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By P. J. Spencer, J. N. Pratt

The vapor pressure of manganese over liquid Mn-Sn alloys has been determined by a high-temperature torsion-effusion technique. Alloys containing from 8 to 100 at. pct Mn were investigated in the temperature range jrom 1280" to 1580" and the measured pressure values were used to calculate the partial and integral thermodynamic properties of the liquid alloys. The activities show small negative departures from ideality while the integral heats and excess entropies of mixing are asymmetric inform, changing from positive to negative with increasing manganese content. The possible contribution of various factors to the observed thermodynamic properties is discussed. COMPARATIVELY few thermodynamic data are available for manganese alloy systems.' Therefore, as part of a continuing program of studies of the thermodynamic properties of transition metal alloys, measurements have been made on various binary alloys involving this element. In recent publications,2~3 investigations of liquid Mn-Cu alloys and of the Mn-Au system in both solid and liquid states have been reported. For the first-mentioned system the work suggested that magnetic interactions may be responsible for the observed form of the thermodynamic properties, while in the second the influence of the electrochemical factor appears to be dominant. The present paper describes a similar study of liquid Mn-Sn alloys. Again the thermodynamic properties have been obtained from vapor pressure measurements made by use of a high-temperature torsion-effusion technique.4 A detailed description of the apparatus and of the experimental procedures used in alloy preparation and pressure measurement may be found elsewhere.2'4 EXPERIMENTAL RESULTS Sixteen alloys, ranging in composition from 8 to 100 at. pct Mn, were prepared from spectroscopically standardized manganese of 99.99 pct purity and tin of 99.999 pct purity, both supplied by Johnson-Matthey and Co., Ltd. One-gram samples of the alloys were obtained by carefully weighing appropriate amounts of the pure components into an effusion cell; this was then suspended in the apparatus and the metals melted together in situ by heating under vacuum to approximately 1550°K. After allowing sufficient time for a homogeneous liquid alloy to be formed, vapor pressure measurements were commenced. These were determined as rapidly as possible at a series of steady temperatures within the range of interest. The duration of experimental runs on individual samples was kept sufficiently short to ensure insignificant varia- tion of alloy composition during investigation. After completing pressure measurements, the alloys were rapidly cooled and their compositions checked by weighing or chemical analysis. All experiments were conducted using effusion cells machined entirely from boron nitride. Measurements were made using a variety of cells with orifice areas ranging from 0.0032 to 0.0075 sq cm and lengths of the order of 0.04 cm; the usual effusion correction factors for orifice geometry and molecular distribution were calculated using Freeman and Searcy's equation5 and had values between 0.6 and 0.75 for the orifices employed here. The vapor pressures of manganese over the alloys were measured at approximately 20°K intervals in the temperature range 1280" to 1580°K. In view of the close approximation of the measured pressures to Clausius-Clapeyron behavior in the experimental temperature range, the data for each alloy have been expressed by equations of the form: logp(atm) =-A/T + B A least-squares computer treatment was applied to the vapor pressure values in order to obtain the coefficients A and B with their associated error. The resulting equations are listed in Table I, together with the equation for pure solid manganese obtained from a previous study.4 To minimize the effect of possible apparatus calibration errors, the activities and partial free energies of manganese in the alloys were calculated by initial reference to the latter equation, obtained from identical torsion-effusion measurements. The immediately resulting thermodynamic quantities, based on a solid manganese reference state, were then converted to refer to the more appropriate supercooled pure liquid manganese standard; tabulated values of the free energies of solid and liquid manganese from Hultgren et al.' were used for this purpose. Partial entropies of solution of manganese were calculated from the temperature coefficients of the free energies and partial heats from the Gibbs-Helmholtz relationship.

Jan 1, 1969

By Edward W. Parsons

To determine in what part and to what extent creep or movement exists in an installed rock-bolt anchor and the effect of the creep or bleedoff on the tension load on a rock bolt, as well as the effect of normal mining activity on the rock-bolt load and anchor, the Bureau of Mines conducted tests in the laboratory and in two mines using four types of expansion-shell anchors, the wedge bolt, and in one mine, explosive-anchored rock bolts. Specially mounted heavy-duty springs, calibrated to obtain load readings from spring movement, were installed on the test bolts to prevent a drastic reduction in tension load caused by slipping or seating of the anchor. A dial gage, reading to 0.0001 inch, measured movement of the anchor shell against the rock and seating of the anchor wedge plug in the shell. While all bolts had electrical resistance strain gages for load reading, only readings from the laboratory tests were used because those from the field tests were erratic. The tests revealed that most load loss resulting from static loading is due to the slippage and seating of the serrations on the anchor shell into the drill-hole wall. Dynamic loading, primarily from sudden shocks from nearby blasting, and possibly from ground movement or mining activity, will cause the anchor wedge plug to seat in the anchor shell.

