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By Carl F. Lutz, Robert F. Mehl

A layer of brass was formed on 0 brass using a vapor-solid reaction technique. The variation in composition with distance within the phase layer and across the a -ß interface was determined by an electron probe. The diffusivities were calculated as a function of concentration; the diffusion coefficient in brass was found to change by a factor of twenty-five over a compositional range of 8 at. pct. An explanation is proposed to account for the large change in the diffusivity, based on possible stresses in the diffusion zone; it is suggested that the concentration dependence of the thermodynamic factor might decrease the activation energy with increasing zinc content. ALTHOUGH the body of data on diffusion coefficients (D- values) in terminal solid solutions and in isomorphous systems is quite large, there is but little quantitative information for intermediate alloy phases. This paper recounts measurements of D-values for the ? phase in the Cu-Zn system. The experimental method employed consists in the forming of alloy layers and determining and analyzing the concentration-distance (c-x) curve. The experimental method employed could have been applied to several systems; the Cu-Zn system was chosen because the D-values in the a and ß phases are well known, and can thus be used for purposes of comparison. The rates of growth, of intermediate alloy layers are known for a number of systems. In nearly all cases the thickness varies parabolically with time, as to be expected if the rate of growth is controlled by the D-values. In a few cases, nonparabolic behavior has been noted.1 Nonparabolic growth may be the result of a) a variation of the composition at the phase interface with time, b) interface-controlled kinetics, c) short diffusion times coupled with long vacancy lifetimes (in vacancy diffusion),2 and d) crystal reorientation during diffusion in anisotropic systems.3 In the Al-Ni system4 and in the Al-u5 system parabolic growth kinetics are slowly approached after an initial transient period. In general those phases stable at a given temperature in a system A-B appear as separate phase layers when A and B are put in contact and given a diffusion he at-treatment. In most cases the compositions at the interface of adjacent phase layers are those read from the phase diagrams for the termini of the respective phases. This discontinuity in concentration is taken as independent of time, and the growth of one phase at the expense of another is assumed to be independent of interface kinetics, i.e., the rate of interface movement is controlled by volume diffusion in the phases concerned. wagner6 has given many solutions for multi-phase diffusion processes, assuming that the chemical diffusion coefficient, D, is independent of concentration, as have others. Jost7 has pointed out that the familiar Boltzmann-Matano solution is as valid for a (c-x) curve exhibiting concentration discontinuities at phase boundaries as for the usual single-phase case. Heumann and associates8,9 have applied this solution to the multi-phase case, but lacking concentration data within the several layers were forced to assume linear concentration gradients, thus calculating only average D-values. The purposes of the present study were: 1) to determine the dependence of D on concentration, 2) to calculate the intrinsic diffusion coefficients D?Cu and D?Zn, 3) to establish the time-law for the movement of the interface, 4) to determine the concentration limits at the ß-? interface, and, 5) to attempt to clarify the mechanism of diffusion. EXPERIMENTAL PROCEDURE The ? phase is exceedingly brittle; conventional solid couples and conventional sectioning methods were thus inapplicable. Vapor-solid couples were accordingly employed. Unfortunately the equilibrium vapor pressures of Zn for the ? phase are unknown; to escape this awkwardness, samples were enclosed in a capsule containing powder of an alloy of known composition, sufficiently large in quantity as to be effectively an infinite source, maintaining the concentration of Zn at the sample surface constant with time. The sources of Zn-vapor which can be employed in reaction with Cu, or a brass, or ß brass, to form a layer of ? brass satisfactorily wide in concentration range, are alloys of the phases ?, or ? + E or ? + d (depending on the temperature).

Jan 1, 1962

By K. Tangri, M. Chaturvedi

On water-quenching from within the (a + ß) phase region Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys can undergo w transfirmation. This transformation has been attributed to the enrichment of ß Zr phase, at the solution-treatment temperature, by the solute atoms. The ß—w transformation is accompanied by a hardening 01 the alloys and in the optimum condition w-bearing specimens of the ternary alloy are about 14 pct harder than those bearing only a' phase. On aging crl reverls thermally to more stable phases, which results in a considerable softening of the specirnens, as follows: w quenched -waged + ß Zr (enriched) waged - a + ß Nb The high-temperature bcc ß Zr in certain zirconium alloys can transform to a metastable w phase which has a primitive hexagonal structure. In the Zr-Nb alloy system w phase has been observed in the composition range of 7 to 15 wt pct Nb by a number of investigators.'-' In this composition range w can form by water-quenching ß Zr or by aging. known as athermal and thermal w, respectively. Hatt and Roberts' have found that thermal w has a constant c/a ratio of 0.622 +0.002 regardless of the composition and tha: the lattice parameters are of the order of a = 5.02A and c = 3.0 ?. The actual lattice parameters depend upon composition and the temperature of aging. The lattice constants of athermal w. however, depend upon alloy composition.3 The presence of w phase in zirconium alloys can cause substantial strengthening which has been attributed as probably due to the coherency between w and parent ß Zr phase. Robinson et al.' found that the presence of w can raise the strength of Zr-5.0 Nb-2.0 Sn and Zr-5.0 Mo-2.0 Sn alloys to about 170.000 psi: however. the ductility is considerably reduced. With the growing use of Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys in the nuclear reactors, and the significant effect of w transformation on the mechanical properties of zirconium alloys studies were undertaken to study the w transformation in Zr-2.5 Nb and Zr-2.5 Nb-0.5 Cu alloys. This paper is concerned with the formation and reversion of w phase in these alloys. 1) MATERIALS AND EXPERIMENTAL TECHNIQUES The following two alloys were studied: 1) Zr-2.5 Nb-0.5 Cu (referred to as the ternary alloy): 2) Zr-2.5 Nb (referred to as the binary alloy). The chemical analysis of the alloys is given in Table I. The alloys, in the form of 3/8-in.-diam rods and 1/8-in.-thick sheets, were supplied by the Chalk River Nuclear Laboratories of the A.E.C.L. Initial fabrication for preparation of various specimens was carried out by cold rolling and swaging with intermediate anneals at 1000°C. All the heat treatments were carried out after the specimens were wrapped in zirconium foils and encapsulated in silica tubes under a vacuum of 5 x 10-% m of Hg. For optical metallography and hardness measurements disc specimens + in. diam by + in. thick were used. After the required heat treatments the specimens were mechanically and then chemically polished in a 45 pct HNO3, 45 pct H20. and 10 pct HF solution. The hardness measurements on the chemically polished specimens were carried out on a Vickers hardness tester using a 10-kg load. For each specimen at least fifteen indentations were made in order to obtain a representative hardness value. The phase identification and structural analysis were carried out using X-rays and electron diffraction techniques. Wires of 1.5 mm diam reduced to 0.12 mm diam by chemical etching were used for making Debye-Scherrer powder patterns using copper K-a radiations in 114.6-mm-diam camera. Thin films for transmission electron microscopy were prepared by electropolishing heat-treated a by 7 by 0.005-in.-thick strips using a modified Bollmann-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol electropolishing bath was kepl at -5OCC and polishing was done at 5 to 10 v. The thinned specimens were washed in ethyl alcohol at -30" to -40°C and dried between filter papers. The thin films

