Search Documents

By J Kelly, C Kelsey

The Autogenous Kelsey Axial Displacement (KAD) mill is essentially a short retention time centrifugal mill, designed to effectively harness extraordinarily high concentrations of input energy for efficient particle size reduction. The mill which has an operating duty band extending from primary to ultra-fine also has the unique capacity to generate fine and ultra-fine products from single stage size reductions of either coarse or fine mineral feeds, in both wet or dry operating modes. This paper covers developments following a recent optimisation programme which confirmed that mill shell speed and the velocity differential between the shell and the rotating internal disc assembly tips could be optimised in relation to specific energy for a given feed rate and feed size providing the product size remains constant. Subsequent test-work further demonstrated the versatility of the mill with duties ranging from the removal of surface contamination associated with heavy mineral concentrate particles to the production of premium ultra-fine talc, from a single stage of size reduction. Open circuit grinding of a rougher gravity concentrate to liberate gold ahead of leaching also highlighted the scope for the application of low cost KAD milling for a range of hydrometallurgical feed preparation duties, which involve fine and ultra-fine grinding.CITATION:Kelly, J and Kelsey, C, 2013. Advanced media-free fine and ultra-fine grinding - The Kelsey axial displacement mill, in Proceedings MetPlant 2013 , pp 152-161 (The Australasian Institute of Mining and Metallurgy: Melbourne).

Jul 15, 2013

A glance at the accompanying table reveals that for the first time in many years the industrial minerals field did not post significant gains in production. Gypsum, cement, sand and gravel, and crushed stone production-which are very sensitive to construction industry performance-barely managed to stay even, or in fact, declined slightly. Even barite production, in spite of increased oil and gas drilling, was off 5% in 1979 from 1978. Domestic fluorspar production continued to decline while imports, primarily from Mexico, South Africa, and Spain, rose 11%. Worldwide production of iodine (which is used primarily as a catalyst, animal feed additive, and in ink and colorants, stabilizers, pharmaceuticals, photographic film, and disinfectants) grew about 9% on a worldwide basis-from 10.75 Tg to 11.16 Tg. This does not include US production. The two US producers increased demand as apparent domestic production increased about 5%. Increasing energy costs and the transportation situation have had a prohibitive effect on the industrial minerals field. The critical shortage of railroad cars, according to consultant Neil Johnson, is due to several factors: • High capital investment rind interest rates needed to build new equipment, • Accidents, obsolescence, and failure to repair equipment reduced car availability, •Reduction in some train schedules due to fuel economy, and •Increased turnaround time as a result of new safety regulations. Railroad rates have increased an average of 7% per year for the past 10 years. Last year's increases were about 11.1% due to fuel costs and surcharges. Truck shipments have increased their penetration in certain areas of the southeast because of the shortage of box cars and the producers' demand for quick shipment. Some producers report a 15-20% increase in truck use. In the feldspar-mica area, which formerly made most shipments by rail, trucks are now used for about half the shipments. Marketing prices, as expected, increased due to higher costs for energy, power, labor, and supplies used in production. Estimated average price increases were about 10% which, in many cases, did not cover increased production costs. These factors-rising energy costs and the critical transportation situation-resulted in an unsatisfactory year for industrial minerals. Stanley J. Lefond Cnntrihuting Editor

Jan 5, 1980

By William M. Knuth

The objectives of this study were the evaluation of the technical and economical feasibility of utilizing color and color infrared aerial photography to satisfy some informational needs for pre-mine planning, mine progress monitoring, and reclamation monitoring. Seven aerial data collection missions were conducted within a two year period over strip mine sites in Pennsylvania. Scales of 1:6,000; 1:12,000; 1:24,000; and 1:48,000 were flown on the first aerial mission. The 1:12,000 was chosen as the optimal scale for the remaining six flights. An extensive literature and legislative review and discussions with strip mine operators identified the informational parameters that could possibly be obtained from aerial photography. Interpretability accuracies were evaluated for parameters in the broad categories of land use, soils, water, vegetation, and geology. Seasonal interpretation characteristics were noted for infrared photography. Man hours for photointerpretation and costs of aerial data collection were monitored closely and incorporated into a cost benefit analysis. Those parameters that could be mapped, identified, or measured with accuracies =75 percent were incorporated into a photointerpretation training manual. An additional analysis was performed to determine the applicability of the interpretation methodologies and results developed for Eastern sites to Western surface coal mining sites. In addition, a separate aerial photographic interpretation manual was produced during this contract. It is intended to be utilized by eastern and midwestern surface mine operators and regulatory personnel.

Jan 1, 1978

By Madan M. Singh

Mining under bodies of water such as oceans, lakes, rivers, streams, ponds, and other impoundments, is not a new technique. However, in the United States it has been fairly restricted to date. In order to recover the maximum amount of the nation's reserves, in the future, it is important that further mining under such bodies of water be planned. However, it is imperative that this is not done at the expense of danger to the nation's miners or appreciable property damage. If all the bodies of water are considered large, and it is assumed that these constitute a hazard when mining in their vicinity irrespective of their size and geometry, a significant percentage of the coal reserves are likely to be sterilized. How- ever, by adopting special mining plans and procedures, it may be possible to release reserves for exploitation that would otherwise be unmineable since the surface of bodies of water below which they lie are relatively small in size. In 1977, the Bureau of Mines issued Information Circular (IC) 8741, entitled "Results of Research to Develop Guidelines for Mining Near Surface and Underground Bodies of Water". This document was based on the work done by Skelly and Loy (under Contract No. H0252083) and K. Wardell and Partners (under Contract No. H0252021). How ever, these guidelines have limited application, because these did not (1) define the bodies of water of concern, and (2) take into account the nature of the strata being encountered. These guidelines assumed that all bodies of water which would be undermined were large. For the purposes of this discussion, water bodies may be divided into three categories, depending upon their potential to cause damage to the underlying mining activity: ? catastrophic ? major ? limited Bodies of water which pose the threat of catastrophic destruction in the mine, should there be an inrush, include oceans, large lakes, big inland reservoirs, and rivers that could flood the mine complete- ly in a short period. Those with major potential for damage would include lakes and streams which might present a danger to both life and property if not considered in the mining plans. Finite bodies of water which have a volume considerably smaller than the mine volume offer a limited potential for damage in the case of an inrush. These include flowing streams which have a relatively small rate of replenishment and which can be dealt with by pumping or isolation in the event of a breakthrough into the mine. Such bodies of water may be inconvenient and cause a temporary disruption of production, and could even cause some damage to equipment, but only in rare instances would personal injuries occur. The scope of this paper is limited to the exploitation of stratified mineral deposits with overlying -face bodies of water. Vein and dessiminate ore deposits are not discussed, nor are breakthroughs into adjacent mines in the same horizon. Overlying, water-logged strata also pose a serious threat to mining, but are not treated in this discussion. Pending further investigation these strata may be considered as bodies of water with catastrophic potential for damage, with the base of the strata being considered the bottom of the water body. When the body of water is large, so that it may cause catastrophic damage to the mine in the event of an inrush, the guidelines suggested in IC 8741 are applicable. In order to develop criteria and procedures for mining under smaller bodies of water, however, it is necessary that the nature of the strata disturbance due to underground mining be examined.

