Intermetallic phases formed during DC-casting

The Al-Fe and Al-Fe-Si particles formed during DC-casting of an Al-0.25 wt pct Fe-0.13 wt pct Si alloy have been examined. The particles were analyzed by transmission electron microscopy (TEM) and energy dispersive spectroscopy of X-rays (EDS). Crystal faults were studied by high resolution electron microscopy (HREM). Samples for electron microscopy were taken at various positions in the ingot,i.e., with different local cooling rates during solidification. At a cooling rate of 6 to 8 K/s the dominating phases were bcc α-AlFeSi and bct Al _m Fe. The space group of bcc α-AlFeSi was verified to be Im3. Superstructure reflections from Al _m Fe were caused by faults on {110}-planes. At a cooling rate of 1 K/s the dominating phases were monoclinic Al₃Fe and the incommensurate structure Al _x Fe. In Al₃Fe, stacking faults on {001} were frequently observed. The structure of Al _x Fe is probably related to Al₆Fe. Some amounts of other phases were detected. For EDS-analysis, extracted particles mounted on holey carbon films were examined. Extracted particles were obtained by dissolving aluminum samples in butanol. Accurate compositions of various Al-Fe-Si phases were determined by EDS-analysis of extracted crystals.     

Intermetallic phases formed during DC-casting of an Al−0.25 Wt Pct Fe−0.13 Wt Pct Si alloy

The Al−Fe and Al−Fe−Si particles formed during DC-casting of an Al-0.25 wt pct Fe-0.13 wt pct Si alloy have been examined. The particles were analyzed by transmission electron microscopy (TEM) and energy dispersive spectroscopy of X-rays (EDS). Crystal faults were studied by high resolution electron microscopy (HREM). Samples for electron microscopy were taken at various positions in the ingot,i.e., with different local cooling rates during solidification. At a cooling rate of 6 to 8 K/s the dominating phases were bcc α-AlFeSi and bct Al _m Fe. The space group of bcc α-AlFeSi was verified to be Im3. Superstructure reflections from Al _m Fe were caused by faults on {110}-planes. At a cooling rate of 1 K/s the dominating phases were monoclinic Al₃Fe and the incommensurate structure Al _x Fe. In Al₃Fe, stacking faults on {001} were frequently observed. The structure of Al _x Fe is probably related to Al₆Fe. Some amounts of other phases were detected. For EDS-analysis, extracted particles mounted on holey carbon films were examined. Extracted particles were obtained by dissolving aluminum samples in butanol. Accurate compositions of various Al−Fe−Si phases were determined by EDS-analysis of extracted crystals.     

Effects of air atomisation,spray deposition and melt overflow techniques on microstructural features of Al–Fe–V–Si alloy

Abstract

Microstructural features of atomised powders, spray deposited preforms and melt overflow strips of 8009 series AlFeVSi alloy were investigated to reveal the microstructural evolution associated with the processing condition variations. X-ray diffraction (XRD), optical microscopy and scanning electron microscopy (SEM) techniques were used to identify the crystal structures of dispersoid phases in the specimens. The primary intermetallic phase was characterised as bcc α-Al(Fe,V)Si having a lattice parameter in the range 12·51–12·53 Å, other phases are identified as icosahedral, Al₃Fe and Al₁₃(Fe,V)₄ by XRD. SEM was used to examine the morphological changes and quantitative element analysis to reveal the chemical composition of these phases in specimens. Various phase morphologies such as starlike, band shape and needlelike intermetallics were observed in these products, whereas the primary phase crystal structure is cubic and stable.     

Processing,microstructure, and properties of laser-clad Ni alloy FP-5 on Al alloy AA333

This article describes the results of laser cladding Ni alloy FP-5 on Al alloy AA333, microstructure and crystal structure characterization, and properties of the clad evaluated by Vickers hardness measurement and wear testing. Direct cladding of Ni alloy on Al alloy creates brittle Ni _x Al _y compounds in the interface, which make the interface very brittle, and result in cracking at the interface. The compound formation is avoided by introducing an intermediate layer of Cu or bronze. The cracking tendency of the clad is prevented by preheating the substrate to 673 K. The microstructure and crystal structure of the clad and interface are investigated by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). Five phases in the clad layer (including three new phases) and two phases in the interface are identified by convergent beam electron diffraction (CBED) and selected area diffraction (SAD) studies. The mechanical properties of the laser-clad Ni alloy are evaluated by Vickers hardness measurements and wear testing, which show superior results over Cu- and Fe-based alloys.     

