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Question 1—1s Structural Deformation of Some Character Always Necessary for the Migration of Mineralizing Solutions. Especially in Limestone or Other Dense Rocks Discussion by: A. M. Bateman.........120 S. F. Kelly........... 125 C. H. Behre. Jr........118.123 G. B. Langford......... 123 T. M. Broderick.........121 H. E. McKinstry......... 116 R. H. Clarke.......... 123 D. H. McLaughlin........ 118 J. D. Forrester for M. H. Gidel . . 11.4 W. H. Newhouse......... 123 G. M. Fowler......112. 116. 121 R. H. Sales.......... 116 L. C. Graton..........121 Edward Sampson......... 122 R. D. Hoffman.........120 J. 8. Stevenson.......... 117 I. B. Joralemon.........121 L.B.Wright........... 112 Question 2—What Are Some of the Chemical and Physical Characteristics in Certain Beds within a Limestone Formation that Determine Their Favorable or Unfavorable Aspects as Reservoirs for Ore and Other Introduced Minerals? Drscussion by: G. W. Bain....... 127 G. M. Fowler....... 131 C. H. Behre, Jr. ..... 131 L. C. Graton....... 128 J. S. Brown. ...... 126 G.. F. Loughlin...... 128 E. L. Bruce ....... 129 R. H. Sales........ 129 R. H. Clarke...... 130 A. K. Snelgrove....... 130 Question 3—Are There Any Criteria by Which One Can Determine When Mining Operations Have Reached the Roots of an Ore Deposit? Discussion by: J. M. Boutwell...... 139 G. B. Langford..... 139 T. M. Broderick..... 138 . G. F. Loughlin...... 135 E. Y. Dougherty..... 137 R. H. Sales...... . 134 G. M. Fowler....... 132 F. A. Thomson..... 136 I. B. Joralcmon...... 132 Edward Wisser:..... 140 General s. F. Kelly................. . ...........145 References...............................146

Jan 1, 1941

By Eugene Guccione

FOREWORD-In an interview published in the May 1975 issue of MINING ENGINEERING, one of the world's most respected scientists, nuclear physicist Hans Bethe discussed the reliability of nuclear power plants and their unmatched safety record. The Nobel laureate also refuted the anti-nuclear lobby's claims that nuclear plants emit harmful radioactive effluents. When asked whether the US has sufficient uranium resources, Dr. Bethe replied that "it depends on the time of introduction of the commercial and advanced breeder reactors which, in effect, can generate more nuclear fuel than they consume." The United States has enough U3O8 to last until the breeders can take over in 1995, said Dr. Bethe. But lie warned that extensive environmental damage would be caused if "environmentalists" succeed in delaying the development of the breeder. Any such delays, he noted, would vastly increase the demand for U3O8 in conventional reactors; this, in turn, would impose an immense burden on the US uranium mining industry, which would then have no alternative but to mine the low-grade Chattanooga deposits.

Jan 4, 1977

By O. N. Carlson, E. D. Gibson, D. W. Bare

SEVERAL investigators have recently contributed to a clarification of the controversy concerning the allotropy of chromium. Ross and Hume-Rothery1 and Vasyutinskii et a1.2 have reported high-temperature X-ray data on unalloyed chromium from which they have concluded that no allotropic transformation occurs in this metal. Bechtoldt and vacher3 in their investigation of the Cr-Ni alloy system found no evidence for a transformation in chromium and Smart and Haynes4 in critically analyzing the data presented to date concerning this system could find no evidence for a eutectoid reaction in this system. Wyder and Hoch5 in research concurrent with our

Jan 1, 1964

By R. E. Maringer, B. A. Wilcox, A. H. Clauer, J. H. Sovik

Internal friction measurements from 25° to 1100°C have been made on specimens machined from TD nickel bar. The use of miniature torsional pendulum specimens alloived damping measurements to be made on specimens whose axes were parallel to the original bar axis, at 45 deg to the bar axis, and at 90 deg to the bar axis. There was a marked difference in damping with varying specimen orientation, and this was related to the pronounced micro structural anisotropy of this material. Varying the microstructure by cold working plus re crystallization annealing also greatly affected the damping behavior. Recrystallization essentially removed the broad peak found at 400° to 650°C in as-received TD nickel, and in addition greatly lowered the background damping. DlSPERSION hardened metals such as SAP-type alloys (Al + AI2O3) and TD nickel (Ni + 2 pet ThO2) have been developed primarily to improve high temperature strength properties. There have been numerous damping studies on SAP-type alloys,1"9 and one of these6 showed that the damping capacity was greater than that of 99.5 pet pure Al from 25° to 600°C. Several preliminary observations on TD nickel10 and an experimental Ni-ThO2 alloy11 indicated that here, too, the damping capacity was quite high. A high damping capacity in TD nickel, combined with its relatively good high temperature creep strength, could provide an additional benefit for some high temperature structural applications. This work was undertaken to study the internal friction behavior of TD nickel bar over the temperature range 25° to 1100°C. In Ni-ThO2 and Ni-Cr-ThOz alloys, other properties such as steady-state creep rate,12 creep rupture life,12'13 modulus of elasticity,14 and ultimate tensile strength,14 are sensitive to orientation with respect to the primary working direction. Thus, in order to examine anisotropy effects in damping, studies were made on specimens machined with their axes parallel, at 45 deg, and transverse to the original bar axis. MATERIAL DESCRIPTION AND EXPERIMENTAL PROCEDURES The material used in this investigation was from the same £ in. diam TD nickel bar that has been studied previously.13'15 The structure typical of the as-received TD nickel is shown in Fig. 1, and additional details of the structure can be found in Refs. 13 and 15. Elongated grains are parallel to the bar axis and are about 10 to 15 u long and about 1 u in diam. This structure is very stable, since no recrys-tallization or grain growth occurred after annealing for 3 hr at 1300°C.13 The total ThO2 content was 2.3 vol pct and the particles ranged in diameter from 100 to 1000. The average particle diameter was 370Å and the mean planar center-to-center particle spacing was 2340Å. To examine the influence of grain and subgrain structure on damping, a section of TD nickel bar was swaged and drawn at room temperature from 1/2 in. diam bar to 0.040 in. diam wire (99.3 pct re-duction in area). The structure in Fig. 2 resulted after the wire was annealed for 3 hr at 1250°C. The recrystallized grains were much larger than the grains in as-received bar, but still showed considerable elongation in the direction of working. A high