Jan 1, 1981

By F. W. Maclennan

The following is a list of members who died in 1929. It is compiled from reports to the Secretary's office. Biographical sketches published in Mining and Metallurgy are indicated in the last two columns. Issue Containing Year of Date of Biography Election Name Death 1929 Page 1922 Adams, Waldo Peck................................ June 2 September 435 1903 Argall, Joseph.................................... Nov. 1 1894 Austin, Leonard S................................. Oct. 30 December 580 1905 Bailed, Melbourne.........................Dec. 12 May* 280 1890 Barringer, Daniel M.............................. Nov. 30 February* 119 1895 Batchelleb, HENRr R.............................. Jan. 15 March 167 BINGHAM Charles E............................. 1923 Bibkett, E. H..................................... Apr. 23 August 394 1913 Bibbell, Robert W................................. Jan. 18 April 214 1926 Blackmar, Don C.................................. June 4 March* 189 1920 Bosworth, T. O.................................... Jan. 18 April 214 1894 Boucher, Arthur S..........................Jan. 6 May 257 1927 Brandenthaler, R. R.............................. Dec. 14 January* 64 1889 Bretherton, S. E............................... Oct. 4 December 581 1911 Brunton, F. K................ .. Aug. 17 1924 Brush, Charles F................................. June 15 September 436 1925 Cameron, Frank N........................May 16 July 349 1891 Chambers, Robert E.............................. May July 348 1908 Clark, J. Murray................................. Dec. 3 August* 407 1883 Clymer, edward T............................. Oct. 13 December 579 1914 Cofpey, George T................................. May 8 July 349 1901 Cooke, L. H....................................... Aug. 23 November 537 1915 Cousins, George W........................ June 30 February* 118 1886 Cowles, A. H...................................... Aug. 13 September 436 1892 Crane, Theron I............................. Nov. 2 December 579 1921 Crowell, Samuel B................................ Dec. 9 January* 64 1907 Curtis, Bracey.................................... Jan. 31 April 215 1895 de Batz, Rene........ ............... January March* 189 1928 Dines, Tyson S.................................... March May 257 1909 Dufourco, R. G............................. January November 537 1920 Ericsson, Eric A.................................. Apr. 4 April* 231 1920 Erlenborn, Willi....................... Nov. 10 March* 189 1888 Farish, John B.................................... Nov. 14 December 580 1891 Ferguson, Vincent........................... Mar. 8 May 257 1917 Finucane, T. R.................................... Mar. 22 February* 118 1921 Foch, Ferdinand................................... Mar. 20 1885 Foucar, Edouard I....................... Jan. 3 March 167 1880 Francklyn, Charles G............................. Jan. 11 April 214 1925 Friedl, Arthur J.............................. July 9 September 435 1894 Fuller, James W. Jr............................... April 4 1876 . Goodale, Charles W.............................. April 11 May 258 1912 Graybill, John..........................Dec. 11 June* 320 1887 Gritzner, F. A.................................. Mar. 26 July 394 1906 Gunther, C. Godfrey.............................. Deo. 26 February* 118 1900 Haggen, Edward A.............................. Apr. 22 October 488 1895 Hall, Benjamin M........................Nov. 19 March* 189 1913 Harper, Walter S................................. Feb. 15 August 394 1921 Hawley, William S................................ Apr. 20 June 298 1920 Herd, Walter ... . . . ..... Aug. 23 September 435 1889 Holt, Marmadtjke B............................... Dec. 26 February* 118 1928 Husk, Albert D .... ................June 9 December 580 1920 Hutchinson, S. Pemberton......................... Feb. 16 April 214 1926 Ivanoff, Constantine L............................ Mar. 7 June 298 1914 Jaeger, Frederick..... .............July 28 October 488 1891 Jennings, Robert E.......................Oct. 30 December 580 1922 Johnson, Homar L................................. Sept. 17 October 483 1927 Johnson, Homer H................................. Mar. 21 May 257 1918 Jones, Jesse L ..... ...............Jan. 22 April 214 1892 Jones, Thomas J................................... June 18 October 488 1883 JoOet, Cavalier H................................. Mar. 29 May 258 1881 Kennedy, John S . . . . ... ..........Apr. 22 August 394 1881 Kennedy,- Hugh.................................... May 23 September 435 1895 Kebvin, James H................................... September October 488 1926 Kkom, L. J........................................ Dec. 24 January* 64 1905 La Follette, Harvey M........................... Sept. 19 November 538 1920 MacDonald, Alexander............................ May August 394 1904 Marengo, Paolo................................... May 257 1925 Maxwell, Allison R............................... May 7 November 538

Jan 1, 1930

By Roland de Wit

Conclusions reached in a paper by weertmanl are amplified in a mathematical and graphical way. It is shown that in a stressed crystal a straight dis-location may be in a position of unstable equilibrium with respect to helix development, and how a small perturbation may start the dislocation off on its helical path. Helix development occurs by diffision; bulk diffusion is necessary for development from a screw, but core diffision is sufficient for development from an edge dislocation. For a mixed dislocation helix development can take place by core diffusion if the helix also climbs; a diffusion reversal takes place when the slope of the helix equals the slope of the Burgers vector. The work done on a crystal by the external forces is graphically shown to be consistent with the direction of motion of the dislocation in helix development. Finally a detailed mechanism for tangle formation from helices is presented. In the course of helix development a glide situation may be reached where a segment of each helix loop lies in a slip plane in which it can expand by glide. Thus is it visualized how a helix can deteriorate into a tangle. A paper by weertmanl proposes dislocation-tangle formation by a helical dislocation mechanism. The present paper amplifies Weertman's arguments in a mathematical and graphical way and presents some further ramifications of the proposed model. The importance of dislocation tangles is discussed in Weertman's paper and the references cited therein. The mechanism of their formation is not yet understood. Weertman proposes that in cold-worked crystals conditions exist suitable for the conversion of straight dislocations into helices and subsequently into tangles. Helix formation occurs by a diffusion process and Weertman proposes that at low temperatures only core diffusion is important. The present paper necessarily repeats some of Weertman's work. Since the mathematical development and notation used are slightly different from his, this was thought necessary for clarity. GEOMETRY OF THE HELIX weertman2 has shown that the equilibrium form of any dislocation line is a helix if it is assumed that the energy per unit length (line tension) of the dis- location is constant. For the special case that no external forces act on the dislocation, the helix reduces to a straight line. We shall assume in this paper that the dislocation has a helical form even when it is not in equilibrium. The parametric equation of a helix is, see Fig. l: where a is the radius of the cylinder tangent to the helix, and 270 the pitch of the helix. The pitch is related to v, the number of turns per unit length (along the axis) by the relation: When a = 0 Eq. [I] gives a straight line in the z direction. The growth of a straight dislocation into a helix is now described mathematically by letting a vary from zero to its final value. The tangent to the helix is given by where is the path length along the helix. This gives the direction of t and since its magnitude is unity we must have So that dcp/ds = (a2 +p2)-"'. The curvature of the helix is given by where n = -(cos + sin) is a unit vector normal to the helix. Note that the vectors k and n always point toward the axis of the helix, Fig. 1. FORCE ON THE HELIX The external force which makes the straight dis-

