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ALABAMA Anniston.-Bretz, J. A. Carrington, F. G. Gerber, A. B. Klugh, B. G. Semple, R. A. Ashland.-Barton, .J. Bessemer.-Abbott, C. E'. Ball, T. L. Beaver, J. J. Ferguson, V. Salmon, H. S. Sehaber, C. F. Whatley, W. J.

Jan 1, 1923

By Herbert Clark Hoover, Lou Henry Hoover

QUESTIONS of assaying were explained in the last Book, and I have now come to a greater task, that is, to the description of how we extract the metals. First of all I will explain the method of preparing the ore1 ; for since Nature usually creates metals in an impure state, mixed with earth, stones, and solidified juices, it is necessary to separate most of these impurities from the ores as far as can be, before they are smelted, and therefore I will now describe the methods by which the ores are sorted, broken with hammers, burnt, crushed with stamps, ground into powder, sifted, washed, roasted, and calcineda.

Jan 1, 1950

By Axel Hultgren, B. Herrlander

THE original aim of the present investigation was to study the mechanism of cracking on hot-deforming red-short steels. During the microscopical examination of hot-deformed soft steels attention was directed to various patterns of so-called veining in the ferrite, as related to variations in the deformation procedure and the heat-treatment of the steels. Later, a hot-short steel of higher carbon content was also studied. Historical Storey1 suggested in 1914 that veining developed during the gamma-alpha transformation from the growing together of several alpha nuclei within the same gamma grain. Rawdon and Berglund2 found that in soft steel, forged somewhat below A3, the ferrite showed profuse veining, while usually there was very little veining after forging at very high temperatures. They emphasized the similarity between the pattern of veining formed by deformation about 600°C and that of slip lines formed by deformation at room temperature. They also suggested a relation between the gamma-alpha transformation and veining. They became convinced from etching and from the slip-line pattern in cold-deformed material, that the orientation within any ferrite grain showing veining was uniform. In recrystallized alpha grains veining was absent. Veining did not appear materially to affect the properties of ferrite. Independent of veining, continuous networks, attributed to preexisting delta and gamma grain boundaries were sometimes found, the former only in cast steel. Those networks were usually connected with small inclusions. Ammermann and Kornfeld3 confirmed that recrystallized ferrite grains showed no veining. In soft steel annealed between At and A3 there was veining only in the ferrite grains formed during cooling. Electrolytic iron deformed at 880° was free from veining. Bannister and Jones4 considered veining in ferrite to be a manifestation of microscopical and submicroscopical inclusions. Northcott,5,6 in studies on veining in steel and other metals and alloys, attributed veining to oxide inclusions. It could generally be removed by annealing in hydrogen, a suggestion earlier made by Rawdon and Berglund. Tritton (discussion in ref. 5) emphasized the importance of perfect polishing for bringing out veining by etching, and thought the veins or subboundaries were produced during the gamma-alpha transformation, when it was rapid, as a result of the volume change and contraction during cooling.

Jan 1, 1946

By W. E. Ruder

It has been pointed out by Stead1 that grains of considerable coarseness may be developed in steels containing from 3 to 5 per cent. of silicon, and in a previous paper2 the present author has shown that under the proper conditions of annealing single grains having an area of 50 sq. cm. may be obtained. It is the purpose of this paper to discuss in more detail the conditions promoting the growth of these grains, in the hope that the experiments described may throw some additional light upon the general problem of grain growth in metals. These silicon steels lend themselves particularly' well to such experiments because of the ease with which the change in structure may be followed and because of the fact that the effect of the carbon present is largely obliterated by the high percentage of silicon. This forms a solid solution with the iron, which acts very much like a pure metal. This alloy has been found to have but one critical point, viz., Ar2. This point, determined by Roberts-Austen for Hadfield's 4 per cent. Si alloy, was found to lie at 703" C. Vigouroux3, in working with pure Fe-Si alloys, finds this point to rise from 750" C. for 0.5 per cent. Si to 860" for 6 per cent. alloy, with a flat region in the curve from 3 to 4 per cent. Si. No accurate determination of this point has been undertaken by the present author, but it has been roughly determined to lie between 730" and 750:. In the preliminary experiments, some of which have been described in a previous paper, it was found that practically all of these steels, when annealed to 1,200" C. or higher, showed grains of comparatively large size. These grains varied in size, however, from about 4 sq. mm. to 50 sq. cm. (see Figs. 1 and 2) under the same conditions of anneal, and it was for the purpose of determining the cause of this difference in size that the experiments herein described were undertaken. From a preliminary study of the chemical composition of the specimens

