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By Sadtler, C. B.

IT is generally assumed that in most metals and alloys a given tensile stress produces a given deformation irrespective of the length of time during which the stress is applied. This assumption is justified in most of the important engineering applications of metallic materials; however, in the case of some metals and alloys used or tested at temperatures over or slightly under that at which recrystallization can occur the assumption is not warranted.1 This paper deals with the variations of deformation produced by a constant tensile stress acting for varied lengths of time in the special case of thin aluminum alloy sheet of the duralumin type, which had been subjected to different thermal and mechanical treatments. The material investigated was in the form of a sheet approximately 0.002 in. thick, which had been cold-rolled from a thickness of 0.020 in. It has been found that at a sufficiently high constant tension such material deforms continuously with the elapse of time and that the rate at which this deformation occurs is greatly influenced by the previous thermal and mechanical treatments.

Jan 1, 1927

By M. Metzger, L. E. Hendrickson

Grain boundary attack in 16 pct HCl was found to be substantially the same at low penetrations in zone-refined aluminum (individual impurities 0.1 at. ppm), superior electrolytically refined aluminum (51 at. ppm), and aluminum with various impurities at much higher levels. It was concluded that impurity atom segregation affecting corrosion would have been detected and that the corrosion susceptibility did not originate in this segregation but in the structure of the boundary. It was pointed out that the significance of most previous studies in this system had been obscured by an unrecognized autocatalytic copper reaction. Although general corrosion rate was also impurity-insensitive , there was shallow pitting attributed to iron seg-regation and a hillocked surface texture associated with copper; these were interpreted as due to cathodic damage affecting- cathode distribution. THE grain boundary corrosion suffered by high-purity aluminum in hydrochloric acid has been the object of some interest (the earlier work is summarized in Ref. 1). The central metallurgical question here is whether the corrosion susceptibility of a boundary originates in its structure or in impurity atom segregation. An attempt to study this question revealed large catalytic effects associated with small quantities of the copper impurity in the aluminum, 220 ppm, or in the corrodent, and these magnified the preferential boundary attack and obscured the intrinsic susceptibility question.' After the catalytic effects had been examined,"' test conditions could be designed to avoid them and methods developed for studying the shallow boundary penetrations prevailing when they were absent.= It then became possible to determine whether intrinsic boundary corrosion in aluminum involves impurity segregation. The older work provided no firm information on grain boundary segregation of specific solutes influencing corrosion although several studies suggested iron segregates.4,5 perryman4 found in 10 pct HC1 that the slowly developing (microns per month) grain boundary grooves were deeper in material of higher iron content (range 10 to 550 ppm, with 5 to 80 pprn Cu) but he did not measure the depths of general corrosion, which were probably several times greater, and his reference surfaces may have varied more than did the groove depths. Metzger and Intrater's results for 20 pct HC~,' which yielded higher time-average rates (mm per month) and deeper penetrations, suggested that boundary segregation of iron (range 4 to 230 ppm, with 22 pprn Cu) decreased the penetration rate. However. in the stronger acid the autocatalytic l.E. HENDRICKSON, Student Member AIME ,and M. METZGER, Member AIME, are Research Assistant and Professor of Physical Metallurgy, respectively, Department of Mining, Metallurgy and Petroleum Engineering, University of Illinois, Urbana, Ill. Manuscript submitted March 11, 1968. IMD effect of the copper impurity is greater1,2 and it is now evident that their corrosion rates had been much magnified by this effect and did not provide a proper basis for the analysis of segregation. In exploratory studies of other solutes (made under the same conditions), 1000 ppm additions of Mg, Mn, or Si or 100 ppm Ca were without effect.1 Montariol6 noted that boundary attack in 22 pct HC1 persisted after zone refining although with fewer deep fissures (ranges 0.06 to 4 ppm Cu, 4 to 23 ppm Fe). Autocatalytic effects may have influenced these results also and those of Perryman.4 The present objective was, as a first step, to see whether quantitative tests designed to exclude autocatalytic influences would indicate the existence of low-level impurity effects on intrinsic boundary corrosion. Comparison of electrolytically refined with zone-refined aluminum of lower copper and iron contents revealed no differences in boundary corrosion, but certain impurity-sensitive differences in general corrosion morphology were noted and investigated further at higher impurity levels. I) EXPERIMENTAL PROCEDURE A) Material. A selection from commercially available material was made with the cooperation of several producers. An electrolytically refined lot (III-A, 1 ppm Cu, 2.4 ppm Fe) studied previously3 provided a starting point. Since material of substantially higher copper content could not be used if the autocatalytic corrosion reaction were to be avoided,2,3 a lot (III-B) of about the same purity was added as a check, two zone-refined lots (I-A and I-B) with copper and iron an order of magnitude lower were selected for comparison, and an intermediate lot (11) was included. Analytical data are given in Table I. For copper in I-B and iron in I-A and I-B, the actual concentration is thought to be near the limit given, i.e., about 0.1 ppm. For titanium, vanadium, and chromium in III-B, the actual amounts are thought to be, like those in III-A, substantially lower than in the zone-refined lots (these elements concentrate in the solid on freezing). Data are given later for some additional lots surveved. B) Specimen Preparation and Procedure. one by 3 cm blanks with a notched stem were cut from 1-mm cold-rolled sheet, annealed 24 hr in air at 650°C, and quenched in an air stream (42°C per sec initial cooling rate, 100 sec to cool to 100°C). The high annealing temperature maximized diffusivities and approach to equilibrium impurity distributions. A water quench, in principle more efficient in preserving the distribution established, was undesirable because the boundaries were almost plane and they tended to shear and migrate during the quench and thus to be separated from any existing impurity atmosphere. Test procedures, previously described,3 involved electropolish-ing, etching 2 min in 10 pct HF at 24.0°C, and expos-

