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By R. Grierson, L. J. Bonis

The high-temperature Strength oF TD nickel has been observed to be dependent upon the previons thermal and mechanical history of the material. Variations in both the level and the anisotropy of strength have been observed. 01 this paper- these variations are correlated with the storing of annealing resistant elastic strain energy in the matrix of the TD nickel. An x-vay line -broadening tecknique is used to measure the maLrTis elastie strain. THE inclusion of a finely dispersed second phase into a ductile matrix has long been recognized as an extremely effective method of strengthening the matrix both at high and at low homologous temperatures. It has been found, however, that the factors which determine the high-temperature strength are not the same as those which are important at low temperatures. Below 0.5 Tm the size and distribution of the second phase particles are of prime importance in determining the strength,')' while above this temperature the strength is mainly dependent upon the previous thermal and mechanical history of the alloy,3-7 This paper is primarily concerned with explaining the response of the high-temperature mechanical strength of one of these alloys (DuPont's TD nickel) to various thermo-mechanical treatments. It will be shown that this response is not associated with the occurrence of any form of dislocation substructure within the matrix of the alloy. It has been found, however, that a correlation does exist between the elastic strain level in the matrix and the previous thermomechanical history of the alloy and that the observed changes in elastic strain level parallel the measured changes in high-temperature strength. It therefore must be concluded that variations in high-temperature strength are a direct result of the variations in elastic strain level. MATERIAL TD nickel contains approximately 2 vol pct of Tho2 in an unalloyed nickel matrix. It is formed, as a powder, by a chemical technique and this powder is compacted to form ingots which are then extruded to give 21/2-in.-diam rod. Rod of smaller diameter is prepared from the as-extruded rod by swaging. In the studies reported in this paper, 1/2-in.-diam rod was used. This rod received an anneal of 1 hr at 1100°C prior to being used in any of these studies. EXPERIMENTAL TECHNIQUES Two methods were used to examine the structure of the nickel matrix of the TD nickel. These were: 1) transmission electron microscopy; 2) the analysis of the position and profile of X-ray diffraction lines obtained using the nickel matrix as the diffracting media. To prepare thin foils for electron-microscopical examination, slices of TD nickel approximately 0.050 in. thick were cut from the as-received 1/2-in.-diam rod. These were then chemically polished down to 0.045 in., rolled to 0.009 in., given a predetermined heat treatment, and thinned, using a modified Bollman technique, to provide the foils for observation. All observations were carried out at 100 kv, using a Hitachi HU-11 electron microscope. Specimens of the undeformed rod were prepared by grinding down the 0.050-in.-thick slices to approximately 0.015 in. and then thinning chemically and electrolytically to give the thin foils. The X-ray specimens were prepared by rolling 0.375-in.-thick rectangular blocks down to 0.075 in. The surfaces of the rolled material were ground flat, chemically polished to remove the layer disturbed by the grinding, and given a predetermined anneal in an inert atmosphere. They were then ground lightly to check their flatness and given a final chemical polish prior to being examined. The X-ray diffraction line profiles were measured using an automated Picker biplane diffractometer. A special specimen holder was built to allow a more accurate and reproducible positioning of the specimen. The line profiles were determined by carrying out intensity measurements at intervals of either 1/30 deg or 1/60 deg over a range of 3 deg on either side of the nickel peaks of interest. A piece of pure nickel which had been recrystallized to give a large grain size was used as a standard to give the X-ray line profile generated by a strain-free matrix. The analysis of the X-ray diffraction line profiles is a modification of that due initially to Warren and Aver-bach8and has been described elsewhere.3 This analysis gives a measurement of two parameters associated with the structure of the nickel matrix. These parameters are: 1) the size of the coherently diffracting domains within the nickel matrix; 2) the magnitude of the elastic strains in these domains. Both of these parameters are first determined in terms of a Fourier series. These series are obtained from other Fourier series which describe the measured profile of the X-ray diffraction lines. Thus, for both the coherently diffracting domain size and the elastic strain level, it is possible to plot Ft (the Fourier coefficient) against t (the term in the Fourier series), where t can be expressed in terms of a distance L and the Fourier coefficient Ft(S) (associated with elastic strain level) can be expressed in terms of the root mean square strain (e2)1/2. Thus a plot of (F 2)1/2 vs L can be obtained. Plots of this type are shown graphically in Figs. 6 and 8. Interpretation

