Two types of S phase precipitates in Al–Cu–Mg alloys
Introduction
The high strength AA2024 alloys (nominal compositions around Al–4.2 Cu–1.5 Mg–0.6 Mn (wt.%)) are primary structural materials used in the aircraft industry. The derivatives AA2x24 (x = 1–5), with stricter limits on Fe and Si impurity levels as well as other recent improved versions of AA2024, have improved toughness, fatigue resistance and strength. These alloys are specified for many aerospace structural applications, such as in fuselage structures, wing tension members, shear webs and ribs. The AA2x24 series alloys are strengthened by microstructure evolution during ageing.
In the early 1950s, Bagaryatsky [1] first proposed a four-stage precipitation sequence for the ageing of Al–Cu–Mg alloys:where SSS stands for supersaturated solid solution and GPB stands for Guinier–Preston–Bagaryatsky (a term popularised by Silcock [2] and subsequent workers). Bagaryatsky [1] considered the GPB zone to be a short-range ordering of Cu and Mg solute atoms. The structure of S″/GPB2 has not been clearly confirmed. Orthorhombic [1], [3], [4], [5], cubic [2], tetragonal [6], [7] and monoclinic [8] structures have been proposed for the S″/GPB2 structure. An S phase with a composition of Al2CuMg has been determined as an orthorhombic Cmcm structure with lattice parameters aS = 0.400 nm, bS = 0.923 nm, cS = 0.714 nm by Perlitz and Westgren [9] (PW) on the basis of X-ray diffraction (XRD) work. The S phase is an equilibrium phase and is incoherent with the Al matrix. Other models for the S phase have been reported [4], but the most accepted structure for the S phase is the PW model. The S′ phase has generally been considered as semi-coherent with the matrix, possessing the same structure as the S phase but with slightly different lattice parameters of , , [1], [10]. As the proposed S′ structures have essentially the same crystal structures as the S phase, with very small differences in lattice parameters, many recent publications make no distinction between the S′ and S phases. On the other hand, Charai et al. [8] suggested that the asymmetric shape of the differential scanning calorimetry (DSC) peaks in the range of 290–370 °C in Al–2.1 Cu–1.3 Mg–0.09 Zr–0.04 Si (wt.%) indicated the presence of both S′ and S phases: orthorhombic semi-coherent S′ was reported to possess lattice parameters , , and the incoherent equilibrium S phase was reported with aS = 0.405 nm, bS = 0.921 ± 0.006 nm,cS = 0.719 ± 0.012 nm.
The S and S′ phases have been reported to precipitate in the form of laths with the {1 2 0}Al habit planes elongated along the 〈1 0 0〉Al direction [1], [11]. Both phases have been suggested to possess the same orientation relationship (OR) with aluminium matrix as follows [1], [8]:where the subscript Al stands for aluminium matrix. The OR between the distorted S″ phase and the matrix was proposed as follows [1], [12]:Based on the matrix transformation by Li and Yan [13], the (OR2) is equivalent to:Since the mid-1990s, atom probe field ion microscopy (APFIM) and three-dimensional atom probe (3DAP) have shown that the first-stage of age-hardening is due to the formation of co-clusters [14], [15], [16], [17]. Ringer et al. [14], [15] proposed the following three-stage precipitation sequences for the ageing of Al–Cu–Mg alloys:where the Cu–Mg co-clusters are responsible for the initial hardening, GPB zone is the dominant precipitate at peak strengthening and the S phase appears in the softening stage.
By a combination of transmission electron microscopy (TEM) and DSC, Wang et al. [18], [19] confirmed the existence of S″ [4] and confirmed that the peak strengthening in quenched and subsequently stretched Al–Cu–Mg alloys corresponds to S phase. Thus the following three-stage precipitation sequence was proposed:where the S phase is responsible for the peak strengthening. S″/GPB2 was considered to provide only limited contribution to strengthening for alloys with substantial Cu content [20].
