A novel Al–Ti–B alloy for grain refining Al–Si foundry alloys

Al–Ti–B refiners with excess-Ti (Ti:B > 2.2) perform adequately for wrought aluminium alloys but they are not as efficient in the case of foundry alloys. Silicon, which is abundant in the latter, forms silicides with Ti and severely impairs the potency of TiB₂ and Al₃Ti particles. Hence, Al–Ti–B alloys with excess-B (Ti:B < 2.2) and binary Al–B alloys are favored to grain refine hypoeutectic Al–Si alloys. These grain refiners rely on the insoluble (Al,Ti)B₂ or AlB₂ particles for grain refinement, and thus do not enjoy the growth restriction provided by solute Ti. It would be very attractive to produce excess-B Al–Ti–B alloys which additionally contain Al₃Ti particles to maximize their grain refining efficiency for aluminium foundry alloys. A powder metallurgy process was employed to produce an experimental Al–3Ti–3B grain refiner which contains both the insoluble AlB₂ and the soluble Al₃Ti particles. Inoculation of a hypoeutectic Al–Si foundry alloy with this grain refiner has produced a fine equiaxed grain structure across the entire section of the test sample which was more or less retained for holding times up to 15 min.     

Effect of strontium on the grain refining efficiency of Mg–3Al alloy refined by carbon inoculation

The effect of Sr on the grain refining efficiency of the Mg–3Al alloy refined by carbon inoculation has been investigated in the present study. A significant grain refinement was obtained for the Mg–3Al alloy treated with either 0.2% C or 0.2% Sr. The Al–C–O particles were found in the sample refined by 0.2% C, and the element O should come from reaction between Al₄C₃ nuclei of Mg grains and water during the process of sample preparation. The grain size of the sample refined by carbon inoculation was further decreased after the combined addition of Sr. The grain size decreased with increasing Sr content. Much higher refining efficiency was obtained when the Sr addition was increased to 0.5%. Sr is an effective element to improve the grain refining efficiency for the Mg–Al alloys refined by carbon inoculation. The number of Al₄C₃ particles in the sample refined by the combination of carbon and Sr was more than that in the sample refined by only carbon. No Al–C–O–Sr-rich particles were obviously found in the sample refined by the combination of carbon and a little (<0.5%) Sr addition.     

Al–Ti–C–Sr master alloy—A melt inoculant for simultaneous grain refinement and modification of hypoeutectic Al–Si alloys

Present article is focused on the microstructural features of Al–Ti–C–Sr master alloy, an inoculant for simultaneous grain refinement and modification of hypoeutectic Al–Si alloys. This master alloy is basically a metal matrix composite consisting of TiC and Al₄Sr phases formed in situ in the Al-matrix. TiC particles initiate the refinement of primary α-Al through heterogeneous nucleation in molten hypoeutectic Al–Si alloy, while Al₄Sr phase dissolves in molten Al–7Si alloy enriching the melt with Sr, which eventually leads to modification of eutectic silicon during solidification of the Al–7Si alloy casting. Thus present master alloy serves in both ways, as a grain refiner and a modifier for hypoeutectic Al–Si alloys.     

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.     

The effects of lanthanum on microstructure and electrochemical properties of Al–Zn–In based sacrificial anode alloys

In order to improve the non-uniform corrosion of Al–0.5Zn–0.03In–1Mg–0.05Ti alloys, Al–5Zn–0.03In–1Mg–0.05Ti–xLa (x = 0.3, 0.5 and 0.7 wt.%) alloys were developed. Microstructures and electrochemical properties of the alloys were investigated. The results show that the optimal microstructures and electrochemical properties are obtained in Al–5Zn–0.03In–1Mg–0.05Ti–0.5La alloy. The main precipitate phase is Al₂LaZn₂ particles. The excellent electrochemical properties of Al–5Zn–0.03In–1Mg–0.05Ti–0.5La alloy is mainly attributed to fine grains and grain boundaries containing fine Al₂LaZn₂ precipitates. At the same time the fine grains can improve the non-uniform corrosion of Al–0.5Zn–0.03In–1Mg–0.05Ti alloy.     

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.     