Jan 1, 1963

By L. J. Kashar

By means of a continuum-mechanics approach, the following equation for calculuting the pressure for axisymmetric cold forging has been derived: where P is the extrusion pressure; K and n are the strength coefficient and the strain-hardening exponent of the steel, respectively; a and b are strain factors, primarily dependent on the die design; and R is the extrusion ratio. Data from compression and backward-extrusion tests performed at a crosshead speed of 0.5 in. per min on a nonwork-hardening material pure lead) and on a work-hardening material (1018 steel) indicate that the predicted extrusion pressures obtained from the derived equation agree very closely with the actual extrusion pressures. Therefore, the equation can be used to predict backward-extrusion pressures from standard mechanical-test data. The decision to cold-forge a part is based on several considerations, most of them related to economics. However, the production feasibility is usually determined by the pressure needed to extrude the part and the availability of tooling that can repeatedly sustain this pressure for over several hundred thousand cycles. A major contribution to the cold forging of steel would result from the development of steels that could be forged at reduced pressures and from the development of techniques that permit the accurate prediction of the pressure necessary to cold-forge the part. The ability to predict extrusion pressures is necessary to properly design the tooling required for the cold-forging process. Attempts to theoretically predict the cold-forging pressure, especially for work-hardening materials, have been unsuccessful in previous investigations of the cold-forging process.1-3 Many empirical relations have, therefore, been proposed relating the extrusion pressure to the chemical composition,' the extrusion ratio,2 the tensile strength,273 and hard- ness, 2 and the yield strength3 of the material being extruded. However, these empirical relations are valid only for the specific conditions used in their development, and, in some cases, the empirical relations had a scatter band greater than * 10 pct. A study of the extrusion process was undertaken with the objective of developing predictor equations for extrusion pressure by using a continuum-mechanics approach to the problem of metal flow. The derivation of such an equation is summarized in this paper, and data are presented that confirm the validity of the derived equation. DERIVATION OF THE EXTRUSION EQUATION If the extrusion of a material through a die is considered in terms of a work balance, it has been shown that it is possible to derive an equation relating the work necessary to move the material through the die to the deformation work undergone by the material. The work done, W, in moving the material is given by where F is the extrusion force, I, is the distance the material was moved, A, is the original cross-sectional area of the material, P is the extrusion pressure acting on A,, Ap is the cross-sectional area of the extrusion punch, and Pp is the punch pressure acting on Ap. The deformation work, WD, is given by where V is the volume of metal deformed, ay is the flow stress of the metal, and E is the total strain undergone in the extrusion. If there are no frictional losses, the work done, W, is equal to the deformation work, WD; therefore, Since the volume of a metal does not change during deformation,

Jan 1, 1968

By W. A. Tiller, W. C. Johnston

The solute distribution ahead of an advancing solid-liquid interface is controlled by varying the momentum boundary layer thickness in the liquid adjacent to the interface. Single pass zone-melting experiments on Pb-Sn alloys are described in which the liquid is caused to rotate in the plane of the solid-liquid interface due to the influence of a magnetic .field rotating in this plane. The variation of the diffusion boundary layer thickness with field strength is determined. DURING the freezing of an alloy a partitioning of solute occurs at the solid-liquid interface, the magnitude depending upon the difference in the solid and liquid solubilities at the interface temperature. Because the diffusion coefficient is so small in the liquid relative to normal freezing rates, either a pileup or a depletion of solute occurs at the interface when the partition coefficient, kO, is either less than or greater than unity, respectively. This solute distribution is completely characterized by i) the solute concentration in the liquid far from the interface, C) the ratio of freezing velocityf, to diffusion coefficient, D, iii) the partition coefficient, ko, and iv) the thickness of the diffusion boundary layer, 6,, at the interface (this assumes that an electric field is not applied across the interface). For a particular system, C ko and D are fixed so that the character of the solute distribution can be changed only by a variation of either f or 6c. As is well known, the freezing velocity can be changed by varying the thermal environment of the alloy. Whereas, the thickness of the solute rich layer can be changed by varying the fluid flow conditions in the liquid. It is the purpose of this paper to describe a technique for controlling the fluid flow conditions in the liquid. During most solidification experiments the driving force for fluid flow arises from either mechanical mixing, thermal convection or electromagnetic mixing. The momentum boundary layer thickness, 6m at the solid-liquid interface, within which negligible movement of the liquid occurs, has been calculated for the special case of mechanical mixing where crystals are solidified by the Czochralski technique.l,2 The calculation of 6m for the case of thermal convection has also been made.3 In some calculations it has been assumed that the momentum boundary layer thickness, 6,, may be equated to the thickness of the diffusion boundary layer 6c, at the interface. However, as may be seen from Appendix I this is only true in very special cases and in general where a is a constant of magnitude determined by the particular liquid alloy. Neither calculations nor experimental data exist for the dependence of upon electromagnetic mixing. In this paper experiments are described in which the liquid is caused to rotate in the plane of the interface due to the torque developed on it by a magnetic field, H, rotating in this plane. The variation of is controlled by the variation of H. A magnetic field rotating in a plane perpendicular to the interface has been used for the purification of aluminum. Experimentally, we determined the variation of 5C with H by using the zone-melting technique of pfann5and a theoretical relationship of Burton et a1.2 pfann5 has shown that, after the passage of one molten zone through a long cylindrical charge of uniform concentration, CO, the solute distribution, C,, in the solid is as shown in Fig. 1. The equation describing this distribution is