Jan 1, 1969

C. A. Stetefeldt, Oakland, Cal.: It has always been assumed by the writer, and also by others, that the silver volatilized by roasting in a Stetefeldt furnace was a minimum as compared with roasting in other furnaces. Now, Mr. Morse records at

Jan 1, 1896

By George Branner

THERE is no record of the commercial production of bentonite as such in Arkansas up to the present although fuller's earth has been. mined intermittently in the state1 from 1891 to 1922. Very nearly all of this was obtained from decomposed basaltic dikes.2 Recently three bentonite deposits of close stratigraphical relationship have been identified in the Coastal Plain region of southern Arkansas, which suggests that there may be a considerable quantity of commercial material available. LOCATION AND ACCESSIBILITY The three deposits are located respectively in the SE. ;1/4 SE. ;1/4 sec. 13, T.- 2 S., R. 13 W., Saline County, the SE. ;1/4 SE. 1/4 sec. 18, T. 5 S., R. 16 W., Hot Spring County, and in sec. 22, T. 12 S., R. 19 W., Ouachita County. The Saline County deposit is 14 miles by state highway 167 south of the city limits of Little Rock and 1 mile west by dirt road. This deposit is 12 miles south of Ward switch on the C. R. I. P. Railroad and 11 miles east of Vimy Ridge by dirt road. Vimy Ridge is also on the C. R. I. P. Railroad. The Hot Spring County deposit located in the SE. ¼ SE. ;1/4 sec. 18, T. 5 S., R. 16 W. is about 100 yd. north of Durian switch on the C. R. I. P. Railroad and on state highway No. 9, 5 miles southeast of Malvern. The Ouachita County deposit has not been accurately located but lies in sec. 22, T. 12 S., R. 19 W. It is approximately 2 miles by road from Chidester on the Missouri Pacific Railroad.

