Phase Equilibria in the Ni - Sc - Ti System in the Region TiNi - Ni - Sc1.13Ni At 900°C

Physicochemical analysis methods have been applied to TiNi - Ni - Sc_1.13Ni alloys at 900°C. It is shown that the phase interactions under these conditions are unaltered by comparison with those at subsolidus temperatures. The phase based on TiNi₃, which is the most thermally and thermodynamically stable, is in equilibrium with all the other phases in the TiNi - Ni - Sc_1.13Ni subsystem. The solubilities of the components in the phases based on Ti_{50– x} NiSc_x, TiNi₃, and the Laves phase are reduced at 900°C. The temperature of the polymorphic transformation in the phase based on ScNi₅ is raised by the dissolution of titanium in it to 930°C. An isothermal section is constructed of the phase diagram for the TiNi - Ni - Sc_1.13Ni subsystem together with two polysections.__________Translated from Poroshkovaya Metallurgiya, Nos. 1–2(441), pp. 39–46, January–February, 2005.     

Projection of the solidus surface of the Sc-Ti-Ni system in the region of TiNi-Ni-ScNi alloys

Metallographic, x-ray diffraction, electron probe, and differential thermal analyses are used to study, the interaction of components of TiNi-Ni-ScNi alloys in the Sc-Ti-Ni system at subsolidus temperatures. The vertical section between TiNi and ScNi is shown to be quasibinary. No ternary compounds are observed in the TiNi-Sc-ScNi system. Six single-phase, nine two-phase, and four three-phase fields are found on the solidus of the TiNi-Ni-ScNi system. A characteristic feature of the structure of alloys of the Sc-Ti-Ni system in the region TiNi-Ni-ScNi ε is that the E phase is in equilibrium with the phases based on TiNi(δ), Ni(γ), ScNi₂(λ₂), Sc₂Ni(v), SCNi₅(ξ), and ScNi(δ).     

Alloy composition and equilibrium phase diagram for the Sc–Ru–Rh system. V. 930°C isothermal section of the Sc–ScRu–ScRh partial system

Microstructure, x-ray diffraction, and electron-microprobe methods have been applied to study an isothermal section at 930°C of the Sc–ScRu–ScRh partial system. As no ternary compounds are formed in the system, the phase equilibria involve intermediate phases based on binary compounds that exist in the bounding Sc–Ru and Sc–Rh systems: Sc₅Ru₃, Sc₂Ru, Sc₁₁Ru₄, Sc₄₄Ru₇, Sc₃Rh, and the δ and η phases (continuous series of solid solutions between the isostructural phases, respectively, ScRu and ScRh (with the CsCl structure type), together with Sc₅₇Ru₁₃ and Sc₅₇Rh₁₃ (Sc₅₇Rh₁₃ structure type)), and a solid solution based on α-scandium. The components of the isothermal section are eight single-phase regions, five three-phase fields (δ+<Sc₅Ru₃>+<Sc₂Ru>, <Sc₂Ru>+<Sc₁₁Ru₄> + η. <Sc₂Ru> + δ+η, >Sc₃Rh> + δ+η, >Sc₄₄Ru₇+<α-Sc>+η, and twelve two-phase regions. There is a difference from the phase equilibria at subsolidus temperatures in that at 930°C there are no equilibria related to the decomposition of the phase based on the compound Sc₂Rh (γ′), and they are replaced by new ones due to the formation of phases based on the compound Sc₄₄Ru₇ (θ) in the solid state.     

The Effect of Minor Addition of Ni on the Microstructural Evolution and Mechanical Properties of Suction Cast Al-0.14 at. pct Sc Binary Alloy

Aluminum scandium binary alloys represent a promising precipitation-hardening alloy system. However, the hardness of the binary alloys decreases with the rapid coarsening of Al₃Sc precipitate during high-temperature aging. In the current study, we report a new approach to compensate for the loss of mechanical properties by combining rapid solidification with very small ternary addition of transition metal Ni. This addition yields dispersion, and at a critical concentration improves the mechanical properties. We explore additions of a maximum of 0.06 at. pct of Nickel to a binary Al-0.14 at. pct Sc alloy, which yield nickel-rich dispersions. We report two kinds of biphasic dispersions containing AlNi₂Sc/Al₉Ni₂ and α-Al/Al₉Ni₂ phase combinations. The maximum improvement in mechanical properties occurs with the addition of 0.045 at. pct Ni with a yield strength of 239 ± 7 MPa for an aging treatment at 583 K (310 °C) for 15 hours.     