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 analysis of alloys Ti–Al, Ti–V, Al–V and Ti–Al–V

Thermodynamic analysis of three binary Ti-based alloys: Ti–Al, Ti–V, and Al–V, as well as ternary alloy Ti–Al–V, is shown in this paper. Thermodynamic analysis involved thermodynamic determination of activities, coefficient of activities, partial and integral values for enthalpies and Gibbs energies of mixing and excess energies at four different temperatures: 2000, 2073, 2200 and 2273 K, as well as calculated phase diagrams for the investigated binary and ternary systems. The FactSage is used for all thermodynamic calculations.     

Effect of Zr, Mn and Sc additions on the grain size of Mg–Gd alloy

The effect of Zr, Mn and Mn + Sc additions on the grain size of Mg–10Gd alloy has been investigated and the grain refinement mechanisms are also suggested. The results reveal that the addition of Zr results in a significant grain refinement of as-cast Mg–10Gd alloy by generating nucleants. However, it cannot restrict grain growth during homogenization treatment at 520 °C, and completely loss the grain refining effect for extruded alloy sample. Mn has a negligible effect on grain size of as-cast Mg–10Gd alloy, but α-Mn particles precipitate during homogenization treatment, which helps to refine the grains of extruded alloy sample due to α-Mn particles restricting recrystallization grain growth during extrusion. Successful grain refinement can be obtained by the addition of Mn + Sc. It is effective to refine microstructure of as-cast Mg–10Gd alloy, inhibit grain growth during homogenization treatment and also have a significant grain refining effect on extruded Mg–10Gd alloy sample, which are ascribed to the precipitation of a large number of Mn₂Sc particles.     

Effect of manganese on the microstructure of Mg–3Al alloy

The effect of manganese on the microstructure of Mg–3Al alloy, especially the nucleation efficiency of Al–Mn particles on primary Mg, has been investigated in this paper. Mg–0.72Mn was used to fabricate Mg–3Al–xMn (x = 0, 0.1, 0.2, 0.3, 0.4 and 0.5) alloys, and the grain sizes of these alloys fluctuate at 390 μm indicating addition of manganese does not evidently influence the grain size of Mg–3Al alloy. Through XRD, FESEM and TEM detection, it is found that Al_0.89Mn_1.11 compound is the dominant Al–Mn phase in Mg–3Al–0.3Mn, Mg–3Al–0.4Mn and Mg–3Al–0.5Mn, and distributes in primary Mg matrix and interdendritic regions with an angular blocky morphology. The number of Al_0.89Mn_1.11 increases gradually with increasing manganese content while the grain sizes of primary Mg are nearly the same in Mg–3Al, Mg–3Al–0.3Mn, Mg–3Al–0.4Mn and Mg–3Al–0.5Mn, indicating Al_0.89Mn_1.11 has low nucleation efficiency on primary Mg.     

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.     

Effects of the addition of Pd on the hydrogen absorption–desorption characteristics of Ti33V33Cr34 alloys

The hydrogen absorption–desorption performance of the body-centered-cubic (bcc) Ti–V–Cr–Pd alloys have been investigated. Ti₃₃V₃₃Cr₃₄ ingots with 0, 0.05, 0.5 at.% Pd were prepared by arc melting. X-ray diffraction (XRD) revealed that all of these alloys were homogeneous bcc solid solutions. Pd-containing (0.05, 0.5 at.% Pd) Ti–V–Cr alloys have better initial activation properties than those without Pd, and the desorption plateau pressure of the (Ti₃₃V₃₃Cr₃₄)_99.5Pd_0.5 alloy was substantially higher than that of the alloy without Pd. It is also found that the hysteresis difference is smaller in these alloys and degradation of hydrogen absorption capacity becomes steady after the 25th cycling test. (Ti₃₃V₃₃Cr₃₄)_99.5Pd_0.5 alloy exhibits large hydrogen absorption and desorption capacity of up to 3.42 and 2.07 mass% at 353 K, respectively.     

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.     

Cu-segregation at the Q/α-Al interface in Al–Mg–Si–Cu alloy

Cu segregation was detected at the Q^′/α-Al interface in an Al–Mg–Si–Cu alloy by energy-filtered transmission electron microscopy. By contrast, in a Cu-free Al–Mg–Si alloy no segregation was observed at the interface between the matrix and Type-C precipitate.     