Jan 1, 1929

By R. W. Smith, K. H. Lee

Bond type simulated closed-circuit grindability tests and batch grindability tests were run on six pure minerals plus a limestone and a taconite ore. The simulated closed-circuit tests were made at five different meshes of grind. The minerals tested varied widely in hardness and breakage pattern. For all simulated closed-circuit tests both Bond and Holmes work indices were calculated and compared. The Bond grindability testlm6 has been widely and quite successfully used for predictions of such items as ball mill power requirements and as an aid in plant comminution equipment selection. The test is interesting in that although run as a series of batch tests it is over-all a simulated closed-circuit test with a circulating load of 250%. The simulated closed-circuit feature is probably the primary reason for the notable success of the test. Perhaps the major criticism that can be brought against the test is that it is time consuming, since it normally takes a trained technician about an eight-hour shift to perform it.4 For many purposes this is an acceptable time requirement. However, if a grindability test is to be used for mill control purposes it should be performed in as short a time interval as possible, preferably under an hour. Sometimes in industry strictly batch tumbling mill tests are run and an attempt is made to correlate such batch tests with plant mill operation. Then an attempt is made to adjust operating tumbling mills based on data from the tests. The tests, however, frequently do not correlate well with operating data. The purpose of the present work was to explore some of the similarities and differences between a simulated closed-circui.t grindability test and a batch type test. Six different minerals plus a limestone and a taconite ore were tested in the experimentation. These materials varied widely in hardness, breakage characteristics, etc. Standard dry Bond grindability tests were mn on all materials at 48, 65, 100, 150 and 200 meshes of grind. Feed material was prepared by stage cmshing to -6 mesh. Initial feed volume was 700 cc. Care was taken to maintain the standard constant weight (about 20125 gm) and number of balls (283) during all the tests. Rotational speed was 70 rpm. Table I lists the materials studied along with pertinent data regarding the materials and Bond grindabilities obtained. In addition, dry batch tests were performed on all materials. These tests were made in the same mill as was used for the Bond tests. An identical ball charge and an identical feed volume were used. After rotating the mill for a predetermined number of revolutions at the same rotational speed as for the Bond test, the material was removed and separated from the ball charge. A screen analysis was performed on this product. Knowing the original size distribution of the feed and the size distribution of the product, it was possible to calculate the grindability or grams of a particular mesh produced per revolution of the mill. After the screen analysis of the product, the various fractions were recombined and placed back in the mill for further grinding. After grinding for additional revolutions of the mill the procedure was repeated. Thus, it was possible to obtain grindabilities for each material at a number of meshes of grind and for various revolutions of the mill. Grind-abilities obtained from batch tests are shown in Table 11. COMPARISON OF BOND AND BATCH GRINDABILITIES Figs. 1A through 1D show relationships between batch and Bond grindabilities. Figs. lA, lB, 1C and 1D show respectively graphs indicating the relationships for batch tests run for 100, 300, 500 and 700 revolutions of the mill. For a particular material on each graph, points are shown for 200, 150, 100, 65, and 48 meshes of grind except for several missing coarse meshes of grind for galena which did not fall conveniently on the figures. For each material and graph the point nearest the origin is for 200 mesh, the next for 150 mesh, the next for 100 mesh, etc. Considering Fig. lA, it is interesting to note that all relationships shown are essentially straight lines. There are, however, considerable differences

Jan 1, 1969

Jan 1, 1967

By A. Eilers

A MARKED peculiarity of most of the smelting-works of the Far West is the looseness with which accounts of the operations are kept. Indeed, probably over half of the works do not keep any detailed accounts at all, the yearly gross statement of profit or loss being considered sufficient for all purposes. The reasons for this must apparently be sought in the as yet unsettled state of all business relations, and in the deplorable fact that only in isolated cases educated metallurgists are in charge of the smelting-works. Continual and regular assays of the ores, by-products, and slags, are almost unknown, so that it is impracticable to ascertain, even approximately, the losses in the smelting processes. Of cases where regular accounts are kept of the quantities of ore smelted, of the composition of charges, and of the fuel consumed, I know only two or three. Under these circumstances it is extremely difficult to collect figures which cover the operations for a considerable length of time, and which are so valuable for the metallurgist, who wishes to get an insight into the economy of smelting operations, as shown by practice. Figures obtained by personal observation, which can of course cover only the brief space of the visit of the travelling metallurgist, must, therefore, be made to replace the more valuable data. The writer has had occasion, during the last and the preceding summer, to visit the larger number of the Western lead-smelting works, and offers, in the following pages, such data relating to the economy of lead-smelting in the blast furnace at various works, as he has been able to obtain. The only object in doing so is to place these figures on record, so that they may be from time to time supplemented with other data, which are now wanting. I. EUREKA CONSOLIDATED COMPANY'S SMELTING-WORKS. These works smelt the ores from the mines of the same company on Ruby Hill, Eureka, Nevada. The supply is almost unlimited,