Interface microstructures in the diffusion bonding of a titanium alloy Ti 6242 to an INCONEL 625

A Ti6242 alloy has been diffusion bonded to a superalloy INCONEL 625. The microstructures of the as-processed products have been analyzed using optical metallography, scanning electron microscope (SEM), and scanning transmission electron microscope (STEM) techniques. The interdiffusion of the different elements through the interface has been determined using energy-dispersive spectroscopy (EDS) microanalysis in both a SEM and a STEM. Several regions around the original interface have been observed. Starting from the superalloy INCONEL 625, first a sigma phase (Cr₄Ni₃Mo₂), followed by several phases like NbNi₃, Ŋ/Ni₃Ti, Cr(20 pct Mo), β Cr₂Ti, NiTi, TiO, TiNi, and Ti₂Ni intermetallics, just before the Ti6242 have been identified. Because the diffusion of Ni in Ti is faster than the diffusion of Ti in the superalloy, a Kirkendall effect was produced. The sequence of formation of the different phases were in agreement with the ternary Ti-Cr-Ni diagram.     

Reactions of TiNi with Oxygen

X-ray diffraction and metallography have been applied to study the formation conditions for the phase in Ti₄Ni₂O in the reaction of TiNi with oxygen. The phase diagram for the Ti Ni O system indicates that the phase (Fe₃W₃C structure type) is a solid solution of oxygen in Ti₂Ni. For TiNi made in an oxygen-bearing medium or from initial materials contaminated with oxygen, one gets the phase and Ni₃Ti. Those phases are formed also in surface oxidation as products from the initial interaction of the alloy with oxygen. They occur when there is a low oxygen partial pressure, as on annealing TiNi powder in a vacuum given by a rotary pump or in a layer under scale formed on a cast TiNi specimen on oxidation in air. The layer under the scale is formed because of the preferential loss of titanium from the TiNi and consists of Ti₄Ni₂O, Ni₃Ti, and an Ni(Ti) solid solution, which in turn occurs because of preferential loss of titanium from Ni₃Ti. In the subsequent oxidation stages, there is selective oxidation of Ti₄Ni₂O to lower titanium oxides. The decomposition of TiNi alloy containing oxygen when the composition is varied near the equiatomic one can be used for practical purposes. For example, the segregation of the hydride-forming phase may improve the hydrogen uptake by the alloy, while the dispersed segregation of Ni₃Ti and the lower oxides may favor hardening.     

Thermodynamic activities and phase boundaries for the alloys of the Ni3Al-Ni3Ti pseudobinary section in the Ni-Al-Ti system

The pseudobinary section Ni₃Al-Ni₃Ti of the ternary Ni-Al-Ti system has been investigated by ther-mal analysis and Knudsen effusion mass spectrometry. The solidus of the pseudobinary section and the thermodynamic activities of Ni and Al have been determined in the alloys Ni_0.75A1_0.25-xTi_x of the compositionsx = 0.00 to 0.21. Moreover, the thermodynamic activities of Ni and Ti in Ni₃Ti (x = 0.25) as well as the Gibbs energy of mixing of the Ni_0.75Ti_0.25 phase resulted. The ionization cross-sectional ratio Σ(Ni)/Σ(Ti) = 0.77 has been evaluated for the electron impact energy of 50 eV.     

