In situ XRD studies on Al–Ni and Al–Ni–Sr alloys prepared by rapid solidification

In this paper the structure and stability of Al–17 wt.%Ni(Al–17Ni) and Al–17 wt.%Ni–2 wt.%Sr alloys prepared by rapid solidification was investigated by means of XRD techniques. Our work demonstrates that both alloys are crystalline and composed of fcc (Al–Ni) solid solution and orthorhombic Al₃Ni phases. The ternary alloy shows in addition the presence of small amount of tetragonal Al₄Sr phase. In situ XRD experiment demonstrates the stability of the solute solution up to 650 °C, Al₃Ni above 750 °C while Al₄Sr overcomes melting of the major phases at 800 °C. High-temperature structure analysis proved strong bindings between Al and Ni atoms in Al₃Ni phase, corroborating its covalent nature, linear and faster increase of the fcc volume with annealing temperature. The linear correlation between constituting atoms decreases with increase of the temperature.The work also documents the applicability of pair distribution function (PDF) analysis to the study of multiphase crystalline systems.     

On the four-phase reactions in the Ti–Ni–Al system

Two four-phase reactions of transition type in the Ti–Ni–Al system were studied on several alloys, which were annealed at carefully set temperatures and quenched. The phase constitution was established by XRD and EPMA analyses. Due to sluggish reaction kinetics, the transition temperatures were defined by annealing and quenching techniques as no DTA signals could be received. For the reaction NiAl + TiNiAl  TiNiAl₂ + TiNi₂Al, the transition temperature was found to be 925 °C ± 15 °C and for the reaction TiNiAl + Ti₃NiAl₈  TiAl₂ + TiNiAl₂, the transition temperature was found to be 990 °C ± 15 °C. Furthermore we confirmed the three-phase field TiNi₂Al + Ti₃Al + Laves phase (TiNiAl), as reported at 900 °C by Huneau et al. in 1999.     

Constitution of the high-Al region of Al–Cu–Rh

Phase equilibria between 540 and 1010 °C were studied in Al–Cu–Rh alloys containing more than 50 at.% Al. Congruent equiatomic AlRh dissolves more than 40 at.% Cu and extends up to 58 at.% Al at the high-Cu part of its compositional range. High-temperature cubic C-Al₅Rh₂ (C-phase) dissolves up to 13 at.% Cu, “Al₃Rh” (₆-phase) up to 15 at.% Cu and Al₉Rh₂ up to 1.5 at.% Cu. The solubility of the third element in other binary Al–Rh and Al–Cu phases is below 0.5 at.%. Close to the high-Cu limit of the C-phase region the fcc C₂-phase structurally related to the C-phase is formed. Stable decagonal phase (D₁-phase) is formed below 1005 °C in a compositional range extending from Al₆₅Cu₁₆Rh₁₉ to Al₆₂Cu₂₃Rh₁₅, which shifts to higher Cu concentrations with decreasing temperature. An additional ternary phase forming around the Al₇₀Cu₂₀Rh₁₀ composition below 660 °C was revealed. Partial 1010, 990, 900, 800, 700, 600 and 540 °C isothermal sections were determined.     

Identification of σ and Ω phases in AA2009/SiC composites

The microstructure evolution during ageing treatment at 170 and 190 °C of AA2009/SiC composites, reinforced with 15 vol.% particulates and whiskers, was studied by transmission electron microscopy. Besides θ′ and S′ phases, the typical hardening precipitates on Al–Cu–Mg alloys, it was found the presence of Ω and σ (Al₅Cu₆Mg₂) phases in the matrix. σ phase was only found in the matrix of particulate composite, while Ω phase appeared in both. This phase has not been previously observed in Al matrix composites based on conventional Al–Cu–Mg alloys.     

