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Influence of Silicon Addition on Precipitation Behavior in an Al-Cu-Mg Alloy
Abstract
Influence of silicon (Si) addition on aging precipitation behavior of an Al-4Cu-1.3Mg alloy was investigated by hardness measurements and TEM analysis. The results indicated that each of the Si-containing alloys exhibits an enhanced age-hardening response, a cubic phase, previously designated σ phase, was obtained. The σ phase was showed to be a (semi-)coherent and coplanar phase with the Al matrix, i.e.,{100}σ//{100}Al and <100>σ//<100>Al. In addition, S’ phase was observed. Si addition played an important role on the formation of σ phase. σ phases were observed in the Si containing alloy for all the aging conditions studied but was absent in Si-free alloy. The content of Si significantly influenced the volume fraction of σ phase in the alloy. Prolongation of aging time will be conducive to precipitation of a greater number of σ phases.
... Therefore, to make sure the precipitation behavior of σ phases becomes a concern. To meet the heat-resistance produ requirements stringent of cts for applications, an novel aerospace Al-Cu-Mg-Si alloy containing a cubic σ phase with excellent high temperature stability was developed [10,11]. ...
To meet the stringent requirements of heat-resistance products for aerospace applications, an Al-4.0Cu-1.3Mg-0.6Si alloy containing a cubic σ phase (Al5Cu6Mg2) with excellent high temperature stability was developed. Influence of deformation on aging precipitation behavior of the alloy was investigated by transmission electron microscope (TEM) and differential scanning calorimetry (DSC). The results indicated that cold deformation had great effect on aging precipitation behavior in the alloy. When the alloy was processed a pre-stretching deformation before aging, the precipitation of σ phase was restrained while the S′ phase (Al2CuMg) was promoted. And the nucleation sites played an important role on them.
Vortex-free high-speed stir casting (VFHSC) methodology can achieve uniform dispersion of particles in melt without air entrapment for fabricating particle reinforced composites, and it has been proved to be feasible for preparing micron-composites by this methodology. In this work, in order to study deeply on particles in nano-size magnitude in composites by this methodology, the preparation of 1.5 vol.% Nano-Al2O3p/Al–Cu–Mg–Si composite is also investigated. The proper stirring parameters for ideal particle dispersion are determined to prepare the materials. Porosity of the composite can be limited to 0.147 % under the VFHSC methodology. The composition and microstructure of ingots, including the incorporation of Al2O3 particles as well as the morphology of precipitated phases, are examined by OM, XRD, SEM, TEM, HRTEM and EDS. The nano-particles are incorporated ideally in the matrix with restricted aggregation and sedimentation, and the well-bounded Al2O3–Al interface possesses semi-coherent interface. Moreover, the VFHSC 1.5 vol.% Nano-Al2O3p/Al–Cu–Mg–Si composite exhibits obvious strengthening, limited ductility reduction, higher hardness as well as better wear resistance than those of matrix, validating the efficacy of the VFHSC methodology on fabricating 1.5 vol.% Nano-Al2O3p/Al–Cu–Mg–Si composite. The work proves that incorporating nano-particles in Al matrix by VFHSC methodology is feasible and efficient. The work presented in this paper proposes a viable approach for the fabrication of nanocomposites using the stir casting method, thereby offering valuable insights for further research on stir casting technology.
The Thermo-Calc software is used for analyzing the Al – Cu –Mg – Mn – Si system as applied to alloys of the 2xxx series including computation of isothermal and polythermal sections. It is shown that the solidus temperature of alloys D16, D1, AK6 and AK8 differs by 22, 38, 40 and 38°C, respectively, within the standardized compositions. The total fraction of excess phases may differ by more than a factor of 5 in alloy AK8 depending on the concentrations of the alloying elements. The data of the computation show that the standardized compositions of alloys of series 2xxx have to be corrected.
The effect of Si additions on vacancy behavior following solution treatment and quenching was studied for an Al-Cu-Mg-Ag alloy system using positron annihilation lifetime spectroscopy (PALS). Both the initial positron lifetime and steady state positron lifetime increase as the Si concentration increases. This behavior is interpreted in terms of Si interacting with vacancies, leading to their retention in the alloy. The ability of PALS to monitor vacancy behavior after solution treatment should allow the improved prediction of phase transformation kinetics and hence the tailoring of ageing treatments.
A century has elapsed since Alfred WiIm made the accidental discovery of age hardening in an aluminium alloy that became known as Duralumin. His work, and the gradual realization that hardening arose because of the presence of fine precipitates which provided barriers to the motion of dislocations, is a good example of the transition of metallurgy from an art to a science. A brief account is given of the development of age hardenable aluminium alloys and the way that modern experimental techniques allow precipitation processes to be understood on an atomic scale. Some contemporary issues in age hardening are then discussed.
Tests at 130 °C and 150 °C have shown that the creep resistance of an Al–Cu–Mg–Ag alloy is significantly increased if it is heat-treated at an elevated temperature to an underaged condition rather than the fully hardened, T6 temper. This beneficial effect of underageing is manifest in reduced rates of secondary creep. Similar results have been obtained for the commercial alloy 2024. Delays at ambient temperature after underageing and before testing lead to secondary precipitation and a progressive decrease in creep performance that eventually reverts to close to that for the T6 condition. This detrimental effect may be overcome by slow cooling from the underageing temperature, which arrests or impedes subsequent secondary precipitation. Microstructural observations suggest that the enhanced creep resistance in the underaged condition is a consequence of the presence of “free” solute in solid solution that is not yet involved in precipitation.
The microstructure evolution during ageing treatment at 170 and 190°C of AA2009/SiC composites, reinforced with 15vol.% 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.
