Reducing the degree of oxidation of low-carbon steel

Krivorozhstal Combine. Translated from Metallurg, No. 8, p. 37, August, 1989.     

Equilibrium Phase Diagrams for the Dy ― Al ― B and Er ― Al ― B Systems

X-ray diffraction has been applied to component interactions in order to construct isothermal sections of the phase diagrams for ternary systems Dy Al B at 600°C (region above 50 at.% Al) and 800°C (region below 50 at.% Al) and Er Al B at 800°C. It is confirmed that there are ternary borides: DyAl₃B> _x (structure type BaPb₃, space group R , a = 0.6156(3) nm, and c = 2.109(1) nm), DyAlB₁₄ (type MgAlB₁₄, Imma, a = 0.5819(2) nm, b = 1.0380(3) nm, c = 0.8176(5) nm), ErAlB₄ (type YCrB₄, Pbam, a = 0.59258(7) nm, b = 1.1515(2) nm, c = 0.35340(6) nm), and ErAlB₁₄ (type MgAlB₁₄, Imma, a = 0.5819(1) nm, b = 1.0401(2) nm, c = 0.8189(1) nm). The {Dy, Er} Al B systems lack signs of solid solutions based on binary and ternary compounds.     

Role of polymorphic transformation in the hot pressing of alumina

Conclusions At temperatures of up to 1400°C hot pressing under isothermal conditions enables higher (by 5%) densities to be obtained in compacts from-Al₂O₃ powders than in compacts from-Al₂O₃ powders. With rise in temperature the difference between the densification rates of- and -Al₂O₃ powders diminshes. The difference in density between- and-Al₂O₃ compacts is preserved up to a temperature of 1600°C in hot pressing at a constant heating rate of 5 deg C/min. In the temperature range 1300–18QO°C the energy of activation for hot pressing Al₂O₃ varies from 50 to 75 kcal/g· form. Estimated values of viscosity and effective diffusion coefficient are in accord with values determined in creep and sintering experiments.Translated from Poroshkovaya Metallurgiya, No. 9(165), pp. 76–80, September, 1976.     

Improving the effectiveness of reheatings conducted to electrically refine metal

Dneprospetsstal Plant. Translated from Metallurg, No. 4, pp. 34–35, April, 1990.     

The Cr ― Ta ― C Melting Diagram in the (Cr) ― (Ta) ― (TaC) Region

It has been found from an experimental study on alloys in the Cr Ta C system in the region of the (Cr) (Ta) (TaC) subsystem at melting (crystallization) temperatures that there are two nonvariant four-phase equilibria of congruent type: L_E (Ta) + (Ta₂C) + (TaCr₂) at 1935°C and L_E (TaC) + (Cr) + (TaCr₂) at 1675°C; there is also one four-phase nonvariant equilibrium of incongruent transition type: L_U + (Ta₂C) (TaC) + (TaCr₂) at 1943°C, and two nonvariant three-phase equilibria of congruent type: L_e (Ta₂C) + (TaCr₂) at 1960°C and L_e (TaC) + (Cr) at 1695°C. The phase diagram for the subsystem has been constructed in the form of projections of the solidus and liquidus surfaces, and a melting diagram has been constructed.     

Static and cyclic strength of a uniaxially reinforced material in bending (part 1)

Conclusions A built-up degree of damage =f/f_o 8% has virtually no effect on the static strength of an AD33-B CM operating in air. The arbitrary fatigue limit corresponding to this volume of damage on a basis of 2–10⁵ cycles may be taken to be equal to (0.5–0.55) _t. Accelerated buildup of damage on a basis of 2–10⁵ cycles is observed at _a > (0.5–0.55) _t.Translated from Poroshkovaya Metallurgiya, No. 7 (259), pp. 65–67, July, 1984.     