Jan 1, 1962

By J. E. Harris

By elimination of other possible nucleation processes, it has been demonstrated, for commercially pure magnesium and a Mg-Al alloy, that at stresses less than that necessary for triple-point cracking creep cavities are most probably nucleated heterogeneously at second-phase particles. Evidence for the existence of such particles is presented and the role of grain boundary sliding in the nucleation process is examined quantitatively. DURING creep deformation of polycrystalline magnesium and its dilute alloys in the temperature range 0.45 to 0.7 Tm (Tm = absolute melting temperature) cavities form mainly along grain boundaries oriented normal to the stress axis and failure eventually occurs by their growth and impingement.' The possibility of low-ductility creep failure of magnesium-alloy reactor cans during service has prompted a number of studies of cavity formation in these materials.2-5 In this paper divergent theories of cavity nucleation6-10 are examined critically and conclusions drawn regarding the most probable mechanism of formation of cavities in magnesium. The paper is divided into two parts. The first section consists of the review and quantitative assessment of the possible role of grain boundary sliding in the nucleation process. In the second part an experimental program is described where an electron-microscope technique has been used to determine if second-phase particles exist in as-received commercially pure magnesium and in a magnesium alloy 1) THEORETICAL REVIEW Possible nucleation sites for creep cavities are: a) high-angle grain boundary triple points,11-14 b) subboundary/grain boundary intersections,15-18 c) ledges in the grain boundary, 19-25 d) second-phase particles. 9,26-30 Each of these possibilities is now considered with special reference to deformation of magnesium. 1.1) Triple-Point Cracking. The stroh31 condition to nucleate fracture is: " > vr [1] where ss = applied shear stress, ? = surface energy, C = shear modulus, L = length of sliding interface. This relationship is independent of the nature of the interface and it is possible to calculate the stress concentration at the end of a sliding high-angle grain boundary and thus deduce, approximately,* the minimum applied stress necessary to nucleate triple-point cracking. Such theoretical predictions for Nimonic alloys have been made by McLean13 and found to agree quite closely with the minimum stress determined by experiment. Stresses capable of nucleating triple-point cracking should also be sufficient to nucleate fracture at a grain boundary due to intragranular slip, for the lengths of sliding interface in the two cases are approximately equal. Thus the minimum stress for triple-point cracking should represent a demarcation between failure due to general "cracking" and that due to cavitation. Hauser, Landon, and Dorn32 have demonstrated that for polycrystalline magnesium in the temperature range 78° to 298°K the fracture stress is inversely proportional to the square root of the mean grain diameter. As pointed out by stroh31 this result indicates that Eq. [I] describes the fracture characteristics of magnesium at low temperatures. Heal3 has reported that for rapid deformation (112 pct hr-1) of polycrystalline magnesium the transition from cracking to cavitation occurs in the temperature range 175" to 275°C. He also presented the fracture-stress values for the same material tested under similar conditions. From Eq. [I] taking G = 1.3 x 10" dynes , ? = 600 ergs . cm-2,33 and L = 0.017 cm, a Stroh-McLean13,31 tensile stress of -266 kg. cm-2 is calculated, corresponding to a transition temperature of -190°C; this is in

Jan 1, 1965

By W. A. Backofen, M. L. Ebner

SINCE the early work of Gough with Hanson and Wright,l-3 the study of fatigue has been characterized by experiments on single crystals only in recent times.9-10 Now, increasing attention is given to this aspect of fatigue research for the insight that it may provide into details of mechanism. The investigations have concentrated to a large extent on the development of deformation markings on fatigued crystals, and have shown the cracks to originate in slip bands possibly preceded or accompanied by slip-band extrusions. Experiments of special interest to the present work were conducted by paterson5 on copper crystals and involved both metallographic examination and measurement of change in flow stress. Crystals were cycled in alternating tens ion-compress ion with a constant plastic shear-strain amplitude of approximately 0.8 pct, and were particularly revealing for their demonstration of hardening with accumulated strain similar to that in unidirectional straining, through an easy-glide stage I followed by a stage II of rapid hardening; deformation was not continued beyond 40 cycles, however, so that the eventual course of the hardening curve could not be decided. For the conditions used by Paterson, surface slip markings were similar to those observed in unidirectional straining, but there were no X-ray asterisms and no deformation bands on the surface. In current thinking about the nature of fracture in fatigue, two views relative to mechanism are generally acknowledged, with the reservation that both could apply simultaneously in some measure. As one possibility, fracturing across a slip plane is regarded as a result of loss of cohesion from the creation of many point defects by dislocation movement under the cyclic loading.'' On the other hand, fracture has also been taken to follow as a consequence only of the geometry of slip at a free surface, consisting of offsets and crevices which eventually become fatigue cracks.12, l3 The work of McCammon and Rosenbergl4 showing fatigue in polycrystals at 4.2oK makes clear that any long-range diffusion of point defects is unnecessary, yet studies such as those of Forsyth and stubington15 show that accelerated diffusion may be a characteristic of deformation under fatigue loading. Interest in obtaining data of possible use for resolving such questions led to the experiments described below. Copper crystals were conveniently loaded in alternating four-point bending at constant deflection, a test condition shown to approximate a constant plastic strain amplitude for a wide range of axial orientations. The method of testing was not readily adapted to extensive study of temperature effects, but an investigation of the geometry of crack formation could simply be made by orienting,the slip direction at different angles to the surface. EXPERIMENTAL PROCEDURES Single crystals were grown by a modified Bridgman method in a stationary gradient furnace under an atmosphere of purified dry nitrogen. Purity of as-grown crystals was 99.999 pct as determined by spectrographic analysis plus vacuum fusion and gravimetric analyses for oxygen and sulfur, respectively. Crystals were of reasonable perfection as evidenced by a critical resolved shear stress at 10 deg from (110) of 60 g per sq mm and half-width of a (400) X-ray diffraction line from one crystal of

Jan 1, 1960

By Xianzhong Hu, Qingbo Yu, Qin Qin

"Oxy-fuel combustion is one promising option among different CO2 capture technologies. The object of this work is to show the effect of CO2 dilution on the laminar flame speed, flame structure and reaction pathway for methane oxidation at O2/CO2 atmosphere. In this paper, kinetic simulations were made using CHEMKIN with GRI-Mech 3.0 mechanism, which was validated by experimental data of laminar flame speeds. The reaction pathway of CH4/O2/CO2 mixture was derived from reaction flow analysis. Results showed that CO2 decreased the flame speeds, lowed concentration of major intermediates and increased flame thickness. The major chemical pathways of CH4 to CO2 have not changed at O2/CO2 atmosphere comparing with O2/N2 atmosphere, but contributions to major intermediates formation from each related elementary reaction have changed a lot. In addition, the secondary chemical pathways of CH3 oxidation became different at O2/CO2 atmosphere.IntroductionNowadays, as the global economy is developing fast and the pollutant levels (especially global warming gases) are increasing, a lot of researches are focusing on new technologies which can cut greenhouse gas emissions [1]. New technologies have been developed, such as high Mild combustion, flameless combustion and Integrated Gasification Combined Cycle (IGCC) technology. Oxy-fuel combustion using carbon dioxide as dilution of oxygen is considered as a promising approach for reducing CO2 emission. The properties of CO2 in flue gas are different from N2 in air. On one side, the transport, thermal and radiative properties of CO2 are different from N2 in air. On the other side, CO2 is not inert but directly participates in chemical reactions and recent researches [2-5] also present the chemical effects of CO2 on oxy-combustion. For these differences above, the characteristics of oxy-combustion are not alike with traditional combustion. It is necessary to develop the investigation of oxy-combustion.In all combustion processes, laminar flame speed is an important parameter, which links to combustion behavior. Analysis of this parameter can provide a better understanding of the effect of CO2 dilution in the process of oxy-combustion. Several investigations of laminar flame speeds of oxy-fuel flame have been improved, including experimental measurements and numerical computations. A.A.Konnov et al. measured the flame speeds of oxy-methane and oxy-ethane mixture using the heat flux method with a flat burner [6-8]. V. Ratna Kishore et al. used heat flux method to measure the adiabatic laminar flame speeds of hydrogen [9] and H2/CO mixture [10]. Adiabatic laminar flame speeds of hydrogen in O2/CO2 atmosphere are much smaller than in O2/N2 or in O2/Ar. The burning velocity of H2/CO mixture decreased with increase in CO2 content in mixture. Jeongseog Oh et al. [11] used Bunsen flame to measure the laminar flame speeds of premixed oxy-methane mixture with angle methods in atmospheric condition (300k, 1atm). The results were SL=29.1cm/s and SL=29.5cm/s respectively derived from oxy-methane from CH* images and Schlieren photos."

