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By P. Niessen, S. J. Bercovici

RECENTLY Aust et al.&apos; proposed a model of non-equilibrium grain boundary segregation based on a vacancy gradient induced uphill diffusion process of solute to grain boundaries. According to their proposal, solute atoms which are attracted to vacancies would, during the cooling of a sample, migrate with vacancies toward grain boundaries (vacancy sinks) and thereby enhance the solute concentration in boundary regions. For solute atoms which are not attracted to vacancies no such boundary solute enrichment would be expected. These concepts were given experimental support by boundary and bulk mi-crohardness measurements of various thermally treated dilute zinc and lead alloys and also doped aluminum oxide. The underlying assumption in their work was that solute enrichment in boundary regions is indicated by higher microhardness values. From thermodynamic considerations1&apos;2 it appears that the solid-liquid solute distribution coefficient (K), defined as the concentration of solute in the solid to the concentration of solute in the liquid at equilibrium, indicates whether attractive or repulsive forces exist between vacancies and solute atoms. K < 1 indicates a solute-vacancy attractions, K > 1 indicates no, or a repulsive, vacancy-solute interaction. This note is to report on observations of different grain boundary etching phenomena in high purity zinc alloys. These phenomena appear to depend on whether a K < 1 or K > 1 solute is present. In experiments designed to measure directly the solute enrichment in grain boundary regions, samples of zone refined zinc with 0.1 at. pct solute additions were cooled from 380°C at a rate of 3°C per min. A surface layer of approximately 50 was then removed from the samples by chemical polishing. The samples&apos; surfaces were then protected by a chromating treatment such that in a subsequent etch only grain boundary regions were dissolved. The dissolved regions were approximately semicircular troughs, 10 µ in width. The grain boundaries were usually in the middle, i.e., at the bottom, of this trough. For zinc containing the K < 1 solutes cadmium and aluminum, the grain boundary was always indicated by a ridge as shown in Fig. 1. For pure zinc, and zinc containing the K > 1 solutes copper and gold, a groove at the bottom of the trough indicated the boundary position, Fig. 2., To test if these different grain boundary etching behaviors were possibly connected with the segregation process discussed above, the Zn-Cd alloy was again heated to 380°C and quenched into brine. After surface protection and etching it was found that this specimen no longer exhibited the grain boundary ridges but rather grain boundary grooves as shown in Fig. 2. It would appear, therefore, that the boundary ridge is a consequence of a process occurring during the slow cooling of the samples. From absorption spectrophotometric analysis it was found that for slowly cooled samples containing K < 1 solutes, the solute/zinc ratios of the material from the boundary trough were consistently considerably higher than the solute/zinc ratios of the bulk alloys. This boundary region solute enrichment was not observed in quenched samples. It is concluded therefore, that the boundary ridge shown in Fig. 1 is directly or indirectly connected with nonequilibrium solute segregation.

Jan 1, 1970

By R. G. Aspden, D. A. Colling

THE {120}(001) texture has been reported in near 50 pct Ni-Fe alloys <0.004 in. thick1,2 and in copper 0.014 in. thick.3 This texture is one of two types of secondary recrystallization textures observed to replace a cube texture matrix and is desirable in Fe-Ni alloys because of its unique magnetic characteristics.1,2 However, the conditions for and mechanism of the {120}(001) formation are not well understood.3 During an investigation of the influence of alloying elements on the magnetic characteristics of 49 pct Ni-Fe alloys with a sheet thickness of 0.014 in., we have observed the growth of {120}(001) secondaries when strong deoxidizers such as aluminum and calcium were present. It is believed that this observation in unique and provides additional insight with regard to the mechanism for {120}(001) texture formation. Three alloys were prepared with electrolytic iron and nickel and contained a nominal 49 wt pct Ni. Alloy A was a straight binary Fe-Ni alloy, whereas intentional additions of 0.013 wt pct Ca to alloy B and of 0.59 wt pct A1 to alloy C were made. Ingots were forged at 1100 C, hot-rolled to 0.150 in. gage, pickled, and cold-rolled to 0.014 in. Standard U-shaped magnetic test specimens were punched, deburred, sprayed with MgO, and annealed in ~50°C dew point hydrogen. Textures were identified with the aid of etch pits, the

Jan 1, 1970

By H. W. Kerr, B. Soutiere

THE possibility of using unidirectionally solidified eutectic alloys as high-strength composite materials has been investigated by several workers. Unidirectionally solidified lamellar and fibrous eutectics have been shown to have good short-time structural stability at temperatures close to the eutectic melting temperature.&apos;-4 Several investigations of both short-time high temperature tensile properties and creep properties of eutectic alloys have been carried out,3-" with promising results. Little work has been reported, however, on the thermal stability of eutectic alloys which have been unidirectionally solidified and subsequently deformed in compression. Since processes such as rolling and grinding introduce compressive stresses, and since the properties depend on the spacing and perfection of the composite structure10,11 this aspect is of definite importance if unidirectionally solidified eutectic alloys are to be considered for commercial use. A Cd-Zn alloy made from 99.999 pct starting materials was cast into a graphite boat and zone levelled12 to obtain a completely lamellar structure at one end of the bar. Sections from this region of the bar were removed, recast into 6.5 mm I.D. Pyrex tubing, and solidified unidirectionally upwards at a rate of 2.8 cm per hr. Specimens were cut from the solidified rods, heated in air at 200°C (0.88 of the eutectic melting temperature) for various lengths of time, and examined metal-lographically. Compression specimens were also cut from the solidified rod carefully machined to a length of & in., and deformed with the compressive axis parallel to the original growth axis to produce a total longitudinal strain of 12 pct. Some of these specimens were examined metallographically; others were heated in air at 200°C for various lengths of time and then examined. Metallographic examination of the deformed specimens showed the two types of deformation previously described; slip lines13 (smoothly bent lamellae) and kinks13,14 (sharply bent lamellae). Very close to the specimen ends there were a large number of deformation bands, but at a distance of about -re- in. the deformation was very localized, as shown in Fig. 1.