The phase equilibria in W-Fe-Co-Ni system alloys. II. Isotherms of the solidus and the phase composition of alloys containing 10 and 20% (Fe + Co + Ni) at temperatures above 1200°C

Conclusions The isostructural intermediate -phases Fe₇W₅ and Co₇W₆ in the W—Fe—Co system form a continuous series of -solid solutions. In the 1640–1630°C range the L + (Fe₇W₆) peritectic equilibrium in this system changes to a similar L + (Co₇W₆) equilibrium, where is the tungsten-base boundary solution.In the W-Fe-Co-Ni polythermal tetrahedron in the 1470–1460°C range conversion of the L +(Fe₇W₆)+, peritectic equilibrium into the similar L + (Co₇W₆) + , where is the nickel-, -ironcobalt-base boundary solution, occurs.Upon completion of crystallization at 1400°C, the W-Fe-Co system alloys with 10–20% (Fe + Co) have a + phase composition, while the W-Fe-Co-Ni system alloys with 10–20% (Fe + Co + Ni) accordingly have + , + + or + . At temperatures below 1215°C in alloys rich in iron, FeW may be formed instead of -phase and therefore the alloys may have an + FeW, + + FeW, + + + FeW and + + FeW phase composition.Translated from Poroshkovaya Metallurgiya, No. 5(281), pp. 86–89, May, 1986.     

Polythermal Ti-∑MoNi (9∶1) section of the ternary system Ti-Mo-Ni

Conclusions As a result of an investigation of the titanium corner of the system Ti-Mo-Ni, a partial phase diagram of the Ti + (0–55%) MoNi (91) section was constructed. The solubility of Mo and Ni in-Ti is 0.65% MoNi. Thea + phase field is bounded by MoNi contents of 0.65 and 17% at 600°C, 0.5 and 14% at 700°C, and 0.4 and 6% at 800°C. At MoNi contents of 55% and higher, there is a + Ti₂Ni field, the existence of which was confirmed by x-ray structural studies.Translated from Poroshkovaya Metallurgiya, No. 2 (110), pp. 33–37, February, 1972.     

Phase relations in the Ti-TiNi-NbNi-Nb region of the ternary system Ti-Nb-Ni

We have used microstructural, differential thermal, x-ray phase, and electron probe analysis to study alloys in the Ti-TiNi-NbNi-Nb region of the ternary Ti-Nb-Ni system, both as-cast and annealed at 900°C. We have located a pseudobinary TiNi- section, where is a Nb-rich (Nb, Ti)-based solidsolution. The pseudobinary eutectic point parameters are 1170°C and 38 Ni 26 Nb (at. %). We have found two invariant peritectic four-phase equilibria for crystallization of the alloys in the region of interest: L + TiNi + Ti₂Ni (950 °C) and L + TiNi + NbNi (1140°C). We have not confirmed the existence of a ternary compound Ti₃(Ni,Nb)₂.Institute of Materials Science Problems, Ukrainian Academy of Sciences, Kiev. Translated from Poroshkovaya Metallurgiya, Nos. 3/4, pp. 48–54, March–April, 1995.     

Phase equilibria in W-Fe-Co-Ni system alloys. I. Alloys containing 10% (Fe + Co + Ni) at 1400–1200°C

Conclusions The isostructural intermediate -phases of Fe₇W₆ and Co₇W₆ in the W-Fe-Co system form a continuous series of solid solutions and transformation of the L + peritectic equilibrium into, the similar L + (Co₇W₆), occurring in a narrow temperature range (1640–1630°C) is observed. In the W-Fe-Co-Ni system in the 1470–1460°C range transition of the L + (Fe₇W₆)+ peritectic equilibrium into the similar. L + (Co₇W₆+ is also observed.Upon completion of crystallization and at temperatures of 1400–1200°C alloys of the primary section with 10% (Fe + Co + Ni) have a two-( + or +) or three-phase (+ +) structure. In alloys rich in iron at temperatures below 1215°C FeW may form instead of (Fe₇W₆) phase.Translated from Poroshkovaya Metallurgiya, No. 4(280), pp. 60–64, April, 1986.     

Automation and provision of remote-control facilities for the industrial water supply of a combine

Magnitogorsk Metallurgical Combine Energostal Scientific-Industrial Association. Translated from Metallurg, No. 11, pp. 26–27, November, 1989.     