Jan 1, 1970

By A. J. Griest, P. D. Frost, H. A. Robinson

The flow characteristics and uniform elongation of alpha-beta titanium alloys in the solution-treated condition were shown to be markedly affected by the solution temperature. Two classes of alpha-beta alloys were distinguished, based on the effect of solution temperature on the uniform elongation. Martensitic alpha-beta alloys, which are relatively weakly beta-stabilized, showed decreasing values of uniform elongation as the solution temperature was initially increased. However, after solution-treatment at still higher temperatures in the alpha-beta field, the uniform elongation of these alloys was again high, and often attained its maximum value. Nonmartensitic alpha-beta alloys, which are more heavily beta-stabilized, showed a continuous decrease of uniform elongation with increasing solution temperature. The behavior was rationalized in terms of the alpha-beta ratio and the alloy content of the beta existing at the solution temperature. The data indicate that in commercial practice close temperature control may be required to get maximum uniform elongotion, and hence, maximum potential formability, in solution-treated sheet. Bend-test data also were obtained for the solution-treated conditions. A MOST important objective of current titanium research is the development of age-hardenable alloys with superior formability in the solution -treated condition. Alpha-beta titanium alloys have been shown to be amenable to high-strength aging treatments after solution-treatment in the alpha-beta field.' However, heat-treatment data on the properties of solution-treated sheet, such as uniform elongation, minimum bend radius, and flow properties, which have a bearing on the potential formability, have not been widely available. This paper presents only the most significant results of a research program conducted at Battelle for the Air Force in an effort to develop sheet alloys that can be formed readily as solution-treated and subsequently aged to very high strengths. Data for 10 compositions representative of the more than 25 alloys studied in this program are presented. A primary objective of the paper is to demonstrate the pronounced dependence of the flow characteristics and uniform elongation on solution-treatment temperature. This effect is believed to be vitally important in the commercial heat-treatment of age-hardenable titanium sheet alloys. Experimental Procedures Melting and Fabrication—The nominal and actual compositions, beta transi temperatures, and rolling temperatures of 10 of the 30 alloys studied are presented in Table I. Stock :€or two of the alloys, E (Ti- 8Mn) and J (Ti-6A1-4V) was obtained from commercial heats. The remaining heats, except D and I, were double melted as 10-lb heats in a consumable-electrode furnace. Melting stock was 120 Brine11 hardness sponge. The oxygen and hydrogen analyses spotted throughout the heats indicate the range of interstitial levels resulting from the melting practice used. Heats D and I were melted as 20-lb ingots with the same grade of sponge. These ingots were cut in half and one portion used in the present evaluation. The ingots, which were about 4 in. diam and 4 in. in height, were upset forged and drawn at 1750°F to sheet bar 5/8 in. thick by 2 1/2 in. wide in such a manner that the top of the ingot corresponded to the top of the sheet bar. Slabs of the sheet bar, 5 1/2 in. in length were scalped to 1/2 in. thickness and rolled at the temperatures indicated in Table I to approximately 0.055 in. sheet. Elongation during rolling took place at 90" to the original forged length. The fabrication procedure used is described in some detail since the experimental heats showed considerable directionality in mechanical properties. The rolling procedures for the commercial heats were unknown, but were such as to produce essentially isotropic properties. All heats, as rolled, showed microstructures typical of alpha-beta alloys worked in the alpha-beta field; that is, the structures consisted of equiaxed alpha particles in a beta matrix. Heat Treatment— The effects of solution and aging heat-treatments on the properties of the alloys were evaluated uSing longitudinal and transverse tensile and bend coupons sheared from the sheet. Two or three solution-treated conditions were evaluated for each alloy. In addition, the alloys were also tested after annealing at 1200°F for 4 hr and furnace cooling. The solution-treated and annealed conditions were those in which the alloys might be expected to

Jan 1, 1959

By R. T. O’Brien, J. M. W. Mackenzie

A microelectrophoresis technique has been used to measure the zeta potential of quartz over a range of pH and Ni (11) and Co (Ilj concentrations. Results have been discussed in terms of adsorption of Ni (11) and Co (11) cations and hydroxide precipitates. Possible relationships between the experimental results and the formation of lateritic nickel deposits have been discussed. The adsorption of metallic cations to the surface of silicate minerals is of significance in many industrial and natural processes. In an earlier publication,' the adsorption of Fe (111) to the quartz surface, a phenomenon of importance in flotation, was studied qualitatively using microelectrophoresis techniques. This paper reports the extension of this work to an investigation of the adsorption from aqueous solution of Co (11) and Ni (11) to the quartz surface. The adsorption of these metals was of interest for two reasons; firstly, the importance of such an adsorption process in the formation of nickel deposits by the lateritization of nickeliferous rocks, and secondly, the comparison of electrophoresis data for these metals, whose hydroxides precipitate in the pH range 7-8, with similar data for Fe (111) where the hydroxide first precipitates at pH 2.8. The development of laterites is essentially a near-surface geological process involving vertical and, in part, horizontal redistribution of elements derived from rock-forming minerals. This redistribution of elements leads to the formation of layers within the lateritic profile which are enriched in one or more of these elements and, depending on the nature of the original rocks, continued lateritization may give rise to iron ores, bauxites or nickeliferous laterites. The mechanisms discussed in this paper deal with the last mentioned type of deposit although the considered factors are applicable to laterites in general. Nickeliferous laterites are derived from basic and ultrabasic rocks containing ferro-magnesian silicates. These rocks may be partly or wholly converted to serpentinites. Besides their high Mg, Si and Fe contents, they also contain small amounts of Mn, Al, Cr, Ni, Co, Cu and Zn. The nickel content is usually near to 0.2-0.3%, but for economic use this must be increased during lateritization to over 1% and preferably over 2%. The zone of nickel enrichment is usually at some depth within the lateritic profile. Unpublished work by Mr. O'Brien on nickel-bearing laterites in eastern Australia indicated four visually different types of material which were separated from crushed samples under a stereo-microscope. Although chemical analyses indicated that all four types of fragments contained 2-4% nickel and smaller amounts of cobalt, X-ray diffraction and optical studies revealed only quartz, hematite and chromite as well as, in one of the types, a serpentine mineral. No distinct, nickeliferous minerals were detected. Furthermore, it was possible to extract some 60% of the nickel content of the laterites with cold 0.5N hydrochloric acid. It was therefore postulated that much of the contained nickel was present as amorphous, adsorbed material. In considering this fact, as well as the problem of the actual method of transfer and fixation of elements within the lateritic profile, it was further postulated that the transfer of these elements probably took place partly as solutions and partly as colloids and that the fixation of these was caused by adsorption and flocculation. The existence of nickel in solution in water draining from serpentinites has been confirmed by Smurov.2 Although we have no evidence for the transportation of nickel in the colloidal state, it is well accepted that colloids are important in the formation of laterites. Previous work by Mr. Mackenzie' has already indicated that iron may be adsorbed on quartz under conditions which could be expected in a lateritic profile, but no information was available for nickel or cobalt. Although adsorption possibly also takes place on minerals other than quartz, the latter is significant because silica is an important constituent of the silicate parent rocks. During lateritization, it is released to form free quartz and it is likely that it is initially extremely fine-grained. This would provide abundant opportunities for adsorption and it was de-