Jan 1, 1968

Jan 1, 1954

By R. A. Vandermeer, J. C. Ogle

The preferred crystallographic orientations (texture) developed in randomly oriented, poly crystalline niobium during rolling were studied by means of X-ray diflraction techniques. The evolution of texture at both the surface and center regions of the rolled strip was carefully examined as a function of increasing defamation throughout the range 43 to 99.5 pct reduction in thickness. Certain aspects of the center texture development in niobium are in agreement with the predictions of a theory by Dillamore and Roberts, but others cannot be explained by the theory in its present form. Above 87 pct reduction by rolling, a distinctly different texture appeared in the surface layers which was unlike the center texture. The present results are compared with previous results obtained from other bcc metals and alloys. RANDOMLY oriented, poly crystalline metal aggregates when plastically deformed to a sufficiently large extent develop preferred orientations or textures. In a recent review article, Dillamore and Roberts1 pointed out that the nature of the developed texture may be influenced by a large number of variables. These include both material variables such as crystal structure and composition and treatment variables such as stress system, amount of deformation, deformation temperature, strain rate, prior thermal-mechanical history, and so forth. From a practical point of view, the control of preferred orientation may often be important for the successful fabrication of metals into usable components. During the past few decades many experiments have been devoted to the study of deformation textures. This work, however, has been confined in large part to metals and alloys that have an fcc crystal lattice. By comparison, bcc metals and alloys have received much less attention, and consequently our understanding of preferred orientations in these materials is only shallow. This state of affairs worsens when it is realized that almost all of our present howledge about this class of materials derives from studies on irons and steels.' The bcc refractory metals, which are relative newcomers to the industrial world, have, on the other hand, been given at best only passing glances in the area of texture development. Our understanding of the evolution of preferred orientations in bcc metals can only remain fairly limited until systematic studies of metals and alloys other than the irons and steels have been carried out and the influence of the many variables has been determined. To that end a program was initiated to investigate in detail texture development in niobium. The present paper reports some of the results of this study. Textures were determined at both the center and surface of strips rolled variously to as much as 99.5 pct reduction in thickness at subzero temperatures. Emphasis in this paper is on texture description and on texture evolution during rolling to progressively heavier deformation. EXPERIMENTAL PROCEDURE The niobium was purchased from the Wah Chang Corp. as a 3-in.-diam electron-beam-melted billet. Chemical analysis indicated the impurities to be less than 300 ppm Ta, 40 ppm C, 10 ppm H, 170 ppm 0, and 110 ppm N. All other impurities were below the limits of detection by spectrochemical analysis. This large-grained billet was fabricated into specimen stock so that a fine-grained randomly oriented grain structure resulted. This was accomplished in three deformation steps alternated with recrystalli-zation anneals of 1 hr at 1200°C in a vacuum of low 10"6 Torr range after each deformation step. The first step was to alternately compress the billet 10 to 20 pct in each of three orthogonal directions. The second step was to compress in only two directions 90 deg apart to produce a 2-in.-sq bar. The final step was to roll this bar 50 pct to give a 1-in. by 2-in. cross section. After the final anneal, metallo-graphic examination showed the material to have an average grain size equivalent to ASTM No. 5 at 100 times (i.e., 0.065 in. diam). Specimens cut from the center and edges of this bar gave no indication of detectable preferred orientation when examined by X-ray diffraction. Samples 1.5 in. long, either 0.625 or 0.750 in. wide, and approximately 0.400 in. thick were machined from this fabricated ingot. The surfaces corresponding to the rolling planes were ground so as to be parallel. The samples were chemically polished in a solution of 60 pct nitric acid and 40 pct hydrofluoric acid (48 pct solution) prior to rolling to remove any cold work introduced in the machining operations. Rolling was accomplished with a 2-high hand-operated laboratory rolling mill that had 2.72-in.-diam rolls. Prior to operation, the rolls were polished with 600 grit paper, cleaned with acetone, and then soaked in a container of liquid nitrogen for several hours. The samples were also soaked in liquid nitrogen prior to rolling and were recooled between each pass. While some slight heating of the samples occurred during rolling, this procedure maintained the sample temperature well below 0°C at all times. The samples were rolled unidirectionally, and the rolling plane surfaces were not inverted during any phase of the operation. The draft per pass averaged between 0.010 to 0.012 in. After 96 or 97 pct reduction the draft was reduced to 0.001 to 0.002 in. per pass. Samples were rolled to various reductions in thickness between 43 and 99.5 pct.