Jan 1, 1970

By S. Floreen, R. M. Pilliar, H. W. Hayden

The fracture behavior of a 50 pct Cu-50 pct Fe mi-croduplex alloy, laminated composites of copper and iron and an extruded 50-50 Cu-Fe elemental powder composite was studied. Very low ductile-brittle transition temperatures were achieved in all cases, but for different reasons. In the microduplex alloy both the initiation and also the propagation of cleavage fractures appeared retarded by the very small in-terphase distances. In the composites, crack propagation through the sumples was prevented in most cases by delamination fractures perpendicular to the advancing cracks. These delaminations occurred at different regions and by different mechanisms in the various composites. In the extruded powder composite, de-lamination appeared to take place along preexisting flaws. In the crack arrest geometry of the laminated plates, delamination took place by localized shear fractures within the copper near the Fe-Cu interfaces. In this case delamination was enhanced by thicker laminate layers, and by having the resistance to shear failure of the copper sufficiently low compared to the toughness of the iron. BRITTLE fracture in engineering materials has long been a problem, and many different ways of preventing it have been considered. One method that has been of growing interest lately is to prevent crack propagation by the introduction of mechanical discontinuities into the structure. These discontinuities may act in several ways. They may simply act as crack stoppers. They may introduce secondary fractures such as de-laminations that deflect the initial crack into new, less damaging directions. Alternatively, they may subdivide a fairly large bulk sample that would have been loaded in plane strain, for example, into a number of subunits that are individually loaded in plane stress and thus are more resistant to fracture. Other mechanisms, or combinations of mechanisms, are also feasible. A number of methods exist for introducing mechanical discontinuities into a structure. Composites by their nature have discontinuities in structure, and numerous studies have shown that fracture propagation in materials of this type can be radically changed by suitable control of the composite parameters. Of particular significance to the present work are recent investigations of layered composites made by joining high strength steel sheets by various means.&apos;-4 These studies have shown that through proper control of the mechanical properties of the bonds joining the sheets it was possible to introduce delamination fractures that markedly improved the overall toughness of the composites and in some cases completely prevented through-the-thickness fractures. Another technique for introducing structural discontinuities is simply to use a two-phase alloy. It has been recognized for many years that a small amount of a second phase may improve toughness either by homogenizing plastic flow and thus preventing localized stress concentrations that nucleate fracture, or by interacting with an advancing crack. In most of these studies of two-phase materials, the decreases in ductile-brittle transition temperatures produced by the second phase were relatively small. More recently, work on two-phase stainless steels having a very fine grain microduplex structure has shown that the presence of on the order of 40 to 50 pct of a tougher second phase may lower the ductile-brittle transition temperature of the brittle phase by approximately 300°F. 5-7 In these alloys delaminations were seldom observed. The tougher second phase appeared to minimize the ease of both the initiation and the propagation of cleavage fractures. These results show that both the composite approach and the microduplex alloy approach are effective methods of preventing brittle fracture. Therefore, it was of interest to compare the fracture behavior of a microduplex alloy with composites made from the two-phases that were present in the alloy. To simplify this comparison the 50 pct Cu-50 pct Fe system was selected for study. At low temperatures the equilibrium tie line phases in this system are essentially pure ferrite and pure copper. A 50-50 alloy was cast and hot worked to produce a microduplex structure. Two types of composites were studied; laminated structures prepared by roll bonding iron and copper sheets of the tie line compositions, and an extruded powder composite made from high purity elemental powders. The fracture behavior of these materials was then compared. EXPERIMENTAL PROCEDURE Alloy Preparation. The 50-50 Fe-Cu alloy and the components for the roll bonded composites were prepared by vacuum induction melting 30-lb heats using electrolytic grades of iron and copper as charge materials. A carbon boil was used to deoxidize the melts. Small additions of copper and iron were made to the iron and copper heats, respectively, to approximate

Jan 1, 1970

By C. S. Giggins, F. S. Pettit

The oxidation of Ni-Cr alloys in 0.1 atm of oxygen has been studied at temperatures between 800" and 1200°C. For alloys with 30 wt pct or more Cr, continuous layers of Cr2O3 are formed during oxidation. In the case of alloys with chromium concentrations between approximately 5 to 30 wt pct, external scales of Cr203 are formed over grain boundaries whereas internal precipitates of Cr2O3 and external layers of NiO are formed at other areas on the alloy surface. When such conditions are present on the alloy surface, chromium diffuses laterally from those areas covered with a continuous layer of Cr2O3 to areas where a Cr2O3 sub scale exists and it is possible for the sub-scale zone to become separated from the alloy by a continuous layer of Cr2O3. Whether such a state will be attained depends upon the initial grain size of the alloy and the oxidation time. When the concentration of chromium in the alloy is less than 5 pct, Cr2O3 is formed internally both at grain boundaries and within the interior of grains and the alloy is covered with an external layer of NiO. MECHANISMS which describe the growth of oxide scales on nickel-base superalloys are complex and the effects produced by the various elements in these alloys on the oxidation behavior of superalloys are not clearly understood. In order to determine the influence of the different elements on the oxidation behavior of superalloys, it is first necessary to examine the oxidation properties of binary nickel-base systems which contain the principal elements present in the superalloys and then progressively more complex systems until compositions typical of the superalloys are attained. Chromium is present in virtually all nickel-base superalloys and the purpose of the present studies was to examine the selective oxidation of chromium in Ni-Cr alloys. The oxidation characteristics of Ni-Cr alloys have been extensively studied1-" to date principally as a result of the high oxidation resistance exhibited by some of these alloys. Ni-20Cr* has long been known *All compositions are given as wcight percent unless specified otherwise. to be oxidation resistant and is commonly used as resistance heating elements for service temperatures up to 1100°C. This alloy cannot be used for extended periods of time at higher temperatures because of the apparent reaction of the external scale with oxygen to form gaseous CrO3. In spite of the considerable work cited above some important aspects of Ni-Cr oxidation still remain unresolved. Virtually all of the previous studies agree that small additions of chromium to nickel, e.g., <10 wt pct Cr, result in increased oxidation rates as compared to that of pure nickel, whereas larger additions, e.g., 20 to 30 wt pct Cr, form alloys with substantially lower oxidation rates. The controversial aspects of the oxidation mechanisms for these alloys that still remain unresolved are as follows: 1) A description of the oxidation mechanism for the low chromium alloys. 2) A description of the oxidation mechanism for the high chromium alloys, particularly with respect to the composition of the external scale which results in the lower oxidation rates. 3) The specific alloy compositions at which the oxidation mechanism changes from that obtained for low chromium contents to that of the high chromium alloys and the reason for this transition. EXPERIMENTAL The Ni-Cr alloys listed in Table I were prepared from high purity metals by nonconsumably arc melting and casting as buttons. These alloys were then given a preliminary annealing treatment in argon at 815°C for 100 hr to promote homogeneity. Each button was cut into 0.250 in. thick sections that were subsequently cold-rolled to 0.050 in. thicknesses and annealed in argon at 815°C for 48 hr to provide a twinned, equi-axed grain structure. The grain size for these alloys was not uniform and the limits, within which the average grain size lies, are given in Table I for the single-phase alloys. All the alloys were single phase with the exception of the Ni4OCr alloy in agreement with the Ni-Cr phase diagram.&apos;&apos; Rectangular specimens were cut from the sheet to provide surface areas of approximately 2.5 sq cm. Exact areas were determined with a micrometer after surface preparation was completed. All of the specimens except the Ni-40Cr alloy and pure chromium were polished through 600-grit Sic abrasive paper, ultrasonically agitated in ethylene trichloride, rinsed with ethyl alcohol, and electro-polished. The specimens were electropolished in a 10 vol pct H2SO4 (conc), 6 vol pct lactic acid, methyl alcohol solution at 70" to 80°C for 2 min at a current density of 0.8 to 1.2 amp per sq cm. This electro-polishing procedure did not produce acceptable surfaces on the Ni-40Cr alloy nor on pure chromium and the oxidation properties of these materials were obtained for specimens polished through 600-grit Sic