Jan 1, 1970

By Nicholas J. Grant, William H. Ferguson, Bill C. Giessen

Ta-lr alloys have been examined over the complete range of compositions using metallographic and X-ray techniques. The terminal solid-solubility limits, solidus temperatures, and intermediate phases were determined. There are four inter-mediate phases: o, tetragonal, analogous to a (Fe-Cr); 0,. orthorhombic; az, tetragonal; and a TaZr,, cubic, AuCu, structure. Of these, o and a, melt peritecticaLly, a, decomposes peritectoidally, and a TaIrs has a congruent melting point. Three peri-tectic and one eutectic reactions occur. 1 HE Ta-Ir phase diagram has not been treated in the literature. Intermetallic phases and terminal solid solubilities have been described and estimated.'-4 The previous findings have been confirmed in this work. A complete diagram with additional phases and solidus curves is presented. EXPERIMENTAL METHODS Starting materials were tantalum powder of 99.8 pct purity, supplied by the National Research Corp., Cambridge, Mass., and iridium powder of 99.9 pct purity, supplied by the Bishop Co., Malvern, Pa. Both powders were -200 mesh. Table I lists the analyses as given by the suppliers. A standard procedure was employed in making the alloy buttons, which was described in detail in Refs. 5 and 6, and which included repeated arc-melting of powder compacts and determination of the composition by the weight-balance method. This is especially justifiable due to the similar vapor pressures of tantalum and iridium. In the critical parts of the diagram, a spacing of 1/2 at. pct was used, based on premelted master alloy material. The observed structures and X-ray spectra were consistent with the composition sequence of the alloys. Table II lists the final alloy atomic percentages with the calculated percent error of composition. Cross sections representative in grain size and composition were selected for the study, as described in Ref. 5. Based on prior experience, no contamination from the tungsten electrode was considered. Temperature measurements were obtained using a Leeds and Northrup optical pyrometer and have been described extensively in Refs. 5 and 6. All heat treatments were made in the high-vacuum tantalum tube furnace mentioned in Ref. 5. Samples annealed in tantalum containers cooled from 1800" to 1000"~ in 5 sec; those annealed in ceramic crucibles cooled over the same range in 15 to 30 sec. The thermal treatments are listed in Table 111. All specimens were homogenized at 1735°C for 24 hr. Tantalum-rich compositions were treated in tantalum buckets. Iridium-rich compositions were annealed in tantalum buckets below the eutectic temperature of 1950°C; up to 2200°C, tungsten-lined tantalum baskets were used; and above 2200°C thoria crucibles were used. It was found that refractory oxides (in early stages of the investigation alumina crucibles were used up to 1800°C and thoria crucibles were used to 2400"~) give off minute amounts of oxygen and thereby influence the microstructure considerably, or react directly with the alloys.

Jan 1, 1963

By Tracy C. Jarrett

In 14-karat gold alloys, as in 10-karat gold alloys,3 the addition of such metals as zinc, nickel, copper and silver produces a wide range of physical and chemical properties such as color, hardness, corrosion resistance and melting points. These important factors made gold alloys desirable in the optical work and for jewelry. In 14-karat gold alloys, as well as in 10-karat or 12-karat, it is possible to have two golds of the same color, one of which will corrode more than the other, although it may be harder. Some of these alloys have been investigated by F. E. Carter,' E. M. Wise,= and L. Sterner Rainer.4 In this group of alloys the gold content of 58.3 per cent (14-karat) was held constant, while the copper varied from 36 to 24 per cent. The balance of the alloy consisted of one, two, or three of the metals silver, zinc, and nickel. All the alloys were melted in an induction furnace and cast into an ingot that measured 3 by 1/4 by 3/4 in. The method of determining the melting points, color determinations and acid test have been discussed in previous papers by the author.2, 3 It is necessary to use small test bars (Fig. I). The tensile bars were annealed at 1300oF. for 20 min. and then air-cooled. There being a very marked drop when the yield point was reached, "the drop of the beam" method was used in determining the yield point. Alloys Gold-copper.zilzc A~loys.—Gold-copper-zinc alloys melt, cast, and labricate easily. The low melting points, however, limit their use to some extent. The physical properties of these alloys are not entirely satisfactory. When the ultimate strength is satisfactory, the elongation and reduction of area are too low (Table I). Gold-copper-sinr-nickel A lloys.—In this group of alloys, in which zinc and nickel replace copper in approximately equal parts (Table I), the yield point, hardness, and ultimate strength increase with the increase of zinc and nickel content, while the reduction of area and elongation decrease. The physical properties are desirable, particularly of the white gold alloys. The alloys of this group are not difficult to cast and fabricate, and most of these colors could be used commercially. Gold-copper-nickel A lloys.—Gold-copper-nickel alloys are similar to the gold-copper-zinc and the gold-copper-zinc-nickel alloys. They do, however, possess high yield point, ultimate strength and hardness, somewhat higher than those of the other two groups mentioned (Table I). The difficulty with these alloys is that they fire-crack readily and heavy rolling is necessary between annealings. Being hard to fabricate, many