The structure of the high-temperature phase MnAl(h) and the displacive transformation from MnAl(h) into Mn5Al8

Splat-cooling investigations, as well as heat treatment at temperatures corresponding to the stability range of the high-temperature phase MnAl(h), show that this phase cannot be retained by quenching. Therefore, high-temperature X-ray powder diffraction was used for the determination of the crystal structure. MnAl(h) is isotypic to tungsten,a _l230K = 0.3063 nm, and Im-3m. The displacive transformation from MnAl(h) into Mn₅Al₈ (Cr₅Al₈ type) was studied by means of X-ray diffraction and optical and scanning electron microscopy (SEM). For the low-temperature phase Mn₅Al₈, the axial ratioc/a, as well as the volume of the unit cell, was investigated in dependence on mole fraction. Crystal chemical parameters of phases isotypic to Cr₅Al₈ are compared.     

Phase equilibria in the Mg?Al?Ca ternary system at 773 and 673 K

The isothermal sections of the Mg−Al−Ca ternary system at 773 and 673 K were determined by phase analysis with electron-probe microanalysis (EPMA) and transmission electron microscopy (TEM). The C36 phase exists between the C14 (Mg₂Ca) and C15 (Al₂Ca) phases, and its stoichiometry is close to Mg₂Al₄Ca₃. The α-Mg phase equilibriates with the C14 and C36 phases at 773 K, but with C14, C15, and β phases at 673 K, due to the decomposition of the C36 phase into C14 and C15 phases. These intermetallic phases have significant solid-solubility in the ternary system.     

Phase equilibria in the Mg?Al?Ca ternary system at 773 and 673 K

The isothermal sections of the Mg−Al−Ca ternary system at 773 and 673 K were determined by phase analysis with electron-probe microanalysis (EPMA) and transmission electron microscopy (TEM). The C36 phase exists between the C14 (Mg₂Ca) and C15 (Al₂Ca) phases, and its stoichiometry is close to Mg₂Al₄Ca₃. The α-Mg phase equilibriates with the C14 and C36 phases at 773 K, but with C14, C15, and β phases at 673 K, due to the decomposition of the C36 phase into C14 and C15 phases. These intermetallic phases have significant solid-solubility in the ternary system.     

Phase equilibria in the Mg-Al-Ca ternary system at 773 and 673 K

The isothermal sections of the Mg-Al-Ca ternary system at 773 and 673 K were determined by phase analysis with electron-probe microanalysis (EPMA) and transmission electron microscopy (TEM). The C36 phase exists between the C14 (Mg₂Ca) and C15 (Al₂Ca) phases, and its stoichiometry is close to Mg₂Al₄Ca₃. The α-Mg phase equilibrates with the C14 and C36 phases at 773 K, but with C14, C15, and β phases at 673 K, due to the decomposition of the C36 phase into C14 and C15 phases. These intermetallic phases have significant solid-solubility in the ternary system.     

Phase Constituents and Microstructure of Interaction Layer Formed in U-Mo Alloys <Emphasis Type="Italic">vs</Emphasis> Al Diffusion Couples Annealed at 873 K (600 °C)