Jan 1, 1970

By R. O. Williams

Analytical representation of the thermodynamics of solutions is highly desirable from the standpoint of accuracy, compactness, and numerical manipulations. In particular, computer calculations are greatly implemented. Mathematical considerations show that previous expressions have one or more serious defects. This investigation shows a Fourier series to be satisfactory but that it is also possible to derive a new series which fits certain additional conditions. Included examples show the value of analytical expressions in giving a simple characterization of each system using some two to five parameters, the elimination of the Gibbs-Duhem integration, and the es timation of the error for the experimental function as well as derived functions. It is further shown that the present characterization provides easy comparison between systems. IN the past, thermodynamic calculations have depended to a considerable extent on tabular and graphical methods. As the volume and precision of such data increase such methods become less satisfactory. Specifically, the selection of the optimum representation and the estimation of errors require statistical methods which in turn require analytical representation. The utilization of such data require further manipulations which are best done analytically for maximum precision. For example, phase equilibria are determined by common tangents to free-energy curves: a graphical determination is normally of low accuracy. As computers are increasingly used analytical representations become almost mandatory. Insufficient mathematical consideration has been given previously to the selection of empirical expressions. Those expressions having some theoretical justification are generally too inflexible and mathematically unattractive. We consider the problem in some detail and show that a Fourier series can be effectively used. Also a new series is defined which has certain advantages. ANALYSIS We wish to consider the analytical representation of the heat of mixing, AH, the excess free energy, ?Gxs, and the excess entropy, ?sXS, as a function of composition, X, for binary solutions relative to the pure components in the same state. When a distinction is not required, we use W to denote any one of the above functions. One may use a Taylor expansion around X = 0 to generate a power series. As the derivatives are un- known we represent the series as W = A + BX + CX2 + DX3 + EX4 + ... [l] where the constants A , B, C , ..- are to be selected to provide some optimum fit. For the extremes of composition W is necessarily zero so it follows that A = 0 [2a] B +C + D + E +••• = 0 [2b] Nonelectrolytes, which we are considering, appear to satisfy the condition that d3W/dx3 = 0 [3] in the terminal regions. This is the basis of the a, ß, and Q functions used by Hultgren et al.&apos; and others. While this condition does not have a strong theoretical basis it does appear desirable that any analytical relation should satisfy this condition. Darken2 and Turk-dogan and Darken3 have shown that many systems exhibit this behavior over an extended range from each terminal region, departure being restricted to a limited intermediate region. Since we have no a priori knowledge as to where this transition occurs we can require that this condition be satisfied only as a limit at the extreme compositions as a general condition. We will show later how more restricted conditions can be included in specific solutions. Darken2 has called this behavior the quadratic formalism; we call our application the limiting quadratic formalism, LQF. This condition applied to the above power series requires that D = 0 [4a] 4-3-2E +5-4-3_F + 6 • 5 . 4G + ••• =0 [4b] The form of the power series normally used, due to Margules,4 is W=X(1-X)(A + BX + CX2 + DX3 + EX4 + •••) [5] where A, B, C, --. are a new set of constants. (Guggenheim5 has given a variation of this expression in a more desirable form. Since, however, it is contained in the above expression it does not require separate consideration.) This form is precisely what results by incorporating the conditions in Eq. [2] into the power series and regrouping the constants. The LQF requires that B =C [6a] and 4.3.2(D-C) +5-4-3(E-D) + ••• =0 [6b] Thus, the correct form of the Margules expression with two adjustable parameters is w =X(1-X)[A + B +X2-2/3x3)] 171 and the EX4 term must be included before three adjustable parameters are permitted.

