. Although many suggestions for the existence of pre-precipitates in Al-Cu-Mg alloys with a Cu/Mg atomic ratio close to 1 have been made, a critical review reveals that evidence exists for only two truly distinct ones. The precipitation sequence is best represented as: supersaturated solid solutionRco-clustersRGPB2/S"RS where clusters are predominantly Cu-Mg co-clusters (also termed GPB or GPB I zones), GPB2/S" is an orthorhombic phase that is coherent with the matrix (probable composition Al 10 Cu 3 Mg 3 ) for which both the term GPB2 and S" have been used, and S phase is the equilibrium Al 2 CuMg phase. GPB2/S" can co-exist with S phase before the completion of the formation of S phase. It is further mostly accepted that the crystal structure of S' (Al 2 CuMg) is identical to the equilibrium S phase (Al 2 CuMg). The Perlitz and Westgren model for S phase is viewed to be the most accepted structure. 3DAP analysis showed that Cu-Mg clusters form within a short time of natural and artificial aging. Cu-Mg clusters and S phase contribute to the first and second stage hardening during aging. In Al-Cu alloys, the h phase (Al 2 Cu) has I4/ mcm structure with a50 . 607 nm and c50 . 487 nm, and h' phase with tetragonal structure and a50 . 404 nm, c50 . 58 nm, the space group is I4 m2. Gerold's model for h" (or GPII) appears to be favourable in terms of free energy, and is consistent with most experimental data. The transformation from GPI to GPII (or h") seems continuous, and as Cu atoms will not tend to cluster together or cluster with vacancies, the precipitation sequence can thus be captured as: supersaturated solid solutionRh" (Al 3 Cu)Rh' (Al 2 Cu)Rh (Al 2 Cu). The V phase (Al 2 Cu) has been variously proposed as monoclinic, orthorhombic, hexagonal and tetragonal distorted h phase structures. It has been shown that V phase forms initially on {111} Al with c50 . 935 nm and on further aging, the c lattice parameter changes continuously to 0 . 848 nm, to become identical to the orthorhombic structure proposed by Knowles and Stobbs (a50 . 496 nm, b50 . 858 nm and c50 . 848 nm). Other models are either wrong (for example, monoclinic and hexagonal) or refer to a transition phase (for example, the Garg and Howe model with c50 . 858 in a converted orthorhombic structure). For Al-Li-Cu-Mg alloys, the L1 2 ordered metastable d' (Al 3 Li) phase has been observed by many researchers. The Huang and Ardell model for T 1 phase (space group P6/ mmm, a50 . 496 nm and c50 . 935 nm), appears more likely than other proposed structures. Other proposed structures are perhaps due to the T 1 phase forming by the dissociation of Kn110m dislocations into 1/6n211m Shockley partials bounding a region of intrinsic stacking fault, in which copper and lithium enrichment of the fault produces a thin layer of the T 1 phase.

Hot rolled Al-6Li-1Cu-1Mg-0.2Mn (at.%) (Al-1.6Li-2.2Cu-0.9Mg-0.4Mn, wt.%) and Al-6Li-1Cu-1Mg-0.03Zr (at.%) (Al-1.6Li-2.3Cu-1Mg-0.1Zr, wt.%) alloys developed for age forming were studied by tensile testing, electron backscatter diffraction (EBSD), three-dimensional atom probe (3DAP), transmission electron microscopy (TEM) and differential scanning calorimetry (DSC). For both alloys, DSC analysis shows that ageing at 150°C leads initially to formation of zones/clusters, which are later gradually replaced by S phase. On ageing at 190°C, S phase formation is completed within 12 h. The precipitates identified by 3D atom probe and TEM can be classified into (a) Li-rich clusters containing Cu and Mg, (b) a plate-shaped metastable precipitate (similar to GPB2 zones/S''), (c) S phase and (d) δ' spherical particles rich in Li. The Zr containing alloy also contains β' (Al 3 Zr) precipitates and composite β'/δ' particles. The β' precipitates reduces recrystallisation and grain growth leading to fine grains and subgrains.