U-Mo dispersion and monolithic fuels are being developed to fulfill the requirements for research reactors, under the Reduced Enrichment for Research and Test Reactors program. In dispersion fuels, particles of U-Mo alloys are embedded in the Al-alloy matrix, while in monolithic fuels, U-Mo monoliths are roll bonded to the Al-alloy matrix. In this study, interdiffusion and microstructural development in the solid-to-solid diffusion couples, namely, U-15.7 at. pct Mo (7 wt pct Mo) vs pure Al, U-21.6 at. pct Mo (10 wt pct Mo) vs pure Al, and U-25.3 at. pct Mo (12 wt pct Mo) vs pure Al, annealed at 873 K (600 °C) for 24 hours, were examined in detail. Scanning electron microscopy (SEM), transmission electron microscopy (TEM), and electron probe microanalysis (EPMA) were employed to examine the development of a very fine multiphase interaction layer with an approximately constant average composition of 80 at. pct Al. Extensive TEM was carried out to identify the constituent phases across the interaction layer based on selected area electron diffraction and convergent beam electron diffraction (CBED). The cubic-UAl₃, orthorhombic-UAl₄, hexagonal-U₆Mo₄Al₄₃, and cubic-UMo₂Al₂₀ phases were identified within the interaction layer that included two- and three-phase layers. Residual stress from large differences in molar volume, evidenced by vertical cracks within the interaction layer, high Al mobility, Mo supersaturation, and partitioning toward equilibrium in the interdiffusion zone were employed to describe the complex microstructure and phase constituents observed. A mechanism by compositional modification of the Al alloy is explored to mitigate the development of the U₆Mo₄Al₄₃ phase, which exhibits poor irradiation behavior that includes void formation and swelling.     

The apparent ’five-fold’ nature of large T2 (AI6Li3Cu) crystals

LargeT ₂ (Al₆Li₃Cu) crystals which display apparent five-fold symmetry have been studied using several different electron microscopy techniques. Electron microprobe X-ray analysis was used to determine the phases present in two different alloys cast with bulk compositions close to stoi-chiometricT ₂. Convergent beam electron diffraction (CBED) of these crystals indicates that they do not display five-fold rotational symmetry, and TEM images and selected area diffraction from thin foil specimens indicate a microcrystalline or twinned structure is responsible for the apparent five-fold symmetry ofT ₂. Microdiffraction patterns obtained from the thinnest regions of the crystals are consistent with a cubic structure having a lattice parameter of ≈0.40 nm.     

Resistance-Spot-Welded AZ31 Magnesium Alloys: Part I. Dependence of Fusion Zone Microstructures on Second-Phase Particles

A comparison of microstructural features in resistance spot welds of two AZ31 magnesium (Mg) alloys, AZ31-SA (from supplier A) and AZ31-SB (from supplier B), with the same sheet thickness and welding conditions, was performed via optical microscopy, scanning electron microscopy (SEM), X-ray diffraction (XRD), and transmission electron microscopy (TEM). These alloys have similar chemical composition but different sizes of second-phase particles due to manufacturing process differences. Both columnar and equiaxed dendritic structures were observed in the weld fusion zones of these AZ31 SA and SB alloys. However, columnar dendritic grains were well developed and the width of the columnar dendritic zone (CDZ) was much larger in the SB alloy. In contrast, columnar grains were restricted within narrow strip regions, and equiaxed grains were promoted in the SA alloy. Microstructural examination showed that the as-received Mg alloys contained two sizes of Al₈Mn₅ second-phase particles. Submicron Al₈Mn₅ particles of 0.09 to 0.4 μm in length occured in both SA and SB alloys; however, larger Al₈Mn₅ particles of 4 to 10 μm in length were observed only in the SA alloy. The welding process did not have a great effect on the populations of Al₈Mn₅ particles in these AZ31 welds. The earlier columnar-equiaxed transition (CET) is believed to be related to the pre-existence of the coarse Al₈Mn₅ intermetallic phases in the SA alloy as an inoculant of α-Mg heterogeneous nucleation. This was revealed by the presence of Al₈Mn₅ particles at the origin of some equiaxed dendrites. Finally, the columnar grains of the SB alloy, which did not contain coarse second-phase particles, were efficiently restrained and equiaxed grains were found to be promoted by adding 10 μm-long Mn particles into the fusion zone during resistance spot welding (RSW).     

