pct W) was obtained in the as-drawn and surface ground condition (0.030 in. diam "ground seal rod"). Doped and undoped tungsten rod (0.075 in. diam) representing an earlier stage of final processing was obtained from the same commercial source (Refractory Metals Division, General Electric Co.). Lengths of both the doped and undoped rod materials were single-pass melted in an electron-beam zone refiner to examine the effect of vacuum melting on the substructure. Annealing was carried out in a tungsten crucible in a tantalum strip resistance furnace under a vacuum of l0-15 mm Hg. Longitudinal sections of the wire and rod materials were examined by light and electron microscopy. The preparation of thin foils suitable for electron transmission from 0.030 in. diam tungsten wire and the rod specimens was carried out by means of a high-precision microjet technique developed to provide lack of jet stability and precise control of the area thinned. The method is described in detail elsewhere.' The foils were examined in a JEM 6A electron microscope using a goniometer stage (±20 deg tilt, 360 deg rotation) and operated at 100 kV. To minimize contamination problems a 200 µ condenser aperture was used in conjunction with a useful beam current of 50 µA. II) RESULTS AND DISCUSSION A) Diffraction Contrast Analysis. In order to determine the optimum conditions for the development of the substructural feature, a series of isochronal 30 min annealing experiments were carried out on specimens of the doped PM tungsten wire. The transmission electron microscopy analysis showed that the as-drawn wire, Fig. 1(a), consists of 'fibers' whose long axis is closely parallel to the wire axis of (110). The fiber width averages some 0.5 µ. Dense disloca-

Jan 1, 1969

By C. S. Simons

Manufacturers of filtering, thickening or clarifying equipment, have been active in developing machines to meet increasingly rigid operator requirements, and the innovations resulting from this effort are moving into commercial applications. In general, the results of these applications are reported to be good. Little has been reported on progress in theory. The growing insistence on the part of pollution control agencies that waste water discharged into streams contain less and less suspended solids is being felt by a broad cross section of industry. A new clarifier, based on the often rediscovered fact that suspended solids settle more rapidly in inclined channels, is being marketed by Neptune MicroFLOC of Corvallis, Ore. This unit employs modules of inclined tubes, mounted in the upper zone of the clarifier, which carry the water flow. These were developed for light solids, such as those encountered in domestic water treatment and in paper mill waste water. In these applications, flow rates up to about three times those permissible in conventional clarifiers are reported to give comparable overflow turbidity. By combining these units with mixed media filters, high clarity may be obtained with minimum capital. Use of such units in the metallurgical industry, particularly where water reuse or effluent quality is important, should grow.

Jan 1, 1970