By B. J. Nelson, F. N. Rhinos

Studies upon the rates of oxidation of copper alloys containing small quantities of the alloying elementsl,2 have shown that steady growth of the scales at predictable rates is limited to a small concentration range and to oxidation at temperatures above 600°C., whereas unpredictable and erratic oxidation behavior is found with alloys containing larger additions and at low temperatures. The latter effect appears to be associated with modifications in the geometric disposition of the oxides. At low concentrations of the baser elements in copper, simple oxidation behavior is typified by the formation of two distinguishable zones of oxidation: (I) a subscale composed of discrete particles of the oxide of the alloying element embedded in substantially pure copper and (2) an external scale composed of cuprous oxide (with a very thin layer of cupric oxide on the outside) containing particles of the alloying element distributed throughout, but most profusely near the metal surface (see Fig. 2, top). Where this type of structure obtains, the process of oxidation is believed to be implemented, and its rate controlled, by the diffusion of the reacting elements through a' continuous matrix phase; i.e., through the alpha (metallic) phase of the subscale and alloy and through the cuprous oxide of the external scale. Copper is presumed to diffuse outward through the cuprous oxide layer to the external surface, where it reacts to form more cuprous oxide.3 Oxygen is delivered at the oxide-metal interface in the form of cuprous oxide, which may either react with the alloying element to deposit its oxide at the inner limit of the external scale or dissolve in the pure copper that remains at the surface of the metal after it has become depleted in its oxidiza-ble alloying elements. Both oxygen and the alloying element are pictured as diffusing through the metal, the oxygen inward and the alloying element outward. Where the two diffusion streams meet in sufficient concentration the oxide of the subscale is precipitated. If the oxygen pressure in the surrounding atmosphere is just below the pressure necessary to maintain cuprous oxide undecomposed, no cuprous oxide can be formed; some of the oxide of the alloying element may form upon the external surface, if its decomposition pressure is below that of cuprous oxide, and a normal sub-scale appears. This condition is achieved experimentally by enclosing the alloy with cuprous oxide together with an excess of copper. At temperatures from 600°C. downward, the precipitated oxides tend to accumulate at the grain boundaries of the alpha phase where a relatively small volume establishes a partial barrier to the diffusion of the reactants- Similarly, large volumes of Precipitated oxides tend to deposit in various patterns that may be expected to

Jan 1, 1944

By F. N. Rhinos, B. J. Nelson

Studies upon the rates of oxidation of copper alloys containing small quantities of the alloying elementsl,2 have shown that steady growth of the scales at predictable rates is limited to a small concentration range and to oxidation at temperatures above 600°C., whereas unpredictable and erratic oxidation behavior is found with alloys containing larger additions and at low temperatures. The latter effect appears to be associated with modifications in the geometric disposition of the oxides. At low concentrations of the baser elements in copper, simple oxidation behavior is typified by the formation of two distinguishable zones of oxidation: (I) a subscale composed of discrete particles of the oxide of the alloying element embedded in substantially pure copper and (2) an external scale composed of cuprous oxide (with a very thin layer of cupric oxide on the outside) containing particles of the alloying element distributed throughout, but most profusely near the metal surface (see Fig. 2, top). Where this type of structure obtains, the process of oxidation is believed to be implemented, and its rate controlled, by the diffusion of the reacting elements through a' continuous matrix phase; i.e., through the alpha (metallic) phase of the subscale and alloy and through the cuprous oxide of the external scale. Copper is presumed to diffuse outward through the cuprous oxide layer to the external surface, where it reacts to form more cuprous oxide.3 Oxygen is delivered at the oxide-metal interface in the form of cuprous oxide, which may either react with the alloying element to deposit its oxide at the inner limit of the external scale or dissolve in the pure copper that remains at the surface of the metal after it has become depleted in its oxidiza-ble alloying elements. Both oxygen and the alloying element are pictured as diffusing through the metal, the oxygen inward and the alloying element outward. Where the two diffusion streams meet in sufficient concentration the oxide of the subscale is precipitated. If the oxygen pressure in the surrounding atmosphere is just below the pressure necessary to maintain cuprous oxide undecomposed, no cuprous oxide can be formed; some of the oxide of the alloying element may form upon the external surface, if its decomposition pressure is below that of cuprous oxide, and a normal sub-scale appears. This condition is achieved experimentally by enclosing the alloy with cuprous oxide together with an excess of copper. At temperatures from 600°C. downward, the precipitated oxides tend to accumulate at the grain boundaries of the alpha phase where a relatively small volume establishes a partial barrier to the diffusion of the reactants- Similarly, large volumes of Precipitated oxides tend to deposit in various patterns that may be expected to