Jan 1, 1969

By John M. Gomes

The preparation of tungsten carbide by molten-salt electrolysis was investigated. Sodium tungstate in the electrolyte and carbon from the graphite anode were the source materials, Optimum results were obtained with an electrolyte consisting of 83 mole-percent sodium chloride and 5.7 mole-percent each of sodium tungstate, sodium metaborate, and sodium hydroxide; a temperature of 1,000° to 1,025° C; an anode current density of 37 amp/dm2 or less; and an initial cathode current density of 150 amp/dm2. The yield was approximately 0.45 g/amp hr. The electrolytic tungsten carbide product contained approximately 5.4 weight-percent combined carbon as compared to 6.1 weight-percent for pure tungsten carbide. Microhardness of the electrolytic tungsten carbide, determined with a 50-gram load, varied from 3,480 to 3,970 dph. The material was friable and could be ground in a laboratory ball mill to less than 400-mesh particle size.

Jan 1, 1969

By J. A. Beesley

This paper reviews contemporary international law concerning the seabed, current developments in the United Nations on the seabed, and the main outstanding problems. It emphasizes the importance that the Canadian government attaches to this area.

Jan 1, 1969

By R. L. Bell, T. G. Langdon

A technique was developed- for determining the grain strain, and hence the grain boundary sliding contribution, occurring during the high- temperature creep of a magnesium alloy, from the distortion of a grid photographically printed on the specimen surface. The results were compared with those obtained from measurements of grain shape, both at the surface and interrwlly, and it was concluded that the grain shape technique may substantially underestimate the grain strain and overestimate the sliding contribution due to the tendency for migration to spheroidize the grains. ALTHOUGH a considerable volume of work has been published on the role of grain boundary sliding in high-temperature creep, many of the estimates of Egb (the contribution of grain boundary sliding to the total strain) have been in error due to the use of incorrect formulas or inadequate averaging procedures.' One of the most easy and convenient measurements from which to compute Egb is that of v, the step normal to the surface where a grain boundary is incident. Unfortunately, this parameter is also the one associated with the treatest number of pitfalls. Values of v have been used to calculate Egb from the equation: egb =knrVr [1] where k is a geometrical averaging factor, n is the number of grains per unit length before deformation, v is the average value of v, and the subscript ,r denotes the procedure of averaging along a number of randomly directed lines. If the dependence of sliding on stress were assumed, it would be possible, in principle, to calculate k from the known distribution of angles between boundaries and the surface. This in itself is difficult because the distribution depends on the history of the surface,' but the problem is even further complicated by the fact that v depends on other factors such as the unbalanced pressure from subsurface grains.3 However, the great simplicity of the measurement procedure for v makes it highly desirable that this problem of k determination should be overcome. In the present experiments, this was achieved by the use of an indirect empirical method in which the grain strain, eg, at the surface was determined by the use of a photographically printed grid. The assumption here is that the total strain, et, is simply the sum of that due to grain boundary sliding, egb, and that due to slip or other processes within the grains, eg. SO that: Et = Eg + Egb [2] Thus k is given by: In practice, it is customary to indicate the importance of sliding by expressing it as a percentage of the total creep strain; this quantity is termed y (= 100Egb/Et). The determination of Eg from a printed grid within the grains avoids the difficulties due to boundary migration which should be considered when the grain strain is calculated from measurements of the average grain shape before and after deformation. As first pointed out by Rachinger,4,5 however, this latter technique has the particular advantage that it can also be applied in the interior of a polycrystal. Recently, several workers have produced evidence on a variety of materials6-'' to support the observation, first made by Rachinger on aluminum,4,5 that 7 can be very high, 70 to 100 pct, in the interior, even when the surface value, determined from boundary offsets, is very much lower.10'11 Although there have been criticisms both of the shortcomings of the grain shape technique'' and of the different procedures used to determine y at the surface,' it seemed important to check whether measurements of sliding by grain shape gave values of y which were truly representative of the material. In the present experiments, grain shape measurements were therefore made both at the surface and in the interior for comparison with one another and with the independent measurements of grain strain using the surface grid technique. EXPERIMENTAL TECHNIQUES The material used in this investigation was Magnox AL80, a Mg-0.78 wt pct A1 alloy supplied by Magnesium Elektron Ltd., Manchester. Tensile specimens, about 7 cm in length, were prepared from a 1.27-cm-diam rod, with two parallel longitudinal flat faces each approximately 3 cm in length. The specimens were annealed for 2 hr in an oxygen-free capsule, at temperatures in the range 430° to 540°C, to give varying grain sizes, and, prior to testing, the grain size of each was carefully determined using the linear intercept method. This revealed that the grains were elongated -0.5 to 5 pct in the longitudinal direction. Testing was carried out in Dennison Model T47E machines under constant load at temperatures in the range 150" to 300°C. At temperatures of 200°C and below, tests were conducted in air with the polished flat faces coated with a thin film of silicone oil to prevent oxidation; at higher temperatures, an argon atmosphere was used. To determine v,, each test was interrupted at regular increments of strain and the specimen removed from the machine. At the lower strains, when v, was less than about 1 pm, measurements were taken on a Zeiss Linnik interference microscope;

