Experimental Investigation in the Quaternary Systems Ti–Ni–Al–N and Ti–Ni–Al–O (original) (raw)
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On the Quaternary System Ti-Fe-Ni-Al
Journal of Phase Equilibria and Diffusion, 2008
The homogeneity ranges of the Laves phases and phase relations concerning the Laves phases in the quaternary system Ti-Fe-Ni-Al at 900°C were defined by x-ray powder diffraction (XPD) data and electron probe microanalysis (EPMA). Although at higher temperatures the Laves phase forms a continuous solid solution, two separate homogeneity fields of TiFe 2-based (denoted by k Fe) and Ti(TiNiAl) 2-based (denoted by k Ni) Laves phases appear at 900°C. The relative locations of Laves phases, G phase, Heusler phase, and CsCl-type phase as well as the associated tie-tetrahedra were experimentally established in the quaternary for 900°C and presented in three-dimensional (3D) view. Furthermore, a partial isothermal section TiFe 2-TiAl 2-TiNi 2 was constructed, and a connectivity scheme, derived for equilibria involving Laves phases in the Ti-Fe-Ni-Al quaternary system at 900°C was derived. As a characteristic feature of the quaternary phase diagram, the solid solubility of fourth elements in both the TiFe 2-based and Ti(NiAl) 2-based Laves phases is limited at 900°C and is dependent on the ternary Laves phase composition. A maximum solubility of about 8 at.% Ni is reached for composition Ti 33.3 Fe 33.3 Al 33.4. Structural details have been evaluated from powder x-ray and neutron diffraction data for (i) the Ti-Fe-Ni ternary and the Ti-Fe-Ni-Al quaternary Laves phases (MgZn 2-type, space group: P6 3 /mmc) and (ii) the quaternary G phase. Atom site occupation behavior for all phases from the quaternary system corresponds to that of the ternary systems. For the quaternary Laves phase, Ti occupies the 4f site and additional Ti (for compositions higher than 33.3 at.%Ti) preferably enters the 6h site. Aluminum and (Fe,Ni) share the 6h and the 2a sites. The compositional dependence of unit cell dimensions, atomic coordinates, and interatomic distances for the Laves phases from the quaternary system is discussed. For the quaternary cubic G phase, a centrosymmetric as well as a noncentrosymmetric variety was observed depending on the composition: from combined x-ray and neutron powder diffraction measurements Ti 33.33 Fe 13.33 Ni 10.67 Al 42.67 was found to adopt the lower symmetry with space groupF " 43m.
Phase relations in the Al-rich corner of the Ti–Ni–Al system
Journal of Alloys and Compounds, 2001
Phase relations in the Al-rich corner of the Ti-Ni-Al system were determined for three partial isothermal sections at 900, 800 and 6008C by diffusion couples and 14 bulk alloy samples employing X-ray powder diffraction, optical microscopy and electron probe microanalysis (EPMA). The equilibrium diagrams are dominated by the ternary compound t with AuCu-type. In some samples, 1 3 annealed at 600 and 8008C, a (non-equilibrium) Ti Al phase with the ZrAl-type was found, which formed during solidification 5 11 3 followed by incomplete decomposition. A new metastable phase t was observed with a composition around Ti Ni Al (at.%). On 5 13.8 20.1 66.1 oxidation a quaternary phase Ti NiAl O with a filled Ti Ni-type is formed. A Rietveld refinement of t with AuCu-type is presented. 3 2 1 2x 2 1 3
Intermetallics, 1999
Phase relations in the ternary system Al±Ni±Ti have been experimentally established for the isothermal section at 900 C for concentrations 0.1 4 x Al 4 0.7. The investigation is based on X-ray powder diraction, metallography, SEM and EMPA-techniques on about 40 ternary alloys, prepared by argon-arc or vacuum-electron beam melting of proper elemental powder blends. The existence of four ternary compounds, t 1 to t 4 , is con®rmed, however, in contrast to earlier investigations at signi®cantly dierent compositions and with dierent shape of the homogeneity regions. This is particularly true for the phase regions of t 3-Al 3 NiTi 2 with the MgZn 2-type structure ranging from Al 30 Ni 28 Ti 42 (composition lowest in Al) to Al 50 Ni 16 Ti 34 (composition richest in Al) and for t 2-Al 2 NiTi. The complex atom site substitution mechanism in t 3 changing from Ti/Al exchange at Al-poor compositions towards Ni/Al replacement for the Al-rich part was monitored in detail by quantitative X-ray powder diraction techniques (Rietveld analyses). In contrast to earlier reports, claiming a two-phase region Ni{Al x Ti 1-x } 2 +t 3 , we observed two closely adjoining three-phase equilibria: a 2-AlTi 3 +Ni{Al x Ti 1-x } 2 + t 4-AlNi 2 Ti and a 2-AlTi 3 +t 3-Al 2 NiTi 2 +t 4-AlNi 2 Ti. The earlier reported ``homogeneous phase at Al 23 Ni 26 Ti 51 H '' was shown by high resolution microprobe and X-ray diraction measurements to be an extremely ®negrained eutectic. The experimental results are in ®ne agreement with the thermodynamic calculation.