Jan 1, 1941

By Tracy C. Jarrett

In 14-karat gold alloys, as in 10-karat gold alloys,3 the addition of such metals as zinc, nickel, copper and silver produces a wide range of physical and chemical properties such as color, hardness, corrosion resistance and melting points. These important factors made gold alloys desirable in the optical work and for jewelry. In 14-karat gold alloys, as well as in 10-karat or 12-karat, it is possible to have two golds of the same color, one of which will corrode more than the other, although it may be harder. Some of these alloys have been investigated by F. E. Carter,' E. M. Wise,= and L. Sterner Rainer.4 In this group of alloys the gold content of 58.3 per cent (14-karat) was held constant, while the copper varied from 36 to 24 per cent. The balance of the alloy consisted of one, two, or three of the metals silver, zinc, and nickel. All the alloys were melted in an induction furnace and cast into an ingot that measured 3 by 1/4 by 3/4 in. The method of determining the melting points, color determinations and acid test have been discussed in previous papers by the author.2, 3 It is necessary to use small test bars (Fig. I). The tensile bars were annealed at 1300oF. for 20 min. and then air-cooled. There being a very marked drop when the yield point was reached, "the drop of the beam" method was used in determining the yield point. Alloys Gold-copper.zilzc A~loys.—Gold-copper-zinc alloys melt, cast, and labricate easily. The low melting points, however, limit their use to some extent. The physical properties of these alloys are not entirely satisfactory. When the ultimate strength is satisfactory, the elongation and reduction of area are too low (Table I). Gold-copper-sinr-nickel A lloys.—In this group of alloys, in which zinc and nickel replace copper in approximately equal parts (Table I), the yield point, hardness, and ultimate strength increase with the increase of zinc and nickel content, while the reduction of area and elongation decrease. The physical properties are desirable, particularly of the white gold alloys. The alloys of this group are not difficult to cast and fabricate, and most of these colors could be used commercially. Gold-copper-nickel A lloys.—Gold-copper-nickel alloys are similar to the gold-copper-zinc and the gold-copper-zinc-nickel alloys. They do, however, possess high yield point, ultimate strength and hardness, somewhat higher than those of the other two groups mentioned (Table I). The difficulty with these alloys is that they fire-crack readily and heavy rolling is necessary between annealings. Being hard to fabricate, many

Jan 1, 1941

By R. E. Boni, G. Derge

Application of surface tension measurements has been made to molten silicates in order to determine the effect of sulphur upon the surface tensions of synthetic blast furnace slags. In melts with the same molar ratios of AI2O3 + SiO2 to the basic oxides BaO, CaO, and MgO, the surface tensions increase in the order of oxides listed. Sulphur lowers the surface tensions of the basic BaO, basic CaO, acid BaO, acid CaO, and acid MgO slags in that order; and the concentration of excess sulphur at the slag surface increases in this same order of slags. The rate and amount of desulphurization of iron by these slags is favored by a low excess surface sulphur in the slag. IN a previous paper,1 the authors discussed the applications of surface tension measurements of molten silicates to problems of metallurgical interest. Of major importance is the aid that such studies lend to forming a working concept of slag structure. Since the ability of a slag to desulphurize metal depends strongly upon its composition,2"4 fundamental information about the structural differences between various slags which contain sulphur should aid the development of better desulphurizing media. For that reason, the present investigation concerns the effect of sulphur upon the surface tensions of synthetic blast furnace slags. Experimental Procedure , The maximum bubble-pressure method was used to measure the surface tensions of synthetic blast furnace slags in the temperature range 1300° to 1600°C. The method is based upon the equation which was developed by Schrödinger5 from a consideration of the forces of hydrostatic pressure and surface tension of a liquid which oppose bubble formation when a gas is forced through a tube immersed in the liquid. where y is the surface tension between argon and slag in dynes per centimeter; g is the gravitational constant, centimeters per second2; H is the manometer reading in centimeters; s2 is the density of manometer fluid in grams per cubic centimeter; S1 is the density of the medium in grams per cubic

Jan 1, 1957

By completing a double-track railroad cut-off, involving a large volume of cut and fill, in addition to a half-mile bridge spanning an important French river, engineers in the zone of the Services of Supply have opened up a route by means of which hours of valuable time are being saved in transporting men and materials to combat areas. Its most valuable feature is that it provides a means for by-passing a certain city where several, French railroad lines converge and where, consequently, congestion of traffic has been responsible for delayed train movements. To avoid a double-track turnout, which would have involved a cross-over cutting one of the French tracks, the two track connections of the cut-off are taken off from opposite sides of the French right-of-way. This involved running one of the new lines of the cut-off out from the main line track on a fill to a point where it could be brought back and over the French track on a steel girder bridge. Beyond this bridge the two tracks of the cut-off converge and straighten out in the standard type of double-track construction. The new cut-off is about 5 3/4 mi. (9 km.) long and has required 160,000 cu. yd. of cut (122,329 cu. m.) and 414,000 cu. yd. (316,527 cu. m.) of fill. In the biggest continuous fill a-yardage of 180,000 has been placed; the road-bed in this section-being-about 40-ft. (12 m.) above the original ground surface. The bridge portion of the cut-off, 2190 ft. (667.5 in.) long, is made up of ninety-nine 14-ft. (4.3 m.) timber spans and sixteen 50-ft. (15.3 m.) steel girder spans. This is the longest bridge which engineers of the American Expeditionary Forces have built in France. The timber spans are supported by pile trestle bents, while the longer steel girders rest on clusters of 43 piles. Two pile drivers of the steam-hammer type worked from each -end of the bridge toward the center of the river, while a drop hammer worked westward from the shore end. In this structure nearly two thousand piles 50 ft. or so in length had to be driven.