Jan 1, 1969

By C. L. Magee, H. W. Paxton

PREVIOUS workersl'% ave found that as-quenched ferrous martensites can be plastically deformed at low stresses. They have also found that the stress to obtain small strains can be significantly increased by aging at temperatures up to 150°C. This behavior has been interpreted in two ways: 1) that the deformation at low applied stress is due to internal stresses and thus the aging phenomenon is due to relaxation of the internal stresses;' 2) that the microdeformation arises through the motion of unpinned dislocations and the aging process is the diffusion of carbon to the dislocations to pin them.293 In a recent paper,4 we have shown that under certain conditions the dominant microdeformation mechanism of as-quenched Fe-31Ni marten-site-austenite aggregates is the transformation of retained austenite. Our results on resistance changes during deformation and temperature dependence of the deformation also established that, at temperatures above 300° K, the microdeformation was not due to this transformation. The purpose of this note is to report results of experiments which indicate that explanation 1 is appropriate for this deformation. All results presented here are from conditions where transformation does not make a measurable contribution to the microplastic deformation, in Fe-31Ni-0.002C, or in Fe-29Ni-0.02C-0.25Ti. The experimental techniques that were used for measurement of the plastic deformation are described4 and analyzed5 elsewhere. Fig. 1 shows the microplastic response of 100 pct austenite at 300" and 350° K. The microplastic response of an austenite-martensite aggregate is also given in Fig. 1 and it is seen that, after formation of martensite, the initial plastic deformation occurs at lower stresses. Thus, the softness is not entirely due to retained austenite—the quenched structure can be deformed at lower stresses despite the fact that the macroscopic strength is significantly higher. A result which allows a distinction to be drawn between the alternative explanations for this behavior is derived from aging experiments on the low-carbon alloy containing titanium. We have previously shown4 that the presence of titanium in this alloy reduced the free carbon to a sufficiently low level that thermal stabilization of the austenite-martensite transformation was not detected. Despite this, Fig. 2 shows that the microplastic response is decreased by aging treatments. A similar aging effect in Fe-ZONi, with some titanium, has previously been reported by . Thus, carbon is not necessary for the aging process and explanation 2 is at best incomplete. Since internal stress relief can occur independently of carbon (any temperature dependence of the flow stress can result in decreased internal

Jan 1, 1969

By W. Jolley

Decarburized iron and Fe-3.28 pct Ni alloys were impact and tension tested in the temperature range of ambient to 4°K. It was found that alloying with nickel improves the fracture properties of the ferrite matrix. The improvement in toughness of iron alloyed with nickel is attributed to the effect of the nickel addition on the temperature and strain rate dependence of the yield stress as well as on slip behavior. The Fe-Ni alloy exhibits wavy slip when strained at temperatures down to at least 50°K. This behavior contrasts with that of the unalloyed iron where a change It is a generally accepted fact that of the common alloying elements of steel only nickel and manganese improve both the room-temperature yield strength and the notch impact properties. It has been argued for a number of years that manganese refines the ferrite grain size, the pearlite spacing, and the grain boundary carbides in medium- and low-carbon steels and thereby increases resistance to brittle fracture.''' This hypothesis recently was confirmed for a 0.04 pct C iron containing manganese, but at the same time from wavy to planar slip occurs at about 120°K. The observed lower temperature and strain rate dependence of the lower yield stress in the Fe-Ni alloy as compared to the unalloyed iron may be attributed to this difference in slip behavior. In contrast to suggestions published previously, transmission electron microscopy showed that the uniform dislocation distributions and the lack of dislocation tangles which are produced by low-temperature straining of bcc materials cannot be attributed to the absence of cross slip at these temperatures. no increase in notch toughness was observed in a decarburized Fe-Mn alloy.3 It thus appears that manganese only increases the toughness of Fe-Mn-C alloys by refining grain boundary carbides and has no toughening effect in carbon-free iron. On the other hand, the mechanism whereby nickel improves toughness has not been established. Hodge et a1.4 have clearly shown that a 3.28 pct Ni addition lowers the impact transition temperature of 0.02 pct C steel approximately 35°C when the grain size is maintained constant. Recently, Wullaert and Tetelman6 stated that nickel refined the grain boundary carbides of an Fe-C alloy and consequently increased the toughness. They also observed the previously reporteda-'