Jan 1, 1969

By AIME AIME

FOR the development of a mine, a shaft of small cross-section is usually sunk, of no larger size than is absolutely necessary. After the mine has been developed and put on a production basis it is a common practice to sink a shaft of larger cross-section for hoisting and use the original shaft for the handling of men and material into the mine. While that ar¬rangement is usually expedient, the small size of the supply cages presents a problem in the handling of ma¬terial. Where a cage is, say, 4 by 4 ft., which is not an uncommon size, all lengthy and bulky material must be handled piece by piece several times from the supply department to the underground stations. Thus drill steel and timbers over 4 ft. long will be loaded on cars at the supply department and wheeled to the shaft, where each piece must be unloaded from the cars and stood up in the cage. They are then lowered to the shaft station, unloaded piece by piece and placed in cars, run to the supply station and again unloaded and stored until called for. In such a case handling becomes a major item in the cost of supplies.

Jan 1, 1930

By Shrikant Lele

INTRINSIC and extrinsic stacking faults may be formed in fcc structures by similar processes with bounding S-and D-Shockley partial dislocations respectively.' The only difference between these two dislocations is that the side on which the fault lies is reversed. The number of nearest and next nearest neighbors of atoms is identical for the two faults, but that of next next nearest neighbors is different. Thus, the energies associated with intrinsic and extrinsic faults may be expected to be different but of the same order of magnitude. An extrinsic fault can be constructed from two intrinsic faults on adjacent close-packed planes. This suggests that the extrinsic fault energy is higher than the intrinsic fault energy. It is, therefore, surprising that extrinsic faults have not been detected in X-ray studies2 of deformed alloys of low stacking fault energy. These have, however, been observed in cerium3 by X-ray diffractometry and in Au-4.8 at. pct Sn4 and Cu-16.6 at. pct A15 alloys by transmission electron microscopy.