Jan 1, 1914

By C. Laird, C. E. Feltner

May different kinds of dislocation structures have been observed in strain-cycled metals and alloys. In order to understand their pattern and causes, an experimental program has been carried out to determine the influence on the dislocation structures of the three variables: 1) slip character of the material, 2) test temperature, and 3) strain amplitude. The results show that at high strain amplitudes cell structures me formed when the slip character is wavy, and that these are progressively replaced by uniform distributions of dislocations as the stacking fault energy is decreased. At lower strains, dislocation debris is formed which consists primarily of dipoles in wavy slip mode materials and multipoles in planar slip mode materials. Temperature merely acts to change the scale of the structure, smaller cells, and clumps of dislocation debris being associated with lower temperatures. It is shown that the results for many metals fit this pattern, which Parallels that occurring in unidirectional deformation. DISLOCATION structures produced by cyclic strain (fatigue) have been examined in a number of metals by transmission electron microscopy. These studies have produced a variety of interesting and often seemingly conflicting results. For example, different investigators have reported such structural features as cells.le4 bands of tangled dislocations,4'5 dense patches or clusters of prismatic dislocation loops, planar arrays,4'10 and various combinations or mixtures of these different structures. Most of these observations have been made on materials which were initially annealed and cyclically strained at low amplitudes resulting in long lives. Recently we have reported observations of the dislocation structures produced in copper and Cu-7.5 pct Al cycled at large amplitudes, resulting in lives of less than 104 cycles.4 These results, examined in combination with those in the literature, have suggested that a common or consistent structural pattern exists. Variations in this pattern appear to be determined chiefly by the three variables, namely, the slip character of the material,4,11 test temperature. and the strain amplitude. To verify this interpretation, we have studied [he influence of the above three variables (in different combinations) on the resultant structures in cyclically strained metals. Copper, fatigued at room temperature, was chosen as a reference state to which all other observations can be compared. The effect of slip character has been investigated by employing fcc metals of different stacking fault energy. Thus aluminum which has a more wavy slip character than copper, and Cu-2.5 pct A1 having a more planar slip char- acter, have been examined. The aluminum samples were fatigued at 210°K thus making their homologous temperature equal to that of copper at room temperature. The influence of temperature has been evaluated by examining the structures in copper at room temperature and 78°K. Finally the effect of strain amplitude was studied by looking at the structures at amplitudes giving lives ranging from 104 to 107 cycles. All of the specimens were examined at the 50 pct life level at which stage the structures have reached a stable configuration.12 I) EXPERIMENTAL PROCEDURE Strip specimens, 0.006 in. in thickness, were prepared from base elements of 99.99 pct purity or greater. Specimens were fatigued by cementing the strips to a lucite substrate which was subjected to reverse plane bending. This method of testing has been described e1sewhere.7 After fatiguing, specimens were thinned and examined in a Philips EM 200 which was equipped with a goniometer stage capable of ±30-deg tilt and 330-deg rotation of the specimen. On the basis of separate calibrations,13 allowances were made for the relative rotation and inversions between the bright-field images and the diffraction patterns. II) RESULTS AND DISCUSSION The life behavior of the materials under different test conditions is shown in Fig. 1 in the form of plots of total strain range vs cycles to failure. Comparisons of structures produced in the different materials were made at amplitudes which produced equal numbers of cycles to failure. The influence of strain amplitude on the structures produced in the reference state material (copper tested at room temperature) is shown in Fig. 2. At the 104 life level the structure produced comprises cells similar to those previously observed.3,4 They are approximately 0.5 p in diam and the cell walls are generally more regular or sharper than those produced by unidirectional deformation.14 At the 10' life level the