Jan 1, 1944

By Ray W. Carpenter, Ralph Hultgren, Raymond L. Orr

Heats of formation of Au-Zn alloys of compositions Au3Zn and AuZn were rneasured at several temperatures by liquid tin solution calorimetry. The data for Au3Zn show that much smaller heat and entropy effects are associated with the a2- a1 transformation than for the a1 - a transformation. This result is consistent with reported X-ray diffraction studies which indicate that the a1 and (a2 phases have ordered structures which only slightly differ, whereas the a phase is disordered. The heat of formation of AuZn(ß') does not change significantly between 322o and 8000 K, confirming that the ß' alloy remains highly ordered to high temperatures. Not many measurements have been made of the energies of order-disorder transformations in alloy phases. In many cases these energies may be conveniently determined by measuring heats of formation as a function of temperature. Many alloy phases approximately obey Kopp's law of additivity of heat capacities; as a consequence their heats of formation do not change with temperature. In ordered alloys, however, a decrease in degree of order with increasing temperature will contribute an endothermic component to the heat of formation. Thus, for ordered phases where other anomalous contributions to heat capacity in the alloy or its pure components are absent, a measured change in the heat of formation may be attributed to the energy of disordering. A portion of the Au-Zn phase diagram1" is shown in Fig. 1. The composition AuZn(ß') has a well-ordered superlattice structure of the CsCl type at room temperature. The alloy is considered to remain ordered to its melting point,1 although, from X-ray diffraction data and electromotive force measurements of zinc activities, Terpilowski3 concluded that considerable disordering began at -673°K. The Au3Zn composition is a disordered fee solid solution (a) at high temperatures. Between 693o and 543oK, the ailoy orders to form the a1 phase, having a structure based on the L12(AuCu3) type, but with a regular distribution of ordering faults in the c direction, making it tetragonal. Below 543oK, the alloy orders further by means of slight atomic shifts, without diffusion, yielding a complex long-period orthorhom-bic superlattice structure, a2.4-7 Heats of formation of Au3Zn(a2) and AuZn(ß') were measured at 363°K by Biltz et al.8 by aqueous solution calorimetry. Heat capacities of several alloys between 14.9 and 38.4 at. pet Zn in the range from 373" to 723°K were measured by Iwasaki ef al.6 Stoichiometric Au3Zn, however, was not included in the alloys measured, the nearest composition being 26.1 at, pet Zn. Also, the heating rate used for the measurements, 2 deg per min, may not have permitted equilibrium to be reached in the neighborhood of the transformations. Heat contents of solid and liquid AuZn were measured between 488o and 1138°K by Kubaschewski.9 It was decided to determine the ordering energies of Au3Zn and to check on the possible disordering of AuZn by measuring the heats of formation of these alloys at several temperatures. The present paper reports measurements by liquid tin solution calorimetry of the heats of formation of the three phases of Au3Zn at temperatures within their regions of stability, and of AuZn at a low (322°K) and a high (800°K) temperature. EXPERIMENTAL Alloy Preparation and X-Ray Examination. Ten-gram ingots of Au3Zn and AuZn were prepared by melting weighed amounts of gold (99.95 pct Au) and zinc (99.99 pet Zn) together in sealed evacuated Vycor tubes at 1060°K, followed by rapid quenching. The resulting ingots were homogenized at 900°K for 10 days. The Au3Zn alloy lost no weight in preparation, the loss from the AuZn alloy was enough to introduce an uncertainty of only 0.1 at. pet in its composition. Filings for X-ray examination were taken from various parts of each ingot and strain-annealed at 723°K for 20 min in vacuo. The Au3Zn filings were quenched rapidly to retain the disordered phase. Diffraction patterns were determined with an X-ray dif-fractometer using copper Ka radiation.

Jan 1, 1968

By D. J. McAdam

As shown in three recent papers by the author, in two papers by McAdam and Mebs, and in a paper by McAdam, Mebs, and Geil," the technical cohesive strength of a metal, in any particular state as regards mechanical treatment and heat-treatment, cannot be represented by a single stress value, but must be represented by a diagram with the principal stressest as coordinates. Each point on the boundary of such a diagram represents a technical cohesion limit. The technical cohesive strength thus comprises an infinite number of technical cohesion limits, each representing fracture under a specific stress combination. Moreover, neither the yield strength1 nor the ultimate strength can be represented completely by a single stress value, but must be represented by a diagram with the principal stresses as coordinates. All combinations of the principal stresses may be classified in three groups as illustrated by Fig. I. Fig. Ia shows three typical stress combinations with no two principal stresses equal. The magnitude of each principal stress is indicated qualitatively by the length of the arrow. Arrows pointing away from the cube indicate tensile stresses; arrows pointing toward the cube indicate compressive stresses. Typical stress combinations with two principal stresses equal are illustrated in Figs. rb and Ic. Such stress combinations would be produced by subjecting a cylinder to combinations of uniform axial and radial stresses. The uniform radial stress may be resolved into two equal principal stresses. with the mutually perpendicular directions rotatable around the axis of the cylinder. Fig. ib shows stress combinations with Sz equal to The greatest principal stress is in the axial direction, and the stress combination tends to cause either absolute or relative increase in length.* Fig. u shows stress combinations with S2 equal to St. The greatest principal stress is in the radial direction, and the combination tends to cause either absolute or relative decrease in length. The complete representation of either technical cohesive strength, yield strength, or ultimate strength generally requires a three-dimensional diagram. When two of the principal stresses are kept equal, however, the strength of a metal may be represented by a two-dimensional diagram with the coordinates representing the axial and radial stresses.6-II In this paper, attention will be confined almost entirely to stress combinations with two of the prin-