Jan 12, 1918

"Located on Missouri River 16 miles N. E. from Helena, 67 miles from Butte and 18 miles below Canyon Ferry. River drainage area 16,000 sq. miles.DAM: Built in 1907; rebuilt of concrete in 1911. Maximum height 129.3 ft., length over all 720 ft.; elevation of crest 3621; elevation of bridge 3640; normal head 60'; normal tail water elevation 3568.5.SPILLWAY: Length over all 491 ft.; clear opening in 24 sections; 445' flashboards, 14 high to elevation 3635, which is 3.5' above Canyon Ferry normal tail . water elevation.PENSTOCKS: 5 of steel, 12 ft. diam., about 75 ft. long from canal to main turbines, and 3 40-inch diam. from canal to exciters.WHEELS: 5 main pairs, horizontal, center discharge, 48-inch bronze turbines, 225 rev., 4,000 h-p. at 60 ft. head, made by S. Morgan Smith Co., direct coupled to generators, and--3 single, horizontal, 25-inch units, direct coupled to exciters.GENERATORS: 5 W. E. & M. Co. water wheel type, revolving field, 2,800 k-w., 2,400 v., 225 r. p. m., 3 phase, 60 cycle alternators.OPERATION: The plant has a larger water storage than Canyon Ferry and can also use stored water from Canyon Ferry storage. Separate transmission lines 13 miles long connect to lines from Canyon Ferry to Butte, and the Plant operates in parallel with the Canyon Ferry plant."

Jan 1, 1913

By Philip D. Wilson

SPEED is essentiaI in this war program and it is hard to keep up with developments. When the title of this paper was chosen, the contemplated magnesium production for which plants were then under construction was some 200,000 tons of metal per year or sixty times the 3350-ton production in prewar 1939. A few- days later an analysis of the President's 185,000-plane pro-gram, incendiary bomb schedules, and the additional requirements of the Army. Navy and Maritime Commission made it apparent that a much greater amount of magnesium must be provided in 1943. The new figures are huge but the censor will not per-mit their disclosure.

Jan 1, 1942

By Paul A. Flinn, Jean-Marie Robert Genin

The MCssbauer effect was used to study the distribution of carbon in Fe-C alloys of high carbon content during room-temperature aging. In freshly formed rnartensite the carbon distribution is approximately random; carbon clustering occurs dun'ng aging. In the retained austete the carbon atoms are isolated, suggesting an efiective repulsion; and no rearrangement during room-temperature aging is evident. The kinetics of the carbon clustering in martensite has been studied in detail at 0" . 10". and 20°C. The time and temperature dependence observed is sinzilar to that reported for a variety of other aging effects, such as resistance change, hardness increase, and recovery of internal friction. The stabilization of the martensite kransformation also appears to be related to the clustering effect. DESPITE extensive efforts since the early days of Reaumur, the mechanism which provides the strengthening of steel remains controversial in many cases for several reasons: 1) The structure of as-quenched martensite is not well-determined. 2) Experimental methods capable of distinguishing between various hardening mechanisms are not well-developed. 3) The temperature of formation of many low-alloy martensites is so high that the martensite formed during quenching may be already tempered to some extent before any observation. Among all the hardening phenomena, the most puzzling one occurs at room temperature, where no structural changes have been detected by the usual methods of X-ray diffraction or electron microscopy. However, a new and more sensitive technique became available with the discovery of the MBssbauer effect. This effect makes possible the detection of changes in the local environment of iron atoms in solid solution. It has been extensively used in the study of substitutional alloy systems, for example, Fe-Al,1 Fe-Si,2,3 Fe-Rh,4 and Fe-Ni.5 Making use of the possibility of obtaining new structural information, we undertook an investigation of the aging of high-carbon Fe-C martensite to follow the structural changes involved in the aging at room temperature. Preliminary results6 showed that a clustering of the carbon atoms occurs with time. A systematic study of the kinetics of the process showed a striking similarity to the kinetics of stabilization of the marten-sitic transformation, thus suggesting further investigation of the relationship. EXPERIMENTAL TECHNIQUE To eliminate difficulties with the aging of martensite prior to the first experimental measurements, we needed an alloy with an M, temperature not above room temperature. Other investigators7 were misled by results obtained on alloys which transformed above room temperature and underwent structural changes prior to the first Mossbauer measurements. Winchell and cohen8 dealt with this difficulty by using a third element, nickel, to depress the MS below room temperature. The presence of nickel did not seriously interfere with hardness or resistivity measurements, but it would complicate the interpretation of the Mossbauer data by increasing the variety of iron environments. We chose, therefore, to obtain the necessary lowering of the M, simply by increasing the carbon content. After the first quench to room temperature, our specimens are retained austenite, which we later quench into liquid nitrogen. The martensitic transformation takes place suddenly, so that the as-quenched condition of the samples is retained, since the temperature is low enough to prevent any carbon diffusion. A Massbauer spectrum at 77°K can be taken and analyzed; after an appropriate up-quenching to the aging temperature, a new Mossbauer measurement at 77°K can be made. Fig. 1 shows a typical cycle of experiments. The samples were foils 25 p thick of high-purity