The phase equilibria of the Al–Ce–Ni system at 500 °C

The phase equilibria at 500 °C in the Al–Ce–Ni system in the composition region of 0–33.3 at.% Ce are investigated using XRD and SEM/EDX techniques applied to equilibrated alloys. The previously reported ternary phases and the variation of the lattice parameters versus the composition for different solid solution phases are investigated. It is confirmed that τ₂(Al₂CeNi) exists at 500 °C, while τ₃(Al₅Ce₂Ni₅) does not exist at 500 °C. A new compound τ₉ with composition of about Al₃₅Ce_16.5Ni_48.5 is found. The solubility of Ni in Al₁₁Ce₃ and αAl₃Ce is generally about 1 at.%, while the solubility of Ni in Al₂Ce is measured to be 2.7 at.%. The solubility of Ce in Al₃Ni, Al₃Ni₂, AlNi and AlNi₃ is all less than 1 at.%. The solubility of Al in CeNi₅, Ce₂Ni₇ and CeNi₃ is measured to be 30.4, 4.8 and 9.2 at.%, respectively, while there is no detectable solubility for Al in CeNi₂. A revised isothermal section at 500 °C in the Al–Ce–Ni system has been presented.     

A contribution to the Al–Ni–Cr phase diagram

The Al–Ni–Cr phase diagram was specified at 1000 °C and partially at 900 °C. The results concerning the region below 60 at.% Al agreed qualitatively with the literature data. The binary Al–Cr phases μ and γ dissolve up to 1 and 3 at.% Ni, respectively, and Al₃Ni₂ up to 2.5 at.% Cr. Two ternary phases were revealed: hexagonal ζ (a ≈ 1.77, c ≈ 1.24 nm) in a wide range between Al₈₁Ni₃Cr₁₆, Al_76.5Ni₃Cr_20.5, Al_76.5Ni₉Cr_14.5 and Al_71.5Ni₉Cr_19.5, and high-temperature orthorhombic (a ≈ 1.26, b ≈ 3.48, c ≈ 2.02 nm) around Al_76.5Ni_2.0Cr_21.5.     

Experimental determination and thermodynamic assessment of the phase diagram in the Co–Zr system

The phase diagram in the Co–Zr system was determined using the electron probe microanalyzer (EPMA), differential thermal analysis (DTA) and X-ray diffraction (XRD) technique. The experimental results indicate that (1) the solubility region of the Co₂Zr phase is from 25 to 34 at.% Zr; (2) the CoZr₃ phase exists and the temperature of the peritectoid reaction (CoZr₂ + (βZr) ↔ CoZr₃) is about 981 °C; (3) the solubilities of Zr in the (αCo) phase and Co in the (βZr) phase are about 0.15 and 2.5 at.%, respectively. A thermodynamic assessment of the Co–Zr binary system was carried out by the CALPHAD (calculation of phase diagrams) method. The Gibbs free energies of the solution phases (liquid, fcc, bcc and hcp) were described by the subregular solution model, and those of the intermetallic compounds (Co₁₁Zr₂, Co₂₃Zr₆, Co₂Zr, CoZr, CoZr₂ and CoZr₃) were described by the sublattice model. A proper set of the thermodynamic parameters has been derived for describing the Gibbs free energies of each phase in the Co–Zr system. An agreement between the calculated results and experimental data is obtained.     

Strong bonding of infrared brazed α2-Ti3Al and Ti–6Al–4V using Ti–Cu–Ni fillers

Infrared dissimilar brazing of α₂-Ti₃Al and Ti–6Al–4V using Ti–15Cu–25Ni and Ti–15Cu–15Ni filler metals has been performed in this study. The brazed joint consists primarily of Ti-rich and Ti₂Ni phases, and there is no interfacial phase among the braze alloy, α₂-Ti₃Al and Ti–6Al–4V substrates. The existence of the Ti₂Ni intermetallic compound is detrimental to the bonding strength of the joint. The amount of Ti₂Ni decreases with increasing brazing temperature and/or time due to the depletion of Ni content from the braze alloy into the Ti–6Al–4V substrate during brazing. The shear strength of the brazed joint free of the blocky Ti₂Ni phase is comparable with that of the α₂-Ti₃Al substrate, and strong bonding can thus be obtained.     

