Al Phase Diagram Research Articles

First-principles prediction of the Co–Al phase diagram including configurational, vibrational and magnetic contributions

Co-based superalloys have attracted attention to replace Ni-based superalloys in high temperature structural applications because of their higher melting temperature and corrosion resistance. Strengthening is provided by Co3(Al, X) (X = W, Cr, Mo, Ni) phases and optimization of their properties requires detailed information about the free energies of the different phases and of the phase diagram of the Co–Al system. This is achieved in this paper through first principles calculations in combination with statistical mechanics. Configurational entropic contributions to the free energy were included through Monte Carlo simulations using the cluster expansion formalism. The vibrational entropy of each phase was determined through the length-stiffness relationship while the magnetic entropy was also included through the Heisenberg Hamiltonian. The computed phase diagram is compared with the currently accepted experimental phase diagram and the different contributions to the stability of each phase are analyzed independently. The potential of this strategy to predict phase diagrams of magnetic systems is clearly established.

The growth behavior and kinetics of intermetallic compounds in Cu–Al interface at 600°C–800 °C

Due to the low production cost and light weight, copper-clad aluminum conductors show great potential value in the field of power transmission, while the phase structure and organizational changes at the Cu–Al interface in high-temperature environments have a great impact on the electrical and mechanical properties. In this paper, Cu–Al diffusion couples are prepared to study the diffusion process, microstructure, and growth kinetics of intermetallic compounds (IMCs) in Cu–Al interface during heating from 600 °C to 800 °C, and the mechanical and electrical properties of Cu–Al diffusion couples are discussed. The results show that after heat treatment, IMC layers are formed at the Cu–Al interface and the Al-side is composed of CuAl and CuAl2, in which the CuAl phase transforms from long to thick dendrites with the increase of temperature, and finally grows into a “popcorn” form. The IMCs at the Cu–Al interface from the Cu side to the Al side are in the order of Cu3Al layer, Cu9Al4 layer, Cu3Al2 layer and CuAl layer, among which the Cu3Al2 layer is the thickest (168.42 μm at 700 °C), while the Cu3Al layer is the thinnest (37.32 μm at 700 °C). The kinetic calculations show that the diffusion layer grows up rapidly, which is caused by the generation of liquid-phase during the heating process, leading to the acceleration of the diffusion of Cu and Al. The kinetic factor of the Cu3Al2 layer is the largest (n = 0.69), causing the fastest growth rate of the Cu3Al2 layer, while the Cu3Al layer shows the largest activation energy for diffusion and the smallest diffusion rate (Q = 9.48 kJ/mol, D0 = 2.06 × 10−1 m2/s), therefore its growth rate is the slowest. Based on the thermodynamic calculations and the Cu–Al phase diagram, a growth mechanism of the Cu–Al interface is proposed, the formation sequence of IMCs at the Cu–Al interface is Cu9Al4, CuAl, Cu3Al2 and Cu3Al. The formation of IMCs layer at the Cu–Al interface leads to a significant increase in the hardness of Cu–Al samples (the hardness of the Cu3Al2 is 10 times higher than that of Cu) and an obvious decrease in electrical conductivity (the electrical resistivity of Cu–Al is 5 orders of magnitude higher than that of Cu), affecting its application in the electrical industry.

Alloying Elements in Intermetallic γ‐TiAl Based Alloys – A Review on Their Influence on Phase Equilibria and Phase Transformations

Intermetallic γ‐TiAl based alloys are innovative structural materials dedicated to applications in the automotive and aeronautic industries. Especially their low density, high specific yield strength, and excellent resistance against creep and oxidation make them a suitable choice for structural high‐temperature components in combustion engines. However, further improvement of their properties and processability is required to conquer new application areas and facilitate cost‐effective production. While the incorporation of additional alloying elements is a promising possibility, their effects on the phase transformations and phase equilibria need to be considered rigorously. In this context, this review provides a detailed survey of the research work on the influence of technically important alloying elements on the Ti–Al phase diagram. First, an introduction to the fundamentals of the phase transformations in γ‐TiAl based alloys and the consequences of changed cooling conditions, relevant for example to the latest developments within the field of processing, is given. Afterward, the alloying elements, categorized with respect to their stabilization effect, are discussed and their particularities are highlighted. Topics covered include established ternary phase diagrams and isoplethal sections, the stabilization of additional phases as well as the influence of alloying elements on the microstructure of modern engineering intermetallic γ‐TiAl based alloys.