Jan 1, 1969

By E. R. Gilbert

Klopp et al. have reported data on tensile and creep properties of are-melted and electron-beam-melted tungsten. We would like to point out some similarities between their creep results and ours on are-melted and powder-metallurgy tungstenand offer some alternate interpretations of the data which lead to different rate-controlling mechanisms. Klopp et al. observed extensive transient creep in the 2250" to 3500°F temperature range where the acthetivation energy is 106,400 ± 3300 cal mole-1. They suggest cross slip as the possible rate-controlling mechanism since "cross slip may occur only by the movement of screw dislocations, and not edge dislocations, thus allowing only partial recovery of the strain hardening and resulting in a continuously decreasing creep rate, characteristic of transient creep.'' Our tests on are-melted tungsten and powder-metallurgy tungsten displayed an extensive transient creep stage in the same temperature range but they also displayed a steady-state creep stage. The stress levels and strain rates of the two investigations, 1.4 to 7 kg mm-2 and 10-7 to 10-4 sec-1 for that of Klopp et al. and 4.5 to 10 kg mm-2 and 10-8 to 10-4 sec-1 for those of Gilbert et al., are overlapping. The activation energy for creep of our powder-metallurgy tungsten is 105,000 cal mole-1. This close agreement with the value found by Klopp et al. suggests that the same mechanism may be common to both types of materials. It may be possible to gain some additional insight into the rate-controlling mechanism by applying the analysis made by conrad22 in terms of the activation volume. The value for the activation volume depends upon the dislocation mechanism and may be used to help differentiate between various dislocation mechanisms. The activation volume for cross slip is 10 to 100 b3. Values determined from our data were 130 to 220 b3 which are in better agreement with the range of 102 to 104 b3 for nonconservative motion of jogs. Values determined from data of Klopp et al. are 240 to 1040 b3. We suggest that the rate-controlling mechanism for creep in this temperature range may be nonconservative motion of jogs assisted by enhanced diffusion along grain boundaries or dislocations. This model requires an activation energy lower than that for self-diffusion as is observed. If this is the controlling mechanism, a higher activation energy which is closer to that for self-diffusion might be expected for tests conducted on specimens with reduced grain boundary area, i.e., larger grain size. In support of this premise we have conducted tests on speci- mens having one to three grains per cross section for which a higher activation energy of 130,000 cal mole-1 was observed. Also, in agreement with results of Klopp et al., a faster creep rate resulted from the increased grain size and probable decreased dislocation density. Further argument against the cross slip mechanism is based on the constant activation energy we obsserved over the range of 5.0 to 10 kg mm-2,21 whereas the activation energy for cross slip displays a significant stress dependence.23 In addition to this, the powe on the stress term for the strain-rate dependence appears too high for cross slip. The values for tungsten fall in the range of 4.5 to 7 (Klopp et al. and Ref. 21) whereas values of approximately 2 are observed in creep by cross slip for hep metals.24 At temperatures from 3300° to 4000°F, Klopp et al. observed an activation energy of 141,000 ± 4000 cal mole-' which is less than that for self-diffusion, 153,000 cal mole-1, 25 and less than that reported for high-temperature creep of powder-metallurgy tungsten by Green:' 160,000 cal mole-1. Klopp et al. suggested that recovery of strain hardening by dislocation climb is the rate-controlling mechanism in this temperature region. Actually an activation energy larger than that for self-diffusion is predicted by this model if the temperature dependence of the elastic modulus27 is taken into consideration. Klopp et al. attributed this lack of correlation to the compositional differences between powder-metallurgy tungsten used by Green and their own arc- and EB-melted tungsten. It is our position that the value of 141,000 ± 4000 cal mole-' represents a transitional value between the higher and lower activation energies and does not represent a discrete mechanism of deformation. If the tests had been extended to higher temperatures the higher value of activation energy would have been observed as has been demonstrated by both our analy sis'' of isothermal test data by GE-NMPO28 and differential temperature tests28 on are-melted tungsten. The activation energies determined (170,000 and 160,000 cal mole-1) were greater than that for self-diffusion and are in better agreement with the value found by Green. The model of dislocation climb is associated with ; activation volume of 1 b3 according to Conrad.22 The activation volume calculated for the creep data of fou independent investigators in the temperature range where the activation energy for creep is near that for self-diffusion (Klopp et al. and Refs. 26, 28, and 30) is 110 to 2800 b3 and is in the range suggested for nonconservative motion of jogs. In the case of tungsten this mechanism may be preferred over that of dislocation climb. Sincere appreciation is expressed to Messrs. J. E. Flinn and F. L. Yaggee for the helpful discussions and suggestions in the preparation of this communication.