Jan 1, 1970

By T. E. Leontis

Data are presented in this paper to show that addition of zirconium to sand-cast Mg-Th alloys effects a marked decrease in the grain size of these alloys which is accompanied by a significant increase in the mechanical properties over the entire range of thorium content investigated. The beneficial effect of zirconium on the strength properties is maintained at elevated temperatures up to 700°F. In addition the alloys exhibit exceptionally high creep resistance at temperatures up to 600°F. Zirconium does not greatly improve the strength or creep characteristics of extruded Mg-'Th alloys. Cerium is not a desirable addition to Mg-Th-Zr alloys. IN a previous paper,' it was shown that addition of thorium to magnesium imparts exceptionally high resistance to creep at elevated temperatures both in the sand-cast and in the extruded states. The large grain size of the sand-cast alloys, however, renders them of questionable value as engineering materials. Sauerwald² showed that the grain size can be reduced effectively by the addition of zirconium. Results of a detailed study of the properties of Mg-Th-Zr alloys over a wide range of thorium content are presented in this paper. In accordance with the preliminary findings of Sauerwald,² it is shown here that Mg-Th-Zr alloys possess the unusual combination of fine grain size and a high level of strength and creep resistance at elevated temperatures. Some information on the effect of cerium on the properties of Mg-Th-Zr alloys is presented also. Effect of Th on Solubility of Zr in Molten Mg Sand-cast test bars and 3 in. diam extrusion billets of Mg-Th-Zr and Mg-Th-Ce-Zr alloys* were pre- pared according to the procedures described in the previous paper.' The zirconium was added in the form of sponge metal as supplied by the Bureau of Mines.' The maximum amount of zirconium that usually can be introduced into magnesium at the customary alloying temperature of 1350° F is about 0.7 pct. The analyses in Table I clearly show that thorium in- creases the solubility of zirconium in liquid magnesium. The tabulated zirconium contents are designated as "acid soluble;" this serves to distinguish these values from the small amount of zirconium that is not dissolved when the alloy is subjected to the action of a dilute acid solution. The "acid-insoluble" zirconium is believed to be ineffective as an alloying ingredient. The acid-insoluble zirconium contents ranged between 0.03 and 0.16 pct in the present alloys, the tendency was toward higher values in alloys with higher thorium content. In contrast to the effect of thorium, cerium decreases the solubility of zirconium in magnesium. This is shown by the three Mg-Th-Ce-Zr alloyst in- cluded in Table I and stands in agreement with the observations of Nelson and Stricter".' on Mg-misch-metal-Zr alloys. Effect of Zr on Grain Size of Mg-Th Alloys The potent effect of zirconium as a grain refiner has been demonstrated for many types of magnesium alloys.3-8 Of greatest importance to date are Mg-Zn-Zr alloys5,6,8 and Mg-mischmetal-Zr alloys.3,4,7 Zirconium has also been found to exert a marked grain-refining action on Mg-Th alloys. All the Mg-Th-Zr alloys cast in the form of ½ in. diam test bars have an exceedingly fine grain size (Table I). Reference to the previous paper' will show that binary Mg-Th alloys containing up to 6 pct Th are almost completely columnar even in the M in. diam reduced section of a test bar, and that 10 to 50 pct Th alloys have an equiaxed grain structure of 0.03 to 0.08 in. The data in Table I also show that addition of cerium to Mg-Th-Zr alloys has no effect on the grain size. In addition to increasing the strength properties, the fine grain size of Mg-Th-Zr alloys

Jan 1, 1953

By A. F. Nylander, J. H. Jensen

Magnesium, in its combined forms, is the sixth most abundant element and the third most abundant metal in the earth's crust, but it is so reactive that it is never found in nature in the elemental state. As part of the rock-forming minerals like magnesite, dolomite, olivine, and serpentine, it is combined in virtually insoluble forms. As part of soluble minerals like kainite (KC1 MgS04 - 3H20), leonite (K2S0, MgSO, 4H20), langbeinite (K,SO, 2MgS04), and carnallite (KC1 . MgCl? 6H20), it is found in evaporite deposits, generally of marine origin, deep within the earth's crust. Since these minerals are soluble, they are found only in deposits protected from the percolation and leaching action of ground waters by an insoluble and impermeable mantle. Usually this mantle is from several hundred to many thousands of feet thick. Finally, as a solute in marine-like brines, magnesium is common in surface and subterranean waters throughout the world. The oceans, which contain 0.13% Mg, illustrate the inexhaustible nature of the reserve. The recovery of magnesium from the rock-forming minerals, usually by thermal decomposition, and from dilute brines such as sea water by precipitation with lime, are frequent industrial practices. But the recovery of magnesium as magnesium chloride from evaporites or natural brines of concentration greater than sea water has been of little industrial significance for economic and technological reasons. The present US capacity is preponderantly based on the sea-water magnesia process with thermal reduction capacity in second place. No present magnesium capacity is dependent on recovery from the stronger natural brines. But projected capacity, although still weighted in favor of the sea-water magnesia process, is partly based on the recovery of magnesium chloride as a byproduct of potash operations.

Jan 11, 1964

By R. A. Bowen

THE Montreal mine, operated by the Montreal Mining Co., is four miles west of Ironwood, Mich., in Iron County, Wisconsin. It is the westernmost producing property on the Gogebic Iron Range of the Lake Superior district. Its production is 4000 tons of hematite per day from an average depth of 2300 ft. This tonnage is hoisted through a five-compartment vertical shaft. Several mining methods are in use but in the larger orebodies the method is sublcvel caving. ORE OCCURRENCE AND CHARATERISTICS The main orebodies of the Montreal mine lie in eastward-pitching troughs formed by the intersections of intrusive dikes with impervious members of the iron formation. The formation has an average width of 425 ft., strikes east and west and dips north at approximately 62°. The dikes dip south at 45° and bear north 20° from the strike of the formation, making the intersections with various members pitch east at 16°. The dikes are not continuous across the formation but are interrupted by a longitudinal bedding fault, which has a throw of 380 ft. vertically and 885 ft. horizontally. North of the fault zone the formation is called "hanging," south it is called " footwall" formation. The faulting has created two ore-bearing horizons; viz., at the inter- section of hanging dikes with the slate members and the foot dikes with the quartzite footwall. In each formation the main ore-bearing member consists of wavy beds of hard chert interlaid with hematite. In the most productive part of the mine, where the sublevel caving method is employed, a hanging-formation dike lies in juxtaposition with a footwall dike and the iron formation carries ore in certain sections from the quartzite footwall to the hanging dike. Here a horizontal section of the orebody may be 300 to 400 ft. wide for a length of 150 ft. or more and tail out lengthwise in both the foot and hanging formations. The footwall ore section is much longer than the hanging wall because of the orebody's greater height above the "pitching dike."

Jan 1, 1938

By W. O. Milligan, J. L. McAtee

THE constitution and structure of the hydrous forms of aluminum oxides and hydroxides have been studied in the laboratory of the Depart¬ment of Chemistry at Rice Institute for almost three decades. Progress reports have been made from time to time 13,17,18 summarizing and evaluating the existing state of knowledge concerning the complex inorganic system A12Oa-H2O. X-ray and electron diffraction analysis, together with dehydra¬tion isobars and isotherms, have been the most useful methods for increasing the understanding of the structure and interrelationships of the various forms of alumina and aluminum hydroxides. The purpose of this paper is to bring up to date the information now available, in the hope that a clearer understanding of the structures of these materials may prove useful in aiding mineralogists and geologists in studies of the genetic significance of the aluminum hydroxide minerals, and possibly of the closely related clay minerals. Early investigators found the system alumina-water to be confusing, since, prior to X-ray diffraction examination, it was not possible to clearly identify and characterize the various crystalline forms. One could not be certain as to whether or not different investigators were working with the same chemical individuals. In later studies, including some from the authors' laboratory, addi¬tional confusion resulted from difficulty in preparing pure samples of the various aluminum hydroxides and in obtaining relatively large crystals of dimensions above the colloidal range of particle size. Even at present, it is difficult to establish sets of interplanar spacings for the several forms of the aluminum hydroxides, as illustrated by the 29 different X-ray diffraction pat¬terns listed in a current publication. The nomenclature of the aluminum hydroxides is also extremely confus¬ing, inasmuch as two different systems are employed by large groups of in¬vestigators. Plans are being made to formulate a uniform system of nomen¬clature but in this report the designations employed will agree with previous papers from the same source.13,17,18 Four forms of crystalline aluminum hydroxides may be considered to be well defined and recognized, which, together with hydrous alumina "amor¬phous to X-rays," comprise an exceedingly complex inorganic system. Several