Jan 1, 1970

By G. R. Speich

The distribution of carbon and the type of substructure in iron-carbon martensites containing 0.02 to 0.57pct C has been studied in the as-quenched condition and after tempering at 25" to 700°C by using electrical resistivity, internal friction, hardness, and light and electron microscope techniques. in marten-sites containing less than 0.2 pct C, almost 90 pct of the carbon segregates to dislocations and to lath boundaries during quenching; in martensites containing greater than 0.20 pct C, appreciable amounts of carbon enter normal interstitial positions located far from defects. Tempering martensites with carbon contents below 0.20 pct at temperatures below 150°C results in additional carbon segregation to dislocations and to lath boundaries but no carbide precipitation whereas -carbide precipitation occurs in martensites with carbon contents exceeding 0.2 pct. Above 150°C, a rod-shaped carbide (either Fe3C or Hagg) is precipitated in all cases. At 400°C, spheroidal Fe3C precipitates at lath boundaries and at former aus-tenite grain boundaries. At 400" to 600"C, recovery of the martensite defect structure occurs. At 600" to 700°C, recrystallization of the martensite and Ost-waW ripening of the Fe3C occur. The effects of the carbon segregation that occurs during quenching and the subsequent substructural changes that occur during tempering on martensite tetragonality, hardness, and precipitation behavior are discussed. A mathematical analysis of carbon segregation during quenching is presented. RECENT studies of the strength of low-carbon martensitel-4 emphasize the importance of carbon segregation to the martensite lath boundaries and to the dislocations contained between them during quenching. Unfortunately, very few studies of the tempering of low-carbon martensites have been conducted, so the exact nature of this segregation is poorly understood. In fact, most early tempering studies5,6 were restricted to carbon contents greater than 0.20 pct. Moreover, these studies did not determine the amount of carbon segregated to the martensite substructure during quenching so that the initial state of the martensite was not established. Aborn7 studied the precipitation of carbide in low-carbon martensite during quenching but did not establish whether carbon segregation occurs prior to carbide precipitation, nor did he study the subsequent tempering sequence in detail. In the present work we have used electrical resistance and internal friction measurements, supplemented by electron transmission microscopy to establish the carbon distribution in as-quenched specimens. Specimens thin enough to avoid carbide precipitation (but not carbon segregation) were employed. The redistribution of carbon on subsequent tempering below 250°C was followed by measurements of elec- trical resistance. Additional studies were made on specimens tempered at 250" to 700°C to elucidate the overall tempering behavior of low-carbon martensites, including the formation of cementite and recrystalli-zation of the martensite. EXPERIMENTAL PROCEDURE Eight iron-carbon alloys with 0.026, 0.057, 0.097, 0.18, 0.20, 0.29, 0.39, and 0.57 wt pct C were prepared as 8-lb ingots by vacuum melting. Typical impurities in wt ppm were 40 Si, 20 Mn, 30 S, 10 P, and 10 N. These alloys were hot rolled to 3 in. plate at 1095°C) (2000°F). The hot-rolled plates were surface ground to remove scale and the decarburized layer, then cold rolled to 0.010 in. sheet. Specimens cut from the sheet were austenitized for 30 min at 1000°C (1830°F) in a vacuum tube furnace in which the pressure did not exceed 2 x 10-3 torr. Chemical analysis of specimens after austenitization indicated no decarburization at this pressure. Immediately before quenching, the furnace was filled with prepurified helium. The specimen was then pushed rapidly through an aluminum foil gasket, which sealed the bottom of the furnace, into an iced-brine bath (10 pct NaC1, 2 pct NaOH). The quenching rate at the M, temperature is about 104&apos;c per sec for 0.010 in thick specimens, as calculated from Newton&apos;s law of heat flow2 using a heat transfer coefficient of 25 ft-&apos;. This quenching rate is sufficiently high so that all the alloys transformed completely to martensite throughout the entire 0.010 in thickness and no carbide precipitation occurred in the martensite. All specimens were immediately transferred to liquid nitrogen after quenching and stored there until needed. Tempering below 250°C (480°F) was done in silicone oil baths thermostatically controlled to *;"C. Tempering above 250°C was done in circulating air furnaces or lead pots with the specimens contained in evacuated silica capsules. Electrical resistance was determined by measurement of the potential drop across both a standard resistance and the specimen, connected in series. All resistance measurements were made in liquid nitrogen (77K, -196°C) to minimize thermal scattering of electrons and thus maximize the contribution of impurity scattering to the resistance. Specimen dimensions were 5.10 by 0.19 by 0.025 cm. Although the precision in the electrical resistance measurements was +0.1 pct, the electrical resistivities could only be measured with an accuracy of +5 pct because of uncertainty in the specimen dimensions. Internal friction measurements were performed in an inverted pendulum apparatus at vibration frequencies of either 1.9 or 66 Hz. The specimen dimensions were 5.10 by 0.375 by 0.025 cm. Hardness measurements were made with a Leitz-Wetzlar microhardness machine with loads of 100 g. Specimens were examined by light microscopy after etching in 2 pct Nital and by electron transmission microscopy after preparation of thin sections by electrolytic thinning in a chromic-acetic acid solution.

Jan 1, 1970

By C. S. Giggins, F. S. Pettit

The oxidation properties of Ni-Cr alloys with fine grains, coarse grains, and deformed surface layers have been studied at temperatures of 900" and 1100°C in 0.1 atm of oxygen. The oxidation rates of alloys containing between 10 and 30 wt pct Cr have been found to be dependent upon the grain size of the alloy. Finegrained alloys had smaller oxidation rates than coarsegvained alloys because of the selective oxidation of chromium at alloy grain boundaries. In this compositional range alloys with deformed surface layers behaved similar to fine-grained alloys due to recrys-tallization of the deformed surface layer. In the preceding paper1 it was found that during the oxidation of Ni-Cr alloys, the volume fraction of precipitated Cr2O3 could be greater at alloy grain boundaries than at other areas of the alloy surface. In the case of alloys with chromium concentrations equal to or greater than 30 pct,* the volume fraction of Cr2O3 *AIL compositions are given as weight percent unless specified otherwise. precipitated at grain boundaries and within grains on the alloy surface both exceeded the critical amount required for lateral growth of the Cr2O3 particles and the surfaces of these alloys were completely covered with a continuous, external layer of Cr203 during oxidation. However, in the case of alloys with chromium concentrations between approximately 5 to 30 pct, the volume fraction of precipitated Cr2O3 exceeded the critical value required for external scale formation only at grain boundaries.but not within the interior of the grains. Consequently, the surfaces of these alloys had external scales of Cr3O3 over the grain boundaries but internal Cr2O3 subscales with external scales of NiO away from the grain boundaries. Under these latter conditions, it was found that chromium could diffuse laterally in the alloy from those areas covered with an external layer of Cr2O3, i.e., grain boundaries, to areas where the Cr2O3 was present as a subscale. This diffusion of chromium resulted in an increase in the volume fraction of Cr2O3 precipitated in the sub-scale zone and continuous layers of Cr2O3 could be formed at the subscale front in these regions. For the alloys used in the previous studies,&apos; continuous layers of Cr2O3 were formed on Ni-20Cr alloys in the subscale regions after approximately 30 hr of oxidation at 900°C. For shorter periods of oxidation, the Cr2O3 layer was semicontinuous with the continuous portion at the subscale front emanating from points where the Cr2O3 had been formed as an external scale over alloy grain boundaries. Some lateral growth of a Cr2O3 layer in the subscale region was observed on Ni-15Cr and even Ni-10Cr alloys but this layer was never continuous after 30 hr of oxidation. These results indicate that the selective oxidation of chromium in Ni-Cr alloys with chromium contents between 5 to 30 pct may be dependent upon the grain size of the alloy. Fine-grained specimens in this compositional range should have a larger fraction of the surface covered with external Cr2O3 than coarsegrained specimens and the subscale areas required to be sealed via lateral diffusion of chromium should be smaller. It is therefore to be expected that a continuous layer of Cr2O3 can be formed on alloys in this compositional range after short periods of oxidation providing the alloy grain-size is sufficiently small. bo studieS2,3 have established that the oxidation behavior of alloys can be significantly influenced by pretreatments which produce mechanically deformed surfaces. It has been found that deformed surfaces usually promote the selective oxidation of elements in alloys and it is believed that these effects are due to rapid diffusion of elements in the deformed layer. In view of the previous results,&apos; which showed that alloy grain boundaries may play an important role in the selective oxidation of chromium in Ni-Cr alloys, deformed surfaces may promote the selective oxidation of elements in alloys as a result of the numerous grain boundaries formed on the alloy surfaces via recrys-tallization during heating to the oxidation temperature. The purpose of the present studies was to determine the effect of alloy grain size and surface deformation on the selective oxidation of chromium in Ni-Cr alloys at temperatures of 900" and 1100°C in 0.1 atm of oxygen. EXPERIMENTAL The average grain diameter of the alloys used in the previous studies1 was not less than 0.04 mm and alloys with compositions between 5 and 30 pct chromium had average grain diameters between 0.04 and 0.14 mm. Since the oxidation kinetics were already available for these relatively coarse-grained alloys, it was desirable to use these same alloys in the present studies. The surfaces of the alloys listed in Table I of the previous paper were deformed by using a Model F S.S. White Industrial Airbrasive Unit, which delivered a controlled mixture of 25 µ Al2O3 particles in a stream of dry air at high velocity against the surface of a specimen. The amount of surface deformation produced by this treatment was not determined but a re-crystallized layer about 15 µ thick was formed upon annealing deformed specimens. The grain size at the surface of the specimens was reduced to an average grain diameter of 0.01 mm by annealing the deformed specimens, i.e., grit-blasted,