Jan 1, 1969

By G. K. Burgess

(New York Meeting, October, 1913.) 1. COOLING AND HEATING CURVES BY G. V. BURGESS AND J. J. CROWE. § 22. THE results published in Professor Howe's paper 10 of our determinations on the Ac3 and Ar3 points for a series of commercial carbon steels 11 containing manganese in varying proportion, represent the average temperatures found by taking several (4 or 5) heating and cooling curves. The accompanying curves, Figs. 6 and 7, are, for each steel, the first heating and cooling curves as obtained by' the inverse-rate method of Osmond, using a cylindrical chronograph sensitive to -0.1 sec. for recording .time. Simultaneous observations were also taken by the differential method with a platinum neutral and plotted by Rosenhain's derived differential method and gave substantially the same results, the curves having sensibly the same contours in both cases.12. The observations were taken in vacuo on cylinders 5/8 by 1 in. in length bored to center with a 1/8-in. hole to receive porcelain insulated thereto-couple, and inclosed in an out- and in-glazed Berlin porcelain tube heated electrically by, platinum foil wound on the tube. The vacuum was usually better than 0.1 mm. The heating and cooling of the furnace was automatically controlled by a salt-water rheostat with shaped electrodes, giving a uniform rate of heating and cooling.

Jan 6, 1913

By G. Howell Clevenger

FORMERLY great divergence of opinion existed in regard to the influence of copper in steel, as affecting its various physical properties. More recently the investigations of Stead,1 Breuil,2 Wigham,3 Burgess and Aston,4 Ball and Wingham,5 Miiller,6 Dillner,7 and others have shown fairly good agreement in the major points. Copper steel has been shown by these investigators to possess special properties which make it valuable for certain uses. The general range of composition of the useful alloys of copper and iron, and of copper, iron, and carbon, has been quite definitely established, for the former by Burgess and Aston,8 and for the latter by Breuil 9 and Stead.10 Stead 11 has removed one difficulty in reconciling certain of the earlier data by calling particular attention to the influence of carbon in copper steels. Another point which caused confusion in the earlier investigations, particularly in regard to " red-shortness," was the presence of sulphur. Failure to note the presence of sulphur often led to attributing " red-shortness " to the presence of small amounts of copper. Turner 12 has called attention to the fact that where copper exists in an iron ore it often occurs as the sulphide, so that the presence of copper in steel made from such ores might also, necessarily, mean the presence of sulphur. It is evident that the effect in this case might be quite different than where pure copper was added.

Jan 10, 1913

By E. P. Klier, Jane Jellison

The course of the transformations on cooling in a series of Cu-A1 alloys has been followed by means of thermal analysis, X-ray diffraction, and optical metallogrAphy. Specimen size was found to have a pronounced effect on the time of completion of the pearlite transformation due to modification in the unmber of nuclei. The ordeving transformation zvas observed for all cooling rates for all alloys studied, and is considered insuppvessible. At intermediate cooling' rates, a transitional struchire for 011 alloy compositions and based ox the bainite reaction was observed. The alloys of aluminum and copper possess a eu-tectoidal region at from about 9.4 to 16 pct Al. Above the critical temperature, the disordered bcc phase, 0, is stable. At subcritical temperatures, the 0 phase decomposes under equilibrium conditions into the stable phases, a and y,. At the eutectoid composition, these phases develop a lamellar pearlite structure, in appearance much like that found in steels. During rapid quenching from supercritical temperatures, the diffusion processes required for the formation of the equilibrium phases are suppressed, and the 0 transforms to an acicular martensitic phase, designated B or y' depending on the aluminum content, or to a mixture of the two. Early investigators of this system concentrated on determining the equilibrium phase diagram, ''' the structure of the equilibrium phases,3 and later the isothermal decomposition of the 0 phase at subcritical temperatures, or the decomposition of the martensite on tempering. Other work which treated the nonequilibrium transformations concentrated largely on the martensite transformations and the structure of the b and y' phases.15"17 Since the pearlite and martensite transformations in this system resemble those in steels, the one being formed by nucleation and growth of adjacent lamellae of the equilibrium phases, and the other by a diffusionless, shear-type transformation, the present study has been undertaken to investigate the structures formed on continuous cooling, with the objective of constructing cooling transformation diagrams analogous to those derived for alloy steel systems. In particular it was desired to determine whether parallels between the A1-Cu and iron-iron carbide systems on continuous cooling extended to encompass the intermediate phase, namely the bainite structures, which form by a combined martensitic-diffusion mechanism,'' but which are supersaturated with respect to the equilibrium phase upon which their crystal structure is based. To this end the transformations in representative alloys of the Cu-A1 B eutectoid system have been followed, using the methods of thermal analysis, X-ray diffraction, and optical microscopy. EXPERIMENTAL PROCEDURE Materials. The materials tested were those used by Klier and Grmko 5 in their study of the isothermal decomposition of these alloys. The alloys were melted in graphite crucibles and chill-cast into 5 by 2 by 1/4 in. ingots. The ingots were then hot-rolled to sheets 1/8, 1/16, and 1/32 in. thick. The compositions are given in Table I. Specimen Preparation. Specimens for thermal analysis were in the form of small squares approximately 1/4 in. on a side. For the 11.9 and 13.5 Al-Cu alloys, the sheet thickness was 1/16 in. and for the 10.5 A1-Cu alloy the 1/32-in. sheet was used. One lead of a chromel-alumel thermocouple was spot-welded to either face of the square with the specimen itself thus forming the hot junction of the thermocouple. Specimens for X-ray analysis were prepared as filings from martensitic sheet stock and were sealed in glass ampoules under 1/2 atm of argon. Metallographic examination of the powder specimens was carried out on a portion of the heat-treated filings mounted in lucite. All specimens were examined metallographic ally after etching