Structurization in the sintering of heterophase materials TiN-Ni

Conclusions As a result of the investigation of the structurization in the quasibinary system TiN-Ni after high-temperature liquid-phase and low-temperature solid-phase sintering it was established that structurization in short-term liquid-phase sintering is accompanied by the denitration of TiN, decrease of its microhardness, change of the compositon of its metal bond and of its amount in consequence of diffusion of Ti from TiN (formation of the eutectic Ti(Ni) + Ni₃Ti).The considerable increase of microhardness of grains of TiN and the linear increase of its lattice period after lengthy isothermal sintering are due to the dissolution of nickel in TiN. In the metal bond there appear excess crystals of Ni₃Ti and of the eutectic Ni₃Ti + TiNi.The growth of TiN grains in the heterophase material TiN-Ni proceeds by the mechanism of coalescence and dissolution-segregation. After solid-phase sintering the phase composition of Ni and TiN did not change, the grain size of TiN remained stable.Translated from Poroshkovaya Metallurgiya, No. 4(328), pp. 25–30, April, 1990.     

Absorption of hydrogen by sintered intermetallics of the Sc — Ni system

The absorption of hydrogen by the intermetallics Sc₂Ni, ScNi, ScNi₂, Sc₂Ni₇, and ScNi₅ produced by sintering was investigated. Hydrogen absorption isobars at the pressures 0.1 and 3 MPa were obtained. The maximum hydrogen absorption capacity of the compounds was determined. The absorption properties were compared with those of specimens produced by arc melting. Translated from Poroshkovaya Metallurgiya, Nos. 5–6(413), pp. 78–82, May–June, 2000.     

TiNi Shape Memory Alloys with High Sintered Densities and Well-Defined Martensitic Transformation Behavior

This study demonstrates that a high density and a high transformation heat can be attained for PM TiNi. With the use of fine elemental powders, a composition of Ti₅₁Ni₄₉, two-step heating, and persistent liquid-phase sintering at 1553 K (1280 °C), a 95.3 pct sintered density was attained for compacts with a green density of 66 pct. A transformation heat, ΔH, of 31.9 J/g was also achieved, which is much higher than reported previously for sintered TiNi and is approaching the highest ΔH reported to date, 35 J/g, for wrought TiNi with low C, O, and N contents. The main reason for having these properties in powder metal TiNi with higher amounts of C, O, and N is that the extra Ti, that over the equiatomic portion in the Ti-rich Ti₅₁Ni₄₉, forms Ti₂Ni compound, which traps most of the C, O, and N. This results in low interstitial contents and a high Ti/Ni ratio of 50.5/49.5 in the TiNi matrix. The tensile strength, elongation, and shape recovery rate after five training cycles were 638 MPa, 14.6, and 99.1 pct, respectively, despite the presence of Ti₂Ni compounds at grain boundaries.     

Metallographic and structural observations in the pseudo-binary section Ni3Si-Ni3Ti of the Ni-Si-Ti system

The metallographic and structural features of the alloys in the pseudo-binary section between Ni₃Si (L1₂) and Ni₃Ti (D0₂₄) were investigated. The L1₂ phase (γ′) extended along a pseudo-binary line between the two phases with Ti solubility of about 11 at.%. Also, the L1₂ phase expanded up to about 80 at.% Ni at high Ti concentration. Consistent with previous observations, the addition of the Ti elements into the Ni₃Si alloy led to congruent melting of L1₂ phase. Based on the observations of X-ray diffraction and an electron channeling technique aided with EDX (Energy Dispersive X-ray spectrometer), it was shown that the alloying element, Ti, substituted for the component element of Si the ternary alloy Ni₃(Si, Ti) was highly ordered. The solubility limit of boron in Ni₃(Si, Ti) was shown to be very low, i.e. less than 50 ppm.     

The Solidus Surface in the Ti ― Ni ― Hf System in the Ti ― TiNi ― HfNi ― Hf Region

Studies have been done on the phase equilibria at subsolidus temperatures in the Ti TiNi HfNi Hf region of the Ti Ni Hf ternary system. The phases based on binary compounds and solid solutions of these components are accompanied in the equilibria by a phase based on an equiatomic ternary compound. This new phase belongs to the family of Laves phases and has a hexagonal crystal structure of MgZn₂ type. The solidus surface in the Ti TiNi HfNi Hf subsystem consists of the surface of the ternary phase alone, the surfaces of the six solid solutions based on the components and binary intermediate phases, the planes of five conode triangles, and the corresponding lineated surfaces.     