Jan 1, 1970

By A. J. Heckler, J. A. Peterson

A capsule technique was successfully employed to investigate the effect of nickel on the activity of nitrogen in Fe-Ni-N austenite in the temperature range 600" to 1200°C. This technique consisted of equilibrating nitrogen among various Fe-Ni alloys within a sealed silica capsule. Nitrogen transfer among the specimens occurred by N, gas at 900°, lOOO? and 1200?C. Nitrogen gas pressures within the capsules were estimated to be as high as 22 atm. The activity coefficient of nitrogen, fN , in Fe-Ni-N austenite is adequately described by the linear interaction equation: log . wt pct Ni where the standard state is chosen such that fN = I as wt pct Napproaches zero in binary Fe-N. This relationship was determined over the temperature range 873" to 1473°K and for nickel contents of 0 to 35 wt pct. ALTHOUGH chemical thermodynamics of liquid iron alloys have been extensively studied, experimental data for the solid state are needed. These thermody-namic data will provide a basis for understanding phase transformations, precipitation reactions, metal-gas equilibria, and so forth. The interaction of sub-stitutional alloying elements with the interstitial elements is of particular interest. In this investigation the thermodynamic behavior of Fe-Ni-N austenite has been studied. The solubility of nitrogen gas in iron austenite is known to obey Sieverts&apos; law up to about 65 atm.1-6 In addition, the solubility of nitrogen in Fe-Ni austenite has been investigated5"8 using the classical method of equilibrating Fe-Ni alloys with nitrogen gas at 1 atm. A capsule technique similar to that used to study the activity of carbon in alloyed austeniteg&apos;&apos;&apos; was employed in the present work to determine the effect of nickel on the activity of nitrogen in Fe-Ni austenite over the temperature range 600" to 1200°C. EXPERIMENTAL PROCEDURE A series of Fe-Ni alloys up to 35 wt pct Ni was vacuum melted and cast into 1 by 3 by 6 in. ingots. Chemical analyses at the top and bottom of each ingot demonstrated that the ingots were homogeneous with respect to nickel content. The nickel contents are given in Table I. Additional chemical analyses showed that wt pct Si < 0.05, s < 0.01, C < 0.01, Al < 0.006, 0 < 0.004, Mn < 0.002, and P < 0.002. A 2 in. section of each ingot was cold rolled to 0.015 in. The material was then decarburized to a carbon content of less than 0.004 wt pct. A portion of the material of each nickel content was nitrided to various levels in a H2-NH3 gas atmosphere to provide a source of nitrogen during subsequent equilibration. The experimental technique consisted of equilibrating the series of Fe-Ni-N alloys in a partially evacuated sealed silica capsule at the temperature of interest. Both Vycor and quartz capsules were used. In general, the final equilibrium nitrogen content for each Fe-Ni alloy was approached from both higher and lower nitrogen levels. The criterion for establishing that equilibrium was attained was that the final nitrogen content for each Fe-Ni alloy was the same irrespective of the initial level. A schematic drawing of the sample configuration in a capsule is shown in Fig. 1. The samples were arranged so that there was a minimum of physical contact. The samples were also dusted with a fine, high purity alumina powder to help prevent sticking. Several different types of furnaces were used in this study. In each case, a thermocouple was placed immediately adjacent to the capsule during equilibration and the temperature was controlled to within *4?C of that reported. At each equilibration temperature, the following times were found to be more than sufficient to attain equilibrium: 600°C-250 hr, 900°C-150 hr, 1000°C-150 hr, and 1200°C-50 hr. After equilibration the capsules were quenched in water and the nitrogen contents of the specimens determined by a Strohlein analyzer. Analyses of samples after equilibration at 1000" and 1200°C showed no silicon pickup from the silica capsules. RESULTS AND DISCUSSION Transfer Mechanism. The mechanism by which nitrogen was transferred among specimens in an initially hydrogen flushed and partially evacuated capsule equilibrated at 1000°C was investigated. After equilibration the gas in the capsule was collected over water and an estimate of the pressure at temperature