Jan 1, 1965

By J. J. Kramer, W. A. Tiller, G. F. Bolling

THE axial orientations of columnar crystals in unidirectionally solidified ingots of zone-refined zinc and tin have been examined using the techniques recently described by us.' Both metals had an initial purity of 99.999+ pct before zone refining and satisfied all qualitative guides for very high purity after zone refining. Fig. 1 presents the orientations of the columnar crystals for tin and zinc, respectively. For tin, there is a preference for grains growing in the [I101 direction, i.e., the dendrite direction.' For zinc, a weak preference for the [0001] direction is seen (a random population in the unit triangle would give equal density on either side of the 60-deg circle centered on [0001]).' There was no apparent evidence for dendritic growth in either metal. These results, and those previously reported,' confirm the recent investigation by Hellawell and Herbert of casting textures observed for zone-refined lead, zinc, and tin.4 The preferred casting orientations for these metals in both the impure and zone-refined states are given in Table I. In the experiments on zone-refined lead,' it was shown that two separate textures existed: a preferred (111) nucleation texture existed at the chill surface and a preferred (111) texture developed during solidification. Either effect is sufficient to produce the result recorded in Table I. Similarly the results recorded for tin and zinc may be attributable to either a preferred nucleation or a preferred growth phenomenon and additional experiments are required to separate these two possible explanations. Hellawell and Herbert' attribute their results to a preferred growth process and account for it on the basis of a "layer-growth" model proposed much earlier by one of the present authors.5 An equally satisfactory growth model would be that based upon anisotropic surface energies. However, both of these growth models would indicate that the preferred growth direction for very pure metals should be normal to the plane of closest packing in the metal. This is consistent for zinc but, contrary to the conclusion of Hellawell and Herbert, it is not consistent for tin since here the (100) is expected to exhibit a greater density of packing than the (110). Although tin is expected to exhibit an anisotropy of thermal conductivity, we would expect the thermal conductivity to be the same in the [1101 as the [l00]

Jan 1, 1963

By W. Vickers

Single-point rotating cantilever fatigue tests have been carried out at 550" to 650°C on beryllium produced by a variety of fabrication routes. All the specimens gave similar plots of stress against number of cycles to failure but failure occurred at a lower stress range in the hot-pressed material than in the extruded material. A discontinuity or "kink" appears to exist in the curves and this is considered. A true endurance limit was not evident up to 10 cycles. When Brush hot-pressed warm-extruded beryllium was tested at 600oC a relationship where Nf = number of cycles to failure, ?ep = plastic strain range, and c = constant, was shown to be applicable for values of ?ep > 6.3 x 10-4. Further for this moterial and temperature and for frequencies between 150 and 4500 cycles per min, the variation of number of cycles to failure with frequency obeyed the expression N = Af0.44 where N =number of cycles to failure, f = frequency, and A = constant. Examination of the fracture surfaces indicated that they were generally intergranular and in all cases the higher the stress the more irregular was the fracture. Microexamination has shown that grain boundary cavities and fine intergranular cracks were associated with the fractures. REVERSING or cyclic stresses set up in fuel elements can lead to failure by fatigue. Fatigue stresses may be developed by instability of the fuel element in the coolant flow leading to rattling against the channel wall or fin vibration. Thermal cycling also sets up reversing stresses due to differential expansion of fuel and can. Fatigue resistance is of some importance therefore in the evaluation of canning materials. At the time the work described in this paper was initiated, beryllium was being considered for the can and endcaps of an advanced gas-cooled reactor fuel element. With this in mind, beryllium obtained from different manufacturing routes was tested on single-point rotating cantilever fatigue machines at temperatures the same as those envisaged in the reactor to obtain appropriate fatigue data. The results obtained indicated a discontinuity in the S-N curves coinciding with the yield stress in the specimens. Although in some cases the data is insufficient to warrant a detailed quantitative analysis, a qualitative analysis of this effect was thought to be justified. 1) EXPERIMENTAL DETAILS The materials used are given in Table I together with the production route and available analyses. The specimen shape is shown in Fig. 1. The fatigue specimens were machined slightly oversize, after which they were given a stress-relieving anneal of 1 hr at '750° C in vacuum. This was followed by an etch in sulfuric-phosphoric-chromic acid mixture down to the finished dimensions given in Fig. 1. Etching served to remove the interference colored oxide film formed on the surface during annealing and very small recrystallized surface grains too. The specimen was mounted in the grips of a cantilever machine and rotated inside a furnace controlled at the testing temperature. The specimen was loaded in increments as the speed was increased until the required load was applied at 5000 cycles per min. A revolution counter was used to record the total number of cycles.