Determination of RuAl phase boundaries in binary Ru–Al phase diagram at room temperature and 1200 °C

Abstract Binary Ru–Al alloys with nominal compositions from 29.39 at.% Al up to 60.47 at.% Al were manufactured by powder-metallurgy processing. Single-phase RuAl with a homogeneous microstructure was prepared which is difficult to achieve by melting metallurgy. The microstructure and lattice parameters of the alloys were investigated by optical and scanning electron microscopy, energy-dispersive spectroscopy, wavelength-dispersive spectroscopy and X-ray diffraction. The parametric method was employed to determine the RuAl phase boundaries. Single-phase RuAl was found to exist from near-stoichiometry up to 53.8 at.% Al at room temperature and up to 54.5 at.% Al at 1200 °C, respectively. In addition, the phase Ru2Al3 was found to be stable at room temperature and to have a very narrow range of stoichiometry.

The (γ + γ′ )/γ′ phase boundary in the Ni–Al phase diagram from 600 to 1200 ° C

Abstract The phase boundary separating the γ + γ′ and γ′ regions of the Ni –Al phase diagram (the γ solvus) was established over the temperature range 600– 1200 °C. The γ phase is the Ni –Al solid solution, and γ′ is (non-stoichiometric) Ni3Al. Cast alloys containing 22.0, 22.2, 22.4, 22.6, 22.8 and 23.0 at.% Al, all compositions being within ± 0.03 at.%, were used in the experiments. At temperatures from 1000 to 1200 °C, the γ solvus was determined using aging treatments designed to eliminate the dendritic structure in the as-received castings. At lower temperatures, from 600 to ≈ 900 °C, dissolution experiments were conducted. The method takes advantage of the absence of a barrier to the dissolution of a dispersed phase. Specimens of the supersaturated γ′ phase were aged at low temperatures, 600 or 650 °C, for long times (up to 633 h), to produce fine dispersions of the coherent γ phase precipitates. These specimens were then re-aged at 10 –20 °C intervals of temperature and re-examined to determine whether the γ phase precipitates had dissolved. The absence of the γ phase precipitates enabled the phase boundary to be located to within ± 10 °C for an alloy of known composition, or ± 0.03 at.% at a given temperature, from 600 to 900 °C. The phase boundary generally agrees with the lower end of Al concentrations established in previous research.

Growth kinetics analysis of Nb–Al intermetallic compounds interfacial layers based on Nb–Al phase diagram

A model for studying the growth kinetics of intermetallic compounds by means of binary phase diagrams was proposed with some assumptions. The parabolic growth relationship of intermetallic compound is deduced mathematically using analytical solution of diffusion equations and the mass balance law. By comparing the solubility of the adjacent phase in intermetallic compound phase, the predominant parameter controlling the growth kinetics is the solute diffusion coefficient. The diffusion coefficient of Al atoms in intermetallic layers is quantitatively described and the smallest diffusion coefficient is 0.373 × 10−21 m2/s in Nb3Al phase. It is demonstrated intermetallic layers formed at interface indeed inhibit the Nb container being corroded further by TiAl melt, and Nb3Al layer is a dominant factor.

Interpretation of Fe-rich part of Fe–Al phase diagram from magnetic properties of A2-, B2-, and DO3-phases

Iron-rich Fe–Al binary phase diagram in the composition range of 23–26 at% Al shows a complex structure partly due to magnetic properties of the disordered A2-phase, the ordered B2-phase and the ordered DO 3 -phase. In this study, the disorder-order transformations between the constituent phases have been investigated through magnetization measurements of Fe- x Al alloys ( x = 22.0, 24.0, 24.3, 24.7, 26.5 at%). It is found that spontaneous magnetization of the B2-phase is larger than that of the DO 3 -phase and smaller than that of the A2-phase. The former relation disagree with previous reports but was further confirmed as the decrease of the B2-DO 3 transformation temperature under magnetic field (−1.58 K/T). It is also found that the Curie temperature of the B2-phase (836 K for Fe-24.7 at%Al alloy) coincides with that of the A2-phase; furthermore, these temperatures agree with the phase boundary between B2-region and (A2+B2)-region. • Temperature dependence of spontaneous magnetization was evaluated for the A2-, B2- and DO3-phases of Fe24.7at%Al alloy, and found that the value decreases in this order at any temperature. • Magnetic field dependence of B2-DO3 transformation temperature was evaluated for several Fe-Al alloys, and confirmed that it decreases as magnetic field increases. • Critical behavior for the magnetic transition was analyzed for several Fe-Al alloys and revealed that Curie temperature of the B2-phase almost coincides with that of the A2-phase.