The S precipitates commonly observed are lath-shaped, having habit planes of {2 1 0}Al//(0 1 0)S and lath axis of {1 0 0}Al//(1 0 0)S, and their orientation relationship with Al matrix satisfies (OR1). But, using high-resolution electron microscopy (HREM), Zhang et al. [21] first observed some S′(S) precipitates with the habit plane between the and in Al–4.43 Cu–2.00 Mg–0.53 Mn (wt.%) after water quenching and aged at 250 °C for 6 h. Radmilovic et al. [22] observed a similar type of precipitates, in the Al–2.01 Cu–1.06 Mg–0.14 Zr–0.08 Fe (wt.%) after quenching into ice brine and ageing for 72 h at 190 °C, which had habit planes of {2 1 0}Al//(0 4 3)S and were designated as Type II precipitates. The precipitates which followed (OR1) were designated as Type I. Both types of S phase precipitates had the same lattice parameters of aS = 0.403 ± 0.001 nm, bS = 0.930 ± 0.013 nm, cS = 0.708 ± 0.01 nm [22]. However, the relationship between the two types of S precipitates has not been given. Majimel et al. [23], studying an Al–2.66 Cu–1.85 Mg–0.2 Fe–0.21 Si–0.34 Mn–0.23 Ni–0.09 Ti (wt.%) (AA2650) alloy, observed that the differences between the two orientation variants ranged over several degrees, with most of the data falling into a range of about 4–6°. The following orientation relationship was satisfied:In recent work on a quenched and aged Al–2.96 Mg–0.42 Cu–0.12 Si–0.25 Mn–0.21 Fe (wt.%) alloy, Kovarik et al. [24] also reported the existence of Type II precipitates. As (OR3) has the same orientation relationship as (OR2a), they referred to the Type II S phase as S″ phase, which is the distorted S structure proposed by Bagaryatsky [1]. Energy-dispersive X-ray spectrometry (EDS) and atom probe tomography analysis indicated that the Type II S phase was slightly enriched in Si [24]. The driving force for the formation of Type II was explained to result from an invariant line transformation strain [22], or related to a competition between the length of the (0 0 1)S and (021)S interfaces [24].
Throughout the various works cited above the sequence in which these S phase variants appeared is unclear and little is know about the factors determining the predominance of either of the variants. Clarification of this is essential to understand the strengthening in Al–Cu–Mg alloys, and is the purpose of this study. To identify the effect of Si on the formation of Type II precipitates, we compared the commercial AA2024 alloy and its high purity variant AA2324 alloy by DSC and TEM.
Section snippets
Experimental procedures
The chemical compositions of the AA2024 and AA2324 alloys studied are given in Table 1. The Si and Fe contents are close to the maximum allowable for the two alloys. The alloys were supplied in the form of 13 mm plate in T351 (solution treated, quenched, stretched 1.5–3%) condition.
Table 2 shows the heat treatments and working performed on the samples. For all treatments slices of initially about 3 mm thickness and diameter about 5 mm were used. All solution treatments were performed at 495 °C.
Results
Fig. 1 shows the DSC curves of the AA2324 alloy in three different conditions. Five main effects may be identified in these thermograms [25], [26], [27]: an exothermic peak, A, between 60 and 110 °C, is due to the formation of co-clusters [17], [20]; an endothermic effect, B, between 160 and 240 °C, may be attributed to Cu–Mg co-cluster dissolution (with possibly some GPB2 dissolution); an exothermic effect, C (containing two overlapping reactions; see below), between about 230 and 340 °C, is
Formation temperatures for Types I and II S phase
TEM analysis of interrupted DSC experiments on the AA2324 alloy show that the peaks at 274 and 300 °C (Fig. 1) are caused by the formation of two distinct types of S phase precipitates, Type I and Type II, respectively, with an ∼4° difference in orientation relationship (Fig. 2, Fig. 3). Zhang et al. [21] first reported the S precipitates with OR different from the common (OR1) in a sample of an Al–4.43 Cu–2.00 Mg–0.53 Mn (wt.%) alloy that was water quenched and aged at 250 °C for 6 h. Radmilovic
Conclusions
TEM and DSC have been used to study the formation of two variants of S phase precipitation in an Al–4.2 Cu–1.5 Mg–0.6 Mn–0.5 Si (AA2024) and an Al–4.2 Cu–1.5 Mg–0.6 Mn–0.08 Si (AA2324) (wt.%) alloy. In DSC experiments on the 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. The combination of DSC and TEM experiments in Al–Cu–Mg alloys confirm the
References (31)
- et al.
Scripta Met
(1984) - et al.
Mater Sci Eng A
(2004) - et al.
Acta Mater
(2004) Acta Mater
(2001)- et al.
Acta Mater
(2000) - et al.
Mater Sci Eng
(1983) - et al.
Acta Mater
(1996) - et al.
Scripta Mater
(1998) - et al.
Mater Sci Eng A
(2004) - et al.
Scripta Mater
(2006)
Acta Mater
Acta Mater
Mater Sci Eng A
Thermochim Acta
Dokl Akad SSSR
Cited by (293)
Inhibiting segregation enabled outstanding combination of mechanical and corrosion properties in precipitation-strengthened aluminum alloys
2024, Journal of Materials Science and TechnologyHeat treatment response and mechanical properties of a Zr-modified AA2618 aluminum alloy fabricated by laser powder bed fusion
2023, Journal of Alloys and CompoundsDevelopment and applications of aluminum alloys for aerospace industry
2023, Journal of Materials Research and TechnologyMicrostructure and mechanical properties of 2524 aluminum alloy with dislocation loops by various quenching rates
2023, Materials Science and Engineering: A