Jan 1, 1967

By Guido Bassi

IT is known'" that secondary recrystallization occurs in copper sheet with at least 90 pct reduction after annealing at high temperatures, 700" to 1000°C. Turkalo and Turnbull4 have found recently that secondary recrystallization might occur even after annealing at such low temperatures as 500°C for a high purity, coarse-grained copper sheet. It will be shown in the following that in an electrolytic tough pitch copper wire, with a high degree of deformation, secondary recrystallization might occur at even lower temperatures. A hot-rolled rod of 9.5 mm diam (temperature at which rolling commenced was 850°C) was drawn in continuous drawing machines to 0.4 mm, corresponding to a cross-section reduction of 99.8 pct. The texture in the center of the wire (the sample for X-ray analysis was etched to 0.15 mm diam) is a mixture of the [111] and the [loo] direction, as can be seen from Fig. 1. This is in accordance with the result of Schmid and Wassermann.' After a few minutes annealing of the wire at 200 °C, there is a noticeable recrystallization of the structure with a preferred orientation in the [l00] direction at the expense of the [Ill] deformation texture, as Figs. 2 and 3 show. Farnham and OINeill" reported a similar result. The final recrystallized structure shows mainly the [loo] orientation. There is however a small amount of [111] and random orientations. If the wire is annealed at 400°C, the recrystallized structure first shows a [loo] texture, Fig. 4, but as the annealing time is increased the fine-grained structure transforms into coarse grains, Fig. 5, giving a new preferred orientation, Fig. 6. The calculations show this to correspond to the [I121 direction. Even if the annealing that gives the primary recrystallization is carried out for six weeks at 20O°C, the same secondary recrystallization texture is obtained if the annealing is continued at 400°C. Schmid and Wassermann7 found the [I121 texture, but only after annealing at 1000°C or higher. The appearance of a secondary recrystallized

Jan 1, 1952

By Henry F. Adams, Guy H. Ruggles

In this paper the authors aim to chronicle the experience and salient points brought out in changing flotation reagents at a concentrator which had probably been using a minimum amount of oil at a minimum cost per ton, to a more complicated system requiring a partial chemical control. The original oils were coal tar and pine oil. The new reagents are lime, xanthate and pine oil. It was found that there were no insurmountable difficulties and eventually the cost of the new reagents per ton of ore proved to be very little more than that of the old. Lime Plant When the question of handling lime came up there were two propositions open: (1) to handle the lime dry; (2) to handle it wet in the form of an emulsion. Granting that a dry lime feeder would function properly, delivering lime uniformly, the chief objection was the dust that would be raised when these feeders were being charged or operated, as the only possible location was on a gallery beneath which men were working. With this objection the handling of lime as an emulsion seemed to be the better scheme. The lime plant which was first installed was for experimental work on one section. This was gradually changed and developed until it was of sufficient size to accommodate 20 sections. The various changes will not be discussed and the following description deals only with the fully developed plant. Crude as it is, and perhaps somewhat expensive to operate, the lime-handling plant has been quite successful. Much has been learned which will pave the way for a better installation in the future. Burnt lime is received loose in standard gage box cars. This is the product of the kilns and has had no screening or sorting. A grab sample is taken from each car as it is unloaded. An average of the analyses of samples from 44 cars gave the following: Available CaO Total CaO, Oxide and Hy- MgO, Fe²O³ + Al²O³, Insoluble, Per Cent. drate, Per Cent. Per Cent. Per Cent. Pek Cent. 88.3 79.2 1.6 ' 1.3 2.2

Jan 1, 1928

LIST OF MEMBERS AND ASSOCIATES ARRANGED ACCORDING TO STATES AND TOWNS ALABAMA Aldrich, Aldrich, W F Anniston, Noble, A E, Smith, N B Aibacoche, Bradley, D II Bessemer, Fergusson, V Birmingham, Aldrich, T H, Aldrich, T H, Jr, Bowron, J, Bowron, C E, Davis, F H, Dimmick, F D, Hillman, McCarthy, M E , McCormack, McCrery, Ramsay, E, Shook, A M Bridge¬port, Wheeler, S K Ensley, McDonald, S, Shook, J W Fanbanks, Hamil¬ton, A L Helena, Squire, J Huntsville, Moss Lehigh, Toulmin Oxmoor, Green¬amyer, A G Plasma, Gage Rock Run, Garvin Sheffileld, Miller, D I Tuscaloesa, Lodge, J University, McCalley Woodward, Bukofzer, Woodward, A H ALASKA Dawson, Buck Douglas Island, Corbus, J P , Macdonald, J Hames, Le Wald, E A Jualin, Hoggatt, W B Juneau, Denby, Gaifield, Garside, Hamp¬ton, W H Rodman Bay, Phillpots Snettisham, Tisdale Sumdum, Thane Tel¬ler, Hirshberg, M R Treadwell City, Batcheller, J H , Kinzie Valdez, Esterly, G M Yukon, Dawson, Elmer, J M ARIZONA Bisbee, Clark, G C, Coombs, H L , Douglas, W, Hoatson, Parnall, S A, Pelton, R T, Ricketts, L D, Sherman, G F G, Smith, F W, Smith, H A Black Diamond, Bagge Blanchard, Cummins, Murray, J J Casa Grande, Rich¬mond, A B Clifton, Colquhoun, Grabill, Murray, H T , Plater, Smith, J D A, Veitch, Watson, W S Congress, Goodloe Crown King, Alexander, G E Douglas, Cornell, Kraemer, E L, Lee, G B, Rutherford, F Globe, Berray, Hoar, Snell Harrisburg, Smith, F C Jerome, Clark, W A, Thomas, C A Kingman, Treadwell, Wallace, J P Kirkland, Bretherton, G K Kofa, Lamp¬shire, J O , Pomeroy, H G Mayer, Haggott, E A Metca j, Blumenkranz, M , Forbes, D L H Morenca, Clark, C D, Fleming, J B, Honey man, Mills, C E, Wallace, L R Morristown, Stevens, H W Nogales, Glove, H G, Sandoval, A , Sandoval, P Octave, Rose, G W Phoenix, Armstrong, T , Butler, C E, Defty, Lindsay, L Prescott, Blauvelt, H , Bunker, C E, Claypool, W M , Drescher, F M , Fairbank, Nissen Tombstone, Gage, Poindexter, C L , Staunton, Walker, E W Troy, Cutting Tucson, Blake, W P, Burgess, De Lashmutt, I, McCleary, Mann, Prout, J W , Jr ARKANSAS Fayetteville, Purdue Fort Smith, Bache, F , Denman, H Little Rock, Fordyce, Scholz CALIFORNIA Alameda, Arents, King, L Arnador City, Downs, Tregloan Angeles Camp, Dig¬gles, R N , Kenney, C E Antioch, Belshaw Auburn, Coe, I J Ballarat, Thorns Berkeley, Abadie,, Christy, Dozier, Hall, F F , Hersam, Landfield, Lind¬say, L , Mansfield, N T, Morley, W S , Sharwood, W J Bag Oak Flat, Dron Bodae, Hamilton, W R, Hoover, T, Hoover, T J, Jackson, B L , Nutter , E, H Bullion, Power, H T Callahan, Scott, E H Campo Seco, Baird Cartero,