Jan 1, 1945

By M. J. Sinnott

WITH the advent of nodular iron as an engineering material, considerable interest has been shown in developing this material for various applications. Generally the strength, hardness, or ductility are the more common properties used for design purposes but occasionally applications involving a knowledge of the thermal conductivity of the material is required. A search of the literature showed, rather surprisingly, that relatively little work has been done on this property for all types of cast iron. Insufficient data are available on thermal and electrical conductivities to develop the Weidemann-Franz ratio, or the Lorentz relationship, for estimation purposes. There are no data on nodular irons and the published data on malleable irons are quite sketchy. For these reasons the conductivity of the group of nodular irons listed in Table I were measured. Irons I to IV were made in such a fashion as to maintain the carbon and alloy content essentially the same while varying the silicon content. Test specimens were cut from the mid-section of 4-in. keel castings. Iron V is a commercial casting of nodular iron. The test specimen of this iron was cut from the center of a large cast billet 3x10x26 in. All alloys were tested in the as-cast condition. A microstructural examination showed that each iron had spheroidal graphite but the size and distribution varied with the silicon content of the iron. The spheroid size decreased while the number of graphite nodules increased as the silicon content of the iron was increased. The matrix structures of each iron consisted of pearlite and ferrite with the proportions varying from predominantly pearlite in the low-silicon iron to mostly ferrite in the high-silicon iron. The commercial heat, iron V, was predominantly pearlite probably because of its higher manganese and phosphorus content. Two methods for measuring the thermal conductivity were used. The first was a modification of the techniques described by Powell and Hickman and Van Dusen and Shelton,² the second method was a simpler technique described by Jakob and Hawkins." Since the agreement of both methods as tested on two sets of duplicate samples was excellent, the simpler experimental technique was used. The results of the measurements are given in Table II along with a summary of the microstructural constituents. These conductivity data indicate that as the percentage of silicon increases there is a corresponding decrease in the thermal conductivity even though the structure is changing from pearlitic to ferritic. Such a change in microstructure is usually accompanied by an increase in conductivity. A third factor involved in the behavior of these irons is the size and distribution of the graphite nodules. The relative amounts of graphite do not change greatly but the size and distribution do. Generally graphite improves the conductivity but in the present case such an effect is probably masked by the alloying effect from the silicon and its dispersion through the matrix. The thermal conductivity of the nodular irons studied is less than that of regular foundry irons by some 10 to 30 pct and is less than that of regular malleable irons.' The cause of the decreased conductivity is probably the more highly alloyed fer-rites obtained in these irons and the changed mode of distribution of the graphite as compared to the regular foundry irons. References R. W. Powell and M. J. Hickman: Journal Iron and Steel Inst. (1946) 154, No. 2, p. 99. 'M. S. Van Dusen and S. M. Shelton: U. S. Bur. Standards Journal of Research, RP668 (1934) 12, p. 429. ³M. Jakob and G. A. Hawkins: Elements of Heat Transfer and Insulation. 2nd Ed. (1950) New York. John Wiley and Sons. ' J. W. Donaldson: Journal Iron and Steel Inst. (1933) 128, No. 2, p. 255.

Jan 1, 1954

By F. Z. Schellenberg

A recent visit with the engineering students of the Western University of Pennsylvania to the Keeling coal-mine on the south side of Pittsburgh furnished interesting matter to communicate, as may be the case, however, with regard to the striking features of any of our well-arranged larger collieries; for every mine surely appears unique as a whole, though description fails to give that impression, except in particulars.