Microstructure evolution of an age hardenable Al–3Mg–0.4Cu–0.12Si (wt%) alloy has been studied during artificial aging at 180 °C prior to the formation of the stable S-phase. The primary investigation method used in this study was high-resolution transmission electron microscopy (HRTEM), coupled with image processing and image simulation. After 1 h of aging, the presence of super-lattice reflections was detected in the Fourier spectra of the HRTEM images, suggesting an L10 type ordering of Mg and Cu atoms in the Al matrix. After 4 and 8 h of aging, coherent particles were observed in the microstructure. These particles give rise to diffraction spots that in previous literature have been considered to be characteristic of the S″-phase in the “Cu-lean” Al–Mg–Cu alloys. It is shown that these diffraction spots can be indexed in terms of a crystal structure that is closely related to the L10 ordering formed at the shorter aging times. The crystal structure is orthorhombic with lattice parameters a=1.2 nm, b=0.4 nm, c=0.4 nm and space group Cmmm. We propose to identify these coherent particles as GPB-II zones, and the ordering that precedes them as GPB zones.
Transmission electron microscopy (TEM) and differential scanning calorimetry (DSC) have been used to study S phase precipitation in an Al-4.2Cu-1.5Mg-0.6Mn-0.5Si (AA2024) and an Al-4.2Cu-1.5Mg-0.6Mn-0.08Si (AA2324) (wt.%) alloy. In DSC experiments on as solution treated samples two distinct exothermic peaks are observed in the range 250-350°C, whereas only one peak is observed in solution treated and subsequently stretched or cold worked samples. Samples heated to 270°C and 400°C at a rate of 10°C/min in the DSC have been studied by TEM. The selected area diffraction patterns show that S phase precipitates with the classic orientation relationship form during the lower temperature peak, and for the solution treated samples, the higher temperature peak is caused by the formation of a second type of S phase precipitates which have an orientation relationship that is rotated by ~4o the classic one. The effects of Si and cold work on the formation of second type of S precipitates have been discussed.
An Al-4.3 wt% Cu-2.0 wt% Mg alloy reinforced with 20 vol% reinforcing fibres was examined after a T7 heat treatment. The expected precipitate phase was equilibrium S (Al2CuMg), which was confirmed to form in the monolithic alloy. However when this Al-Cu-Mg alloy was squeeze-cast into a fibre preform and given an identical T7 heat treatment a number of other phases also nucleated; these included (Al2Cu), (Mg2Si) and the cubic phase (Al5Cu6Mg2). These additional phases were determined to nucleate and grow rapidly during the water-quench following solution treatment. The existence of excess Si (approximately 0.5 wt%) in the matrix was determined to be responsible for nucleation of these additional phases. This extra Si entered the composite matrix during squeeze-casting through breakdown of an SiO2 layer which existed at the fibre interfaces. During quenching Si clusters rapidly form and provide nucleation sites for the and phases. The Si clusters apparently created a compressive strain in the matrix which attracted a high concentration of small Cu atoms to their interface. The phase nucleated in this high-Cu region since, on a localized scale, became the equilibrium phase. This type of nucleation process may also explain the enhanced precipitate nucleation which occasionally takes place in other alloy systems when trace amounts of certain elements are added.
The role of silicon in the precipitation of the phase (Al5Cu6Mg2) has been investigated through comparative studies on Al-3.63Cu-1.67Mg (wt%) and Al-3.63Cu-1.67Mg-0.5Si alloys. Both alloys were extensively examined after solution treating at 525C for 2.5 h followed by ageing at 265C for times up to 650 h. Limited studies were also undertaken on both alloys after ageing at 200 and 305C. Precipitation of was observed in Al-3.62%Cu-1.66%Mg-0.5%Si for all ageing conditions studied but was absent in Si-free Al-3.62%Cu-1.66%Mg. In addition, S and phases were observed in both alloys. The volume fraction of phase in the Si containing alloy was substantially reduced by a pre-age stretch followed by ageing for 24 h at 265C with S being the dominant precipitate type. The volume fraction of phase in the Si containing alloy was lower after ageing 24 h at 200C than after 24 h at 265 and 305C. Peak hardness was higher for the Si free alloy on ageing at 200 and 265C, but the Si free alloy softened more rapidly, reflecting the more rapid coarsening kinetics of S compared with .
The individual influence of small additions of Ag and Si on the nucleation of the Ω (i.e., a chemically modified coherent form of ϑ-Al2Cu) and σ (Al5Cu6Mg2) phases, respectively, in Al-Cu-Mg alloys has been known. These phases nucleate directly and exhibit reduced rates of coarsening
at the commercial aging temperatures. Alloys containing a uniform and fine distribution of these phases may, therefore, be
of interest for further investigation for applications at temperatures below 200 °C. In a recent study, using an Al-Cu-Mg-Mn-Ag-Si
alloy aged at 180 °C, it was shown that the Ω phase formed as a major precipitate nucleated predominantly upon the Mn-bearing
dispersoids, while the σ phase was present as a minor one. This article describes the conditions under which widespread nucleation of σ phase may occur at unidentified sites in the matrix as well as upon the Mn-bearing dispersoids in the alloy. Widespread nucleation
of σ phase begins in the alloy following the onset of dissolution of Guinier-Preston-Bagaryatski (G-P-B) zones that form early
in the aging cycle as the major precipitate. It is established that nucleation of σ as a major phase precipitate requires a critical minimum supersaturation of Si in the solid solution. This article further
points out that several constituent phases (implying those which form as solidification products and survive the homogenization
treatment) together with the Mn-bearing dispersoids dissolve Si, thereby considerably reducing the Si supersaturation in the
solid solution. The implications of these observations are discussed in view of the available data on the nucleation of σ phase in such alloys.