Evolution and coarsening of Al2O3 dispersoids at 500 °C to 600 °C in a mechanically alloyed Al-Ti-O based material

The formation and coarsening of Al₂O₃ dispersoids have been investigated at 500 °C, 550 °C, and 600 °C in a mechanically alloyed (MA) extrusion of composition Al-0.35wt pct Li-1wt pct Mg-0.25wt pct C-10vol pct TiO₂ for times up to 1500 hours. In the as-extruded condition, the dispersed phases included Al₃Ti, Al₄C₃, MgO, cubic TiO (C-TiO), monoclinic TiO (M-TiO), TiO₂, and a small amount of Al₂O₃. However, numerous Al₂O₃ dispersoids (various polymorphs: η, γ, α, and δ) with “block-shaped” morphology were formed after heat treatment due to reduction of C-TiO, M-TiO, and TiO₂. Transmission electron microscopy (TEM) and X-ray diffraction (XRD) showed conclusively the transformation of these phases to additional Al₂O₃ and Al₃Ti. High resolution TEM showed that the α-Al₂O₃ dispersoids exhibited some lattice matching with the α-Al matrix. Coalescence of the block-shaped Al₂O₃ dispersoids occurred after heat treatment, and Al₄C₃ also became attached to them. The length and width of the block-shaped Al₂O₃ dispersoids increased by a factor of ∼1.55 between 340 and 1500 hours at 600 °C.     

On the nature of the Fe-bearing particles influencing hard anodizing behavior of AA 7075 extrusion products

The deleterious effects of Fe-bearing constituent particles on the fracture toughness of wrought Al alloys have been known. Recent studies have shown that the presence of Fe-bearing constituent particles is also detrimental to the nature and growth of the hard anodic oxide coating formed on such materials. The present study, using a combination of scanning electron microscopy (SEM), transmission electron microscopy (TEM), and electron probe microanalysis (EPMA), was made to examine the influence of the nature of the Fe-bearing particles on the hard anodizing behavior of AA 7075 extrusion products containing varying amounts of Si, Mn, and Fe impurities. It was found that, in the alloy containing 0.25 wt pct Si, 0.27 wt pct Mn, and 0.25 wt pct Fe, the Fe-bearing constituent particles are based on the Al₁₂(FeMn)₃Si phase (bcc with a=1.260 nm). These particles survive the hard anodizing treatment, add resistance to the electrical path, causing a rapid rise in the bath voltage with time, and cause a nonuniform growth of the anodic oxide film. In the materials containing 0.05 wt pct Si, 0.04 wt pct Mn, and 0.18 wt pct Fe, on the other hand, the formation of the Al₁₂(FeMn)₃Si-based phase is suppressed, and two different Fe-bearing phases, based on Al-Fe-Cu-Mn (simple cubic with a=1.265 nm) and Al₇Cu₂Fe, respectively, form. Neither the Al-Fe-Cu-Mn-based phase nor the Al₇Cu₂Fe-based phase survive the hard anodizing treatment, and this results in a steady rise in the bath voltage with time and a relatively uniform growth of the anodic oxide film. Consideration of the size of the Fe-bearing particles reveals that the smaller the particle, the more uniform the growth of the anodic oxide film.     

An analytical electron microscopy study of constituent particles in commercial 7075-T6 and 2024-T3 alloys

To better understand the role of constituent particles in pitting corrosion, analytical electron microscopic studies were performed on the constituent particles in commercial 7075-T6 and 2024-T3 alloys. Five phases, namely, Al₂₃CuFe₄ and amorphous SiO₂ in 7075-T6 and Al₂CuMg, Al₂Cu, and (Fe,Mn) _x Si(Al,Cu) _y in 2024-T3, were identified. The crystal structure and chemistry of the Al₂₃CuFe₄, Al₂CuMg, and Al₂Cu phases in these alloys are in good agreement with the published data. Small deviations from their stoichiometric compositions were observed and are attributed to the influence of alloy composition on the phase chemistry. For the (Fe,Mn) _x Si(Al,Cu) _y (approximately, x=3 and y=11) phase, a rhombohedral structure, with lattice parameter a=b=c=1.598 nm and α=β=γ=75 deg, was identified and is believed to be a modified form of either Al₈Fe₂Si or Al₁₀Mn₃Si. Information from this study provided technical support for studying the electrochemical interactions between the individual particles (or phases) and the matrix. The corrosion results are reported in a companion article.     