Jan 1, 1969

By D. Atkinson, C. Bodsworth, I. M. Davidson

A geometric model of interstitial solid solutions, which has been used previously as a basis for the prediction of carbon activities in Fe-C austenite, is shown to serve also for the calculation of nitrogen activities in Fe-N austenite. The model has been developed to enable predictions to be made of the activities of an interstitial element in the presence of two host atom species. The activities calculated via the model are shown to be in satisfactory agreement with the measured values in the austenite phase for carbon in Fe-C-Co, Fe-C-Cr, Fe-C-Ni, Fe-C-~n, Fe-C-Si, and Fe-C-V alloys and for nitrogen in Fe-N-Ni alloys. The effect of the second substitu-tional solute on the logarithm of the activity of the interstitial element is expressed as the product of a constant mad the atomic concentration of that solute. The constants so derived we related to the thermo-dynamic interaction coefficients which describe the effect on the activity coefficient of carbon of an added solute element. In recent years the thermodynamic activities of carbon and nitrogen in the single-phase austenite field have been determined for iron binary alloys and for several iron-base ternary alloys. In order to extend the use of these measurements, it is desirable to be able to predict with reasonable accuracy the activities of the interstitials at compositions and temperatures other than those which have been measured experimentally. In all the systems studied to date, the interstitial elements do not conform to ideal behavior. Hence, the available data cannot be extrapolated or interpolated using the simple thermodynamic concepts of solutions. Several models have, therefore, been formulated for the purpose of predicting the activity of an interstitial element in the presence of one species of host atom. These models can be divided into the geometric1"5 and energetic6-' types. The former group is based on the assumption that at low concentrations the activity of the interstitial species is determined by a composition-dependent configurational entropy term and an excess free-energy term which is temperature-dependent but independent of composition. The purpose of this paper is to show that the treatment, based on a geometric model, can be extended to enable predictions to be made of interstitial activities in the presence of two substitutional host atom species. THE CONFIGURATIONAL ENTROPY OF MIXING ICaufman5 has shown that the configurational entropy, S,, for a binary solution comprising of a host atom species, A, and an interstitial species, I, can be expressed as: where NI is the atom fraction of the interstitial species, R is the gas constant, and (2 - 1) is the number of interstitial sites excluded from occupancy by the strain field around each added interstitial atom. The number of interstitial sites per host atom, p, is unityg for the fcc austenite solutions considered here. The configurational entropy of mixing for a ternary solution comprising two substitutional atom species, A and B, and one interstitial species, I, can be derived similarly. Let the number of atoms per mole of each of these species in the solution be represented by «a, ng, and nI. From geometric considerations, it is improbable that the addition of a few atom percent of a second host atom species will change the type of sites (i.e., octahedral) in which the interstitial atom can be accommodated in the austenite lattice. At higher concentrations (determined largely by the relative atomic radii of the atomic species present and any tendency to nonrandom occupancy of the host lattice sites) other types of interstitial sites may become energetically favorable. Restricting consideration to compositions below this limit, for 1 = 1 the number of suitable interstitial sites is given by (n + nB). However, if each interstitial atom excludes from occupancy (Z - 1) additional sites, the total number of sites available for occupation is reduced to (n + ng)/Z. The number of vacant interstitial sites is given by: The total number of recognizable permutations of the atoms must include the recognizable, different configurations of the A and B atoms on the host lattice. Assuming that these arrangements are purely random, and are not affected by the presence of the interstitial species, the total number of recognizable permutations in the ternary alloy is given by: The configurational entropy is obtained by expanding, using Stirling's approximation, and collecting like items, as:

Jan 1, 1969

By Martin Landwehr

"Dust samplers are based on different principles. The konimeter is based on the number of particles counted; mining applications are based on routine sampling; and the photometrical evaluation of samples is used after the pre-separation of particles above 5 micron by the mining konimeter. In special cases, the Draeger sampler is used for quantitative analysis. General trends are toward larger samples, concentration on a weight basis, fractionation into size groups by different techniques and an analysis of composition. Standards or TLVs already exist in all industries. Surface industries base their criteria on concentrations and quartz contents in particles below 5 micron. The mining industry uses the tyndalloscope routinely to determine concentrations; concurrent] y, the rock content, as deter-mined by ashing samples, is obtained by konimeter.Threshold limits are related to the rock content in respirable fractions and tyndallometrically determined concentrations. The TL Vs are based empirically on Landwehr's ""dust values,"" as agreed upon with the Mines Inspectorates. Recently, and in accurate sampling, dust samplers have been used that count the dust in two fractions -above and below 5 microns. Mineral composition is determined either for all minerais or just for quartz. In assessing the dust hazard by Landwehr's ""dust values,"" all minerais are taken into consideration; in coal dust sampling, however, only the quartz content is taken into account. Efforts are now being made to co-ordinate the different assessment methods, arriving at uniform MAC values in agreement with those of other countries."

Jan 1, 1969

By W. E. Wallace

Volume-rate-of-flow measurements were made of the flow of mercury through porous samples under conditions of no field, magnetic field transverse to the direction of flow, and transverse magnetic field and electric current to determine if a magnetic method might be potentially useful in the study of pore-size distribution. No change was observed in the flow rate from the application of a magnetic field alone. The application of both field and current produced an electromagnetic pumping action with or against the direction of flow, depending upon the direction of the electric current. This pumping effect was observed to be greater in a rectangular porcelain sample than the calculated value for an equal-Sized open channel.