Facile synthesis of micro-sized Ni–Al alloy powders through low-temperature chemical alloying

Micro-sized Ni–Al alloy powders were synthesized using different quantities of Ni and Al powders as well as AlCl3 as an activator in a quartz batch reactor at ≤773 K. This method, named low-temperature chemical alloying (LTCA), which allows Al atoms to diffuse into micro-sized nickel particles mediated by aluminum chloride (chemical promotor), is distinct from conventional processes, such as cast-and-crush, and gas atomization. As-synthesized Ni–Al alloys were characterized using different analytical techniques that included X-ray diffraction (XRD), particle size analysis, and field emission scanning electron microscopy in conjunction with energy dispersive X-ray spectroscopy (FESEM-EDS), to confirm the formation of alloy phases, such as Ni3Al, NiAl, Ni2Al3, and NiAl3. The analytical results showed that the crystalline phase compositions of the Ni–Al alloys were highly dependent upon the initial amounts of Ni and Al powders employed at the given alloying conditions (alloying temperature, 773 K; alloying time, 20 h; amount of AlCl3, 1.2 wt%). As a result of the thermal treatment of Ni–Al powder mixtures with the Al contents of (5, 15, 30, and 50) wt.% under continuous powder mixing by rotation, each powder was found to have (i) Ni solid solution + Ni3Al, (ii) Ni3Al, (iii) NiAl, and (iv) Ni2Al3 + NiAl3 phases, respectively, corresponding to the equilibrium states of the Ni–Al phase diagram. The cross-sectional analyses showed that the alloy structures of the heat-treated powders exist in a single-phase or core-shell form, depending on the number of crystalline phase compositions predicted from the phase diagram. In particular, the Ni–50 wt% Al powder has a unique Ni2Al3@NiAl3 core-shell structure. We further evaluate the performance of the as-developed Ni–Al alloy powders as oxidation-resistant materials and template materials for high-surface area (∼60 m2/gcat) nickel catalysts.

Mechanisms on Superfine Alumina Inclusions Formation by Al‐Deoxidation Reaction for liquid Iron

In order to investigate the formation mechanism of alumina particles in early stages of Al‐deoxidation reaction, the samples are prepared by Al‐deoxidation for Fe–O (90 ppm) melt in confocal laser scanning microscope (CLSM). The shapes of alumina inclusions are irregular in various samples, and their equivalent radius is between 15 and 150 nm based on the observation of field emission scanning electron microscope (FESEM) with energy dispersive spectrometer (EDS). The size and number density of alumina inclusions increase with the increase of temperature of deoxidation while decrease with the increase of cooling rate. Based on the EDS results and Fe–Al phase diagram, it is found that the Al‐deoxidation reaction zone is mainly composed of the FeAl, Fe2Al5, and FeAl3 phase. The average size of alumina particles in the FeAl phase are larger than that of in the Fe2Al5 phase. Alumina particle is not observed in the FeAl3 phase by FESEM. The observation and estimation based on the growth mechanisms of inclusions indicate that the growth of nanoscale alumina with large size is controlled by the diffusion growth and Brownian collision in high temperature melt, while nanoscale alumina with small size are secondary precipitate phase in the process of solidification.