Jan 1, 1964

By David Sinclair

THE dimensions and shapes of mine structures may at present be determined by (1) field experience, (2) structural calculations, and (3) barodynamic tests.§ None of these, however, provide information as to the stresses, their direction and distribution in a mine structure. For the design of pillars, the placing of openings relative to working places such as gangways, haulage levels, etc., this information is desirable. Photoelastic methods provide this and, while relatively new, their appli-cation in all engineering fields is assuming increasing importance. This paper, therefore, will present the method with a few applications to mining, and with sufficient reference for further study by the reader. PART I THE PHOTOELASTIC METHOD The photoelastic method of stress measurement determines the stress distribution in a two-dimensional model by optical means. It makes use of the fact that when a piece of isotropic transparent material, such as glass, celluloid, or Bakelite, is stressed and viewed in polarized** light a picture is seen which indicates by its color bands of light, or fringes, the magnitude of the stresses, their direction and their distribution.

Jan 1, 1940

By Charles W. Eichrodt

So much has already been wimitten on this vast subject of ground movement and subsidence, and so many data collected and commented upon, that in this paper the author proposes to confine himself to the submission of some plans showing what has actually occurred in the mining of several of the orebodies which have come under his direction during the last 20 years or so at Rio Tinto, Spain. Some notes are added for the purpose of describing the plans. The author does not attempt to give his opinion as to the reasons why the subsidences have occurred in the manner in which they are found, but confines himself to giving as true a picture as possible of the movements that have taken place. Effects of Excavation on Ground Movement Broadly, the effects of excavations on the overhead or adjacent ground can be classified under three headings, as follows: 1. Ground movements due to excavations made from surface only, such as those made in opencasts where no excavations underlie them. The stability and strength of the ground to resist movement must determine the slopes. 2. Ground movements due to underground excavations only; i. e., where the surface is not removed. 3. Ground movements due to a combination of both; i. e., from excavations taking place from the surface for the purpose of stripping an orebody and recovery of the ore to a certain horizon combined with excavations being carried out below, at the same time, for the winning of the ore below the horizon at which mining by stripping or opencasting under the specific conditions is no longer economical. Nos. 1 and 2 represent independent and distinct problems and No. 3 is a combination, which presents results much more difficult of explanation. The accompanying plans and sections give examples of what has occurred under the three conditions. Unfortunately, the class of ground is not exactly the same in them all, but the difference is not greater than is generally encountered over districts of equal area. Excavations from Surface To illustrate No. 1 three cross-sections are shown, of excavations made for the stripping and excavation from three different orebodies