Jan 1, 1968

By R. A. Graham, R. W. Rohde

The effect of hydrostatic pressure upon the austenite start temperature of a commercial Fe-28.4 at. pct Ni-0.5 at. pct C alloy has been determined. For pressures to 20 kbar, the austenite start temperature decreased from its atmospheric pressure value of 380°C at the rate of about 4°C per kbar. These data are analyzed by two different thermodynamic approaches; first, considering the transformation as an isothermal process, and second, considering the transformation as an isentropic process. It was found that both these approaches fit the experimental data equally well. The effect of hydrostatic pressure upon the austenite start temperature is best described by considering the mechanical work done during the transformation as that work obtained by multiplying the applied pressure with the gross volume change of the transformation. It is widely recognized1 that strain has an important effect on the initiation of martensitic transformations.* For example, the martensite start tempera- *In this paper, use of the term martensitic transformation implies the reversal of martensite to austenite as wen as the formation of martensite from austenite. ture, M,, may be increased by plastic deformation. Similarly, plastic deformation is observed to lower the austenite start temperature, A,. The effect of uniaxial stress on the M, of iron-nickel alloys has been studied by Kulin, Cohen, and Averbach.2 They found that the martensite start temperature was significantly changed by stresses well within the elastic region. Moreover, the effect of tensile and compres-sive stresses differed. These effects were explained in terms of the interaction of the applied stress with both the dilational and shear components of the transformation strain. The magnitudes of the influence of uniaxial tension, compression and hydrostatic pressure on Ms were measured in 30 pct Ni 70 pct Fe by Pate1 and Cohen.3 Their thermodynamic calculations and similar calculations by Fisher and Turnbull4 predicted the experimental results when the transformation was assumed to occur isothermally at some fixed driving force. This driving force was assumed to be supplied by a combination of the chemical free energy difference between the austenitic and martensitic phases and the work performed during transformation by the applied stress. More recently, Russell and winchel15 reported the effect of rapidly applied shear stress on the reversal of martensite to austenite in iron-nickel-carbon alloys. They performed a thermodynamic analysis of this transformation based upon the assumption that the re- versal occurred adiabatically. They concluded that the applied shear stress did not significantly interact with the transformation strain and thus did not assist in inducing the reversal. Rather they concluded that the reversal was effected by localized strain heating which resulted from the gross local shear deformation of the experiment. In either the adiabatic or isothermal analysis it is necessary to compute the work performed by the interaction of the applied stress and the transformation strains. In the case of hydrostatic pressure this interaction has been treated by two different methods. In either case the applied pressure is assumed to remain constant during the transformation. In one treatment the applied pressure is assumed to interact directly with the dilatational strain associated with the formation of an individual martensite plate.3'4 This local strain has been measured at atmospheric pressure in iron-nickel alloys by Machlin and Cohen.6 In the above treatment this local strain is assumed invariant with temperature and pressure changes. In the other treatment the applied pressure is assumed to interact with the gross volume change of the transformation.7,8 The usefulness of this latter treatment has been demonstrated by Kaufman, Leyenaar, and Harvey7 who calculated the effects of pressure upon the martensite and austenite start temperatures of Fe-10 at. pct Ni and Fe-25 at. pct Ni alloys. Excellent agreement was obtained between their calculations and their experimental data on an Fe-9.5 at. pct Ni alloy. However, this treatment suffers from the fact that the data required to calculate the volume change of the transformation (i.e., the initial specific volumes, the thermal expansion and compressibility data for both the austenitic and martensitic phases) is, in general, not available for any material except pure iron. Thus the calculations of Kaufman et al.7 were necessarily performed by assuming that the volume change of the martensitic transformation in the iron-nickel alloys was that same volume change occurring during the a-? transformation in pure iron. While this approximation may suffice for very dilute alloys it is likely to be inaccurate in high nickel alloys. We have performed measurements of the effect of hydrostatic pressure to 20 kbar on the A, temperature of an Fe-28.4 at. pct Ni-0.5 at. pct C alloy. The composition is similar to the alloy used by Pate1 and Cohen3 to determine the effect of pressure upon the M, temperature. The present measurements permit calculation of the interaction between the applied pressure and the transformation strain. Additionally, measurements have been made which allow precise determination of the gross volume change of the transformation. The data allow direct comparison between the alternate hypotheses of the interaction between the applied pressure and a dilatational transformation strain characterized by either the formation