Jan 1, 2014

By Clark L. Corey, Bogdan Lisowsky

Electrical resistivity, X-ray line positions, degree of order, and microstructures have been investigated for Ni-A1 alloys near the Ni3Al composition. The results indicate that Ni3Al undergoes disordering between 1250oC and the mp; also supersaturated y, produced in y + y' alloys by gas quenching, is extensively ordered prior to y' (Ni3Al) phase separation. The latter reaction, if not both, must follow the cooperative phenomena model. Shifts in line position as a function of quench-indulced super satulration are shown. Lattice pavarneters of equilibriunz and nonequilibrium phases are given. THE objective of this investigation was to determine the mechanism and kinetics of the reactions which occur in Ni3A1 and hypostoichiometric Ni3A1 alloys. The results are of particular significance relative to pre-precipitation phenomena. Early work on Ni-A1 binary compositions near the Ni3A1 composition (13.6 wt pct Al) has been adequately reviewed by Hanssen,1 whose phase diagram is presented in Fig. 1. williams2 verified the suggestion made by Taylor and Floyd3 and others4 that y', Ni3A1, may precipitate coherently on (100) planes of the y phase (see Fig. 1 for phase nomenclature); he further suggested a two-step precipitation mechanism consisting initially of probably short-range ordering followed by actual precipitation of y'. The precipitation stage in low-aluminum hypo-eutectic compositions has been studied by a number of investigators;*"7 short-range order or "zone" formation have been suggested as explanations for various diffuse X-ray scattering effects. The ordering reaction in Ni3A1 alloys is characterized by a) a very rapid rate of ordering in nonstoi-chiometric compositions, b) absence of a disordered phase region for the stoichiometric compound, y', or Ni3A1, c) precipitation phenomena as a dominant aspect in hypoeutectic alloys, and d) an increase in elec-trical resistivity on ordering.' The latter two factors were found by Newkirk to characterize ordering in Co-Pt, for which an extensive and detailed analysis was made.' An electrical-resistivity increase on ordering has also been reported for cuAu3.10 PROCEDURE Ni-A1 alloys in the form of 1300-g chill-cast ingots were prepared by vacuum melting electrolytic nickel and high-purity (99.99 pct) aluminum. Very small heats prepared from high-purity carbonyl nickel, supplied through the courtesy of Dr. Goff of the Naval Research Laboratory, were used for comparison with the larger ingots and no differences were noted. The cast material was homogenized in purified argon for a minimum of 72 hr at 1200°C and water-quenched from 1350°C. Compositions studied are given in Table I. The 10 pct A1 alloy was both hot- and cold-rolled. X-ray superlattice line intensities and peak positions were obtained from -270 mesh filings, using CuK, radiation with a LiF crystal monochromat adjusted to receive both 01 and a2. One-degree slits and a 1/8 deg per min scan rate were used to record line traces. A standard tungsten powder, obtained through the courtesy of G.E. Research Laboratory, was used to calibrate line positions.

Jan 1, 1968

By Jim D. Mote, Eugenia M. Howard, Willis L. Barmore, John E. Dorn

The effect of stress and temperature on the creep rate of Ag-33 at. pct A1 and Ag-33 at. pct Al-1 at. pct Zn for prismatic slip was determined over a range of temperatures from 440º to 700°K. The alloys exhibited a brief initial inverted transient creep over which the creep rate increased slightly to a steady-state value which depended only on the stress and temperature, regardless of the previous strain history. The steady-state creep rates were found to be given by: where the activation energy is independent of the stress and strain and agrees fairly well with that estimated for diffusion. The observed creep behavior was shown to be in qualitative agreement with a model based on the stress-induced local disordering in the vicinity of a dislocation. OVER the past fifteen years, considerable progress has been made in correlating the experimen-entally observed plastic behavior of pure metals over each temperature range with one or more of a series of theoretical strain-rate controlling dislocation mechanisms. Despite the current interest in the plastic behavior of intermediate phases and inter metallic compounds, however, very little has been done in identifying the operative dislocation mechanisms in these materials. It is anticipated that the plastic behavior of intermediate phases can be rationalized in terms of the same dislocation mechanisms that are appropriate to pure metals and a solid solutions, provided the additional complexities arising from crystal structure, bonding, long-range order, short-range order, and defect lattices are taken into account. As part of a more extensive investigation on the plastic behavior of intermediate phases, Mote, Tanaka, and Dorn1 recently reported on the effect of temperature on the yield strength for prismatic slip in the hexagonal Ag-33 at. pct A1 intermediate phase as recorded in Fig. 1: Over Region I the deformation was observed to take place as a result of the thermally-activated Peierls mechanism, and over Region II the athermal disordering mechanism applied. The experimental data at that time, however, were not sufficiently extensive to identify the thermally activated mechanisms operative over Region III. It is the purpose of this investigation to attempt to identify the mechanisms that are operative over this high-temperature region. It will be shown that the same mechanism operates all over Region Ill, although the activation energy for this mechanism is in good agreement with that for diffusion, the process cannot be attributed either to the motion of jogged screw dislocations or the climb of edge dislocations. The results, however, are in good qualitative agreement with stress-induced disordering localized at dislocations. EXPERIMENTAL PROCEDURE The alloys studied in this investigation were prepared from silver, aluminum, and zinc (each of purity greater than 99.995 pet) by induction melting under an argon atmosphere and chill casting in a copper mold. Rectangular bar single crystal specimens (1/4 by 1/8 in. in cross-sect ion) of Ag-33 at. pct Al were grown in vacuum by a modified Bridgman technique. Similar Ag-33 at. pct Zn specimens were produced by an identical technique employing an argon atmosphere. The specimens were so oriented that the angle 90-Xo between the tensile axis and the normal to the prismatic plane and the angle A, between the tensile axis and the slip direction were 45 ± 1 deg. The effect of a change in temperature on the creep rate and a change in stress on the creep rate were determined over the temperature range