Jan 1, 1969

By R. P. Laforce, ZJ. R. Low, A. M. Turkalo, D. F. Stein

The interaction of the crlloying eletnenls, nickel and chromium, with the impurity elements, antimony, pIzosphorus, tin, and arsenic, to producse reversible temper brittleness in a series of high-purity steels containing 0.40 wt pct C has been investigated. The alloyed steels contained approximately 3.5 pcl Ni, 1.7 pct Cr, and 0.05 to 0.08 pct of the particular irnpurity to be investigated. Susceptibility to teirlper embrittlement was measured by comparing the notched-bar transition temperature of each steel after quenching from the final temper and after very slow cooling (step cooling;) following the final temper. A plain carbon steel without alloying elements, bu/ ud/h 0.08 pel Sh, does not embrittle when step-cooled through the emzbrittling range of temperatures. The same embrittling treatment, applied to a steel with about the same antinzony content but with nickel and chvonziunz added, causes a 700°C increase in transition temperature. If chromium or nickel is the only alloying element, the increase in transition temperature is only 50%, again with antimony present. A carbon-free iron containing nickel, chromium, and antimony shou~s a 200°C shift in transition temperature for the same thermal treatment. Specific alloy-impurily interactions are also observed for the other impurity elements, phosphorus, tin, and arsenic. Additional investigations involving electron microscopy, trzicrohard-ness tests of vain boundaries, minor additions of zirconiutn and the rare earth and noble metals, nzainly with negative results, are also described. HE particular type of embrittlement investigated is that which is encountered in alloy steels tempered in the temperature range from about 350" to 525'C or slowly cooled through this range of temperatures when tempered above this range. This type of embrittlement is sometimes called reversible temper brittleness to distinguish it from the embrittlement indicated by a minimum in the room-temperature V -notch Charpy energy vs tempering-temperature curve encountered in the range 28 0" to 350°C. Temper brittle-ness seriously restricts the use of many alloy steels since it precludes tempering or use in the embrittling range of temperatures and may significantly raise the ductile-brittle transition temperature of heavy-section forgings and castings tempered above the embrittling range, since such sections cannot be sufficiently rapidly cooled after tempering to avoid embrittlement. The very voluminous literature of temper brittle-ness up to about 1960 has been reviewed by woodfine' and LOW.' Of particular significance to the present investigation was the demonstration by Balajiva, Cook, and worn3 that high-purity Ni-Cr steel does not exhibit temper brittleness and the subsequent detailed and systematic study by Steven and Balajiva~ of the effect of impurity additions on the susceptibility to embrittlement of Ni-Cr steels. Steven and Balajiva showed that, of the impurities which may be found in commercial steels, Sb, As, P, Sn, Mn, and Si could all produce temper brittleness in a high-purity Ni-Cr steel. The principal purpose of the present investigation was to study the effects of particular alloy-impurity combinations on susceptibility to temper embrittlement. The steels used were high-purity 0.30 to 0.40 wt pct C steels containing 3.5 wt pct Ni and 1.7 wt pct Cr, separately or in combination. The susceptibility of these steels was then determined when approximately 500 ppm by weight of antimony, arsenic, phosphorus, or tin were added as an impurity. The melting, casting, and forging practices used in the preparation of the materials investigated are described in Appendix A. Table A-I in this appendix shows the analysis of all steels to be discussed. The steels were produced as 20- or 2-lb heats. The smaller heats were used after it had been demonstrated (see Appendix B) that a small, round, notched test specimen could be used to measure the shift in the ductile-brittle transition temperature caused by temper brittleness with about the same result as that obtained by Charpy testing. HEAT TREATMENT Unless otherwise noted, all steels were tested for embrittlement in the tempered martensitic condition. A typical heat treatment for a 0.40 C, 3.5 Ni, 1.7 Cr steel was: 1 hr at 870"C, in argon, quench into oil at 100"C, quench into liquid nitrogen, temper 1 hr at 625"C, and water-quench. The warm oil quench was used where quench-cracking was encountered; otherwise the initial quench was into room-temperature oil or water. For other compositions austenitizing temperatures were 50°C above Acs with the remainder of the thermal cycle the same. Steels in this condition, with no further heat treatment, are designated as non-embrittled. The above quenching and tempering cycle for the 0.40 pct C steels resulted in as-quenched hardnesses of 48 to 53 RC and as-tempered hardnesses of 24 to 31 Rc except in the case of the plain nickel or plain carbon steels. In these, the as-tempered hardness was as low as 80 to 90 Rg. No attempt was made to adjust the tempering temperature to obtain the same hardness in ali steels since it was felt that a uniform thermal cycle was more important than exactly equivalent hardness values. Pro- the standard quench and temper described above, the standard embrittling treatment was "step-cooling". For this the thermal cycle was: 593"C, 1 hr; furnace-cool to 538"C, hold 15 hr; cool to 524"C, hold 24 hr; cool to 496"C, hold 48 hr; cool to 468'C, hold 72

Jan 1, 1969

By J. C. White

This paper, after a brief introduction giving the history and geology of the area, deals with mining operations at the Copper Mountain mine of Gaspe Copper Mines, Limited. It discusses development operations, and then goes into the mining methods employed, covering drilling, grade control, blasting practice, loading and hauling, and the road and service machines used.