Jan 1, 1970

By R. W. Carpenter, J. C. Ogle

It is possible, by using pack-rolling instead of conventional rolling, to reduce a number of metals to thicknesses of 2µm or less. Such thinfoils are generally made at room temperature without intermediate annealing. In addition, pack-rolled foils fail by developing pinholes at thicknesses near 2µm instead of developing the shear cracks usually observed in cold-rolled ductile metals. This paper presents the results of a general investigation of the deformation substructure and texture developed in copper and iron pack -rolled from 130 to about 2µm thickness. Electron microscopy showed that in both metals a fine (0.2 to 0.5?µ m) deformation subgrain structure formed during pack-rolling; in neither case was this substructure grossly different from substructures formed during conventional rolling. The deformation texture formed in pack-rolled iron was quite similar to usual bcc textures; however, in the case of copper, the cube texture was stable during pack-rolling and the normal copper deformation texture was unstable. It is shown analytically that the constraining pack induced a large hydrostatic pressure in the foils during pack-rolling. The pinhole failure mechanism is attributed to the presence of the large hydrostatic pressure during pack-rolling; this strongly suppressed the growth of shear cracks. The stability of the cube texture in copper is also probably due to the unusuul stress distribution developed during pack-rolling. EXPERIMENTS at several laboratories have shown that very thin foils of the common structural metals and many of the rare earths can be made by "pack-rolling". 1-3 The technique was originally developed to make specimens for nuclear scattering experiments and foils for X-ray filters. It is also useful for making experimental laminar metallic composite bodies and foils thin enough for direct examination by ultra-high voltage electron microscopy without the need for special thinning techniques. Pack-rolling in the present context means a three-layer pack, with the material to be rolled into foil comprising the center layer. The outer two layers, which constrain the foil during reduction, are ordinarily austenitic stainless steel. Typically, a 130 µm (0.005 in.) metal strip can be reduced to a final thickness of 2 µm or less by this process. This is accomplished at room temperature, without intermediate annealing. It has been observed that foils produced by this process do not exhibit at any stage of their reduction the severe work-hardening found in strip rolled by conventional cold-rolling methods. Neither is the failure characteristic the same."&apos; Conventionally cold-rolled ductile metal strip fails by developing shear cracks on planes whose normals nearly bisect the angle between the rolling direction and normal to the rolling plane; these are planes of maximum shear stress. In pack-rolling this mechanism has not been observed; failure occurs by the formation of pinholes on the foil surface (penetrating the foil). If pack-rolling is continued the hole density increases. These differences in behavior imply the existence of appreciably different substructure in pack-rolled foils compared to substructure in conventionally rolled material, or perhaps that the geometry of pack-rolling has an effect on the foil behavior. This paper describes an investigation of deformation substructure and texture in some specimens of pack-rolled copper and iron, and some considerations of the stress distribution in the foils during rolling that result from the geometry of pack-rolling. EXPERIMENTAL DETAILS Three different materials were used for pack-rolling in the present work: soft copper sheet (99.8 pct Cu, 0.03 pct 0, electrolytic tough pitch) and two types of iron, Ferrovac E* and Armco iron. Each was "Crucible Stccl Co. initially in the form of 130 µm annealed strip with grain size ranges of approximately 10 to 40 µm. The initial texture of the copper (determined as noted below) was the normally observed cube type (001)[100]; there was evidence of a small amount of material in the cube-twin orientation reported by Beck and Hu.4 The initial texture of the Ferrovac E was similar to that reported for recrystallized iron by Kurdjumov and sachs,5 who list the principal orientations as {111}<112>, {001}<110> 15degfrom RD and a weak component {112}(110) 15 deg from RD. The starting texture of the Armco iron was not determined. Pack-Rolling Procedure. A four-high mill was used for all specimens. The work roll and backing roll diameters were 1.625 and 5.25 in., respectively. The peripheral roll speed of the work rolls was about 2.5 in. per sec. All foils were initially reduced from 130 to 100 µm by conventional straight rolling and then inserted into a pack, without any intermediate annealing, for further reduction. The pack consisted of an 0.033 in. (838 µm) thick 3 by 6 in. polished sheet of austenitic stainless steel, folded to make a 3 by 3 in. jacket. After folding, the jacket was given a small reduction to close the fold tightly before insertion of the foil. During pack-rolling a constant change in roll spacing was made every third pass. The roll-spacing change corresponded to a 5 pct reduction in thickness for a new pack. This approached a 10 pct reduction when the pack had decreased to about half its original thickness. At this point the deformed pack was discarded and a new one

Jan 1, 1970

By C. A. Edwards

Jan 1, 1935

Jan 1, 1884

By C. A. Edwards

Jan 1, 1935

As reported in Drift last month demolition of old buildings on the 47th St. site has been completed, and detailed plans and studies for the new structure are well underway. As this issue was being sent to press the $5 million industry- business phase of the fund campaign had passed the half- way mark toward its goal. Target date for the building program is 1960.

Jan 5, 1958

By R. E. Rowlands, I. M. Daniel

An experimental study is discussed of the state of strain, deformation and fracture around lined and unlined cylindrical cavities in rock under biaxial compressive static loading. The specimens were limestone plates 36 in. high, 24 in. wide and 3 in. thick with a 4 in. diameter hole unlined or lined with hydrostone and aluminum liners. The ratio of horizontal to vertical applied stress was one-third. Experimental techniques used were strain gages, differential transformers and photoelastic coatings. Strain distributions were plotted along the horizontal axis of symmetry for various levels of applied stress. The agreement with theory was good at low stress levels but appreciable deviations occurred at higher levels and near the hole. The experimental strain concentration factor was nearly the same, slightly higher and appreciably higher than the theoretical value for the hydrostone-lined, unlined and aluminum-lined specimens, respectively. Failure modes were splitting along a vertical plane for the unlined and hydrostone-fined specimens, compressive (shear) failure for one unlined specimen and splitting with radial cracking for the aluminum-lined specimen. In the latter, a maximum tensile stress of more than three times the uniaxial tensile strength existed at the interface with the liner at the time of fracture,

Jan 1, 1971

By P. J. Connors, William J. McCormick, F. B. Olender, Jerome F. Sheldon, Donald T. Mylar, Edgar T. Gregory, Owen S. Roberts, H. Colijn

A quiet revolution has taken place in the conveyor belt industry. Man-made materials, improved production facilities, better splicing techniques and emphasis on repair and maintenance have all contributed substantially to lowering material handling costs for the mining industry. Formerly the rubber industry had only cotton for a belt carcass and tree grown rubber for the protective cover. Monofilament rayon, with desirable elongation and elastic properties, was the first synthetic fiber to be used in place of cotton. But wet strength and other limitations restricted its use, and the conveyor belt industry turned to nylon.