Jan 1, 1930

By J. B. Austin

It is now well established, chiefly through the work of Davenport and Bain,' that the influence of temperature upon the rate of transformation of austenite to ferrite at constant temperature is represented by a curve of the type shown in Fig. 5 on page 295 and the reasons for this general shape have been discussed by them in a qualitative way. Thus, the upper part, or nose, of the curve is the resultant of two factors of which the rate of change with temperature is not the same, and by analogy with other like cases it is almost certain that these factors are: (1) the chemical potential, or driving force, of the transformation, which is zero at the equilibrium temperature and increases rapidly as the temperature falls below this point; and (2) a factor analogous to viscosity, or resistance of the crystal lattice to a change in structure, which also increases as the temperature is lowered. Just below the equilibrium temperature this virtual "viscosity" of the metal is relatively low—that is, the metal offers little resistance to a change in lattice—but the driving force is also small, and the net result is that the reaction proceeds slowly. As the temperature is lowered the magnitude of both factors increases; but at first the driving force increases more rapidly than the resistance and the reaction is observed to proceed more rapidly. At lower temperatures, however, the resistance begins to increase faster than the driving force, and eventually, at the nose of the curve, overbalances it, causing the reaction to proceed more and more slowly thereafter. This concept is somewhat, though not completely, analogous to that employed in certain electrical cases where the precise factors are known and are separately measurable. The current I (in amperes or coulombs per second) flowing through a given wire is proportional to the impressed potential, or driving force E, and is inversely proportional to the resistance R of the wire; that is, I = E/R. Stated in a slightly different way, the time required for one coulomb to pass a given point (1/I) is equal to R/E. If the temperature of the whole circuit is changed, the resultant change in 1/I will depend on the relative rates of change of R and of E,

Jan 1, 1935

By J. B. Austin

It is now well established, chiefly through the work of Davenport and Bain,' that the influence of temperature upon the rate of transformation of austenite to ferrite at constant temperature is represented by a curve of the type shown in Fig. 5 on page 295 and the reasons for this general shape have been discussed by them in a qualitative way. Thus, the upper part, or nose, of the curve is the resultant of two factors of which the rate of change with temperature is not the same, and by analogy with other like cases it is almost certain that these factors are: (1) the chemical potential, or driving force, of the transformation, which is zero at the equilibrium temperature and increases rapidly as the temperature falls below this point; and (2) a factor analogous to viscosity, or resistance of the crystal lattice to a change in structure, which also increases as the temperature is lowered. Just below the equilibrium temperature this virtual "viscosity" of the metal is relatively low—that is, the metal offers little resistance to a change in lattice—but the driving force is also small, and the net result is that the reaction proceeds slowly. As the temperature is lowered the magnitude of both factors increases; but at first the driving force increases more rapidly than the resistance and the reaction is observed to proceed more rapidly. At lower temperatures, however, the resistance begins to increase faster than the driving force, and eventually, at the nose of the curve, overbalances it, causing the reaction to proceed more and more slowly thereafter. This concept is somewhat, though not completely, analogous to that employed in certain electrical cases where the precise factors are known and are separately measurable. The current I (in amperes or coulombs per second) flowing through a given wire is proportional to the impressed potential, or driving force E, and is inversely proportional to the resistance R of the wire; that is, I = E/R. Stated in a slightly different way, the time required for one coulomb to pass a given point (1/I) is equal to R/E. If the temperature of the whole circuit is changed, the resultant change in 1/I will depend on the relative rates of change of R and of E,

Jan 1, 1935

By J. Pohlig

It is with more than ordinary pleasure that I have complied with the request of the President of the Verein Deutscher Eisenh?tten leute, to read before this meeting and in this country a paper on a system of transport by wire ropeways which, during the past twenty years, has advanced with rapid strides in Europe, and more partic ularly in Germany. America, with its immense network of railways and water communications, should offer peculiarly favorable oppor tunities for the introduction of these ropeways as feeders, more par ticularly in the mining districts, where, as in Europe, a hilly or mountainous region frequently opposes serious obstacles to the em ployment of surface railroads. The importance of the aerial ropeway as a factor in the develop ment of mineral resources deserves the consideration of every mining engineer. As a case in point, the rich iron-ore of the sierra de Bedar in southern Spain, would probably remain untouched to this day but for an aerial ropeway, 15.6 km. (9 3/4 miles) long, which con nects the mines with the shore of the Mediterranean near the town of Garrucha, affording cheap carriage to the nearest shipping point. A railway would have cost about $480,000 (£100,000), but such a heavy outlay would have more than absorbed all prospective profits on the sale of ore. As the aërial ropeway could be built for a little more than one-fourth of that amount, or $124,800 (£26,000), its erection furnished the successful solution of a profitable mining enterprise. This line has been in satisfactory operation for over two yearn; its details will be referred to later. But beside their importance as feeders, ropeways are valuable as a means of conveyance between works and distant warehouses, and between separate buildings in the same works. Their rapidly spreading introduction for such purposes indicates how keenly the necessity of effecting economies is felt under the stress of an everaccentuating competitiop.