Jan 1, 1900

By Albert H. Chester

During the summer of 1875 I visited some of the iron mines of Houghton County, Mich., and was quite interested to observe the progress made there within a few years. Working in open pits is gradually giving place to underground work in most of the mines where this system is practicable. In some cases, as in the Lake Superior Mine, the body of ore is so irregular, and the deposit so liable at any time to be pinched out, or nearly so, that it. has been found impractical, if not impossible, to lay out any plan for systematic underground work. But in many other mines the deposit is so regular in character, that the work can be laid out in shafts at regular intervals, and a series of 60 feet levels. The Michigammi Mine, opened in 1872, and first worked in 1873, is arranged on this plan, though some quarrying is still done. But the men are being put at overhead stoping as fast as it can be got ready for them, and before long it will all be worked in this way. The New York Mine is undergoing a similar change to the pillar system, and several others have either been changed, or it is about to be done. My principal object in visiting this region was to obtain samples from some of the important mines, so as to be able to answer the question: "What is the average percentage of iron in the Lake Superior ores?" Most iron furnaces work on mixtures of ore, so that it is impossible to tell, from their records, just the percentage in each one.

By Hugh Roberts

During 1916 and 1917, the E. J. Longyear Co. of Minneapolis, Minn., carried out a campaign of exploration for nickel-copper ore in the Sudbury District of Ontario. The work was initiated by W. E. Smith, a resident of Sudbury, who called the attention of W. J. Mead, Chief geologist. The E. J. Longyear Co., to the fact that explorable lands in the Sudbury District were still available. Five diamond drills were employed, and holes were drilled in the townships of Levack, Trill, Denison, Blezard; Garson, Falconbridge and MacLennan. The general map of the District, Fig. 1, shows the distribution of the lands chosen. As a result of this exploration, a large body of nickel and copper ore has been found in the western part of the Township of Falconbridge. This property lies in the eastern part of the Sudbury District, east of the Garson mine.

Jan 2, 1918

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 Larry Dale

An approach for predicting the development of mineral resources on public lands has been developed as an aid to resource agencies. It consists of new techniques for evaluating the economic potential of the resource, the uncertainty about mine cost estimates, and the uncertainty about the occurrance of mineral deposits. New mineral sup- plies from inferred and hypothetical resources can be predicted by using these techniques. A case study illustrates their application in predicting development of copper resources in western Montana.

Jan 1, 1982

By N. A. Gjostein

The formal descnption of the decay of an isolated scratch can be written in terms of an appropriate Fourier integral. With the application of certain approximations, this description leads to the second-moment analysis developed by King and Mullins. It is shown that the second-moment analysis does not adequately describe the behavior of a decaying, isolated scratch, and that it is difficult to obtain internally consistent values of the surface self-diffusion coefficient, D,,from this analysis. This inconsistency arises from contributions of higher-order moments to the decay process. Two new analytical methods, based on the Fourier integval representation of a scratch profile, have been developed for determining values of D, from the decay kinetics of isolated scratches. In the profile-matching method, a given initial profile is resolved into its Fourier components, the components are allowed to decay exponentially over an annealing time interval, At, and finally they are synthesized to form a decayed profile. The predicted profile is adjusted to fit the experimental profile by varying only the parameter, D,At; the contribution from volume diffusion is directly taken into account and the integrals involved in the Fourier analysis are evaluated by numerical integration on an electronic computer. In the kinetic method, the widths, W, of a series of predicted profiles are determined for various volume-diffusion contributions, and the results are plotted in the form W vs DSAt. For a fixed ratio of surface to volume diffusivities, R, these plots are linear; the slope of the lines increases as R decreases. A simi-lar analysis can be performed on the depths of the predicted profiles. In this case, plots of (l/df3 vs D,At are found to be nonlinear. The proper value of Ds is found by superimposing the data for the experimental widths and depths on Plots of this kind and adjusting Ds until a fit is obtained with the line having the appropriate value oj R. An analysis of the slopes of these plots for R —= (surface diffusion only) reveals that they do not give the correct value, 2276, for width tf second-moment standard profile. This finding allows one to modify the second-moment analysis in such a way as to obtain internally consistent values of Ds. ThERE are several processes, involving a flow of matter along a surface, due to gradients in curvature of the surface, that may be used to determine the surface self-diffusion coefficient, Ds. These processes are commonly known as 1) thermal grooving, 2) single and 3) multiple scratch decay, 4) blunting of field-emitter tips, and 5) sintering. In this series of papers, the first three of these processes are examined in detail in an attempt to subject the phenomenological theories governing these processes to a critical test and, furthermore, to make a critical comparison of the values of D, obtained from these techniques. Although the thermal grooving technique has been established'' as a reliable method for determining D, , this has not been the case for the process of scratch flattening. Mullins and coworkers374 have proposed phenomenological laws governing the growth of the width and the decay of the amplitude of an isolated scratch, but only in one study5 has the width-growth law been supported by experimental evidence. In paper I of this series detailed experimental evidence is presented that points up the deficiencies of the second-moment analysis; the origins of these deficiencies are also discussed. Following this development, some new analytical procedures, based on representing the scratch profile by Fourier integrals, are presented as a means of obtaining reliable values of D, from the kinetics of the scratch-flattening process. Further applications of the Fourier integral approach are reported in papers 11' and 111,~ where the method has been useful in interpreting the anisotropy of flattening rates of isolated scratches (paper 11) and in describing the multiple scratch decay process (paper 111). THEORY OF SCRATCH DECAY Following ~ullins" approach we consider an arbitrarily shaped scratch profile, Fig. 6, at some time, t. The profile is designated by a function z(x,t) which can be expressed in terms of a Fourier integral as follows:

Jan 1, 1967

By Samuel Hoyt

THE writer has long been interested in seeking an explanation of the upper heat effect in the copper-tin alloys over the a + ß range, first described in 1913. These notes are offered, not at all as the final explanation of this heat effect, but rather to indicate certain progress which has been made toward establishing what happens over this temperature interval. While working on the thermal analysis of the copper-rich kalchoids (copper-tin-zinc alloys), it was noted that those alloys containing major portions of tin and minor portions of zinc exhibited two marked heat effects, one at about 520° C. and the other at about 600° C., instead of the single effect which is generally observed in either of the two binary series. This matter was discussed with Dr. Guertler, who suggested that a heat effect at about 600° C. in the pure copper-tin alloys might be expected. This led to a more searching examination of the thermal critical points in the pure copper-tin alloys, with the result that a marked, although somewhat weak, heat effect was located at about 600°. Somewhat later, this upper heat effect was discussed with Dr. Burgess, and it was requested that the Bureau of Standards make heating and cooling curves of one of these alloys for the purpose of removing any possible doubt as to the actual presence of the heat effect. Results obtained by the Bureau are given in Plate I. In an earlier paper1 differential curves were published showing that the heat effect occurred at constant temperature over the a + ß range and increased in magnitude with the amount of a up to a point somewhat above 20 per cent. tin. No heat effect was noticed at this temperature in the eutectoid alloy. The existence of this heat effect seems to require a modification of the present copper-tin diagram in the a + ß field, and it has been the object of the work here presented to secure the evidence upon which such a change should be based.

Jan 12, 1918

By C. H. Green

. RECENTLY two papers on the structure of eutectics were read before the British Institute of Metals, one by F. L. Brady1 and the other by A. Portevin.2 In the preparation of photomicrographs of laboratory specimens as illustrations for student work, and in connection with various researches; the writer has been working along the same line for some, time and, therefore, was much interested in the results given, especially those of Portevin. He has correlated the information on the subject and has outlined a broad scientific basis for future work along this line. It is largely with the idea-of adding something by way of illustration to the work already done that this paper is written. The eutectic is generally understood to mean the alloy having the lowest melting point of any alloy of the system, and in which the melt becomes supersaturated with reference to the two or more phases which, separate out simultaneously (or alternately). From the surface pattern, the eutectic means; a structure with no primary crystals. It is not always easy to get the exact eutectic composition in making any alloy; in fact, a slight range in composition may be expected and some of the coarser particles may be richer in either of the composing metals. Even when chemical analysis shows the right composition, differences in density of the pure metals may cause segregation, or. supercooling resulting from different crystallizing properties may give a surface pattern the appearance of being on one side of the eutectic point. However, the alloys on either side of the eutectic composition, showing primary crystals of one of the phases and the effect of the pro-eutectic form on the surface pattern of the eutectic make a very interesting study, which is slightly touched upon in this paper. The phases, or microconstituents, of a eutectic may be pure metals (Pb-Sb, Tl-Au), solid solutions (Cd-Zn), a pure metal and a solid solution; (Ag-Sn), or a pure metal or solid solution with an intermetallic compound (Sb-Mg, Sb-Cu, As-Cu, Bi-Te, etc.).

Jan 2, 1925