Jan 1, 1910

By B. F. Oliver, F. Garofalo

Gas-metal heterogeneous reactions and zone-lrelting were sinultarneously employed to produce several high-purity irons with low interstitial contents in a levitating- zone melter. Successive zone-tneltitzg pusses were made in carious sequences of atmospheves which included prepuvified tank hy-drogeta, water vapor hydrogen, palladiuril-purified hydrogen, and vacuunm of 2 x 10-6 Torr. Different irons ranging in purity from 99.89 to 99.996 pct with residual curbon contents of 17, 8 to 10, and 3 to 5 ppm have been Drepared. No measuvable gradient in the residual carbon content was observed along the entire zone-melted length (6-10 in.) in I-ilz.-diameter bars. The absence of a carbon-concentration gradient in the solid indicates an apparent distvihution coefficient of 1, an unusual condition. Internal friction, solid-state extraction, and hot-hardness results indicate that the level of mobile carbon in a iron depends on the gas-metal heterogeneous reactions employed. Purification under an oxidizing atmosphere resulted in the immobilization of nearly all of the carbon present at leuels of 10 ppm or less. Continuation of purification in a reducing atmosphere veturued the carbon to a state of mobility which was aguin reversed by further purification in all oxidizing atmosphere. The immobilization of curbon at "vesidual levels " is believed to be associated with a strong interaction between carbon and oxygen atoms. frons with low carbon and nitrogen contetzts which did not exhibit a hot-hardness peak were produced after zone ynelting in vacuum. CARBON, nitrogen, and oxygen have strong and well-known effects on many properties of iron. These effects persist even at concentrations of 5 to 20 ppm. The intent of the present investigation was to establish procedures for the preparation of high-purity iron with concentrations of these elements in the range of 5 to 20 ppm and lower. The results reported here deal with the procedures employed to prepare high-purity iron, and with observations on the behavior of certain elements at "residual levels": particularly carbon. The details of the (levitating) zone melter in which the final purification steps are performed and the necessary control of the pertinent solidification variables have been described previously.1,2 MATERIALS AND PROCEDURES Commercial electrolytic iron was induction vacuum-melted and cast into 30-lb ingots. Melt additions and ingot analyses for carbon, oxygen, and nitrogen are noted in Table I. All ingots were melted in stabilized zirconia crucibles conditioned with one iron wash heat. Middle sections of the ingots were forged to 1-1/4-in. rounds, machined to 1-in.-diameter bars, then cut to 36-in. lengths. The zone-melting operations were performed on six iron bars (A to F) at a freezing velocity of 2 in. per hr with the freezing interface rotating2 at 200, 300, or 400 rpm. Liquid zones were passed through the bars in various sequences of atmospheres, Table 11, including prepurified tank H2, palladium-purified Hz, wet H2, and system vacuum of 2 x 1018 Torr. Wet-Hz atmospheres were obtained by bubbling palladium-purified Hz through water exhibiting a resistance of 3 megohms. This water was heated with a platinum-immersion heater and was degassed prior to use. Zone melting in a wet-HZ atmosphere resulted in oxidation of the elements which are oxi-dizable with respect to iron. Fine oxide particles were observed on the surface of the liquid zone, and were encased in the bar surface with the passage of the molten zone. Between each zone pass, the surface of the bar was removed by filing and deep etching; this was continued until the oxide formation ceased. The filings were magnetically separated and the residues analyzed by optical spectroscopy. These were especially high in silicon and manganese. Samples were removed from the iron bars, with few exceptions, following each zone-melting operation. A 0.08-in.-thick disc was cut with an Al2O3 cut-off wheel at a point 1/2 in. above the beginning of freezing. Material from each disc was used for hot-hardness measurements, resistivity-ratio measurements, and carbon, nitrogen, and oxygen analyses. Analyses for most other elements were