The microstructures and mechanical properties of cast Mg–Zn–Al–RE alloys

The microstructures and mechanical properties of cast Mg–Zn–Al–RE alloys with 4 wt.% RE and variable Zn and Al contents were investigated. The results show that the alloys mainly consist of α-Mg, Al₂REZn₂, Al₄RE and τ-Mg₃₂(Al,Zn)₄₉ phases, and a little amount of the β-Mg₁₇Al₁₂ phase will also be formed with certain Zn and Al contents. When increasing the Zn or Al content, the distribution of the Al₂REZn₂ and Al₄RE phases will be changed from cluster to dispersed, and the content of τ-Mg₃₂(Al,Zn)₄₉ phase increased gradually. The distribution of the Al₂REZn₂ and Al₄RE phases, and the content of β- or τ-phase are critical to the mechanical properties of Mg–Zn–Al–RE alloys. The Mg–6Zn–5Al–4RE alloy with cluster Al₂REZn₂ phase and low content of β-phase, exhibits the optimal mechanical properties, and the ultimate tensile strength, yield strength and elongation are 242 MPa, 140 MPa and 6.4% at room temperature, respectively.     

Microstructure and mechanical properties of a directionally solidified and aged intermetallic Ni–Al–Cr–Ti alloy with β-γ′-γ-α structure

Microstructure and mechanical properties were investigated in a directionally solidified (DS) Ni–21.7Al–7.5Cr–6.5Ti (at.%) alloy. The dendrites of the as-grown alloy were composed of β(B2)-matrix (NiAl), coarse γ′(L1₂)-particles (Ni₃Al), fine γ′-needles and spherical α(A2)-precipitates (Cr-based solid solution). The majority of fine γ′-precipitates was found to be twinned. The interdendritic region contained γ(A1)-matrix (Ni-based solid solution) separating ordered domains of γ′-phase and fine lath-shaped α-precipitates. Ageing in the temperature range 973–1373 K decreased the volume fraction of dendrites from about 50 vol.% measured in the as-grown material to about 38 vol.% in the material aged at 1373 K for 300 h. During ageing in the temperature range 973–1273 K the γ-phase transformed to the γ′-phase in the interdendritic region. This transformation was connected with precipitation of lath-shaped α-precipitates. Ageing at higher temperatures of 1373 and 1473 K resulted in stabilisation of the γ-phase and precipitation of spherical γ′-particles in the interdendritic region. Ageing at 973 K significantly increased the microhardness, hardness and decreased room-temperature tensile ductility. Neither ageing nor finer dendritic microstructure were found to be effective in increasing the ductility of the alloy. The measured tensile ductility up to 1.1% can be attributed to the effect of extrinsic toughening mechanisms operating in the β-phase such as blunting and bridging of cracks by the α- and γ′-precipitates.     

In situ impedance spectroscopy study of the electrochemical corrosion of Ti and Ti–6Al–4V in simulated body fluid at 25 °C and 37 °C

The corrosion resistance of Ti and Ti–6Al–4V was investigated through electrochemical impedance spectroscopy, EIS, potentiodynamic polarisation curves and UV–Vis spectrophotometry. The tests were done in Hank solution at 25 °C and 37 °C. The EIS measurements were done at the open circuit potential at specific immersion times. An increase of the resistance as a function of the immersion time was observed, for Ti (at 25 °C and 37 °C), and for Ti–6Al–4V (at 25 °C), which was interpreted as the formation and growth of a passive film on the metallic surfaces.     

The third-element effect in the oxidation of Ni–xCr–7Al (x = 0, 5, 10, 15 at.%) alloys in 1 atm O2 at 900–1000 °C

The oxidation in 1 atm of pure oxygen of Ni–Cr–Al alloys with a constant aluminum content of 7 at.% and containing 5, 10 and 15 at.% Cr was studied at 900 and 1000 °C and compared to the behavior of the corresponding binary Ni–Al alloy (Ni–7Al). A dense external scale of NiO overlying a zone of internal oxide precipitates formed on Ni–7Al and Ni–5Cr–7Al at both temperatures. Conversely, an external Al₂O₃ layer formed on Ni–10Cr–7Al at both temperatures and on Ni–15Cr–7Al at 900 °C, while the scales grown initially on Ni–15Cr–7Al at 1000 °C were more complex, but eventually developed an innermost protective alumina layer. Thus, the addition of sufficient chromium levels to Ni–7Al produced a classical third-element effect, inducing the transition between internal and external oxidation of aluminum. This effect is interpreted on the basis of an extension to ternary alloys of a criterion first proposed by Wagner for the transition between internal and external oxidation of the most reactive component in binary alloys.     