Effects of temperature on pressure infiltration of liquid Al and Al–12wt.%Si alloy into packed SiC particles

The effects of temperature on pressure infiltration of packed SiC particles by liquid Al and the eutectic Al–12wt.%Si alloy have been investigated over the temperature range 923–1273 K. Below 1173 K, the threshold pressure P₀ for infiltration of pure Al slightly (and linearly) decreases with temperature. At that temperature a sharp drop occurs, followed by a faster decrease of P₀. The contact angle θ derived from the threshold pressure data shows a similar behavior. Results for the Al–Si eutectic alloy are very similar for temperatures below 1173 K, but, above this temperature, the threshold pressure decreases faster than for pure Al. These results are consistent with contact angle data derived from sessile drop experiments. In addition compact permeability is significantly reduced for T > 1173 K, a result likely related to the reaction between aluminum and SiC.     

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.     

Passive films and corrosion resistance of Ti–Hf alloys in 5% HCl solution

Ti–Hf alloys with Hf contents of 10, 20, 30, and 40 mass% were prepared by a tri-arc furnace and homogenized at 1273 K for 21.6 ks, and then these alloys were cold rolled into 3-mm-thick plates. The alloy specimens were subjected to a solution treatment in a vacuum at 1223 K for 3.6 ks and then rapidly quenched in ice water before corrosion tests. The corrosion resistance of these alloys was investigated by studying the anodic polarization curves at 310 K in 5% HCl solution to determine the potential use of these alloys in biomedical applications. The passive films formed on the surfaces of the alloys were examined by scanning electron microscopy (SEM) and X-ray photoelectron spectroscopy (XPS) analysis. The results reveal that all the Ti–Hf alloys exhibit a passive behavior in 5% HCl solution, which is attributed to the passive film formation of a mixture of both HfO₂ and TiO₂. The corrosion resistance of the Ti–Hf alloys gently increases with the Hf content and the Ti–Hf alloys exhibit better corrosion resistance than commercial pure (CP) Ti—the currently-used metallic biomaterial.     

Precipitation kinetic of Al3(Sc,Zr) dispersoids in aluminium

The non-isothermal formation of Sc- and Zr-containing dispersoids in Al–Zr–Sc ternary alloys has been investigated by atom probe tomography (APT). In the early stages of precipitation, a high number density of Sc-rich clusters form. These clusters subsequently transform into Al₃Sc particles with a L1₂ structure. When Zr diffusion becomes significant, Zr atoms are found to segregate to Al₃Sc/α-Al matrix interfaces. Further annealing at 748 K gives rise to a duplex core/shell structure.     

Structural properties of Li–borate glasses doped with Sm and Eu ions

A Li–borate glass system doped with samarium and europium has been prepared by a conventional melt quenching technique. Europium content was kept constant at 0.01 mol%. The general formula was

xSm₂O₃ + (100 − x) [0.84B₂O₃ + 0.15Li₂O + 0.01Eu₂O₃]

where x = 0.1, 0.2, 0.5, 0.6 and 0.7 mol%. The density was measured and the corresponding molar volume was evaluated .The former was found to increase by increasing Sm while the later exhibits opposite trend. The average optical basicity, Λ_th, electron negativity χ_2av and electron polarizability α²⁻ were calculated for the prepared compositions. Infrared spectra were obtained at room temperature for the prepared glasses before and after γ irradiation. The results showed that the three main appeared bands are most likely due to the bending and/or stretching vibration of both tetrahedral BO₄ and triagonal BO₃ borate units. ESR spectra were recorded at room temperature before and after γ-irradiation. It was found that the oxygen atoms of BO₃ units are responsible for the formation of the hole paramagnetic centers after irradiation in the glass matrix.     

Isothermal section of the Gd–Co–V ternary system at 773 K

The isothermal section of the phase diagram of the Gd–Co–V ternary system at 773 K was investigated by X-ray powder diffraction (XRD), metallographic analysis, electron probe microanalysis, and differential thermal analysis (DTA) techniques. The isothermal section consists of 14 single-phase regions, 26 two-phase regions and 13 three-phase regions. The solid solubilities of V in the compounds Co₁₇Gd₂, Co₃Gd, Co₂Gd, Co₇Gd₁₂ and CoGd₃ were about 10.0, 2.0, 6.0, 1.2 and 5.3 at.% V, respectively. It was found that there are some homogeneity range in the only ternary compound of GdCo_12−xV_x with x = 2.6–3.7 at 773 K. No solubility of Gd in compounds Co₃V, σCoV or CoV₃ was observed. There is no solubility of V in Co₇Gd₂ or Co₃Gd₄ observed at 773 K.