Jan 1, 1965

By Frances Hurd, Clark

Iv 1897, Roberts-Austen(l)$ found an arrest in the thermal curves of alloys of 60 per cent. copper and 40 per cent. zinc. This break occurred from 450° to 470° C. Shepherd, (2) working in 1904, was unable to detect any change in the direction of the thermal curve. In 1911, Carpenter and Edwards(3) confirmed the work of Roberts-Austen and called this a eutectoid transformation of beta to alpha plus gamma. Hudson') in 1914, showed that the beta constituent could be produced synthetically at temperatures below the transformation point. He believed the change in the beta was polymorphic in nature rather than a eutectoid. The work of Matsuda(32) in 1922 demonstrated this break by thermal analysis, microscopic investigations, dilatation, electrical resistance measurements, and a thermoelectric study. He believed the transformation to be continuous from 300° to 480° C., and comparable to the A 2 transformation in iron. Later, in 1924, Andrew and Hay, with electrical resistance measurements, confirmed Matsuda's work. Imai(28) showed the transformation to be progressive and therefore a function of the temperature rather than the time. He stated that it was neither a eutectoid nor an allotropic change of phase. Atomistically expressed, it was a change, not of atomic configuration, but of atomic energy of some kind; and again it was compared to the A 2 transformation in iron. Imai believed that his analyses by X-rays led to the same conclusion. In order to study this change, at 470'C., in homogeneous 60:40 brass, Homerberg and Williams «0> and Homerberg and Shaw(41) quenched their material in iced brine to obtain the beta solid solution and reheated at various temperatures. At 200° C., a change in the microstructure and physical properties occurred. A fine-grained structure, which they termed alpha and which had the appearance of a sorbitic steel, appeared, with an increase in the tensile strength to 82,750 lb. per sq. in., and a Brinell ball hardness of 240 with 3000 kg. load. At 470° C., the alpha

Jan 1, 1927

By E. N. Silverman

Mixtures of' MnS, ZFeO-SiO,, ZMnO.SiO, FeO, and MnO were prepared synthetically and used to study the planes ZFeO. Si0,-ZMnO.SiO,-MnS, FeO-MnO-MnS, and Fe0-MnO.Si0,-MnS in the quaternary system. Because the oxides and silicates studied flux MnS, inclusions in this system are partly liquid during rolling. Also, oxygen can modify subsurface inclusions through the reaction MnS + FeO—FeS + MnO during the soaking of ingots and the reheating of blooms and billets. NONMETALLIC inclusions in steel have long been a subject of study and conjecture because the properties of many steels are known to be affected by the the kind and amount of inclusions present. In free-machining steels, the type of inclusion present has a direct bearing on their machinability; in other steels, cleanliness is important and metallurgists direct their efforts at reducing the amount of non-metallic inclusions. Thus, the accumulation of knowledge about the formation and behavior of inclusions during all stages of steelmaking and processing should aid in improving the quality of steels and decreasing the production costs resulting from rejections. Previous studies of inclusions have contributed greatly to our understanding. The investigators have generally examined inclusions in finished steel. Boulger, Moorehead, and Garvey1 found that in Bessemer steels, low silicon contents results in globular sulfide inclusions with attendant high machin- ability ratings. Van vlack2 also found that low silicon contents favor inclusion morphology that contributes to high machinability ratings. Boquist and Stoll* found that silicon content affected sulfide inclusions and machinability of blooms, billets and bars more than did the finishing temperature in the range between 1775° and 2075°F. Carney and Rudolphy3 studied the distribution of sulfur, manganese, silicon, and carbon in a commercial ingot of resulfur-ized steel and found that silicon was concentrated near the bottom. This helped to explain why bottom cuts had poorer machinability. Klinger and Koch4 developed techniques for extracting inclusions from steel, and the technique devised by Gurry, Chris-takos, and stricker5 for extracting carbides can also be adapted to the study of inclusions. A different approach to the study of inclusions in steel consists of synthesizing the constituents that form inclusions and determining the phase eauilibria in the system that contains the constituents. The resulting phase diagrams can then be used to determine the effect of the various constituents upon each other. In this way, it becomes possible to under-

Jan 1, 1962

Secretary's Note.—M. Beutter's introductory remarks, being simply a resume of Prof. Hone's paper, with a disclaimer of the intention to criticise it, have been omitted, and the remainder of this contribution lias been somewhat condensed, without omitting any essential statements. M. Beutter, St. Etienne, France* (communication to the Secretary†):—Professor Home observes in $ 47 of his paper that the fluid-compression of steel diminishes piping by forcing the liquid and viscous metal, while the pipe is forming, or after it has formed, into the space so produced, and proceeds to study the different systems of compression, alluding to the four principal ones: Whitworth, Illingmorth, Williams, and Harmet. Having now for many years had in use the last-named process, that of M. Henri Harmet, I shall attempt to explain the difference between it and the three other processes named. The Harmet process does not diminish or close the pipe after it has formed, but prevents it from forming; and I may add that it is the only one of the four which produces this result. It is true that the Whitworth process diminishes the interior hollow by localizing it around the central axis of the ingot, which is afterwards removed by trepanning the hollow forg-ings for which this process has been, I believe, always exclusively used. The Illingworth and Williams processes, which, in my opinion, are applicable only to small ingots, of necessity allow the pipe to form, or (to use the expression of Professor Howe) they allow the surface of the lake of liquid steel to sink into the interior of the ingot, since they involve waiting a certain time before compression, so that (in the case of the Williams process) the ingot-mold can be opened without danger, or (in the case of the Illingworth process) so as to avoid the risk that the steel may enter the grooves produced by the removal of the distance-bars. During the time required for

Jan 1, 1908