Jan 1, 1952

By Paul A. Farrar, Sanford Adler

The Tz-Zr system was reinvestigated using both metallographic and X-ray diffraction techniques. It mas found that titanium and zirconium are soluble in all proportions in both the a and 0 phases. The minimum in the transformation was found at 50 at. pct and 535°C. THE Ti-Zr system has previously been the subject of several investigations.1-6 A study by Hayes et al.5 using magnesium-reduced titanium and zirconium melted in graphite crucibles indicated that the system was a continuous series of solid solutions in both a and 0 phases with a minimum at approximately 545°C and 65 pct* Zr. These conclusions were substantiated by the work of Duwez4 as well as by the limited data of Craighead et al.,3 Fast,2 and deBoer.1 However, a more recent investigation of the titanium-rich region by Ence and Margolin6 indicated that the solubility of zirconium in a, titanium at 500°C is approximately 22 pct with the a + 0 field extending to approximately 47 pct Zr. Therefore in order to resolve this discrepancy the following investigation was initiated. EXPERIMENTAL PROCEDURES The alloys used in this investigation were prepared as 20- to 30-g buttons by nonconsumable electrode melting in a helium atmosphere. The starting materials were Bureau of Mines titanium BHN 71 and Foot Mineral iodide crystal bar zirconium. The titanium was premelted before the alloys were prepared to avoid excessive weight losses in the final melting. The compositions of the alloys prepared were as follows: Ti-12.0, 22.2, 29.0, 39.6, 50.0, 58.0, 72.4, 83.5. and 90.8 pct Zr. Prior to heat treatment the alloys were cold-worked 10 to 20 pct, stopping at the first signs of cracking. Specimens for heat treatment were wrapped in molybdenum sheet and annealed in argon-filled quartz capsules for the following times and temperatures: 750°C-40 days, 675°C-61 days, 600°C—109 days, 575°C-112 days, 550°C-112 days, 525°C—127 days, and 450°C —153 days. Quenching was accomplished by breaking the capsules in an iced-brine solution. The standard techniques used for polishing the specimens involved belt grinding, grinding on emery paper, polishing electrolytically, and etching with Remington "A" etch7 or "R" etch 8 Debye-Scherrer X-ray powder photograms were obtained for a number of samples using a 114.6-mm camera and CuK, radiation. Exposure times were from 19 to 23 hr. The powder samples were obtained by filing to 270 mesh size. After filing the samples were wrapped in molybdenum foil capsuled in quartz and reannealed at the temperature of original heat treatments; following the heat treatment the samples were quenched into iced brine without breaking the capsules. RESULTS AND CONCLUSIONS On the basis of the microstructural examinations of the heat-treated samples and of the X-ray diffraction data, the Ti-Zr system shown in Fig. 1 was constructed. The data indicate that a, and 0 titanium form a complete series of solid solutions with a, and ß zirconium. The minimum was found to occur at 65 pct (50 at. pct) Zr and 535°C, in excellent agreement with the earlier investigations,1-5 see Fig. 2. Typical microstructures are shown in Figs. 3 to 7. Figs. 3, 4, and 5 show the 22.2, 39.6, and 50.0 pct Zr alloys after heat treatment at 675°C for 61 days with the 22.2 pct Zr showing equiaxed a, and the 39.6 pct Zr alloy a + transformed ß, while the 50 pct Zr alloy shows only transformed B. Figs. 6 and 7 show the 29.0 and 39.6 pct Zr alloys after heat treatment for 109 days at 600°C. The 29 pct Zr alloy has an equiaxed a structure while the 39.6 pct alloy has a partially recrystallized a+ß structure. X-ray diffraction data obtained from the 50 pct Zr alloys after heat treatment at 525° and 450°C as well as from the as-cold-rolled material showed only those lineos which couid be indexed a: a titanium, good agreement with the previously published values. 2,4,5,9 Chemical analysis* of the 50 pct Zr alloy

Jan 1, 1967

By N. J. Grant, D. Carne, J. B. Seabrook

IT is unnecessary to review much of the literature on hydrogen embrittlement of steel since several excellent reviews and bibliographies exist.1-3 Hot acid pickling and cathodic charging have been known to cause hydrogen embrittlement of steel and have been used as laboratory methods to charge steels with hydrogen.4-8 Experimental Procedure The problem of prime concern in all instances, except in the work of Sims and coworkers,10 is that at no time was there a definite known hydrogen value reported which could be related to a particular value of the degree of embrittlement. It was the purpose of this research to determine the relationship between hydrogen content and degree of em-brittlement quantitatively. The 1020 stock used for the specimens was hot-rolled 7/16 in. diam rod of the following analysis: 0.19 pct C, 0.40 pct Mn, 0.012 pct P, 0.036 pct S' Vacuum-fusion analyses for nitrogen and oxygen gave 0.014 pct O2 and 0.0011 pct N2. Cathodic charging was used to cause embrittle-ment. The advantages of this method were that no quenching was necessary, the impregnation was reasonably controllable, and representative samples for analysis could be prepared simultaneously with the charging of the tensile bars. Analyses for hydrogen were made in a vacuum-fusion type of apparatus which extracted gas from the samples by fusion in a molten tin-iron bath' This apparatus, which has recently been described in detail elsewhere, will not be discussed here.". " Consecutive samples can be analyzed every 15 min with an accuracy of 0'03 ppm. A A from an electrolytic or pickling bath can be charged for analysis in about 3 min, and analysis completed in 15 min. Samples of as-received 1020 stock were analyzed and found to contain a negligible amount of hydrogen (0.06 ppm). Preliminary pickling and electrolytic charging gave highly variable results. Cold-worked samples showed higher values but ranged from 0.6 to 6.0 ppm of hydrogen. Charging at 90°C gave higher values of charged hydrogen, but troublesome films on many of the specimens resulted. It was not until a few drops of a catalytic poison, prepared from 2 g of yellow phosphorus dissolved in 40 ml of carbon disulphide,8 was added to the electrolytic cell that fairly consistent high values of hydrogen could be charged into the as-received 1020 steel at room temperature. The cell used is sketched in fig. 1. The anode and fixture for the tensile specimens were fastened to the cell cover as shown, and the specimen was screwed in place, centrally located with respect to the anode. At first, anodes of lead and graphite were tried, but they contaminated the electrolyte and gave a slight deposit on the specimen. Later, a platinum wire was used in the form of a uniform helix. A current of 2 amp, corresponding to about 0.5 amp per sq in. of cathode, was used for most specimens. The electrolyte was 4 pct H2SO4 solution. Fig. 1 shows the form of the special tensile bar, the main portion of which was an ordinary 1/4 in. diam, 1 in. gauge length test bar, except that two to four small cylinders of stock were left attached by small links to the bottom end of the tensile bar for the purpose of providing samples for hydrogen analysis. Thus the bars and samples were charged with hydrogen simultaneously and an analysis could be made while the bar was being tested mechanically. Experiments were performed to determine how representative the test cylinders were of the middle of the test bar proper. While the bar center was being analyzed, the small cylinder was stored in dry ice, since it had been shown that such storage prevents loss of hydrogen for many hours.12 The results of these checks are shown in table 1. The checks are very good (0.3 ppm) and indicate that it is permissible to use the small end cylinders to establish the hydrogen content of the tensile test bar proper-