Intermetallics, 1999
The Al±Ni±Ti phase diagram has been thermodynamically assessed and a consistent set of thermodynamic functions has been developed. The thermodynamic modeling is based on an experimental investigation of the phase equilibria in the composition range of 0.14x Al 40.7. Alloys were prepared by argon-arc or vacuum-electron beam melting of elemental powder blends. X-ray powder diraction, metallography, SEM and EMPA-techniques were employed to analyze the samples in the as-cast state as well as after annealing at 800, 900 and 1000 C. The existence of the four ternary compounds, (1 to (4 , has been con®rmed, although homogeneity regions dier signi®cantly from reports in the literature. The homogeneous phase, previously claimed at ``Al 23 Ni 26 Ti 51 '', is shown by high resolution microprobe and X-ray diraction measurements to be an extremely ®ne-grained eutectic structure. The congruent melting behavior of (4 elxi 2 i is con®rmed, but, in contrast to earlier reports, primary crystallization and congruent melting have been observed for (1 el 13 xi 2 i 5 and (3 el 3 xii 2. In contrast to earlier assessments, (1 Y(2 and (3 are experimentally found to be stable at 800, 900 and 1000 C. The thermodynamic modeling of the ternary phases (2 and (3 is done with sim-pli®ed sublattice models, considering their crystal structure and homogeneity ranges. The sublattice model for (4 is taken from an earlier asessment of the nickel-rich ternary phase equilibria. The present assessment covers the entire composition range. An application to the solidi®cation behavior of ternary alloys is also exempli®ed.
Non-equilibrium solidification of intermetallic phases in Ni-Al and Ti-Al alloy systems
Ni-Al alloy system exhibits particular solidification behavior due to chemical long-range order of the intermetallic compounds Ni 3 Al (L1 2) and NiAl (B2). Using rapid solidification technique their slow solidification kinetics leads to the entrapment of solute atoms and/or disorder, respectively. In the present work containerless processing by applying electromagnetic levitation technique is used to analyze the aspects of dendrite growth, as observed in solidification of undercooled Ni-Al melts. Sharp-interface modeling is used to illustrate the effect of chemical order on the solidification kinetics of intermetallic phases. Both experimental results and theoretical work denote the role of chemical order in solidification of Ni-Al intermetallics.
Investigation of phase equilibria in the Ti-Al-Si-Nb system at low Nb contents
Acta Materialia, 2003
Phase equilibria in the a/a 2 phase region of the Ti-Al-Si-Nb system at Nb content 2.5, 3.5 and 5 at.% were studied in alloys as-cast and heat-treated at 800°C. Samples were prepared by arc-melting technique, homogenized at 1350°C and then heat-treated at 800°C, followed by ice water cooling. The structure of the alloys was characterized by means of X-Ray diffraction, differential thermal analysis, electron probe microanalysis, scanning electron microscopy and transmission electron microscopy. The continuous solid solutions with variable compositions (Ti 1Ϫx ,Nb x) 3 (Si 1Ϫy ,Al y) (h) (0.05ՅxՅ0.07, 10 Ϫ3 ՅyՅ 0.02) was detected at 800°C for the first time in the multi-component alloys based upon the Ti-Si system. It was stabilized by Nb additions in the alloys with low Al content. A peritectoid reaction b + a→h was observed. Additions of Al neutralized the stabilizing effect of Nb resulting in an a + Ti 5 Si 3 (z) equilibrium.
Influence of Nb on the phase stability of Ti–Al–N
Scripta Materialia, 2010
Metastable solid-solution Ti 1Àx Al x N thin films synthesized by plasma-assisted vapour deposition crystallize in the cubic NaCl structure with AlN mole fractions x 6 0.7. Based on X-ray diffraction and ab initio calculations, we reveal that this critical value only slightly decreases to 0.69 with the addition of up to n = 0.125 mol fraction NbN to form cubic Ti 1ÀxÀn Al x Nb n N. During annealing in vacuum to 1450°C, the as-deposited single-phase cubic coatings decompose to form cubic Ti 1Àn Nb n N and wurtzite AlN.
Experimental study of the Fe–Ni–Ti system
Intermetallics, 2010
In this investigation phase relations in the Fe-Ni-Ti system were studied and two isothermal sections at 800 C and 1000 C as well as a revised liquidus projection were established. Microstructural characterisation of the as-cast alloys and of samples equilibrated at 800 and 1000 C was performed by scanning electron microscopy (SEM), chemical compositions of the phases were analysed by electron probe microanalysis (EPMA), and liquidus temperatures were examined by differential thermal analysis (DTA). The experimental results clarify some uncertainties concerning the melting behaviour and the solid-state phase equilibria between the phases (Ni,Fe)Ti 2 , (Fe,Ni)Ti and b-Ti. The present data also confirm that the solid solubility of Ti in g-(Fe,Ni) varies in dependence on the Fe:Ni ratio and decreases with decreasing temperature. The liquidus projection as well as the reaction scheme in the Ti-lean part are modified because two ternaries eutectic E 1 : L4 g-(Fe,Ni) þ Fe 2 Ti þ Ni 3 Ti and E 2 : L 4 Fe 2 Ti þ Ni 3 Ti þ (Fe,Ni)Ti are found at 1108 and 1099 C, respectively. Intermetallics j o u r n a l h o m e pa g e : w w w . e l s e v i e r . c o m / l o c a t e / i n t e r m e t 0966-9795/$ -see front matter Ó