Jan 1, 1970

By J. J. Jonas, G. Gagnon

SHG zinc was extruded in the temperature range 110" to 350°C and the strain rate range 0.05 to 5 sec-1 The strain rate/flow stress/temperature results were analyzed using a power sinh stress relationship. When temperature changes due to the heat of deformation and to heat losses are neglected, the exponent of the sinh function of the stress is 5.6, and the apparent activation energy of deformation is 28 kcal per mole. When these changes are taken into account, the exponent decreases to 4.7 and the activation energy to 23 + 2 kcal per mole. The corrected stress exponent and activation energy are in very good agreement with published values obtained from creep experiments, and suggest that the hot extrusion of zinc is controlled by a mechanism involving self-diffusion. When the extrusion and creep data are compared using a Zener-Hollomon parameter and either a sinh or an exponential stress term, an appreciable offset is observed, which may be due to the difference in impurity content. For a given set of extrusion conditions the ram speed, maximum pressure, and initial temperature can also be correlated using a Zener-Hollomon parameter and a sinh pressure term. THE deformation of metals at temperatures over about one-half the absolute melting temperature has been extensively studied at creep strain rates. By contrast, relatively little work has been carried out on the behavior of metals at hot working strain rates. Most of the latter investigations have been performed using simulative methods, such as hot torsion and hot compression, in which the friction conditions and temperature rise during deformation may differ appreciably from those existing under industrial conditions. Recently, however, Wong and Jonas1 used a scaled-down industrial process to determine the stress and temperature dependence of the strain rate during the extrusion of aluminum. In such tests, the effects of friction and adiabatic heating are closer to those produced in industrial operations. Also, with regard to the testing of materials of limited ductility such as zinc, hot compression and hot torsion do not permit the attainment of true strains as large as the deformations usually applied commercially. The present study was undertaken to investigate the extrusion behavior of Special High Grade (SHG) zinc. The detailed objectives were: 1) to establish the stress and temperature dependence of the strain rate with and without a consideration of adiabatic heating, 2) to study the pressure and temperature dependence of the ram speed, and 3) to investigate the microstruc- tural changes occurring during the deformation. The last aspect of the investigation will be covered in a separate paper. The treatment described below differs from that of Wong and Jonas in that the adiabatic temperature rise during deformation is taken into account, and the calculation of the mean equivalent strain rate is based on the redundant as well as on the homogeneous work. EXPERIMENTAL PROCEDURE Rods from two shipments of continuously cast SHG zinc* were used in the investigation. The composition *Supplied by courtesy of Cominco Ltd. range of the impurities present, as given by the supplier, was: Pb: 0.0013 to 0.0023 pct, Fe: 0.001 pct max Cd: 0.0001 to 0.0006 pct, Cu: 0.0002 to 0.0005 pct, Ti: 0.0001 pct max; thus, by balance, zinc valued from 99.9963 to 99.9966 pct. The as-received rods were machined into billets having a nominal diameter of 1.56 in. and a nominal length of 1½ in.; longer billets up to 4 in. in length, were also used to investigate special aspects. The as-machined rods were annealed at 400°C for 24 hr and slowly cooled. This treatment produced a columnar grain structure, with a grain size of about $ by 2½ cm which was appreciably larger than the as-cast one. A 150-ton, direct extrusion, vertical press was used. Ram displacement and force were recorded continuously against time. A constant flow control valve permitted the maintenance of a range of preselected ram speeds up to in. per sec. The selected speed was held constant, irrespective of the required force, as long as the load developed was below the maximum available. Strain gages were used to determine the force; the gages were calibrated before and after each testing period with a 200,000-lb capacity load cell. Further details of the experimental equipment can be found in Ref. 2. The billets were preheated for 90 min in the extrusion container, which was well insulated so as to minimize temperature gradients. This period was sufficient for the billet to reach a uniform temperature at all temperatures between 110" and 350°C. A square-shoulder die having a 0.290-in. diam central hole was used. The extrusion ratio was 30 to 1. This is equivalent to a reduction in area of 96.7 pct, an elongation of 2900 pct, and a true strain of 3.4. The ram speed was varied over two orders of magnitude from 0.0027 to 0.26 in. per sec. The ram was water-cooled during most of the tests, although some experiments were conducted with a preheated, uncooled ram. All extrusions were run without lubricant, which resulted in conditions of sticking friction. EXPERIMENTAL RESULTS Stress Dependence of the Strain Rate Neglecting Adiabatic Heating. The maximum force required to extrude is given in Table 1 for each of the five initial

Jan 1, 1970

By K. Lundholm, M. Andersson, L. Nevatalo, Å. Holmström, L. Hedström

"Several copper deposits have the arsenic bearing minerals enargite (Cu3AsS4) and tennantite (Cu12As4S12). Since this mineral contain both arsenic and copper, it is not possible to remove the arsenic by selective flotation. Partial roasting is a method to clean the concentrate from arsenic and thereby produce a suitable raw material for the copper smelter. Also mercury is commonly found in copper bearing minerals. In the partial roasting process both arsenic and mercury are volatilized, while part of the sulphur is used as energy during the roasting. This gives a calcine suitable as feed to a smelter, much like an ordinary copper concentrate. A typical analysis of calcine would be: S content 18–22%, As content below 0.3%, Cu content – not affected, grain size – not affected. However, if the concentrate contains mercury, the mercury will be volatilized during the roasting process. This will lead to very high mercury contents in the subsequent process of sulphuric acid, unless the SO2-rich process gas is pre-treated by the Boliden/Norzink process that recovers mercury as calomel, Hg2Cl2. The volatilized arsenic in the gas phase can be recovered either by dry or wet methods. However, one problem is that arsenic also contaminates the fine grained calcine dust that is captured downstream in the gas cleaning system. The calcine dust contains valuable metals such as copper, gold and silver that needs to be separated from the arsenic. The ultimate aim is to stabilize the recovered arsenic as a compound with minimum impact on the environment and recover the valuable metals. By using wet recovery methods it is possible to selectively precipitate copper, gold and silver as sulphides while arsenic is kept dissolved in the liquid. The precipitated metal sulphides can be mixed with the calcine and sent to the copper smelter. The dissolved arsenic is sent to effluent treatment and is further treated in Outotec ferric arsenate precipitation process, where arsenic is precipitated as the stable compound basic ferric arsenate, formula FeAsO4·xFe(OH)3. All concentrates containing arsenic are unique and the process design of a roaster plant must be based on relevant process data. The process data is collected during tests in lab scale and pilot scale, where the influence of critical process parameters on the quality of calcine is studied. The ability to run tests both in lab scale and pilot scale is crucial to the development of the correct roasting process. Outotec is well equipped for this."

Jan 1, 2014

By W. L. O?Connell

[Patent Issue No. Date Issued To Subject 386,504 July 24, G. Sweanor The separation of quartz from a 1588 metal with a liquid such as CS2 or CHC13 under water by introducing mixture into water layer. Quartz will stay in water layer while metal sinks through organic.]