Jan 1, 1891

By W. J. Taylor

SOME eight years ago I became interested to a considerable extent in one of the well-known deposits of sulphury iron ore in New Jersey, and, as a consequence, soon became interested in the subject of the elimination of the sulphur. On investigation I found that no department of pig-iron manufacture had received less attention than that of roasting or calcining ores for the elimination of sulphur, carbonic acid, water, etc. My first devices to eliminate sulphur were naturally in the lime-kiln line, with solid fuel mixed with the ore, in the same way as common lime is burnt with small coal. While this process will succeed tolerably well for carbonates or calcareous ores, even though they contain sulphur as sulphates, I soon found it would not answer for our magnetites, which are very hard and dense, carrying from 2 to 5 per cent. of sulphur as pyrites, and become very sticky as soon as any attempt is made to roast them at a high heat. Some laboratory experiments developed the fact that the nearer the gaseous current surrounding the hot ore approached in composition pure atmospheric air the more rapidly the sulphur was oxidized, while in an atmosphere composed largely of carbonic oxide, carbonic acid, and nitrogen, it seemed almost impossible to reduce the sulphur more than one-half, which was a confirmation of the well-known fact that just one-half of the sulphur in pyrites volatilizes very easily, and that the trouble is mainly to remove the remainder. It was thus apparent that the atmosphere surrounding the ore was of as much importance as the temperature. This led to the trial of a kiln to heat the ore by conduction through the brick. It was divided into ore-chambers, which were surrounded with flues, through which the products of combustion of gaseous fuel passed without coming in contact with the ore. We did not, however, succeed by this arrangement in getting more than a cherry-red heat, which, of course, was too low to accomplish anything. I had at this point become convinced – 1st. That gaseous fuel only would accomplish the work successfully. 2d. That since a high temperature was required to insure a rapid oxidation of the sulphur, and also to make the ore open and porous,

Jan 1, 1881

The following is a, partial list of members and guests who called at Institute headquarters during the period Aug. 10, 1917, to Sept. 10, 1917. L. P. Barrett, Ann Arbor, Mich. Milo W. Krejci, Great Falls, Mont. H. J. Brown, Ridley Park, Pa. Chas. N. La Porte, Peoria, Ill. George M. Cheney, Guayaquil, Ecuador. Askin Nicholas, Bismuth, Australia. W. Lorrain Cook, Chicago, Ill. C. Edwin Nighman, Butte, Mont. John J. Croston, Atlanta, Ga. W. L. Remick, Norfolk, Va. A. F. Duggleby, Santiago de Cuba. T. C. Roberts, Jerome, Ariz. Geo. A. Easley, La Paz, Bolivia. Burr A. Robinson, New Haven, Conn. Burr J. French, Magdalena, Jal., Mex. J. L. Schueler, Peoria, Ill. A. D. R. Galloway, Komfrodua, W. Af. Walter L. Steele, Van Horn, Tex. P. R. Hines, Milwaukee, Wis. F. J. S. Sur, Denver, Colo. T. H. Jenks, Wickenburg, Ariz. C. A. Thomas, Dawson, Y. T., Canada. Paul S. King, Wilmington, Del. E. E. White, Ann Arbor, Mich. Harry Wolf, Golden, Colo. L. W. Bahney has left the Sheffield Scientific School, Yale, to become metallurgical engineer with the Scovill Manufacturing Co., Waterbury, Conn. H. B. Barling, formerly chief engineer of Breitung & Co., of New York, has been appointed manager of the mining department.

Jan 10, 1917

By P. Payson

INTERGRANULAR corrosion in corrosion-resistant chromium-nickel steels has been widely discussed in the last few years. So far as the author knows, nothing has been published which definitely shows that this tendency can be eliminated, although suggestions have been made that a remedy is possible if certain specifications are satisfied. Strauss, Schottky and Hinueber1 suggested an increase in chromium content to an extent that would prevent the impoverishment of chromium on carbide precipitation sufficient to cause impaired corrosion resistance in the matrix. Pilling2 suggested an increase in nickel, so that the stability of the austenite matrix would be increased. Aborn and Bain,3 say that "intergranular attack-may be eliminated by preventing carbide precipitation (e. g., by lowering the carbon content or somehow raising its solubility) or by precipitating it permanently under conditions which will not remove chromium from solid solution so drastically at any points." In a later paper Aborn and Bain4 state that many different metallic elements, tungsten, vanadium, copper, aluminum, molybdenum, titanium and zirconium, have been added in small per-centages to increase the stability of these alloys, but that it has been reported that according to the copper sulfate test these added elements act only to retard the intergranular corrosion, and that increased nickel and chromium contents have similar effects. Newell5 showed that the lower the quenching temperature before reheating, the less severe was the intergranular attack.