Jan 3, 1972

By K. Tangri, M. Chaturvedi

The addition of 0.5 wt pct Cu to Zr-2.5 Cb alloy increases the as -quenched hardness of the hexagonal martensitic a&apos; phase, produced by water-quenching bccß-Zr phase, by about 35 pct. This strengthening has been attributed to the solid -solution hardening of the matrix. On aging ternary martensite, a&apos; phase reverts to equilibrium a and Zr2Cu and ß-Cb precipitate out, mainly at the twin and grain boundaries, causing a secondary hardening of the matrix. COLD-worked Zircaloy-2 pressure tubes have been in use in power reactors for a considerable period of time. The search for a better material led to the development of Zr-2.5 wt pct Cb alloy which in the quench-aged condition develops 50 pct more strength than that of cold-worked Zircaloy-2, however, its corrosion resistance in water and steam in the temperature range of 316" to 400°C, in absence of neutron flux, is inferior to that of zircaloy-2.&apos; Work carried out by Ells et al.1 and Dalgaard2 has shown that the corrosion properties of Zr-2.5 wt pct Cb alloy can be considerably improved by the ternary addition of 0.5 wt pct Cu. This paper is concerned with the effect of 0.5 wt pct Cu on the formation of martensitic a and its aging characteristics in a Zr-2.5 wt pct Cb alloy. MATERIALS AND EXPERIMENTAL TECHNIQUES Zr-2.5 Cb-0.5 Cu (referred to as the ternary alloy) and Zr-2.5 Cb (referred to as the binary alloy) alloys, supplied by the Chalk River Nuclear Laboratories of the AECL were used. The detailed chemical analysis is given in Table I. Cold rolling and swagging with frequent intermediate anneal of 1000°C were used for the initial fabrication of the alloys. 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 10-6 mm of Hg. For optical metallography and hardness measurements specimens were mechanically and then chemically polished in a 45 pct HNOj, 45 pct HzO, and 10 pct HF solution. Hardness was measured 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 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 Cu Ka radiation in a 114.6 mm diam camera. Carbon extraction replicas were prepared by etching the specimens, after depositing a layer of carbon on the metallographic specimen, in one part HF and thirty parts ethyl alcohol. Thin films were prepared by electropolishing heat-treated 3/4 by 1/2 by 0.005 in. thick strips using a modified Bollman-Window technique. The 10 pct perchloric acid-90 pct methyl alcohol bath was kept at -50°C and polishing was done at 5 to 10 V. The thinned specimens were washed in ethyl alcohol at -30º to -40°C and dried between filter papers. Replicas and thin films were examined in a Phillips 300 G electron microscope. For resistivity measurements thin strip specimens 0.02 by 0.3 by 10.0 cm long were used. The potential leads were spot welded to the specimens in order to maintain a fixed length for the initial and the final resistivity measurements. The resistivity was measured by a Kelvin bridge in a temperature controlled room. The temperature was maintained at 72º ±1°F and the accuracy of the resistivity measurements was 0.03 µa-cm. RESULTS As-Quenched Structures. In order to produce a homogeneous matrix to study the precipitation reaction the solution-treatments of both the alloys were carried out in the -field region. From the Zr-Cb phase diagram due to Lundin and cox3 ß/a + ß phase boundary for Zr-2.5 wt pct Cb alloy is 820°C. Ells et al.1 have reported this boundary for Zr-2.5 Cb alloy containing 1100 ppm 0 to be at 920°C. Also, the addition of 0.5 wt pct Cu reduces this temperature by 50°C. Therefore, the solution-treatments were carried out at 1000°C to ensure that the alloys were in ß-phase region. The soaking time was 1 hr and the specimens were water-quenched. The as-quenched hardness of the binary alloy was 245 Vpn whereas, that of the ternary alloy was 330 Vpn. The X-ray diffraction studies indicated that the as-quenched structure of both the alloys consists of martensitic hexagonal phase a&apos;, with a c/a ratio of 1.591, and some retained ß-Zr. The presence of a&apos; phase was further confirmed by thin film electron microscopy. Electron micrographs of typical ß-quenched structures of the ternary and the binary alloys are shown in Figs. 1 and 2, respectively. Fig. 3 shows the diffraction pattern from an area similar to that shown in Fig. 1. Although, the as-quenched hardness of the ternary alloy is about 35 pct greater than that of the binary alloy, the structure of both the alloys seems to be the same. The matrix of both alloys is heavily twinned and shows very few dislocations. Furthermore, there is no evidence of any precipitation taking place in either of the two specimens during quenching from the solution-treatment temperature. Aging Behavior of Martensitic a&apos;. The aging kinetics of the ternary alloy were followed by resistivity and hardness measurements. The as-quenched values