Jan 1, 1963

By Ken Musgrave

"What Most Effective Management Does?It must satisfy the demands, of three groups of people: owners, employees and the public.OwnersOwners want profitable operations that will pay them a return on their investment, plus operations managed for long life. Many owners are also concerned with employee&apos;s well being and their company&apos;s public image.EmployeesJob security and a good income are probably still the most important demands, however, as time goes by, employees are also demanding more involvement in safety issues, working conditions, and even how work is done.The PublicThe public requires acceptable environmental standards, as well as enjoying the prosperity supplied by a successful operation.How are Mines Managed?Most mines are managed very traditionally. The basic credo is that the managers make the decisions, and the hourly, and junior salaried employees do what they are told. Clearly this is the extreme case, epitomized by the long service shift boss who might say ""It took twenty-years to become a shift boss, and no miner is going to tell me what to do"" .There are some exceptions, two notable ones being Campbell-Chibougamau, and T.G.S. (now Kidd Creek Mines) which, in very different circumstances, decided on a more participative style.Campbell-Chibougamau, having closed down due to unprofitable operations, reopened, using a team structure with working leaders, and a system for sharing gains against cost targets with employees. Although the system has some problems, and development miners are bonused conventionally, this has been a basically successful program.Texasgulf Sulphur, a company with a history of non-union operations, and a participative management style, made a corporate decision to operate their Kidd Creek Mine in the same way. The orebody was rich enough to allow them the luxury to do this, but they have never had a strike, so perhaps they have more than saved the cost of the benefits and salaries they needed to pay. They also reward employees with a stock distribution scheme based on service and earnings."

Jan 1, 1985

By M. D. Flavel

Research programs conducted by Allis-Chalmers Corporation are showing that the most efficient energy use in crushers is achieved when the forces applied to particles undergoing breakage are increased. Laboratory bench scale tests have been developed to confirm this phenomena, which was first observed in the field using HYDROCOME crushers with automatic setting controls. A new testing machine is described and methods of maximizing such benefits are postulated. FOREWORD Allis-Chalmers Corporation was founded in Milwaukee, Wisconsin, in 1847 as the E. P. Allis Company, which had an initial interest in manufacturing components in the grinding of flour. As such, its roots were started in reducing material size, which is covered by the generic term, comminution.

Jan 1, 1981

By Sunil K. Tripathy

Particle characteristics such as size, density, liberation and magnetic susceptibility play vital role in magnetic separation technique. To understand the performance of magnetic separation process for chromite beneficiation industry, particle characteristics need to be analysed. However, to date, this has been an expensive and time consuming task, thus only minimal amounts of this type of data are available in the literature. Several new methods are being developed, that will produce the required mineralogical and liberation data to provide processing information rapidly and inexpensively. Among these, QEMSCAN is a unique and advanced characterisation equipment used for liberation analysis for different ore whereas S.G. Franz Isodynamic Separator for determining the magnetic susceptibility. In Indian chromite deposits, the content of iron bearing minerals in the ore is typical which controls the Cr:Fe ratio in the product while treating in gravity based beneficiation plant. In the present investigation, detailed characterisation of two distinct chromite ore samples from Sukinda chromite deposit, India, was investigated which includes advanced mineral characterisation using QEMSCANand magnetic susceptibility measurement by using Frantz Isodynamic separator, etc.. Presence of haematite and gibbsite along with chromite mineral were identified as major gangue minerals in the high grade ore from XRD pattern whereas Goethite, Quartz, Haematite, Kaolinite, Ringwoodite identified in the low grade ore. The mass magnetic susceptibility of low grade ore is very high 5.2 cm3/g compared to 2.7 cm3/g for high grade ore, due to the presence of abundant quantity of iron bearing minerals. Subsequently, these two types of ore were subjected for magnetic separation for improving Cr:Fe ratio by using different magnetic separators viz. induced roll magnetic separator, wet high intensity magnetic separator, rare earth roll magnetic separator, rare earth drum magnetic separator. The results of magnetic separation studies indicated that, Cr: Fe ratio of the chromite ores can be improved to 3.99 and 1.97 from the feed assaying 3.09 and 1.29 respectively. Studies on rare earth roll magnetic separator indicated that both the ores were upgraded to a maximum Cr:Fe ratio level at optimised condition. Keywords: characterization, QEMSCAN, magnetic separation, rare earth roll magnetic separator, industrial mineral, chromite

Sep 1, 2012

By S. M. Bradshaw, E. Charikinya

"Microwave induced cracks have the potential to enhance metal recovery from coarse sphalerite particles in heap leaching operations by creating new crack surface areas for lixiviant and enhancing value mineral exposure. The mode of fracture and the extent of mineral exposure due to microwave induced cracks have not yet been investigated. The objective of this study was to determine the extent of particle damage and mineral exposure in microwave treated particles. This was done using X-ray computed tomography (XCT) and image analysis techniques. XCT analysis was conducted on small (–5+4.75 mm), medium (–16+9.5 mm), and large (–25+19 mm) sphalerite ore particles before and after microwave treatment (MT). Geometrical properties, such as volume and surface area of each mineral phase in each particle was determined from the XCT reconstructed images using VGStudioMax image analysis software. These geometrical properties were used to calculate the interfacial areas of sulphide mineral grains in microwave treated particles before and after MT. The ratio of specific interfacial area after MT to specific interfacial area before treatment was used as an indicator of the degree of preferential grain boundary fracture. The results showed significant grain boundary fracture after MT at all sizes. For small sized particles, about 31% of the interfacial area was lost compared to 23 and 16% for medium and large particles, respectively. The results also showed an increase in the degree of mineral grain exposure of 28, 26, and 15% for small, medium, and large particles, respectively, after MT. These results provide experimental evidence supporting improved mineral exposure due to MT. INTRODUCTIONFor some years it has been suggested that microwave treatment (MT) of ores might result in preferential grain boundary fracture that would have benefits such as improved grain exposure and particle weakening. (Kingman, 2006). Most of the experimental work investigating the application of microwaves to mineral processing flow sheets has focused on assessing the benefits of applying microwave technology to flow sheets consisting of size reduction and mechanical flotation cells (Kingman et al., 2004). Experimental studies have shown that coarse crushing during tertiary grinding of microwave treated ore resulted in a better liberated coarser flotation feed (Scott, Bradshaw, & Eksteen, 2008). It has been suggested that improved liberation in milled flotation feed, was due to the presence of grain boundary fracture in microwave treated particles. However, this improvement in coarse particle liberation does not necessarily translate to a significant improvement in downstream flotation recovery (Kobusheshe, 2010). This is as a result of the size limit to the flotation of coarse material, due to detachment forces acting on particle-bubble aggregates, which increase as the particle size increases (Jameson, 2012). Current conventional, mechanically agitated flotation cells are not designed to float coarse material (Jameson, 2012). Hence it has been difficult to show the downstream processing benefits of applying microwaves to ores processed by comminution followed by flotation (Kobusheshe, 2010). It is clear from investigations carried out so far that inappropriate downstream exploitation of microwave treated ore would fail to realise a meaningful process benefit (Kobusheshe, 2010; Ali & Bradshaw, 2011)."