Jan 1, 1963

By Mark J. Klein

An internal-friction profile induced by nitrogen in Cr-35 at. pet Re was studied as a function of nitrogen concentration and heat treatment. From these studies, the solubility of nitrogen in this alloy was deduced. The equilibrium solubility of nitrogen in chromium does not appear to be substantially changed by the addition of 35 at. pet Re to chromium. However the precipitation kinetics of nitrogen from solution are significantly decreased. IT has been found that rhenium has a beneficial effect on the low-temperature ductility of Group VI-A metals. As part of a study to investigate the mechanism of this increased ductility, it was of interest to compare the concentration of interstitials retained in solution in one of these ductile bee alloys with the concentration retained in solution in its base metal. In this regard, a study of the solubility of nitrogen in ductile Cr-35 at. pet Re was selected for this phase of the investigation because the solubility of nitrogen in the base metal, chromium, has been determined over a fairly large temperature span.1-3 In addition, nitrogen is known to have an embrittling effect on chromium,4 but does not appear to have a comparable embrittling effect on Cr-35 at. pet Re.5 The experimental method selected for this study was the measurement of internal friction. This technique allows interstitial-concentration measurements to be made in bee metals that are sensitive to small changes in the concentration of a particular interstitial, even when other interstitials are present. When subjected to a cyclic stress, interstitial atoms in solid solution give rise to an internal-friction peak whose origin is the stress-induced ordering of interstitials into preferred lat- tice sites. These internal-friction peaks signify a dissipation of elastic energy which is greatest when the period of oscillation is comparable with the relaxation time for the attainment of an equilibrium interstitial distribution. For dilute interstitial binary alloys, the height of the internal-friction peak is proportional to the concentration of interstitial atoms in solid solutions. snoek8 has shown that the internal friction, Q-1, may be expressed by the equation where A is the relaxation strength, ? is the angular frequency, and 7 is the relaxation time. The relaxation strength, A, is proportional to the concentration of interstitials taking part in the relaxation if no interaction occurs between solute atoms. Thus values of peak height, which can be determined as a function of temperature, give a measure of the relative concentration of interstitial atoms in solution at various temperatures. If equilibrium has been attained, the peak heights will be proportional to the interstitial solubilities. The constant that relates A to the interstitial concentration can be determined experimentally from peak-height measurements as a function of concentration for specimens of constant crystallographic orientation. When this has been done, Q-1 measurements can be used to determine solubilities according to the equation C - KQ-1 where C is the soluble interstitial concentration and K is a constant that can be experimentally determined. It is possible, however, that lattice imperfections may bind interstitials to certain lattice sites. Internal friction does not measure these interstitial~, but does measure interstitials which are free to move in the normal lattice sites when subjected to an applied oscillating stress. On alloying the pure metal with a substitutional

Jan 1, 1965

By John J. Simmons

There is considerable interest in the iron and steel industry in the development of technologies that can reduce carbon dioxide emissions, as the iron and steel industry is the third largest producer of carbon emissions in the world, after automobiles and the power generation industry. The E-Iron Iron Nugget Making Process, invented by Dr. Komar Kawatra and Dr. Tim Eisele, provided a major breakthrough in the iron making process as the process uses biomass in the form of hard wood or soft wood as the reductant in the place of coal. Dr. Kawatra recognized that biomass is rich in both carbon monoxide and hydrogen, two important reducing agents. This resulted in initial experiments conducted by Dr. Kawatra and Dr. Eisele in the laboratories of the Chemical Engineering Department of Michigan Technological University (Michigan Tech) that prepared self-reducing pellets with wheat flour as the reductant and limestone/dolomite as the flux. Initial results were encouraging so they extended their work to include hardwood, which is in huge supply in the forests of Michigan and Minnesota, as the reductant. Dr. Kawatra has also applied the E-Iron Technology to other iron bearing minerals to produce pig iron grade iron nuggets. He used the E-Iron Process to convert red mud, a high iron waste product of the aluminum industry, into pig iron nuggets with the rare earth elements reporting to the slag. This work resulted in a patent with Dr. Kawatra and Dr. Michael Archambeau named as the inventors. Dr. Kawatra then extracted the rare earth elements as a separate high value bi-product using oxalic acid, which he produced from carbon dioxide. Dr. Kawatra also conducted a test program that produced pig iron nuggets from New Zealand beach sands. The vanadium in the beach sands was retained in the iron nuggets, recoverable as a valuable bi- product in a separate process. The titanium dioxide in the beach sands reported to the slag. This demonstrates the potential use of the E-Iron Process to convert the large ilmenite reserves in Northern Minnesota into pig iron and titanium dioxide. Another use of the E-Iron Process discovered by Dr. Kawatra was its ability to produce iron nuggets from manganiferous ores. Dr. Kawatra was able to direct the manganese into either the metal or slag by adjusting process temperature. The above discoveries are included in several patents that were assigned to Carbontec Energy. This presentation will only address the use of the E-Iron Process to produce pig iron nuggets from steel mill waste and un-pelletized, mid-grade iron ore. Carbontec Energy Corporation agreed to sponsor additional research directed at improving this new technology that would later be trade named, The E-Iron Process. Carbontec Energy conducted further work that proved that the E-Iron technology could also convert steel mill waste into metallurgically acceptable iron nuggets. Carbontec also collaborated with the Lawrence Livermore National Laboratory and Purdue University Northwest to conduct a Department of Energy (DOE) sponsored program that would use their high performance computer capability to both verify and model the E-Iron Process. Purdue University Northwest reported that the E-Iron Process can produce pig iron nuggets with 85% less carbon emissions than a blast furnace. The E-Iron Process is similar to the ITmK3 process, developed by Kobe Steel, that also produces iron nuggets but uses coal as the reductant.