Jan 1, 1970

By C. A. Heiland

THE investigations described were made on a portion of Clear Creek basin near Golden, Colo. (A portion of the area under survey is shown in Fig. 1. The photograph was taken in the vicinity of station No. 2 of Line II.) Three principal purposes were in the minds of the investigators: 1. To repeat with modern, more sensitive and more reliable instruments magnetometric placer investigations previously made. 2. To disprove the recently advanced theory that deposits of the glacial period have no pronounced magnetic effects. 3. To develop in detail the theory of the magnetic effects produced by placer deposits. Six years ago, the first magnetic observations were published, which proved that it is possible to locate gold-bearing channels with relatively crude instruments.1 Two years ago, K. C. Laylander published results obtained with similar instruments.2 The measurements of the latter were discussed by the senior author of this paper.3 At that time, however, the material available for such a discussion was somewhat scarce and not made with the most modern magnetic instruments, although very brilliant results were obtained. Therefore, the senior author felt that magnetic investigations should be made on very clear geologic conditions with the most modern instruments.

Jan 1, 1928

By K. J. Brondyke, P. D. Hess

Lack of correlation between densities of aluminum alloy samples solidified under reduced pressure (vacuum gas test) and hydrogen content of the metal is explained on the basis of inclusions serving as nuclei for hydrogen bubble formation. As a result of these findings, the vacuum gas test is suggested as a practical tool to indicate metal cleanliness and the potential effects of hydrogen. Visual obseroation of the sample during solidification under vacuum provides more useful information than that obtained either from density meas -urements or sectioning of the solidified sample. A rapid, practical method of detecting and measuring gas in molten-aluminum alloys has been sought for many years. The development of the vacuum gas test appeared to solve this problem. However, the vacuum gas test is subject to certain limitations, and the results must be interpreted mindful of these limitations, as described herein. Since the only gas which is appreciably soluble in aluminum is hydrogen, the terms gas and hydrogen will be used interchangeably in this paper. In the early days of the aluminum industry, when techniques were not as refined as they are today, it was customary to pour small "pancakes" of alumi- num on a metal plate. These were allowed to solidify under atmospheric pressure. Excessive gas in the metal was evidenced by the appearance of small bubbles on the surface of the "pancake".&apos; It is obvious that such a method would only show very high gas levels far above any which could be tolerated by today&apos;s standards of quality required for most applications. Various attempts were made to develop methods which would be more sensitive. Among these were measurements of density of aluminum alloys solidified under controlled conditions, radiography of cast slabs, and examination of fractures. The method which showed the greatest promise and appeared most practical was the solidification of a sample of the molten aluminum alloy under vacuum. The samples so solidified are evaluated either by observation for bubble emission during solidification, by sectioning the solidified sample and examining for porosity, or by determining the density of the solidified sample. This test is known internationally as the Straube-Pfeiffer test or the reduced-pressure solidification test, but in America it is commonly referred to as the vacuum gas test2-&apos; and will be so designated in this paper. The determination of density appeared an attractive and convenient means of providing a numerical rating for the samples. Attempts were made in some instances to correlate densities to hydrogen contents. Unfortunately, the visible loss of gas during solidification of samples with high gas contents precluded a correlation of vacuum gas test densities with the observed gas evolution. This led to Alcoa&apos;s use of two different pressures; 50 mm of