Jan 1, 1951

By A. R. Troiano, A. B. Greninger

There is need for an adequate working hypothesis that would describe at least qualitatively the crystallographic mechanism for the transformation from austenite to martensite in steel. A general theory would not only be of great assistance in the correlation of existing data and in planning future experimentation, but would also form the basis for an eventual explanation of the property changes that accompany lattice transformations. The mechanism of martensite formation described in the following pages is the outcome of two new experimental determinations: (1) the accurate evaluation of the lattice relationship between austenite and individual crystals of martensite, and (2) the measurement and analysis of the change in position that a volume of austenite undergoes when it transforms into a crystal of martensite. For this latter study, the stereographic analysis of shear was employed; this method greatly simplifies the solution of lattice-shear problems, and some space is devoted to its description. Lattiice Relationships and Habits PREVIOUS WORK The lattice relationships between austenite and martensite in carbon steels have been studied by Kurdjurnow-and Sachs,' and by Wassermann;2 and in iron-nickel alloys by Nishiyama.3 Wassermann,2 and Mehl and Derge.4 The studies by Kurdjumow and Sachs and by Wassermann were made on single (austenite) crystals of quenched 1.4 pct carbon steel, and the results of these two independent studies agree closely with the postulated lattice relationships: (III)r||(0ll)a [101]r [111]a [1] Both Nishiyama and Wassermann worked with an alloy of iron plus 30 pct nickel (all austenite at room temperature), and "martensitic alpha" was formed by cooling in liquid air. They agree that the lattice relationships for these conditions are: (lll)r||(011)a [112]r[011]a [2] Mehl and Derge worked with a range of iron-nickel alloy compositions centered about 30 pct nickel. They concluded that when the transformation takes place near room temperature or above, relationship [l] holds, but if the same specimen is made to transform at -195°C, relationship [2] obtains;intermediate temperatures evidently produced a combination of these two relationships. Sachs and co-worker~,1,7 as well as Nishiyama,3 have proposed transformation mechanisms for martensite formation.* Both of the mechanisms picture the transformation as initiated by shearing movements along the octahedral [Ill] plane of austenite, followed by suitable expansion and contraction to complete the transformation. These mechanisms have been described and commented upon by Mehl, Barrett, and Smith,8 Mehl and Derge,4 Wassermann,2 and others, and generally accepted by investigators in this field. Greninger and Troiano9 have recently published the results of a detailed study of the shapes and orientation habits of martensite crystals in plain-carbon and nickel steels. They found that, contrary to previous assumptions, martensite plates form not along the octahedral planes but along austenite planes of very high indices; these results were in line with those of previous studies of martensite reactions in nonferrous alloy systems by Greninger and coworkers.6,10 Greninger and Troiano concluded that the transformation theories of Kurdjumow and Sachsl and of Nishiyama3 are untenable. Phragmén11 reached a similar conclusion. EXPERlMENTAL METHODS The object of this particular study was the evaluation of (1) lattice relationship between austenite and martensite, and (2) relationship between the martensite lattice and the martensite plate. The technique used for solving these problems was similar to that used by Greninger6 on martensitic structures in 8 copper-aluminum alloys. Briefly, the method consists of grinding and polishing a specimen on a surface parallel to an individual martensite plate and thus exposing a single martensite crystal for an area of l to 3 mm. The orientation of this exposed martensite crystal is then determined by means of a back-reflection Laue X ray pattern.12 Another pattern is obtained from the matrix crystal, and a stereo-graphic plot of data from these two patterns then provides the solution to the problem. It is necessary to repeat this process on several crystals in order to determine whether or not the solution is unique.

Jan 1, 1950

By M. Cohen, D. Caplan

The scaling characteristics of three Fe-Cr alloys have been investigated by determining their weight gain vs. time curves at 1600° to 2000° F. The scales formed thereby have been examined using the techniques of X-ray diffraction and spectrographic and metal-lographic analyses in an attempt to explain the discontinuities in the curves and to elucidate the mechanism of scaling. DESPITE the considerable number of investigations that have been carried out on heat resistant alloys, the characteristics of the scales formed at high temperatures are not fully known. The research reported here was undertaken in an attempt to ascertain the mechanism of scaling of the stainless steels. Scaling experiments were carried out first, the weight increase of the specimens being followed continuously with time. It was observed that, as well as showing the expected decrease in oxidation rate with time, the oxidation curves showed breaks corresponding to intermediate periods of accelerated oxidation, after which protectiveness again increased. This phenomenon was observed with austenitic stainless steels (types 302, 309, and 330) and with Fe-Cr alloys (types 410, 430, and 446), but only the latter are treated in this report. An examination of the scales was made using the techniques of X-ray diffraction and spectrographic and metallographic analyses in an attempt to obtain a correlation between the nature of the scales and the oxidation curves. A search through the literature revealed only a very few previous reports of such periods of accelerated oxidation. Dunn' found breaks in the oxidation-time curves of some Cu-Si alloys but saw no rational explanation of the phenomenon. Heindlhofer and Larsen2 attributed a discontinuity in the weight gain-time curve of iron at 1290°F to the formation of blisters, the subsequent cracking of which exposed an unprotected surface and permitted rapid oxidation until a new protective scale had been reestablished. They advanced no explanation, however, for what they termed the peculiar behavior of a 27 pct Fe-Cr alloy at 2000°F which gained weight very rapidly in between two periods of very slow weight gain. Portevin, Pretet, and Jolivet3 in describing breaks in the weight gain-time curves of Fe-A1 alloys suggested that they might be associated with the occurrence of localized and deeply oxidized areas on the specimens. Bandel4 in a general discussion of oxidation curves of heat resistant alloys considered that the discontinuities were due to a local disruption of the protective layer by the growth of iron-rich oxides. Day and Smith" in their report on the scaling of a large number of iron alloys noted but did not explain occasional relatively rapid changes in oxidation rate at higher temperatures. Chevenard and Wache6 found breaks, often two per specimen, in the oxidation curves of an 18-8 type alloy. They suggested that the cause might be a depletion in chromium of the surface layer of metal due to its selective oxidation, the resultant high concentration of iron and nickel in the scale leading to a poorly protective scale. McCullough, Fontana, and Beck' explained the breaks in the oxidation curves of types 304, 430, and 410 alloys as due to mechanical ruptures. Experimental Work Table I lists the chemical compositions of the materials used. Cylindrical specimens 1/4 in. in diam and 11/2 in. long were machined from cold rolled % in. rod. After a fine finish cut with a sharp tool, the specimens were abraded while still mounted on the lathe with Nos. 2, 1, 0, and 00 metallographic grade emery papers. A 3/64 in. hole was drilled at a distance of 1/8 in. from one end to permit suspension in the furnace. Specimen Nos. 1, 2, and 3 were tested with this surface preparation. All others, after being similarly prepared, were electropolished in a perchloric-acetic electrolyte, electrical contact being made by pressing a tapered platinum hook into the drilled hole. The specimens were then washed in hot water, rinsed with distilled water, rinsed with methanol, dried at 120°F, and weighed. Thereafter,