Jan 1, 2016

By R Nicholson, N D. Cox, N W. A Lambert

"This paper reports on the long-term operation of a particle density meter (PDM) at the Bloomfield coal handling and preparation plant (CHPP). This unit has been operating continuously for over three years, providing real-time data to the operators to allow the monitoring of flotation circuit performance. PDM outputs include slurry density, solids concentration on both mass and volumetric basis, and particle density. Knowing the average particle density of a stream allows for the optimisation of the process in real-time.The calibration of the solids concentration measurement device proved very difficult as the factory calibration techniques and existing laboratory analysis did not give accurate or reliable results. The solids concentration device had to be re-engineered and a new calibration method developed that gave accurate and repeatable results. Independent verification of samples taken from the PDM was undertaken as part of the trial process.The operational reliability of the PDM was a major point of investigation. The initial installation of the unit revealed issues with large, misplaced material of up to 20 mm periodically entering the unit, leading to feed blockages. The presence of these misplaced, oversized particles in the flotation circuit was resolved by installing innovative automatic back flush systems on the PDM.From a plant perspective the PDM has been an invaluable instrument that has enabled irregularities in the flotation circuit performance to be identified, which sampling had not and could not readily identify. This was achieved by close monitoring of real-time tailings composition on a minute by minute basis over short and extended periods. Variations in the tailings stream with respect to particle density and solids concentration were traced back to other issues in the plant. In addition, the PDM has allowed cell optimisation with respect to reagent dose rate to be improved. The PDM with real-time feedback can negate the need for large sampling campaigns and analysis during plant optimisation programs.Citation:Lambert, N, Cox, N and Nicholson, R, 2016. Particle Density Meter, in Proceedings 13th AusIMM Mill Operators’ Conference 2016, pp 245&ndash;250 (The Australasian Institute of Mining and Metallurgy: Melbourne)."

Oct 10, 2016

By Yutaka Yasuda

The particle dispersion phenomena in a turbulent gas jet of the flash smelting process has been investigated. The particle number density in a non- reacting jet was measured experimentally by taking photographs of the particles in flight and counting them using the image analysis. Numerical computations were also performed using a mathematical model for a turbulent particle laden gas jet The experimental particle number density data were then compared with the predicted results of the model. Good agreement was obtained between the experimental and predicted data. The effects of the parameters such as the injector design, particle loading and air flow rate are discussed. Optimization of the concentrate burner design can be achieved with the help of experimental and theoretical simulation as performed in this work.

Jan 1, 1994

By M. Nilmani, H. Pham

"A mathematical model was developed for the dissolution of particles in gas-stirred liquid baths, linking the hydrodynamic flow field of the bath, the average particle trajectory and liquid-side mass transfer controlled dissolution kinetics. The model was validated by comparison with experimental data for the dissolution rates of urea particles in gas-stirred aqueous CaCl2 solutions measured during this study. The dissolution rates of urea. particles were measured using an image analysis system.The effects of operating conditions on the dissolution rates of urea particles were investigated. It was found that the key operating parameters influencing the utilisation efficiency of particles were the gas flowrate, particle diameter and lance submergence. An optimisation of the operating conditions for the model urea-aqueous CaCl2 system indicated that the utilisation efficiency increased with decreasing gas flowrate and decreasing particle size for a given lance submergence.IntroductionThe utilisation solid reagents in molten baths varies from one process to another, depending upon the rates and mechanisms of the reactions between the particles and melts, as well as the trajectory and residence time of the particles in the melts. Previous studies have not attempted to predict the utilisation efficiency of injected solids by coupling the motion of the particle with its dissolution rate. Instead, the particles have usually been assumed to remain suspended in the bath until completely dissolved, and the mass transfer coefficients were based on the average specific mechanical energy dissipation rates in the bath. Investigations of this nature relevant to the present study have been undertaken by Apelian et al. [1], Brucato et al. [2] and Langberg and Nilmani[3]. The objective of the present study was to develop a mathematical model, incorporating both the reaction kinetics between the particles and liquid and the particle motion simultaneously, to predict the dissolution rates and times of the particles. The mathematical model was validated by comparison with experimental dissolution rates of a cold model analogue system measured using an image analysis technique, and then applied to optimise the operating conditions in the model system."

Jan 1, 1996

By J. D. Miller

The recent non-equilibrium electro-phoretic mobility measurements by laser-Doppler electrophoresis coupled with the stability and prevalence of collector colloids in soluble salt flotation systems suggest that the selective flotation of alkali halides is due to the adsorption of oppositely charged collector colloids by hetero-coagulation. Previously reported flotation results confirm this surface charge/collector colloid adsorption model. However, the nature of the inter-particle forces responsible for hetero-coagulation of such oppositely charged particles at high ionic strengths remains to be determined. In this regard, atomic force microscopy was used for inter-particle force measurements at high ionic strengths. Model systems such as polystyrene/quartz and silica/sapphire were studied to measure the particle interaction forces prevalent at high ionic strengths. Results from this study indicate that while repulsive hydration forces exist between similarly charged hydrophilic particles, attractive forces exist between oppositely charged hydrophilic particles in high ionic strength solutions. The repulsive forces between similarly charged hydrophilic surfaces at high ionic strengths (2-4 M) are described by a double exponential function with decay lengths of 0.17 and 1.4 nm depending upon the surfaces involved. On the other hand, attractive forces were observed between oppositely charged hydrophilic surfaces at high ionic strengths (2-4M) and were described by a single exponential function with decay lengths of up to 9 nm depending upon the surfaces involved.

Jan 1, 1996