Preparation of Ti–Al–Si alloys by reactive sintering

The isothermal section of the Dy–Co–Ti system at 500 °C has been investigated in the whole composition range by means of X-ray diffraction, thermal analysis, scanning electron microscopy and energy dispersive X-ray spectroscopy. The only ternary phase DyTi_xCo_12−x is of ThMn₁₂-type structure, space group I4/mmm, and shows a small homogeneity range of 1 ≤ x ≤ 1.6. The lattice parameters for DyTi_xCo_12−x with 1 ≤ x ≤ 1.56 are a = 0.8336(4)–0.8402(1) nm and c = 0.4691(3)–0.4727(1) nm. Along a constant Dy concentration, the solid solubilities of Ti in the compounds Dy₂Co₁₇, DyCo₃, DyCo₂ and Dy₃Co are about 2.0, 2.0, 3.0 and 4.0 at.%, respectively. The TiCo phase has a homogeneity range of 50–54 at.% Co at 500 °C and dissolves up to 2.0 at.% Dy.     

Phase equilibria on Ni–Al interface under low oxygen pressure

Seventeen phases of the Ni–Al–O system at high temperatures were analyzed using thermodynamic calculations. An Ni–Al–O isothermal stability diagram was obtained from the thermochemical data. The diagram describes the interface equations for Ni/Al intermetallic compounds, Al/Al₂O₃, and Al₂O₃/Al_XNi_Y compounds, and their corresponding regions. Four univariant equilibria points and ten bivariant equilibria lines below 1126 K were obtained. The equations for the coexistence points and interface lines were also obtained. A three-domain diagram of Ni–Al–O phase arrangement at temperatures between 900 and 1191 K is shown. Thermodynamic calculations confirmed that the formation of nickel aluminate spinel (NiAl₂O₄) requires a threshold NiO activity (log a_NiO = −205.3/T − 0.347) and the partial pressure of oxygen (log P_O2=−24622/T+8 atm). In the Ni–Al–O system, a_NiO < 0.266 at 900 K, the compounds in the Ni/Al interface are formed in the order Al₃Ni_(s) → Al₃Ni_2(s) → AlNi_(s) → AlNi_3(s) → Al₂O_3(α). When a_NiO < 0.351 at 1911 K, the compounds in the Ni/Al interface are formed in the order AlNi_(s) → Al₂O_3(α).     

Thermodynamic properties of the ternary system InSb–NiSb–Sb in the temperature range 640–860 K

The ternary InSb–NiSb–Sb system has been studied by X-ray diffraction and by potentiometry. The electromotive forces (EMF) have been measured in the temperature range 640<T/K<860 by using the following galvanic cell:

with x (0.075<x<0.498) and y (0<y<0.359). The investigated samples are located on the following lines of the Gibbs triangle: InSb–Ni_0.33Sb_0.66, InSb–Ni_0.48Sb_0.52, InSb–NiSb, Sb–(InSb)_0.75(NiSb)_0.25, Sb–(InSb)_{0. 5}(NiSb)_0.5, Sb–(InSb)_0.25(NiSb)_0.75. From these measurements, the values of the partial molar thermodynamic functions (Δμ°_m,In, ΔH°_m,In, ΔS°_m,In) (data at reference pressure p⁰=10⁵ Pa), for the liquid InSb alloy, for the three solid heterogeneous regions InSb–NiSb₂–Sb, InSb–NiSb_1±δ?–NiSb₂, InSb–NiSb_1±δ, for six ternary liquid–solid alloys, have been calculated.     

Al–Ti–B grain refiners via powder metallurgy processing of Al/K2TiF6/KBF4 powder blends

The response to thermal exposure of ball-milled Al/K₂TiF₆/KBF₄ powder blends was investigated to explore the potential of PM processing for the manufacture of Al–Ti–B alloys. K₂TiF₆ starts to be reduced by aluminium as early as 220 °C when ball-milled Al/K₂TiF₆/KBF₄ powder blends are heated. The reaction of KBF₄ with aluminium follows soon after. The Ti and B thus produced are both solutionized in aluminium before precipitating out as Al₃Ti and TiB₂. All these reactions take place below the melting point of aluminium. The ball-milled Al/K₂TiF₆/KBF₄ powder blends heat treated at approximately 525 °C can be compacted to produce Al–Ti–B pellets with in situ formed Al₃Ti and TiB₂ particles. These pellets are shown to be adequate grain refiners for aluminium alloys.     