Jan 1, 1953

By S. J. M. van Leeuwen, R. Feenstra

The effect of jets on bit penetration has been investigated by means of a 50-ton drilling machine and 8½-in. commercial jet bits, drilling under representative bottom-hole conditions. The conclusions apply to the drilling of impermeable rock only. Penetration rate is in all cases primarily hampered by a differential pressure effect, known as dynamic hold-down. Jet action can reduce this slightly. Major gain in penetration rate, up to 25 or 50 per cent, is already ohtained at medium-high jet power. At high bit load, penetration by soft to medium-hard-formation bits is further hampered by bit-balling. Its alleviation requires intensive tooth scavenging, which is best performed by slanted nozzles with trailing, high-velocity, high-volume jets, hitting bottom in front of the teeth. Penetration may then improve by a 100 per cent. At high bit load, bottom-balling limits the penetration of insert-type bits. In this case bottom scavenging is required, which is most effectively performed with high velocity jets from nozzles extended close to bottom. As a result of this measure, penetration may improve 70 per cent in non-friable rock, and less in friable rock. This will make the insert-type bit more versatile. Toothed hard-formation hits suffer at high bit load from both bit- and bottom-balling. Slanted nozzles, possibly much extended, with high-velocity, high-volume jets then have best prospects. INTRODUCTION For years the use of jet bits in oil well drilling has been increasing, in conjunction with ever greater hydraulic horsepower. To the same extent, the need for efficient use of the horsepower has grown.' This paper, it is hoped, will contribute to an understanding of how effective jet action may be obtained. It deals with the effect of jets on phenomena that are known' " to hamper bit penetration. To study this effect, laboratory drilling tests were performed with 8½-in. commercial bits under bottom-hole pressures representative of deep wells. The investigation has been restricted to the drilling of practically impermeable rock. In the following sections, the increase in bit penetration rates due to jetting is discussed for conditions which made the troublesome phenomena stand out one at a time. A brief definition and interpretation of these phenomena is given in Appendix A. Appendix B gives the results of measurements of fluid velocities along the hole bottom, performed in stationary flow tests. Descriptions of equipment, bits, rocks and mud are collected in Appendix C. Unless otherwise mentioned, drilling conditions were kept standard as given in Table 1. CUTTING TRANSPORT TO THE ANNULUS Jet action was for some time believed to merely improve the transport of cuttings from bottom to annulus, resulting in a clean hole bottom. Flow tests performed in 1954 at the Battelle Memorial Institute6 showed, however, that jet velocities of the order of 40 ft/sec sufficed to clear even large (0.4-in.) chips from bottom in the time between two successive tooth actions. In the field, optimum jet velocities are appreciably higher, indicating that other effects must be involved. This much is confirmed by the results of our drilling experiments in soft rocks performed under atmospheric pressure, i.e., in absence of pressure-differential effects. Fig. 1 shows that variation in the nozzle velocity from 40 to 330 ft/sec had no effect on penetration rate, unless the bit penetrated more than 7 mm per revolution, which is about half the tooth height. At such high rates—150 ft/hr at 100 rpm—cuttings would be caught in the tooth cavities, leading to bit-balling, which was noticed on inspection of the bit. The jets thus assisted in keeping the tooth cavities clean. This is just perceptibly reflected in the penetration-rate curves of Fig. 1: Much the same results were obtained in hard rocks with hard-formation three-cone bits. The conclusion is that transport of loose cuttings towards the annulus does not require a powerful jet action. In the following sections we shall see that in deep-well drilling other phenomena are of greater importance. DYNAMIC HOLD-DOWN It is generally accepted that, in deep-well drilling, bit penetration is hampered primarily by the existence of a difference in pressure between the fluid in the hole and that in the rock at the depth of cutting.2-5 In the drilling

Jan 1, 1965

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

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

Jan 1, 1969

By W. A. Hockings, M. E. Volin, A. L. Mular

Batch grinding tests at short times were made in a laboratory rod mill with 10 x 14 mesh quartzite in corn syrup-water mixtures of varying viscosity. The weight fraction broken and size modulus were found to be independent of viscosity up to 20 cp, but at higher viscosities breakage of the feed decreased while the size modulus increased. The mechanisms of the effect of viscosity are discussed on the basis of particle dynamics. The effect of viscosity of the suspending medium on grinding has not been accorded much attention, although both wet and dry grinding are used in bene-ficiation processes and the viscosities of water and air differ greatly. schweyerl investigated the rates of grinding of quartzite in a pebble mill with air, water and glycerol as the suspending media and found the rate in terms of new surface developed per thousand revolutions of the mill to be constant and dependent on viscosity up to 20,000 rev, but to decrease and become independent of viscosity for longer grinding times. Based on considerations of the drag forces exerted on suspended particles and on the grinding media under turbulent and laminar flow conditions, viscosity can be expected to inhibit the rate of grinding in a given system as it changes the flow conditions from turbulent to laminar. The objective of this study was to determine the effect of viscosity of the suspending fluid in the batch grinding of a homogeneous feed of uniform size for short residence times, a region of more practical interest than the long times investigated by Schweyer. In actual grinding operations, the viscosity of the fluid does not change, but the consistency (apparent viscosity) of the pulp increases with additions of fines. In this study the consistency of the pulp was not measured because it was thought that the grinding times were too short to alter the apparent viscosity appreciably.* *Minus 200 mesh quartz at 65% solids has an apparent viscosity of about 25 cp as measured by a simple consistometer. Personal communication from D. F. Kelsall to A. L. Mular. For the purposes of this study, the effect of viscosity on grinding can be shown sufficiently well on the basis of the cumulative weight fraction finer than the feed size at any time, and the size modulus, ku, of the finer than feed sizes. The parameter, ku, is obtained from a form of the Gaudin-Schuhmann equation. METHOD AND MATERIAL The grinding tests were accomplished in a laboratory rod mill 10-1/2 in. long and 8 in. bore with a 20-lb rod charge consisting of two 1-in., eight 3/4-in., eight 1/2-in., and eight 1/4-in. diam rods. The mill speed was 43.2 rpm. Charges of 400 g 10 x 14 mesh Wisconsin quartzite** were ground for times of 60, **Courtesy of Minnesota Mining and Mfg. Co. 120, 180 and 240 sec in 500 cc of fluid made up of corn syrup and water in measured proportions. The temperature of the pulp was measured immediately upon completion of grinding, and the size distribution of the quartzite was determined by a wash-wet-dry screening technique. Corn syrup was chosen because of its high viscosity and ideal viscous behavior, and also because its density is not greatly different from that of water. The viscosity of each fluid mixture was measured with an Ostwald viscometer at 25 C. and a correction was made for the temperature of the pulp according to tabulated values for sucrose solutions. The density differences were so small that they probably did not significantly affect breakage. RESULTS AND DISCUSSION The conditions of the grinding tests, computed viscosities, weight fractions finer than the feed size and values of a and ku are listed in Table I. The values of a and ku were taken from the linear portions of the log-log plots of cumulative weight fraction finer vs size for the various grinding times. The plots for tests in fluids of average viscosity 8.6 cp are shown in Fig. 1. In general all of the sets of curves were linear throughout most of the size