Isothermal section of the ternary Sn-Cu-Ni system and interfacial reactions in the Sn-Cu/Ni couples at 800 °C

The isothermal section of the Sn-Cu-Ni system at 800 °C has been experimentally determined. There is no ternary compound. A solid solution with a very wide compositional range, the γ phase is formed between the Ni₃Sn(H) phase and Cu₄Sn(H) phase; however, both of these two binary phases are not stable at 800 °C. The binary Ni₃Sn₂ phase also has extensive ternary solubility. The homogeneity ranges of both the γ and Ni₃Sn₂ phases are very large in parallel to the Cu-Ni side, but relatively narrow along the Sn direction. This phenomenon indicates that Cu and Ni are exchangeable in both phases. Three kinds of reaction couples, Sn-55 at. pct Cu/Ni, Sn-65 at. pct Cu/Ni, and Sn-75 at. pct Cu/Ni, were prepared and reacted at 800 °C for 5 to 20 minutes. The reaction paths are liquid/Ni₃Sn₂/γ/Ni₃Sn(L)/Ni for the Sn-55 at. pct Cu/Ni and Sn-65 at. pct Cu/Ni couples, and the reaction path is liquid/γ/Ni₃Sn(L)/Ni for the Sn-75 at. pct Ni couples.     

Phase equilibria of the ternary Ni-Cr-Zr system and interfacial reactions in the Ni-Cr/Zr couples

The phase equilibria of the ternary Ni-Cr-Zr system and interfacial reactions in the Ni-Cr/Zr couples at 900 °C were determined. Fifty alloys were prepared from pure Ni, Cr, and Zr. The alloys were metallographically analyzed. Both X-ray diffraction and electron-probe microanalysis (EPMA) were carried out for structural identification and compositional analysis of phases formed in these alloys. At 900 °C, the Cr-Ni₁₀Zr₇ two-phase region divides the system into two halves. ZrCr₂(C14) exists in the Zr-Cr-Ni₁₀Zr₇ half, and the ZrCr₂ (C15) phase has an extensive ternary solubility. In the Cr-Ni₁₀Zr₇-Ni half, except for the Ni₇Zr₂ phase, most of the binary Zr-Ni compounds are in equilibrium with either Cr or Ni phase. Reaction-couple techniques were used for the interfacial reaction study. The reaction path was Zr/NiZr₂/NiZr/Ni₁₀Zr₇/Ni₂₁Zr₈/Cr/Ni-Cr alloy in the Ni-Cr/Zr couples examined in this study. The results indicate that Ni is the fastest-diffusing species, while Cr is the slowest one.     

Alloy compositions and phase diagram for the Sc−Ru−Rh system. IV. Solidus surface of the Sc−ScRu−ScRh partial system

A projection has been constructed for the solidus surface of the partial, Sc−ScRu−ScRh system on the concentration triangle from data obtained by microstructure analysis, x-ray diffraction, microprobe analysis, differential thermal analysis, and the measurement of temperatures for the start of melting by the Pirani-Alterthum method. It is found that ternary compounds are not formed in this system. The solidus surface has regions representing solid solutions based on β-scandium and phases based on the binary compounds Sc₅Ru₃, Sc₂Ru, Sc₁₁Ru₄, Sc₂Rh, and Sc₃Rh, as well as ones for the δ and η phases (continous solid solution series formed between isostructural phases of CsCl type based on ScRu and ScRh as well as Sc₅₇Rh₁₃ based on Sc₅₇Ru₁₃ and Sc₅₇Rh₁₃). Other components of the solidus surface are twelve lineated surfaces, which bound two-phase volumes, and five isothermal surfaces, which correspond to nonvaiant four-phase equilibria involving the liquid, which occur at temperatures of 1150, 1055, 1045 (two equilibria), and 1000°C. Institute for Problems of Materials Science, Ukraine National Academy of Sciences, Kiev. Translated from Poroshkovaya Metallurgiya, Nos. 7–8(408), pp. 55–63, July–August, 1999.     