Jan 1, 1932

By A. H. Geisler, R. F. Mehl, S. Barrett

The studies of the mechanism of pre= cipitation and of the resulting property changes in aluminum-silver alloysl-3 have presented some new concepts of the aging reactioxl—concepts that may be fundamental in establishing a complete quantitative theory of age-hardening. With both aluminum-copper and aluminum-silver alloys, thin platelike aggregates form parallel to the Widmanstätten planes of precipitation.l,4 These platelets grow in lateral dimensions and in thickness with aging time until they arc thick enough to produce an X-ray diffraction line pattern. Their structure is that of a transition lattice which is coherent with the matrix, since the atomic spacings in the matrix and the precipitate are the same on each side of their common interface.2,4 The equilibrium precipitate forms later from and at the expense of the transition lattice. Since age-hardening occurs during the formation and growth of the transition lattice, the major contributor to the hardening no doubt is the stress condition in the matrix resulting from the matrix holding the precipitate in the coherent condition. The development of these stresses also appears to be responsible for an increase in the electrical resistance superimposed on the normal decrease resulting from a change in the concentration of the matrix during aging.3 The two consecutive reactions constituting the aging process are accompanied by characteristic microstructural alterations. The initial precipitation process results in the formation of a Widmanstatten figure, while the later transformation of the transition lattice to the equilibrium lattice is accompanied by a coalescence, in the case of aluminum-copper alloys,5 and by the appearance of light areas at the grain boundaries with aluminum-silver alloys showing a structural rearrangement of matrix and precipitate.3 This latter reaction is similar microscopically to the so-called "discontinuoils precipitation." Possibly discontinuous precipitation is a process that often if not always follows true precipitation of a coherent lattice and during which the transformation of the coherent transition lattice to the equilibrium lattice takes place accompanied by matrix recrystallization.3 Expermental Procedure The alloys were made using 99.97 per cent A1 or distilled magnesium, with

Jan 1, 1943

By R. F. Mehl, A. H. Geisler, S. Barrett

The studies of the mechanism of pre= cipitation and of the resulting property changes in aluminum-silver alloysl-3 have presented some new concepts of the aging reactioxl—concepts that may be fundamental in establishing a complete quantitative theory of age-hardening. With both aluminum-copper and aluminum-silver alloys, thin platelike aggregates form parallel to the Widmanstätten planes of precipitation.l,4 These platelets grow in lateral dimensions and in thickness with aging time until they arc thick enough to produce an X-ray diffraction line pattern. Their structure is that of a transition lattice which is coherent with the matrix, since the atomic spacings in the matrix and the precipitate are the same on each side of their common interface.2,4 The equilibrium precipitate forms later from and at the expense of the transition lattice. Since age-hardening occurs during the formation and growth of the transition lattice, the major contributor to the hardening no doubt is the stress condition in the matrix resulting from the matrix holding the precipitate in the coherent condition. The development of these stresses also appears to be responsible for an increase in the electrical resistance superimposed on the normal decrease resulting from a change in the concentration of the matrix during aging.3 The two consecutive reactions constituting the aging process are accompanied by characteristic microstructural alterations. The initial precipitation process results in the formation of a Widmanstatten figure, while the later transformation of the transition lattice to the equilibrium lattice is accompanied by a coalescence, in the case of aluminum-copper alloys,5 and by the appearance of light areas at the grain boundaries with aluminum-silver alloys showing a structural rearrangement of matrix and precipitate.3 This latter reaction is similar microscopically to the so-called "discontinuoils precipitation." Possibly discontinuous precipitation is a process that often if not always follows true precipitation of a coherent lattice and during which the transformation of the coherent transition lattice to the equilibrium lattice takes place accompanied by matrix recrystallization.3 Expermental Procedure The alloys were made using 99.97 per cent A1 or distilled magnesium, with

Jan 1, 1943