Crystal structure, homogeneity range and electronic structure of rhombohedral γ-Mn5Al8

Abstract The γ-region of the Mn–Al phase diagram between 45 and 70 at.% Al was re-investigated by a combination of powder and single crystal X-ray diffraction as well as EDS analysis to establish the distribution of Mn and Al atoms. Single crystals of γ-Mn5–x Al8+x were grown using Sn-flux at 650 °C. The crystal structure, atomic coordinates and site occupancy parameters of γ-Mn5−x Al8+x phases were refined from single crystal X-ray data. The γ-Mn5-x Al8+x phase adopts the rhombohedral Cr5Al8-type structure rather than a cubic γ-brass structure. The refined compositions from two crystals extracted from the Al-rich and Mn-rich sides are, respectively, Mn4.76Al8.24(2) (I) and Mn6.32Al6.68(2) (II). The structure was refined in the acentric R3m space group (No.160, Z=6), in order to compare with other reported rhombohedral γ-brasses. In addition, according to X-ray powder diffraction analysis, at the Al-rich side the γ-phase coexists with LT–Mn4Al11 and, at the Mn-rich side, with a hitherto unknown phase. The refined lattice parameters from powder patterns fall in the range a=12.6814(7)−12.6012(5) Å and c=7.9444(2)−7.9311(2) Å from Al-rich to Mn-rich loadings, and the corresponding rhombohedral angles distorted from a pseudo-cubic cell were found to be 89.1(1)°−88.9(1)°. Magnetic susceptibility and magnetization studies of Mn4.92Al8.08(2) are consistent with moment bearing Mn and suggest a spin glass state below 27 K. Tight-binding electronic structure calculations (LMTO-ASA with LSDA) showed that the calculated Fermi level for γ-“Mn5Al8” falls within a pseudogap of the density of states, a result which is in accordance with a Hume-Rothery stabilization mechanism γ-brass type phases.

The Intermetallic Compound Formation for the Wire Bond Between an Al Pad and Ag-xPd Alloy Wire

Silver-palladium alloy wire has been shown as an economical and reliable substitute for gold wire in various applications in the electronic packaging industry. The success of wire bonding relies on the formation of an interfacial intermetallic compound (IMC). This study is aimed to investigate the formation behavior of IMCs between an Al pad and Ag-Pd alloy wire with various Pd concentrations of 1.0–6.0% for the as-bonded commercial Ag/Al joint. The interfacial IMCs were investigated with scanning electron microscopy and energy-dispersive x-ray spectroscopy. The IMCs formed are separate (Ag, Pd)2Al and (Ag, Pd)3Al2 for a Ag6Pd wire bond, while (Ag, Pd)2Al and (Ag, Pd)3Al2 are mixed for the other Ag(1–4.5)Pd alloy wire bonds. The thickness of the total IMC layer varies from 0.65 μm for Ag1Pd to 0.91 μm for Ag6Pd, yet a minimum of 0.44 μm exists for Ag3.5Pd. The compound formation behavior was found to correspond with the Ag–Al phase diagram. After pressure cooker tests, a less stable IMC (Ag, Pd)3Al formed at the AgxPd/Al interface.

Study of the formation of niobium aluminides in the hydride cycle

Current technology requires that newly created materials have not only a set of functional properties but also be solid, lightweight, plastic, and easily machined. Intermetallic compounds based on the Nb–Al system occupy a prominent place among the advanced materials due to their unique properties. High-temperature intermetallides based on Nb and Al (Nb3Al, Nb2Al, NbAl3) are of great interest to aerospace engineering, power industry, etc. Therefore, the studies aimed at developing new methods for producing these materials performed in the present work are highly relevant. It is shown that the hydride cycle (HC) method for synthesizing refractory metal alloys developed in our laboratory can be successfully used to produce niobium aluminides. The main regularities and mechanism of the formation of niobium aluminide in the HC are investigated. All the phases known from the Nb–Al phase diagram are obtained. The advantages of the HC method compared to the conventional are demonstrated. In addition, the combustion of the produced niobium aluminide in a hydrogen atmosphere in the SHS mode is studied. The samples are analyzed using chemical methods, differential thermal analysis, and X-ray phase analysis.

Experimental investigation of phase equilibria in the Ru–Fe–Al system at 1473 K

The low-Al part of the ternary Ru–Fe–Al phase diagram at 1473 K is established in this work. Due to the very promising properties of B2 ruthenium aluminide, the investigation of the B2 region of this system is of special interest. The experimental work includes diffusion methods, as well as quenching of annealed single-phase and two-phase alloys. The results of the different methods are in good agreement. Optical and scanning electron microscopy, energy dispersive X-ray spectroscopy and X-ray diffraction are used to investigate the samples. It is shown in this work that a three-component B2 phase exists over a wide composition range.