Jan 1, 1965

By H. H. Klepfer, M. E. Snyder

UNTIL very recently, the commercial availability of the rare earths as metals has been very limited. Fabrication of mill products from these metals has not been studied in most cases. This note reports the results of the development of fabrication techniques for thulium. Thulium has a melting point of about 1550°C and a hexagonal close-packed crystal structure. It oxidizes in air to give a black oxide (TmzO,). The procedures for producing thin thulium foil were developed on one ingot weighing about 220 g and were subsequently applied in processing about a pound of metal to foil. Several alternates for the various fabrication steps were investigated and will be discussed. The metal fabricated was in the form of commercial chill-cast ingots 1 in. in diam and 2 in. long weighing approximately 220 g. Impurities in the ingots were reported by the vendor to be 4000 ppni tantalum, 2000 ppm calcium, 200 ppm nickel, 100 to 200 ppm iron, 100 ppm europium, and less than 100 ppm copper, lutetium, and ytterbium. In addition to these impurities, several salt-like inclusions as large as l/8 in. in diam were revealed along the center line of the one ingot sectioned. Preliminary tests indicated that small wafers cut from the as-cast ingot would not fabricate readily by rolling. Forging of copper-jacketed wafers was therefore attempted. At 1550"F forging was satisfactory but an apparent reaction of copper with thulium demanded investigation of lower temperatures. Therefore, the remainder of the test ingot was forged at 1450°F—with only minor cracking. All ingots forged were inserted into copper tubes of 1 in. 1D and 0.125-in. wall thickness. The jackets were sealed by flattening the ends of the tubes and welding under helium. Heating time at 1450°F was 30 min. Press forgings of 1/8 in. per pass were used, followed by 10 min reheats. When the ingots had been squared and reduced to 0.250 in. in thickness, the original copper jacket was stripped off and replaced by a new jacket in preparation for hot rolling. Hot rolling at 1450" F without edge cracking was readily accomplished after forging. Excellent results were obtained with 10 pct reductions of thickness followed by 5 to 10 min reheats. After reduction of thickness from 0.250 to 0.100 in. the copper jacket was removed. It was found, in fact, that hot rolling in air was possible. A tenacious black oxide similar to that seen on zirconium was formed during 3 min reheats at 1450°F. Reduction in air to 0.010 in. foil was possible taking 10 pct reductions per pass. The oxide coat formed during hot rolling in air could best be removed by sand blasting and pickling. Common pickling solutions containing polar solvents were found to attack the metal too rapidly and a concentrated nitric-hydrofluoric acid mixture attacked neither the oxide nor the metal. The most satisfactory pickling solution was 52 vol pct concentrated nitric acid-48 vol pct glacial acetic acid. After forging to 0.250 in., vacuum annealing and cold rolling was found to be another satisfactory alternate to hot rolling in copper jackets. After forging. a hardness of Rockwell B76 was found. Annealing in vacuum (2 X l0-5 mm of Hg) for 1 hr at 1200"F did not alter this value. Annealing at 1470"F for 1 hr brought the hardness down to R;]63. With cold rolling (5 pct per pass) the hardness returned to about R,1:76 after 20 pct reduction and edge cracking became noticeable. However, cold rolling to a total of 40 pct reduction in thickness (RI,,83) was possible before edge crack propagation became serious. Good surface finish was obtained, and the metal loss due to oxidation was minimized by cold rolling and vacuum annealing. Using this procedure the yield of 1.25 in. wide by 0.010 in. thick foil from a 1-in. diam ingot was about 40 pct.

Jan 1, 1961

By Marston G. Fleming

E. C. Peterson (Anaconda Copper Mining CO., Darwin, Calif.)—A study of this quite comprehensible and interesting paper by Dr. Fleming brings to mind several observations in the practical application of alkalinity and related factors to the actual practice of lead mineral flotation. Soda ash has long been a widely used and a very helpful alkaline conditioning agent in the flotation of galena from the usual run of lead-zinc ores. Soda ash is one of the common "standard" conditioning agents tried in any laboratory investigation of lead-zinc ores because it has so often proved helpful to galena flotation. However, the use of soda ash in the actual flotation of oxidized lead ores is certainly not widespread. In the flotation of certain lead-zinc ores from the Park City district of Utah, it was found in the usual cyanide-zinc sulphate-xanthate circuit that soda ash had an effect of producing a very condensed and flat froth regardless of the many frothers tried and that substitution of lime for soda ash to produce the same pH (8.0 to 8.4) improved the froth condition. However, the flotation of coarse particles of galena became so critical that some would pass through into the tailing, but caustic soda as a substitute for either soda ash or lime produced a very desirable froth condition in the same pH range and greatly improved the metallurgical outcome. Milling was carried out in typically "hard water" from watersheds of limestone and other calcareous rocks and the ore also apparently contributed magnesium and calcium salts to the pulp. Certain lead-zinc ores of Mexico have shown their greatest flotation response in circuits conditioned with sodium bicarbonate, and such actual mill use was also believed to be related to the water necessarily used in milling. On other lead-zinc ores the optimum has been obtained in actual milling treatment by use of caustic soda in both circuits of the operation. In flotation of oxidized lead ores in which sulphidi-zation is employed by addition of sodium sulphide, very high pH's exist in the flotation circuit. With the usual oxide lead ores, demanding 5 to 15 lb of sodium sulphide per ton of ore, pH's in the circuit (closed and open water-circuits) may range from 9.5 to 12, and for ores that contain much anglesite or a high percentage of iron oxide minerals, sodium sulphide consumption may reach as much as 30 lb per ton of ore and the resultant pH will be correspondingly higher. Yet in such flotation treatment employing either xanthate or oil (high-sulphur crude oil or diesel oils) as the collector, satisfactory to excellent grades of concentrate and lead recovery are obtained. Although there have been great studies and accomplishments in the investigations of flotation fundamentals and flotation theory, and although an investigation to determine the effect of all factors entering into the actual plant flotation of sulphide and oxide ores must become very complicated, as Mr. Fleming has mentioned in his caution regarding the drawing of general conclusions for other ores, it seems that it would be of greater interest if the techniques and investigations could, with time, approach the conditions of actual practice and lead the way to improved and more efficient flotation plant performance. Marston G. Fleming (author's reply)—The gulf between fundamental flotation research and plant operation is, perhaps, less wide than Mr. Peterson suggests. For example, the work described in my paper developed from an extensive investigation of a complex lead-vanadium ore from southwest Africa. This investigation started as a standard ore-testing problem but, at almost every stage, results were obtained which could not be interpreted in terms of previous experience. A program of fundamental research was therefore undertaken and was closely interlocked with the ore testing. This coordinated investigation resulted in a flotation process which has now been proved by two years of successful plant operation, and although the grade and mineralogical constitution of ore as well as smelter requirements have altered more than could have been anticipated, our knowledge of the fundamental character of the problem has made it possible to meet each change in conditions with much more confidence than would have been the case had we neglected this aspect of the investigation.