Microstructure and properties of duplex δ-Al3(Ti, V)/β-(Ti, V) alloys

The microstructures of multiphase intermetallic alloys with compositions Al₇₀Ti₁₀V₂₀ and Al₆₂Ti₁₀V₂₈ based on the trialuminide Al₃Ti have been characterized, following chill casting and postsolidification heat treatment, using a combination of scanning electron microscopy and transmission electron microscopy (TEM). Evidence of a eutectic reaction of the form L → δ-Al₃(Ti, V)+ζ-Al₈V₅, not previously reported in the Al-Ti-V system, has been observed in both alloys solidified at sufficient levels of undercooling. The ζ phase is replaced by metallic β-(Ti, V) phase during subsequent heat treatment in the range 1073 to 1273 K, and differential thermal analysis (DTA) of samples prean-nealed at 1173 K revealed an endothermic peak at ∼1560 K, consistent with equilibrium eutectic melting of the form (δ+β) → L. Although the chill-cast alloys retained metastable intermediate high-temperature phases, duplex metallic-intermetallic microstructures, containing uniform fine-scale distributions of metallic β-(Ti, V) solid solution in a δ-Al₃(Ti, V) intermetallic matrix, have been produced in both alloys during isothermal heat treatments at temperatures in the range 1073 to 1273 K. For both alloys, the bulk Vickers hardness of such microstructures remained in excess of that of binary Al₃Ti, while in the Al₆₂Ti₁₀V₂₈ alloy, where the increased volume fraction of β phase took the form of a near-continuous network within δ matrix, there was evidence arising from indentation tests of a substantial improvement in the cracking resistance compared to both chill-cast ternary alloy and binary Al₃Ti.     

Identification of precipitate phases in a mechanically alloyed rapidly solidified Al-Fe-Ce alloy

The stable and metastable precipitate phases which form during mechanical alloying of a rapidly solidified (RS) Al-8.4Fe-3.4Ce alloy have been unambiguously identified using X-ray diffraction, transmission electron microscopy (TEM), and energy dispersive spectroscopy techniques. The metastable Al₁₀Fe₂Ce and stable Al₁₃Fe₃Ce and Al₁₃Fe₄ intermetallic phases, with crystal structures and lattice parameters as reported in the literature, have been identified. It is shown that the metastable Al₁₀Fe₂Ce intermetallic phase particles have elongated shapes and their sizes range between 100 and 200 ran and are free of any localized faults, whereas the equilibrium Al₁₃Fe₃Ce and Al₁₃Fe₄ intermetallic phases are equiaxed in shape and have particle sizes ranging from 200 to 500 nm. It is suggested that the presence of the metastable Al₁₀Fe₂Ce in this material is due to its incomplete transformation to the equilibrium Al₁₃Fe₃Ce phase.     

Microstructural characterization of the dispersed phases in Al-Ce-Fe system

Analytical electron microscopy studies were conducted on a rapidly solidified Al-8.8Fe-3.7Ce alloy and arc melted buttons of aluminum rich Al-Fe-Ce alloys to determine the characteristics of the metastable and equilibrium phases. The rapidly solidified alloy consisted of binary and ternary metastable phases in the as-extruded condition. The binary metastable phase was identified to be Al₆Fe, while the ternary metastable phases were identified to be Al₁₀Fe₂Ce and Al₂₀Fe₅Ce. The Al₂₀Fe₅Ce was a decagonal quasicrystal while the Al₁₀Fe₂Ce phase was determined to have an orthorhombic crystal structure belonging to space group Cmmm, Cmm2, or C222. Microscopy studies of RS alloy and cast buttons annealed at 700 K established the equilibrium phases to be Al₁₃Fe₄, Al₄Ce, and an Al₁₃Fe₃Ce ternary phase which was first identified in the present study. The crystal structure of the equilibrium ternary phase was determined to be orthorhombic with a Cmcm or Cmc2 space group. The details of X-ray microanalysis and convergent beam electron diffraction analysis are described.