Observation and modeling of α-NiPtAl and Kirkendall void formations during interdiffusion of a Pt coating with a γ-(Ni-13Al) alloy at high temperature

During the last 15years, Pt-rich γ–γ′ bond-coatings have been studied extensively for their corrosion and oxidation resistance, and as a lower cost alternative to β-(Ni,Pt)Al bond-coatings in thermal barrier coating systems. To optimize their fabrication and durability, it is essential to investigate their interdiffusion with Ni-based superalloys. This study reports on experimental results and modeling of the interdiffusion of the model Pt/γ-(Ni-13Al) alloy system. Pt coatings were deposited either by electroplating or by spark plasma sintering using a Pt foil. Heat treatments at 1100°C for 15min to 10h were performed either in a high-temperature X-ray diffraction device under primary vacuum or in a furnace under argon secondary vacuum. The α-NiPtAl phase with L10 crystal structure formed very rapidly, implying fast uphill Al diffusion toward the surface. For Pt electroplating, α-phase transformed to γ′-(Ni,Pt)3Al after only 45min–1h at 1100°C. The resulting two-phased γ–γ′ microstructure remained up to 10h. When using a Pt foil coating, the continuous layer of α-NiPtAl phase disappeared after 10h and the γ′-(Ni,Pt)3Al or γ-(Ni,Pt,Al) phase appeared, resulting in two different diffusion paths in the Ni–Pt–Al phase diagram. Voids also formed at the interdiffusion zone/substrate interface for both systems after 1h or more. Composition analyses confirmed that voids were located at the Pt diffusion front corresponding to the Al-depleted zone. Experiments performed with the samples coated with a Pt foil confirmed that voids are due to a Kirkendall effect and not to the Pt deposition process. Numerical simulations including the cross-term diffusion coefficients in the diffusion flux equations reproduced the experimental concentration profiles for the γ-phased systems.

The effect of native point defect thermodynamics on off-stoichiometry in β-Mg17Al12

The mechanical strength of Mg–Al–Zn alloys can be affected by a fine spatial dispersion of β-Mg17Al12 precipitates in the Mg matrix. In an effort to understand the phase stability and the unusual asymmetric off-stoichiometry observed in β-Mg17Al12, we have performed a series of first-principles density functional theory (DFT) calculations of bulk and defect properties of Mg17Al12. Specifically, we consider native point defects (i.e. vacancies and anti-sites) in all four sublattices of Mg17Al12, i.e. 2a, 8c, 24g (Mg) and 24g (Al). The T=0K static energies of defect Mg17Al12 supercells indicate that anti-site defects are energetically favored over vacancies, and the lowest anti-site defect formation energies are in 24g sites for both AlMg and MgAl. These Al-rich and Mg-rich anti-site defect formation energies are similar in magnitude, and thus do not explain the asymmetric off-stoichiometry of Mg17Al12. We also investigate the effect of atomic vibrations via DFT phonon calculations on native point defect free energies of Mg17Al12 and combine these entropic contributions with the point defect formation energies to evaluate the thermodynamics of off-stoichiometry in this phase. We find that the formation of the AlMg anti-site is not strongly stabilized by vibrational entropy. Thus, we conclude that the observed asymmetry in the off-stoichiometry of the β-Mg17Al12 phase in the Mg–Al phase diagram is not explained by simple native point defect thermodynamics, and must involve a more complicated defect formation mechanism, such as multi-defect clustering.

Re-determination of γ/(γ + α-Mg) phase boundary and experimental evidence of R intermetallic compound existing at lower temperatures in the Mg–Al binary system

Equilibrated phase compositions at 250–400°C in the Mg–Al binary system have been determined by means of both diffusion couple technique and alloy method. The results show that the solubility of Al in the Mg17Al12 (γ) intermetallic compound is ∼36.6at.% when the Mg17Al12 phase is in equilibrium with the α-Mg solid solution, i.e., γ/(γ+α-Mg) phase boundary remarkably shifts to the Mg-rich side. The R intermetallic compound (R3¯) does stably exist at 250°C, and its composition range is 51.2–54.5at.% Al, which is completely different from the presently accepted Mg–Al phase diagram. Moreover, the composition region of β phase (Fd3¯) is determined in the range of 57.5–60.5at.% Al below 400°C.