Jan 1, 1969

By J. D. Verhoeven

In the past 15 years a considerable amount of experimental and theoretical work has been done concerning the onset of convection in liquids as a result of interm1 density gradients. This work, which has been doue in many different fields, is reviewed here and extended slightly to give a rrlore quantitative understanding to the probletrz of conzection in liquid metal dlffusion experinletzts. In liquid metal systems the capillary reservoir technique is currently used, almost exclusively, to measure diffusion coefficients. In this technique it is necessary that the liquid be stagnant in order to avoid mixing by means of convection currents. Convective mixing may result from: 1) convection produced as a result of the initial immersion of the capillary; 2) convection produced in the region of the capillary mouth as the result of the stirring frequency used to avoid solute buildup in the reservoir near the capillary mouth; 3) convection produced during solidification as a result of the volume change; and 4) convection produced as a result of local density differences within the liquid in the capillary. The first three types of convection have been discussed elsewhere1-a and are only mentioned for completeness here. This work is concerned only with the fourth type of convection. Local density differences will arise within the liquid as a result of either a temperature gradient or a concentration gradient. It is usually, but not always, recognized by those employing the capillary reservoir technique that the top of the capillary should be kept slightly hotter than the bottom and that the light element should be made to migrate downward in order to avoid convection. In the past 15 years a considerable amount of work, both theoretical and experimental, has been done in a number of different fields which bear on this problem. This work is reviewed here and extended slightly in an effort to give a more quantitative understanding of the convective motion produced in vertical capillaries by local density differences. The Stokes-Navier equations for an incompressible fluid of constant viscosity in a gravitational field may be written as: %L + (v?)v = - ?£ + Wv - g£ [1] where F is the velocity, t the time, P the pressure, p the density, v the kinematic viscosity, g the gravitational acceleration, and k a unit vector in the vertical direction. A successful diffusion experiment requires the liquid to be motionless, and under this condition Eq. [I] becomes: where a is the thermal expansion coefficient [a =-(l/po)(dp/d)], a' is a solute expansion coefficient [a' = -(l/po)(dp/d)], and the solute is taken as that component which makes a' a positive number. Combining with Eq. [3] the following restriction is obtained: Since there is no fixed relation between VT and VC in a binary diffusion experiment, Eq. [5] shows that the condition of fluid motionlessness requires both the temperature gradient and the concentration gradient to be vertically directed. Given this condition of a density gradient in the vertical direction only, it is obvious that, as this vertical density gradient increases from negative to positive values, the motionless liquid will eventually become unstable and convective movement will begin. The classical treatment of this type of instability problem was given by aleih' in 1916 for the case of a thin fluid film of infinite horizontal extent; and a very comprehensive text has recently been written on the subject by handrasekhar.' It is found that convective motion does not begin until a dimensionless number involving the density gradient exceeds a certain critical value. This dimensionless number is generally referred to as the Rayleigh number, R, and it is equal to the product of the Prandtl and Grashof numbers. For the sake of clarity a distinction will be made between two types of free convection produced by internal density gradients. In the first case a density gradient is present in the vertical direction only, and, since the convection begins only after a critical gradient is attained, this case will be called threshold convection. In the second case a horizontal density gradient is present and in this case a finite convection velocity is produced by a finite density gradient so that it will be termed thresholdless convection. Some experimentalists have performed diffusion experiments using capillaries which were placed in a horizontal or inclined position in order to avoid convection. These positions do put the small capillary dimension in the vertical direction and, consequently, they would be less prone to threshold convection than the vertical position. However, if the diffusion process produced a density variation, as it usually does, it would not be theoretically possible to avoid thresh-

Jan 1, 1969

By J. D. Wright

This paper describes current milling practice at Endako Mines Ltd., a molybdenum producer in central British Columbia. The various steps in the crushing, grinding and flotation of the ore are outlined, with emphasis on the metallurgy of the latter operation. Tailings disposal and concentrate assaying methods are also described.

Jan 1, 1969

This list is part of the long-range conservation program of the Bureau of Mines and is intended for use by those interested in the Bureau's publications dealing with oil shale and shale oil. The items included represent 51 years of research by Bureau of Mines employees on oil shale and shale oil and are available at any U.S. depository library. DISTRIBUTION U.S. Bureau of Mines report series contain both free and sales publications. An asterisk (*) designates publications for sale. Such publications can be obtained from Superintendent of Documents U.S. Government Printing Office Washington, D.C. 20402

Jan 1, 1969

By A. J. Barry

This report presents observations made by Bureau engineers regarding hydraulic-powered roof supports that are currently being used or have been used on nearly all operating longwalls in the United States, A comparative summary of specifications for various types of powered supports with a brief description of the mine conditions has been compiled to enable mining engineers and mine operators to readily evaluate the differences in design data. Some of the support data furnished in the field may vary from those listed in manufacturers' brochures because of the trend to customize various support components to meet specific mine conditions. A method of estimating support requirements for different assumed operating conditions is presented. Similar computations for specific mine conditions may be used to make preliminary estimates of support requirements.