Microstructural instability and embrittlement behaviour of an Al-lean, high-Nb γ-TiAl-based alloy subjected to a long-term thermal exposure in air

Both ingot-cast and forged Ti–44Al–8Nb–1B alloys were exposed at 700 °C in air for up to 10,000 h. The α₂ lamellae in the two conditions are found to be thermodynamically unstable and readily decompose through phase transformations of α₂ → γ, α₂ → B2(ω) and α₂ + γ → B2(ω). Widespread B2(ω) forms throughout the lamellar structure, resulting in a significant increase in volume fraction after 10,000-h exposure. This is attributed to the composition similarity between the transformed and parent phases. The partition coefficients for Ti/Al/Nb between B2(ω) and α₂ and between B2(ω) and α₂ + γ are all measured to be close to 1. The long-term exposure has induced embrittlement owing to oxygen releasing from α₂ decomposition. Room-temperature ductility is only 1/5 and 1/3 of the original value for the two conditions, respectively. However, no clear decreasing trend in S–N fatigue strength is observed, suggesting that the embrittlement effect of B2(ω) on the surface crack initiation is difficult to detect.     

Nanocrystalline Al–Fe intermetallics – light weight alloys with high hardness

Nanocrystalline Al–Fe alloys containing 60–85 at.% Al were produced by consolidation of mechanically alloyed nanocrystalline or amorphous (Al85Fe15 composition) powders at 1000 °C under a pressure of 7.7 GPa. The hardness of the alloys varied between 5.8 and 9.5 GPa, depending on the Al content. The specific strength, calculated using an approximation of the yield strength according to the Tabor relation, was between 544 and 714 kNm/kg. Based on the results obtained, we infer that application of high pressure affected crystallisation of amorphous Al₈₅Fe₁₅ alloy, influencing the phase composition of the crystallisation product, and phase changes in nanocrystalline Al₈₀Fe₂₀ alloy, inhibiting them.     

The Al–R–Mg (R=Gd, Dy, Ho) systems. Part I: experimental investigation

Key experiments were carried out on the three Al–R–Mg (R=Gd,Dy,Ho) systems and the results obtained used for the thermodynamic optimisation reported in a separate paper in this issue [Caccasmani G, De Negri S, Saccone A, Ferro R. Intermetallics this issue.]. The samples were characterized by differential thermal analysis (DTA), X-ray powder diffraction (XRD), light optical microscopy (LOM), scanning electron microscopy (SEM) and quantitative electron probe microanalysis (EPMA). The isothermal sections at 400 °C are all characterized by extended homogeneity regions at a constant rare earth content. The extension of the (Mg,Al)R solid solution, cP2-CsCl type, varies with the R atomic number. Ternary compounds (τ) of Al₂(R,Mg) stoichiometry (hexagonal Laves phases with MgNi₂-type structure) have been found to exist at 400 °C in all the systems. Their temperatures of formation were detected by DTA measurements.     

The corrosion behavior of Cu–Al and Cu–Al–Be shape-memory alloys in 0.5 M H2SO4 solution

The corrosion behavior of Cu–Al and Cu–Al–Be (0.55–1.0 wt%) shape-memory alloys in 0.5 M H₂SO₄ solution at 25 °C was studied by means of anodic polarization, cyclic voltammetry, and alternative current impedance measurements. The results of anodic polarization test show that anodic dissolution rates of alloys decreased slightly with increasing the concentrations of aluminum or beryllium. Severe intergranular corrosion of Cu–Al alloy was observed after alternative current impedance measurement performed at the anodic potential of 0.6 V. However, the addition of a small amount of beryllium was effective to prevent the intergranular corrosion. The effect of beryllium addition on the prevention of intergranular corrosion is possibly attributed to the diffusion of beryllium atoms into grain boundaries, which in turn deactivates the grain boundaries.