Diffusion in the titanium-nickel system: II. calculations of chemical and intrinsic diffusion coefficients

Diffusion coefficients in the Ti-Ni system have been calculated by the aid of equations given by Sauer and Freise, and Wagner. Values for the TiNi (50 at. pct Ni) phase were found to be:D ^u (cm²/s) = 0.0020 exp - 142,000/R for the temperature range between 650 and 940°C. The heat of activation, expressed in J/mol, has an accuracy of ±6000. For the β-Ti(Ni) phase containing 6 at. pct Ni the temperature dependence of the diffusion coefficient is expressed by:D ^u (cm²/s) = 0.0688 exp - 141,000/RT. The uncertainty in the energy of activation is ±12000 J/mol. No clear variation of the diffusion coefficient with concentration could be detected. It was found that Ni is by far the fastest moving component in β-Ti(Ni), Ti₂Ni and TiNi (at least in the composition range between 50 and 53 at. pct Ni). Values ofD _Ni/D _Ti have been calculated with an equation derived by van Loo. The significance of the calculated values is critically examined. By means of a practical example it is shown that the calculated ratio of the intrinsic diffusion coefficients can be extremely sensitive to slight variations in the position of the marker interface.Diffusion coefficients in the Ti-Ni system have been calculated by the aid of equations given by Sauer and Freise, and Wagner. Values for the TiNi (50 at. pct Ni) phase were found to be:D ^u (cm²/s) = 0.0020 exp - 142,000/R for the temperature range between 650 and 940°C. The heat of activation, expressed in J/mol, has an accuracy of ±6000. For the β-Ti(Ni) phase containing 6 at. pct Ni the temperature dependence of the diffusion coefficient is expressed by:D ^u (cm²/s) = 0.0688 exp - 141,000/RT. The uncertainty in the energy of activation is ±12000 J/mol. No clear variation of the diffusion coefficient with concentration could be detected. It was found that Ni is by far the fastest moving component in β-Ti(Ni), Ti₂Ni and TiNi (at least in the composition range between 50 and 53 at. pct Ni). Values ofD _Ni/D _Ti have been calculated with an equation derived by van Loo. The significance of the calculated values is critically examined. By means of a practical example it is shown that the calculated ratio of the intrinsic diffusion coefficients can be extremely sensitive to slight variations in the position of the marker interface. This paper is based on a Thesis submitted by G. F. BASTIN in fulfillment of requirements for the degree of Doctor in Technological Sciences.     

Precipitation processes in near-equiatomic TiNi shape memory alloys

Metallographic studies have been made of precipitation processes in Ti-50 pct Ni and Ti-52 pct Ni (at. pct) shape memory alloys. The eutectoid and peritectoid reactions previously reported for near-equiatomic and Ni-rich TiNi alloys were not observed for either composition. In the Ti-52Ni alloy, diffusional transformations take place, similar to those in supersaturated alloys. The precipitation sequence can be written asβ ₀ → Ti₁₁Ni₁₄ → Ti₂Ni₃ → TiNi₃. The solidus line of the TiNi phase in the Ti-52Ni alloy lies at 812 ± 22 °C. Morphological characteristics of the various precipitate phases are described in detail. M. NISHIDA, formerly with the University of Illinois     

Stabilizing the L1<Subscript>2</Subscript> structure of Pt<Subscript>3</Subscript>Al(r) in the Pt-Al-Sc system

The Pt-Al system has high potential to act as alloy base for so-called refractory superalloys. Although the envisaged strengthening phase Pt₃Al(r) has favorable L1₂ crystal structure only at high temperatures, even small amounts of Sc stabilized L1₂ crystal structure at low temperature. Pt-Al-Sc alloys were arc melted, heat treated, and examined by means of scanning electron microscopy and X-ray diffraction (XRD). Pt₃Al_1−x Sc _x (r) forms a continuous phase field from the Al-rich side to the Sc-rich side of the Pt-Al-Sc ternary system. The absolute value of the lattice misfit between cubic Pt₃Al_1−x Sc _x (r) and the matrix decreases with increasing Sc content.