Jan 1, 1954

By H. Habenicht, J. J. Scott

An experimental method is described to compare strain data for rock and rock bolts when exposed to shock waves released by blastings. Static and dynamic strains were measured. Bolt gage response indicated a static decay of the bolt anchorage coupled with a reaction due to blasting vibrations. A distinct loss of bolt tension was indicated during vibrations and this was accompanied by a smaller but steady loss of tension due to static loading. The strain losses were found to relate to the shock source distance and the vibrational amplitude in the rock and bolts. Millions of roof bolts are used each year to stabilize mine openings and prevent their failure. Frequently, bolts become loose and lose their effectiveness as reinforcement to the mine structure. Many attempts have been made to determine the cause of loosening, including observation of the bolt tension or bolt torque through the use of torque wrenches, compression pads, and load cells. Pull tests have been used extensively to determine anchorage characteristics.102 These investigations have shown that even when bolts are loaded below the nominal anchorage capacity of the rock under 'static' conditions they exhibit a decay of load through adjustments and slippage at the bolt anchor. Since bolts are quite often found with the plate hanging loose from the rock surface, this condition suggests that some additional mechanism may have played a part in loosening the anchorage. Vibration from blasting or shock loading have been suggested as one cause in this type of anchorage deterioration. This paper gives the results of an investigation carried on at the University of Missouri at Rolla to determine the influence of shock waves on rock-bolt anchorage stability. ROCK BOLT ANCHORAGE UNDER DYNAMIC LOADING CONDITIONS In this report dynamic loading conditions are defined as conditions under which rock and bolts are exposed to vibrations occurring as a consequence of blasting, machine vibrations, earthquakes, rock bursts, etc. Whatever the location of the shock source, the waves created suffer extensive alterations along their travel paths, depending upon the vibrational properties and geometric parameters of the transmitting media. Together with reflection, refraction and diffraction at material boundaries, waves suffer changes in their nature to form tensile, compressive and shear waves. They become polarized, retracted, accelerated, dispersed, so that the waves incident at a point in the material can be the result of either one or a combination of this wide range of processes. A certain fraction of the wave energy will be lost due to friction, deformation, and material failure at the anchor-rock contact. Thus, the physical nature of wave propagation is such that it inhibits an exact analytical solution. For this reason this investigation should be considered qualitative rather than quantitative even though numerical data have been used as a base. CONCEPT OF THE INVESTIGATION Experiments were designed to find, from a series of rock bolts installed in natural mine rock, whether or not, and to what extent, bolt tensions would be influenced when bolts were exposed repeatedly to shock waves such as those produced by blasting. Electronic instrumentation available in the rock mechanics laboratory of the University of Missouri at Rolla was used to observe wave development in the rock and vibrational reactions in the bolts. EXPERIMENTAL PROCEDURE Fig. 1 illustrates schematically the experimental equipment. Electric resistance wire strain gages were used as sensors for the strain response of the rock and the bolts. Bolts were 5 ft long and 5/8 in. in diam, equipped with two-leaf type expansion shells for 1 3/8 in. boreholes. The bolt gages were mounted by use of Du Pont Household Cement, covered with an epoxy insulation and then enclosed in a copper foil for electronic shielding. For details of placement and the wiring diagram see Fig. 2 and 3. The output of both the rock gages and the bolt gages was fed into cathode ray oscilloscopes from which the vibrational traces could be photographed. A static strain indicator was also used to read the output of all gages.

Jan 1, 1968

By G. P. Conard, R. C. Hall, J. F. Libsch

The 50 pct Fe-50 pct Co alloy undergoes a transformation from disorder to an ordered structure of the CsCl type reportedly in the vicinity of 732OC. During this process, the coercive force goes through a maximum, apparently as a result of strains associated with the coherent nucleation and growth reaction. This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction, which is associated with annealing to a high degree of order with the continuous application of a magnetic field. The increase in ratio can be explained on the basis of a decrease in 90' domain boundaries and, perhaps, by an increase in anisotropy resulting from lattice distortion. THE 50 pct Fe-50 pct Co alloy undergoes a disorder-order transformation which has been reported to occur in the vicinity of 732°C1,2 The ordered structure is the CsCl type.' This magnetic alloy also shows a marked increase in the ratio of residual to saturation induction as a result of heat treatment in a magnetic field, sometimes called a response to magnetic anneal.'-' The purpose of this investigation was to study the course of the ordering reaction, the nature of the response to .heat treatment in a magnetic field, and the relation, if any, between ordering and the response. Procedure The method of approach in this investigation was to produce an initial structure as completely disordered as possible and then gradually to order the alloy by isothermal anneals at various temperatures under different conditions of the applied magnetic field. Magnetic, magnetostriction, and X-ray analyses were of primary importance in determining the property and structural changes resulting from the isothermal anneals. Rings of the 50 pct Fe-50 pct Co alloy were prepared from the elemental powders by a powder metallurgy technique, further details of which may be found in ref. 7. The initial structure was produced by annealing the specimens for ½ hr at 1000°C, cooling to and holding for ½ hr at 900°C (in the a range above the ordering temperature), and water quenching. Isothermal anneals were performed at 600°, 675°, 720°, and 740°C. For example, rings were heated to 600°C, held for a predetermined period of time, and cooled by natural cooling at a rate slightly slower than an air cool (average of 20" to 25°C per min). The tests (magnetic, etc.) were made after each heat treatment. All high temperature treatments were performed in a purified hydrogen atmosphere. The treatments at the various temperatures were carried out under one or more conditions of an applied field including 1—no field, 2—field of 20 oersteds applied on cooling only, and 3—field of 20 oersteds applied continuously during heating, holding, and cooling. Magnetic measurements were made using the standard Rowland ring technique8 with a maximum field strength of 100 oersteds. The magnetization curve, induction at 100 oersteds (B.), residual induction (Bt), and coercive force (Hc) were determined. All magnetic analysis data were based on an average of the results from three rings. A strain gage technique9 as used for the measurement of magnetostriction. The X-ray determination of the relative amount of ordered phase present was made on the ring specimen used for magnetic measurement. This was done by the back-reflection method using a rotating specimen (because of the large grain size) with unfiltered CoKa radiation and a 7 hr exposure time. Intensity measurements of the ordered line (300) were made by comparing visually the films so obtained with standard films prepared by exposing for different lengths of time a specimen given a long time anneal (high degree of order). Results In all instances the saturation induction (induction at 100 oersteds) was found to increase slightly with annealing time. This effect was small and appears to be the increase in saturation induction to be expected on ordering.10-13 The residual induction behavior was markedly influenced by the field condition during annealing, Figs. 1, 2. For the condition of no applied field, the ratio of residual to saturation induction remained essentially constant for short annealing times but showed a significant increase at longer times. With increasing annealing temperature, less time was required to produce this increase in the ratio. In the case of the 600°C anneals, the increase did not occur until approximately 20 hr, Fig. I, while on annealing at 740°C the increase was immediate, Fig. 2. Slight decreases in the ratio may be observed at 100 hr for specimens treated at 720°C and at 1 hr for those treated at 740°C. Specimens annealed in a field of 20 oersteds showed a residual to saturation induction ratio consistently higher than that for the specimens annealed without the field. The first anneal with the field (¼ hr) caused an abrupt increase in the ratio at all temperatures; thereafter, the increase in the ratio was generally similar for specimens annealed

Jan 1, 1956

By Fredrick Laist, J. P. Cooper

In the spring of 1931 an experiment was made at Anaconda with a small reverberatory furnace, with which was combined a roasting shaft. The arrangement is shown in the accompanying drawing (Fig. 1) and requires but little explanation. The general idea was to roast flotation concentrates, mixed with the proper fluxing materials, by dropping them down a heated shaft, from which the calcines would enter the smelting zone of the reverberatory furnace through an opening in the roof. The hearth of the reverberatory was 3 ft. wide by 21 ft. long. The roasting shaft was 3 ft. square inside and had an effective height of about 20 ft. The reverberatory furnace was arranged so that it could be charged either directly through the roof with regular hot calcines, in the usual manner, or by means of the shaft, at the top of which were a feeding arrangement by which cold flotation concentrates, suitably fluxed, could be fed continuously and a pipe through which cold air was blown in. Means were provided for distributing the feed over the cross-section of the roasting shaft. At first a mechanical arrangement was used but later it was found that a simple cone served the purpose equally well and gave less operating trouble. On account of the small size of the equipment, it was realized that absolute capacity and fuel consumption figures would mean but little. It was not to be expected that ratios comparable to those obtained in large furnaces could be gotten, and the only way in which any idea could be obtained as to the possibilities of the combination furnace was to run it in parallel with the reverberatory furnace charged in the usual manner and operating without the shaft. The results of such parallel runs are given in Table 1. The amount of oil burned in the reverberatory was purposely maintained at substantially the same rate during both runs in order that the benefits resulting from the shaft might be gaged. It was expected that the shaft would considerably increase the capacity of the reverberatory and that a better fuel ratio would be obtained on account of the high temperature at which the calcines would enter the

Jan 1, 1934