In Situ Synchrotron Study of B19 Phase Formation in an Intermetallic γ‐TiAl Alloy

AbstractA multitude of phases exists in the binary Ti–Al phase diagram and even greater numbers are formed in structural TiAl alloys, which contain additional alloying elements to improve their properties. In the current study, a Ti–45 Al–3 Mo–0.1 B (in at%) alloy was investigated with respect to the phases occurring in chemical non‐equilibrium. In situ high‐energy X‐ray diffraction experiments enabled to identify a transient phase to be of the B19 type and to determine its temperatures of formation and dissolution.

Phase transition temperatures of Sn–Zn–Al system and their comparison with calculated phase diagrams

The Sn–Zn–Al system was studied in connection with the possible substitution of lead-based solders for temperatures up to 350 °C. Ternary alloys with up to 3 wt% of aluminium were prepared. The investigated alloys lie close to the monovariant line (eutectic valley) of the Sn–Zn–Al system. The temperatures of phase transitions of six binary Sn–Zn reference alloys and fourteen ternary Sn–Zn–Al alloys using DTA method were investigated in this paper. DTA experiments were performed at the heating/cooling rate of 4 °C min−1 using Setaram SETSYS 18TM experimental equipment. The temperatures of phase transitions in the ternary Sn–Zn–Al system were obtained, namely, the temperature of ternary eutectic reaction T E1 (197.7 ± 0.7 °C), temperature of ternary transition reaction T U1 (278.6 ± 0.7 °C), temperatures of liquidus and other transition temperatures for studied alloys. Temperatures obtained during DTA heating runs were used as authoritative. DTA curves obtained during cooling enabled realising better differentiation of the obtained overlapped heat effects (peaks) during heating. Theoretical isopleths of the Sn–Zn–Al phase diagram were calculated using the Thermocalc software and MP0602 thermodynamic database. Experimental data were compared with the calculated temperatures, and a good agreement was obtained.

Effect of Al content on porous Ni–Al alloys

Porous Ni–Al alloys with different nominal compositions ranging from Ni–10 wt.% Al to Ni–40 wt.% Al were fabricated through reactive synthesis of Ni and Al elemental powders. It had been found that the volume expansion, maximum pore size and permeability of porous Ni–Al alloys increased with increasing the Al content, indicating that the nature of the pores can be manipulated through changing the Al content. In addition, detailed structural characterizations showed that the fabricated porous Ni–Al alloys can have three crystalline phases (i.e., Ni 3Al, NiAl, and Ni 2Al 3) when using different compositions, in good agreement with the Ni–Al phase diagram. Mechanical strength tests of the fabricated porous Ni–Al alloys were also performed. Both of the bending strength and tensile strength of porous Ni–Al alloys were in the range of 5–80 MPa, and decreased with increasing the Al content.

Destructive hydrogenation and recombination of α2-Ti3Al alloy: production of solid nanocomposites and coatings with improved properties. I. Mechanism of destructive hydrogenation of Ti3Al

It is shown that α2-Ti3Al alloy and hydrogen interact at 298–1533 K in accordance with the destructive hydrogenation reaction whose products are titanium hydride and intermetallic compounds rich in aluminum. The structural mechanism of α2-Ti3Al destructive hydrogenation is as follows. Hydrogen dissolves in Ti3Al, thus leading to disordering of its crystalline structure. Then hydrogen interacts with titanium, but not with the disordered alloy, according to the following scheme: α-Ti (hcp) → α-TiHx (hcp) → β-TiHy (bcc) → TiH1.9 (fcc) → TiH2 (bct). Intermetallic compounds rich in aluminum show up when new stoichiometric concentrations of titanium and aluminum form in accordance with the Ti–Al phase diagram. Solid composites with nanosized components result from the destructive hydrogenation of Ti3A at 973 K for 1 and 4 h. Their phase compositions are Ti3Al–(α-TiH–Ti3–xAl) and ((α-Ti–Hx)–(β-Ti-Hy)–TiH1.9 (fcc)–TiH2 (bct))–(Ti3Al–TiAl–TiAl2), respectively.