Feb 1, 2025

By T Kojovic

"The index A*b, determined from the Julius Kruttschnitt Mineral Research Centre (JKMRC) drop weight tester (DWT) results, has become well-known in the mining industry as a reliable measure of ore hardness in impact or crushing, having the inferred units of t/kWh. This index evolved from a long history of comminution studies at the JKMRC, including the development of the now superseded JKTech twin-pendulum, the industry standard JK Drop Weight Tester (JKDWT) and the faster alterative JKMRC rotary breakage tester (JKRBT). A reduced version of the DWT, the semi-autogenous grinding mill comminution test (SMC test®), also provides estimates of the A*b though the testing of only one particle size at five specific energies. Experimental determination of A*b using the above devices/tests requires samples, with sufficient mass, to be sent to designated laboratories, registered to conduct such tests. This protocol is fine for bankable and commercial testing, where material characterisation data is used in conjunction with machine specific data in modelling and simulation, and power based calculations. However, for comparative testing, where results are required at the time of sampling (eg for plant troubleshooting or spot-surveys), or for rapid ranking of geometallurgical samples, these existing tests simply do not provide a timely, viable and cost-effective solution. In order to address this critical void in the realm of comparative testing, a new device has been developed, exploiting a central feature of single particle impact testing – that the A*b can be reliably estimated using one precise low energy test. This device, for which a patent is currently pending, called the HIT (Hardness Index Tester), has been precision engineered to allow users to break narrowly sized fragments at a set specific energy, in a safe and easy manner. The manufacturing materials have been chosen carefully to ensure compatibility with the intended use and portability to any site, whilst minimising cost. The device is supplied with a novel quality assurance / quality control (QA/QC) feature, to ensure that the targeted potential energy is consistently delivered. The product from the HIT breakage test can be quickly sized using a single screen to determine the mass per cent undersize, from which the A*b index is calculated directly at site – eliminating the need to send the raw data off-site for analysis. On presized fragments, a single test using 20 fragments takes less than five minutes to complete, and the results are immediately available. The HIT is currently being tested at a major mining laboratory in the US and the results to date have proved it can deliver on all fronts. This paper describes the new prototype device, which offers users a low cost in-house mechanism for estimating the A*b at any mine site, with potential applications on fragments up to 25 mm from drill core, blast chip rejects or the grinding circuit feed. It also outlines the results of field tests carried out by an objective third party mining group and compares the results with those from current industry standard tests conducted on the same material.Citation:Kojovic, T, 2016. HIT - a Portable Field Device for Rapid Testing at Site, in Proceedings 13th AusIMM Mill Operators’ Conference 2016, pp 9–16 (The Australasian Institute of Mining and Metallurgy: Melbourne)."

Oct 10, 2016

By I. R. Kramer

The delay time for single crystals of iron, zinc, and pre-strained aluminum was measured under conditions of high-speed deformation. The delay time of aluminum was found to be affected by the orientation of the crystal. Under the dynamic conditions employed, tile calculated "activated volume" term appears to be very small for the three metals studied. IN a previous paper,1 it was shown that under dynamic testing conditions at 78° K, there was a delayed yield in single crystals of aluminum and copper which had been cold-worked prior to testing. The experimental results also indicated that the delay-time effect was not due to the interaction of dislocations with impurity atoms or vacancies, but rather to a dislocation-dislocation interaction since aluminum which has not been prestrained shows no delay-time phenomenon. In this paper, some data are reported on the effects of orientation on the delay time of aluminum single crystals; the delay time of zinc single crystals saturated with nitrogen and the delay time of iron single crystals also were measured. EXPERIMENTAL PROCEDURE Aluminum (99.997 pct) and zinc (99.999 pct) crystals, 1/2 in. in diam and 8 in. long were grown by the Bridgman technique. The iron single crystal was supplied by Dr. R. Maddin. The zinc crystals containing nitrogen were prepared by melting high-purity zinc in a graphite crucible and the nitrogen was introduced by additions of ammonium chloride. The liquid zinc was then drawn up into glass tubes of 1/2 in. diam and allowed to solidify. The rods were cooled to room temperature before the transfer to the mold for conversion into single crystals. The orientation of all specimens was determined by the Laue' back-reflection technique, Fig. 2 and Table I. To prepare for the delay-time measurements, the crystals were cut with a fine jewelers saw and the ends carefully machined. The specimens then were electrolytic ally polished to remove the cold-worked metal, mounted in a special fixture and hand-lapped. As a final step, the crystals were annealed in evacuated tubes. Specimens of aluminum crystals, Nos. 24 and 26, were compressed 0.5 and 2 pct, respectively, in the direction of the later impact. The zinc specimens were aged at 50°C for 15 min, In the case of iron, only two specimens could be obtained. To obtain sufficient data, it was necessary to resort to a reaging treatment. After each delay-time test, the specimens were aged at 65° C for 100 min. This treatment, according to Clark,2 is sufficient to restore the original delay time. It is believed that although the specimens were plastically deformed approximately 10-4 pct, after each test the delay-time values are reliable. The delay-time measurements were made in an apparatus similar to that described previously.' Single crystals approximately 1 in. long and having a diameter of 1/2 in. were placed in a pendulum which consisted of a bar 1/2 in. in diam and 8 ft long, designed with a crystal holder to accommodate the specimen at low temperatures. The crystal was held at a position 2 ft from the front part of the pendulum. This portion of the apparatus was supported on a fine molybdenum wire. A projectile bar of the same diameter and length comprised the other portion of the apparatus. This bar was supported on

Jan 1, 1962

By C. S. Samis, J. C. Mitchell

Sodion activity data have been obtained for the Na-A1 system for sodium concentrations in aluminum between NN, = 25 . 10-and 300 x 10'. This concentration intert-a2 encompasses those sodium concentrations nor1)zally encountered in aluminum reduction cell operation. The sodium activity data were measured by a two-phase equilibration of sodium between a lead and an aluminam rich phase. The Na-Pb system was emplojed as a reference for the determination of the actiz7itj 0.f sodium. The activity of sodiin in aluminuwz exhibits large positi2.e deviations frorn ideality and a strong dependence, on concentration below NN, = 100 x 1K6. The inteated data indicate that the activity of aluminium is equal to its mole fraction over the range of miscibility of the Na-A1 sgstem. The data on the Na-A1 binary have been combined with experimental data on the cryolite-A1 system at 1025°C to obtain the activities of Na F and AlF, in liquid cryolite oz2er the cornposition 1-ange encountered in aluminum cell operation. The results are in agreement with activities calculated from the NaF-AlF, phase diagram. ALUMINUM is produced commercially by the electrolytic reduction of aluminum oxide dissolved in molten cryolite. There is considerable controversy over the actual mechanism of the discharge processes, but sodium potentials in the systems appear to play an important role in the reduction mechanism. The metal-salt reaction involving NaF and aluminum to produce A1F3 and sodium is responsible for the presence of sodium in the aluminum metal phase. Additional sodium may be produced from the simultaneous deposition of sodium and aluminum at the cathode as a result of a cathodic overvoltage. runert' has suggested that sodium gas at atmospheric pressure would be the result of reaction between cryolite and aluminum at 1000°C. Subsequent investigations have shown that production of sodium at atmospheric pressure does not occur. These later investigations can be conveniently separated into two categories. The first involves sodium, aluminum, and cryolite, in which the sodium activity was measured. The second includes a lead phase for sodium activity measurement. Jander and errmann,' studied the reaction: at 1090°C in alumina crucibles under an atmosphere of dry, oxygen-free nitrogen. The equilibrium constant for the reaction was expressed in terms of concentration. Their results in Fig. 1 show an increase in sodium content of the aluminum with increasing NaF/A1F3 weight ratio. Pearson and waddington3 performed similar experiments at 1000°C. Their results are included in Fig. 1. Dewing and ollinshead have measured the sodium content of aluminum as a function of the NF/AF, ratio at 1025"C, Fig. 1, and ewin' measured the sodium content of the aluminum in a system involving the solid reactants NaF, cryolite, and liquid aluminum. This reaction is different from those discussed above. Two solid phases, NaF and Na3A1F6, and a liquid aluminum phase were present. The sodium concentration in the aluminum was determined by the temperature. Dewing's measurements covered the temperature range 679" to 896°C. Omo, Matsushima, and 1t02" investigated the equilibrium 6NaF ential thermal analysis. Stokes and Frank6 determined sodium activities by