Jan 1, 1964

The following list comprises the names of those persons who became members during the period Sept. 10, 1919, to Oct. 10, 1919. ADAMS, R. L Chief Min. Engr., Old Ben Coal Corpn., Mine 12, Christopher, Ill. Arizumi, Hiroshi The Kuhara Min. Co., Ltd., Tokyo, Japan. BAILY, T. F Pres. & Gen&apos;l Mgr., Electric Furnace Co., Alliance, 0. BALDWIN, A. J., Min. Engr Box 145, Pikeville, Kentucky. BALTZELL, WILL H., Chief Engr., Canadian Steel Corpn., Ltd., Ojibway, Essex Co., Ont., Canada. BANON, E. MAGAWLY, Examining Engr., Dawes Bros., Inc. 108 So. LaSalle St., Chicago, Ill. BARNES, W., Mech. Engr., Excavating Dept., Ruston & Hornsby, Ltd., Lincoln, England. BASHORE, E. GORGAS, Head Chem., The Babcock & Wilcox Co., Bayonne, N. J. BASS, A. L Supt., Hobson Silver Lead Co., Ltd., Ymir, B. C., Canada. BORN, SIDNEY, Chief Chem., Empire Gas & Fuel Co. & Associated Companies, 1100 Cherokee Ave., Bartlesville, Okla. BRADLEY, J. G,, Coal Operator, Pres.,& Gen&apos;l Mgr., Elk River Coal & Lumber Co., Dundon, W. Va. BENNETT, G. W., Min. Engr., Sales Engr., Taylor-Wharton Iron &- Steel Co., High Bridge,. N. J. BRUFF, CHARLES E., Member firm Bradley, Bruff & Labarthe, 85 Second St., San Francisco, Cal. COLLINS, FRANK W., Chief Engr., Bradley, Bruff & Labarthe, 85 Second St., San Francisco, Cal. CONNELL, F. M., Min. Engr., Eastern Alin. & Mill. Co., Ltd., 905 Bank of Hamilton Bldg., Toronto, Ont., Canada. CooK, HARRY H., Vice-pres., Valley Mould & Iron Corgn., 137 Shenango St., Sharpsville, Pa.

Jan 11, 1919

By Joel Watkins

THE terms kaolin, china clay, ball clay, and paper clay are more or less loosely and interchangeably applied to a large class of white-burning clays. These clays are made up chiefly of hydrous amorphous (colloidal) aluminum silicates with variable amounts of free silica and other impurities in small quantities. Occasionally also the term kaolinite is wrongly applied to a white-burning clay, for kaolinite, is a mineral of definite chemical composition and definite crystalline form which occurs but sparingly in nature. It is not the purpose of this paper, however, to define or classify clays, or to enter into a discussion of the chemistry of clays. It is rather a general description of the occurrence and methods of mining of a kindred group of clays, and a discussion of their future economic aspect. In the Southern Appalachian States, the mining of white-burning clays has in recent years become an important industry. Although these clays are of several distinct types, which occur in as many localities under different geologic conditions, they are for the most part. primarily of the same origin. They are essentially the same in elementary constituents, and are largely consumed in the same industrial arts. From the standpoint of a mining engineer, therefore, and for convenience in comparison and correlation, it should be interesting to discuss them under one heading.

Jan 2, 1915

By Robert C. Stephenson

In 1954, a year when general industrial production declined, it is significant that industrial mineral products continued in high demand. Phenomenal growth of the cement industry, increase in filler-fibre markets such as for asbestos and diatomite, expansion of the gypsum industry, and the increasing production of agricultural raw materials all illustrate the vitality of industrial minerals fields. Pioneer development of new industrial raw materials-wollastonite as a current example-requires keen foresight, hard work, and tremendous intestinal fortitude. The present-day requirements of industry for rigid quality control and uniformity of raw materials have resulted in highly specialized

Jan 3, 1955

By Henry W. Turner

The following notes were taken on a trip through northern Korea in the fall of 1910. We started with about 19 Korean ponies, and as many Koreans, from Shin Anju, on the railway from Seoul to Antung. We passed through the ancient walled city of Anju, and thence up the valley of the Anju river. We then crossed over to the north into the Kalsan drainage, thence over another divide to the Chasung river, and finally reached the old Korean copper-smelting camp of Tong Jom— in all a distance by trail of about 200 miles. The return trip was made with ponies to Mampochin on the Yalu river, and thence&apos; down the river in a junk to Antung. The route of travel is shown on the sketch map, Fig. 1.

Jan 1, 1913