Jan 1, 1969

By H. F. Feldmann

This Bureau of Mines report presents a method for the computer solution of the equilibrium compositions of any complex, multiphase, chemical system for which appropriate thermodynamic data are available. The method describes in detail the computation of systems in which solid or liquid species can form by heterogeneous chemical reactions where these species do not occur in any equilibrium equations. Also described are the calculation of phases formed by condensation (pure or solutions) and the inclusion of other system restrictions such as electroneutrality. The method is being used to compute the equilibrium composition of mixtures of an electrically conductive plasma and a liquid slag phase arising from the high-temperature combustion of coal for magnetohydrodynamic power generation. These mixtures can contain up to 65 species including gases, liquids, free electrons, and ions.

Jan 1, 1969

By P. Bolsaitis, L. Skolnick

A model of metallic solutions is postulated which explains the energy of formation of alloys on the basis of changes in electron density around solute and solvent atoms and changes in pairwise interactions between atoms. The model is based on the Wigner-Seitz cell model of metals, with the electronegativities of the metals as the parameter for the electron-attracting power of the constituent atoms. In contrast to models based on the description of cohesive energy in terms of pairwise interactions only, which require an empirical, adjustable correction parameter to the Berthe-lot relations in order to predict negative energies of formation, the electron cell model accounts for both negative and positive energies on the basis of parameters of the pure components. Calculated heats of mixing based on published parameters of the pure metals are found to compare favorably to experimental values, despite the absence of any adjustable parameters in the model. THE formulation of an "exact" quantum mechanical model for the energy of cohesion of metals reduces ultimately to the many-body problem and for a complete solution would require the knowledge of wave functions and energy levels of all the particles in the system. The case of alloys is further complicated by the fact that the potential fields around the constituent atoms are different and the matching of boundary conditions in terms of the resulting wave functions is still an unresolved mathematical problem. It appears that in spite of the numerous investigations that have led to an elucidation of many properties of the metallic state, the formulation of the Hume-Rothery rules in 1931 marks the lone accomplishment toward simple, qualitative prediction of the energy of formation of alloys. More recently Friedel'' and Arafa have utilized the modern concepts of the electronic structure in metals in combination with the Hume-Rothery rules to arrive at quantitative estimates of the energy of formation of alloys with low solute contents. The relative success of the Hume-Rothery rules and Friedel's and Arafa's studies encouraged broaching the subject of formulation of a model based on realistic physical parameters and of conceptual and mathematical simplicity that could be applied to alloys regardless of concentration. On the basis of the free-electron theory, the cohesive energy of pure metals can be expressed as a function of "valence" (number of free electrons per atom), lattice parameter (or density), and the ioniza-tion energy. In addition to the electronic contributions to cohesion, one has to consider the pairwise exchange interaction energy between neighboring ion cores and the contribution from van der Waals forces. The energy changes upon alloying are viewed as resulting primarily from a redistribution of electron densities around the ions (a simplification of the more basic concept of a change of the wave functions of the electrons in the presence of a perturbing potential field). In addition there is a change in the pairwise interactions between neighboring ions which is usually the only factor taken into account in commonly used adaptations of the quasichemical theory. The electronegativities of the metals forming the alloy are used as the measure of the electron-attracting power of the ions in the alloy and the difference in electronegativities as the driving force for changes in electron energy on alloying. The interplay of pairwise interactions and electronic forces results in positive or negative energies of formation of alloys. Results obtained for systems to which the free electron model is best applicable are in good agreement with experimental values. 1) COHESIVE ENERGY OF METALS IN TERMS OF THE ELECTRON CELL MODEL The simplest physical model of a metal is that of a lattice of positive charges immersed in a sea of uniform, negative charge distribution represented by the

Jan 1, 1969

By M. E. Wadsworth, T. D. Chatwin

The kinetics of sulfidization of hematite (Fe203) by H2S gas at various partial pressures have been determined over the temperature range of 170° to 500°C. The process was controlled by thin films of pyrite (FeS2) and troilite (FeS) which were protective and formed diffusional layers. At higher temperatures the thermodynamic drive to form FeS2 decreased, and consequently the rate of formation of FeS, decreased. Conditions for the development of thin layers of FeSz consisting of 10 to 100 atomic diameters are presented. Laboratory flotation results on synethetic ores indicate floatability identical to that of pyrite of normal xanthate flotation. The prospect of high-temperature sulfidization of metal oxides using H2S, S, or other sulfur bearing compounds is intriguing as a potential activation process for flotation since the thermodynamics are very favorable for a large number of systems. If the kinetic parameters are sufficiently well understood it may be possible to provide surface sulfide layers a few angstroms thick, thus imparting normal sulfide floatability characteristics to the oxide. For thin layers the total sulfur pick-up would be small except for very high surface area materials but in most cases this sulfur could be readily removed by roasting if necessary. The general reactions for the sul-fidization of metal oxide MxO, to a metal sulfide MpSq using H2S may be written for the three conditions py > xq; py < xq; and py = xq: A large thermodynamic advantage is obtained in many systems because of the relatively large negative free energy of formation of H,O compared to H,S. For the temperature range 2S° to 1500°C &apos; where AFOis the standard free energy and T is absolute temperature. As is evident from Eqs. 4 and 5, the free energy of formation of H ,0 is approximately 35 to 37 kilocalories per mole greater than that for H,S over broad variations of temperature. For oxide-sulfide systems as outlined above, sulfidization may be expected to occur for the conditions: Groups I and 111: that is, if the terms on the right are more negative than those on the left the reaction is expected to occur thermodynamically. The value hFg,(,) is included in Eq. 6 since hFLzs is based upon gaseous S2 rather than the normal reference state, orthorhombic sulfur. The