Phase equilibria in the tungsten-rich field of the W-Fe-Ni system in the 800–575°C

Conclusions In the W-Fe-Ni system subjected to annealing at 800 and 575°C, isostructural phases FeW and NiW form a continuous series of (Fe, Ni)W-type solid solutions.At 800 and 575°C, the phases NiW₂ and (Fe, Ni)W or a solid solution of nickel in the FeW phase are in equilibrium with the tungsten-based -solid solution and (Ni, Fe)W or a solid solution of iron in the NiW phase exists in equilibrium with the NiW₂ phase.It is assumed that an invariant eutectoid equilibrium + + (Fe, Ni)W exists in the system under study at a temperature of approximately 1130°C ( is the Fe₇W₆ phase; and and are the tungsten and the nickel-based solid solutions).Depending on the temperature, the phase composition of the binder present in the tungsten-based heavy alloys subjected to long-term annealing can undergo changes: at temperatures exceeding 1110°C, the binder consists of a nickel-based solid solution; in the 1100–1000°C range, an intermediate (Fe, Ni)W phase is present along with this phase; and at temperatures below 1000°C, it consists of the intermediate phases (Fe, Ni)W and NiW₂.Translated from Poroshkovaya Metallurgiya, No. 8(344), pp. 61–67, August, 1991.     

Error in the determination of the particle-size distribution of a powder with a laser analyzer

Conclusions Simple expressions have been obtained from known parameters of a flow-type chamber for estimating errors in laser analyzer measurements of the main numerical particle size distribution characteristics of powders obeying the logarithmic normal law. It is shown that, in spite of marked errors (Fig. 1) in size measurements on single particles whose trajectories are deflected from the laser beam center, it is possible to attain high accuracy in the measurement of numerical particle distribution characteristics by suitable choice of aero or hydrodynamic particle stream focusing [5] (Table 1: maximum error in the determination of the amount of the main fraction in a powder of particle size 3/2 _f = 1.9% at 2_j = 15 and _f = 4.3% at 2_j = 72 m). Without good particle stream focusing, measurements may be very inaccurate (Table 1: _f = 24% at 2_j = 120 m and _f = 55.5% at 2_j = 170 m). The accuracy of laser analyzer measurements grows with increasing curvature (coefficient n) of analyzer calibration characteristics and vice versa. The q/qt relationships obtained may find application in the assessment of errors in particle size analyses of powders with particle distributions differing from the logarithmic normal law.Translated from Poroshkovaya Metallurgiya, No. 12(288), pp. 15–20, December, 1986.     

Deformation structure ofα-SiC single crystals

Conclusions The deformation of-SiC single crystals in bending with the direction of extension and compression approximately parallel to the basal face (the axis of bending parallel to (0001), T = 2000–2100°C) involves slip over the basal plane; the directions of slip are <11¯20> and <10¯10>. Accordingly, crystallographic planes rotate about parallel axes, <10¯10> and <11¯20>. Slip in the <10¯10> direction may be a consequence of migration of dislocations with Burgers vectors equal to b=2/3 <10¯10> (a full dislocation and <10¯10> (a partial dislocation). X-ray microdiffraction observations show [12] that partial dislocations with Burgers vector components parallel to [101¯1] frequently form in-SiC; travel of partial dislocations apparently plays an important part in the deformation of -SiC. The deformation of-SiC is accompanied by kink band formation [1].During the deformation process, dislocations pile up in slip planes, and their density increases by two to three orders; their distribution is very uneven, typical values of being 10⁸–10¹⁰ cm^–2 and values of _max of 10¹¹cm^–2 being recorded in zones with a large local curvature.Because of the changing specimen structure, two stages may be detected in the deformation process: 1) formation of elastically bent regions and 2) comminution (formation of microcracks within the specimen).The crystals investigated exhibited only micropolygonization, when the block size in the basal plane was 30 and the block disorientation 10. Neither postdeformation annealing for 10–30 h at 2000–2100°C nor prolonged (4–8 h) holding of specimens under load produced any macroscopic polygonization. The structure of naturally deformed specimens, too, was found to contain no polygonization macroblocks. It would appear that, because of the appreciable activation energy for diffusion and strong covalent linkage in-SiC, dislocations cannot readily climb over large (>10) distances in this compound; in this connection, the relaxation properties of-SiC would be expected to be weak.The methods of investigation employed in our work are comparatively simple to use. They do not give any information about elementary deformation acts, but throw a light on the character of lattice variation in various volumes (V ranging from 0.1t to 6.0t mm³, where t=200–500 is the crystal thickness) and are particularly useful in topographic x-ray photography when there is no contrast on individual linear defects in crystals.Translated from Poroshkovaya Metallurgiya, No. 8 (128), pp. 63–72, August, 1973.