Jan 1, 1970

By R. D. Young, E. A. Harre

Nitrogen exists in two broad categories commonly designated as elemental nitrogen and fixed nitrogen. Elemental nitrogen is found in nature as a diatomic molecule and constitutes about 78%, by volume, of the earth's atmosphere. Fixed nitrogen is a common term used for nitrogen that is chemically bound to other elements. Although elemental nitrogen is very abundant, nonleguminous plants cannot utilize nitrogen unless it is in the chemically combined form. The content of this chapter is confined to fixed nitrogen since the uses and properties of elemental nitrogen and the fixed nitrogen compounds are different. Certain bacteria associated with legumes are able to fix atmospheric nitrogen, but not to the extent required by succeeding nonleguminous crops. The need to use fertilizers containing fixed nitrogen compounds to grow food crops has been recognized for hundreds of years, and from antiquity farmers have used organic wastes such as animal manures, vegetable wastes, sewage sludge, fish scraps, etc., containing these compounds for fertilizers. These organic materials have a low nitrogen content, however, and cannot support large-scale food production. Known commercial-size deposits of natural nitrates containing much higher nitrogen content exist only in Chile, and the possibility of the discovery of other significant deposits is remote. The Chilean deposits were the major source of nitrogenous fertilizers until the Haber-Bosch process for the production of synthetic ammonia was developed. In the Haber-Bosch process, nitrogen reacts with hydrogen at high temperatures and pressures and in the presence of a catalyst to form ammonia. Modern plants use air, water, and natural gas or other hydrogen sources to produce the ammonia that is now the key building block of the fixed nitrogen industry. Several major problems that developed in the 1970s continue with the nitrogen industry into the 1980s. These are (1) the continuing sharp escalation in the cost of feedstocks for ammonia production as well as all energy sources, and (2) the very steep increase in the investment cost of production facilities. More than 75% of the fixed nitrogen produced domestically is consumed as fertilizers. The types of materials used as fertilizers have changed significantly during the past years. There has been a trend from the use of low-analysis fertilizers to types which contain high percentages of plant nutrients. The average plant nutrient content (N + P2O5 + K2O) of fertilizers used in the United States has increased from 18% in 1930 to 32% in 1960 and today approaches an average of 44%. In the case of nitrogen fertilizers, the low analysis natural nitrates (16% N) gave way to ammonium sulfate (20% N) after about 1920. After World War II, ammonium nitrate (33.5% N) previously used in munitions, became established as the leading nitrogen fertilizer. Urea use as a fertilizer (46% N) began in the 1960s, and in the past 10 years has become the world's leading solid nitrogen fertilizer. Nitrogen Materials Anhydrous Ammonia-Anhydrous ammonia is a colorless gas at ambient temperatures and pressures and has a sharp pungent odor. Anhydrous ammonia is stored and transported as a liquid in high-pressure cylinders and tanks at ambient temperature, and in low-pressure tanks, barges, and ships at -28°F. Anhydrous ammonia is applied directly at 4 to 7-in. depths in the soil as a high-pressure liquid by special applicators, after which it expands to a gas that dissolves in the soil moisture or is absorbed on soil colloids. In other systems ammonia is metered into irrigation water and is thereby introduced into the soil.

Jan 1, 1983

By D. S. Yang, Y. D. Wang, G. Watanabe, A. P. Notaro

"Deep soil mixing (DSM) walls reinforced with steel soldier piles were installed for excavation support and groundwater control for the construction of a grade separation structure at the intersection of Warren Avenue and the Union Pacific Railroad in Fremont, California to mitigate traffic conditions during the busy commute hours and for the extension of San Francisco Bay Area Rapid Transit (BART) system. Instead of treating the DSM wall as a temporary structure, the reinforced DSM walls were integrated in the design of the permanent structures including retaining walls, and the abutments and pier of the railroad and BART bridges. The integration of steel members inside the DSM wall as part of the permanent structures is an innovation which facilitates the use of top down construction techniques to speed construction and utilizes the steel soldier piles that would have been disregarded after the use as a temporary shoring wall. This paper will review the background of the project and DSM method and present the design and construction of the DSM wall composite structure for this grade separation project.INTRODUCTIONUsing the concept of mixed-in-place piles initiated in the 1950’s in the United States, Japanese contractors in the late 1960’s and early 1970’s installed single mixed-in-place soil-cement piles along a row to create a wall for excavation support and groundwater control. Due to the high groundwater conditions in most of the coastal cities in Japan, wall continuity and uniformity became the requirements of soil-cement walls for reliable groundwater control to minimize ground movement and impacts to adjacent buildings and facilities. In late 1970s, Seiko Kogyo, Co. Ltd. developed the Soil Mix Wall (SMW) method, which uses a triple auger mixing tool to improve the uniformity and continuity of soilcement wall. The SMW method is considered a wall installation method in contrast to the Cement Deep Mixing (CDM) method, also developed in Japan, for ground stabilization in soft soils, especially in the marine environment. In the U.S., generic terms like DMM (Deep Mixing Method), DSM (Deep Soil Mixing) or CDSM (Cement Deep Soil Mixing) are used by various government agencies and private organizations to describe the soil mixing technologies without differentiating the origin, procedure and application. The method used in this project is called DSM method by the project owner, Santa Clara Valley Transportation Authority (VTA), to produce a single continuous soil-cement wall for use as shoring wall and composite structure. The mixing tool for DSM wall installation can vary from 3 to 6 mixing augers. Due to the need to install steel soldier piles into the DSM wall for resisting lateral forces, high water cement ratio slurry with bentonite is used to produce low strength soil-cement, generally in the range of 70 to 100 psi (500 to 690 kPa), with high fluidity to facilitate the installation of steel soldier piles. To maintain the wall continuity for groundwater control, overlapping of one full DSM column is made between neighboring elements using two types of installation procedures as illustrated in Figures 1a and 1b. Steel soldier piles, as shown in Figure 2, are used to reinforce the soil-cement wall to resist the lateral shear force and bending moment induced during the excavation."

Jan 1, 2015