Jan 1, 1969

By G. R. Speich

The interlarnmelm spacing, growth rate, and degree of segregation that accompany cellular precipitation in four Fe-Zn alloys containing 9.7, 15.2, 23.5, and 30.5 at. pct Zn have been determined in the temperature range 400" to 600°C. The chemical free-energy change for the reaction was calculated from the available thermodynamic data and the known compositions of the phases. The fraction of the chemical free-energy change for equilibrium segregation that is converted into interfacial free energy decreases from 0.43 to 0.08 as the magnitude of this free-energy change increases from 35 to 270 cal per mole. At constant temperature the cellular growth rate is proportional to the cube of the dissipated free energy. At 600°C newly 100 pct of the equilibrium segregation is achieved during cellulm precipitation whereas at 400°C only 85 pct of the equilibrium segregation is attained. During cellular growth, mass transport of zinc occurs by grain boundary diffusion; excess zinc remaining in the a! phase after the completion of growth is removed slowly by volume diffusion. A modified Cahn theory of cellular precipitation predicts the observed interlamellar spacing within a factor of two. In cellular precipitation reactions such as pearlite formation or discontinuous precipitation, the basic problem is to predict the variation of growth rate G, interlamellar spacing S, and degree of segregation P with composition and temperature. To accomplish this we need three independent equations relating these quantities. One of these equations comes from the diffusion solution. To obtain two additional independent equations, some assumptions must be made. cahnl has suggested recently that two plausible assumptions are 1) that growth rate is proportional to the dissipated free energy and 2) that the spacing which occurs is that which maximizes the dissipated free energy. According to the first assumption, this spacing also maximizes the growth rate and the rate of decrease of free energy per unit area of cell boundary. The present work was undertaken to test these assumptions. To test the first assumption it is necessary to study a cellular reaction over a wide range of supersatura-tions to establish a relationship between G and the dissipated free energy at constant temperature. This is possible only in discontinuous precipitation reactions since in pearlite reactions constituents other than pearlite form if the composition of the parent phase deviates even slightly from the eutectoid composition. The Fe-Zn system was chosen for study because 1) discontinuous precipitation proceeds to completion over a wide temperature and concentration range, 2) the degree of segregation within the cell can be measured by lattice parameter measurements,2 and 3) the thermodynamics of this system have recently been determined by Wriedt.3 In this system the cells consisting d alternate lamellae of a and r phases form from supercooled iron-rich a phase. The a phase within the cells is bcc as is the original a phase, cia, but has a different orientation and a slightly lower zinc content than the original a phase. The r phase has a zinc content of about 70 at. pct and a crystal structure isomor-phous with T brass. EXPERIMENTAL PROCEDURE Four Fe-Zn alloys with 9.7, 15.2, 23.5, and 30.5 at. pct Zn were prepared from carbonyl-iron powder (400 mesh, 99.8 wt pct Fe) and zinc powder (200 mesh, 99.99 wt pct Zn). The powders were ball-milled together and cold-pressed under 60,000 psi to discs $ in. thick by $ in. diam. The cold-pressed discs of the alloys with 9.7 and 15.2 at. pct Zn were sealed in evacuated silica capsules and heated slowly to 1100°C over a period of 1 week (3 days at 600°C, then 3 days at 80O°C, then 1 day at 1100°C). The alloys with 23.5 and 30.5 at. pct Zn were treated similarly except that the final homogenization temperatures were 1000" and 85O°C, respectively, to prevent melting. The alloys were quenched in iced brine from the final homogenization temperature. Specimens of each alloy were subsequently aged in salt pots at temperatures of 400°, 450°, 500°, 550°, 600°, and 650°C for times that varied from a few minutes to several hundred hours. At a late stage of this work, an alloy containing 11.2 at. pct Zn was prepared by vapor-impregnation of iron foil with zinc vapor at 890°C. This alloy proved useful for electron microscope studies because it was free of porosity. The homogenization and aging conditions were based on the recent Fe-Zn phase diagram of Stadelmaier and Bridgers4 rather than the earlier diagram of ansen.5 They consist of a homogenization heat treatment in the homogeneous a field followed by an aging treatment in the two-phase a + r field. The aged specimens were metallographically polished and etched in 2 pct nital and the radius of the largest cell in the microstructure determined. This radius plotted vs time gave a straight line whose slope is the boundary migration rate or growth rate G of the cell. To determine the interlamellar spacing, specimens were examined by surface-replica and thin-section electron microscope techniques. Because of the irregular nature of the lamellae within the cell, the average interlamellar spacing S .of the cell was measured by the method of Cahn and Hagel,6 where S is defined by:

Jan 1, 1969