Al-Cu-Mg-Ag Alloys

The Al-Cu-Mg system with a high Cu-to-Mg ratio and a minor addition of Ag is the base system of the new high-strength aluminum alloys that are currently receiving major research interest for potential aerospace applications. Since the minor amount of Ag atoms has a significant influence on the microstructure and the mechanical properties, the role of Ag atoms in the microstructural evolution has been the subject of many investigations.

The Al-Cu-Mg system with a high Cu-to-Mg ratio and a minor addition of Ag is the base system of the new high-strength aluminum alloys that are currently receiving major research interest for potential aerospace applications. Since the minor amount of Ag atoms has a significant influence on the microstructure and the mechanical properties, the role of Ag atoms in the microstructural evolution has been the subject of many investigations.

The aluminum alloys of Al-Cu-Mg type have been known for decades. Repeated attempts have been made to improve this classic precipitation-hardening alloy by further additions and to optimize its properties for the particular application. To improve the strength properties, alloying of casting alloys of this type with silver has been proposed. Similar proposals have also been made in the field of wrought alloys. To improve the microstructure, the alloys also contain further additions, for example manganese, titanium and the like.

These alloys resulted in most cases from corresponding casting alloys with added nickel. However, since they suffer a rather substantial decrease in strength above 150°C, they cannot really be described as "high-temperature" aluminum alloys in the contemporary sense. In particular, they do not meet the requirements at relatively high temperatures (up to, for example, 250°C.), which are necessary for numerous industrial uses. There is therefore a great demand for a further improvement in wrought aluminum alloys, in particular in their strength properties at elevated temperature.

An objective was to provide a wrought aluminum alloy which can be produced by fusion metallurgy in simple conventional processes and which, in the temperature range from 0° to 250°C and in precipitation-hardened state has markedly higher strength properties than conventional alloys.

A wrought Al-Cu-Mg-type aluminum alloy of high strength in the temperature range between 0° and 250°C, consists essentially of the composition given in Table 1.

Table 1: Chemical composition Al-Cu-Mg-Ag.

Alloy Cu % Mg % Ag % Mn % Zr % V % Si % Al %
1 5.0-7.0 0.3-0.8 0.2-1.0 0.3-1.0 0.1-0.25 0.05-0.15 <0.10 Rem.
2 5.5-6.5 0.4-0.6 0.2-0.8 0.3-0.8 0.1-0.2 0.05-0.15 <0.05 Rem.
3 6.0 0.5 0.4 0.5 0.15 0.10 <0.10 Rem.

In an Al-Cu-Mg-Ag alloy, thin, hexagonally shaped plate-like precipitates, designated as Ω, precipitate uniformly on the {111} planes. The uniform dispersion of this precipitate as well as θ’ plates on the {001} planes is thought to contribute to age hardening in this alloy.

The structure and the chemical composition of the phase have been investigated by selected-area electron diffraction (SAED), high-resolution electron microscopy and convergent-beam electron diffraction, and it has been found that the Ω phase has a composition of Al2Cu with a structure similar to the θ phase. As to the pre-precipitation stage of the phase, evidence for co-clustering of Ag and Mg atoms has been found by the conventional atom-probe field ion microscopy.

More recent investigation using three-dimensional atom probe (3DAP) has clarified the evolution process of the phase Ω as follows. At the initial stage of isothermal ageing at 180°C, co-clusters of Mg and Ag atoms form without the incorporation of Cu atoms. By subsequent ageing, the co-clusters evolve to {111} plate when Cu atoms are incorporated in the clusters. These plates eventually evolve to a distinct phase by rejecting Ag and Mg atoms from the interior of the Ω to the Ω/α interface. Although the evolution process from the co-clusters to the Ω phase was clarified as described above, the mechanism of the formation of initial Ag-Mg co-clusters is still unknown.

It has been known for some time that the addition of low concentrations of silver to Al-Cu-Mg alloys can result in an increased hardening on ageing between 150 and 250°C. It has also been shown that on ageing silver-containing alloys with high copper to magnesium ratios a fine distribution of plate-like precipitates is formed on { 111} planes, and it is conjectured that it is these precipitates which lead to the increased hardness of the alloys. This precipitated phase has been designated n by Chester and Polmear, and Polmear and Cooper have recently described how Al-Cu-Mg alloys with compositions in the α+θ+S or α+θ regions of the phase diagram can be modified by addition of silver to give alloys which contain strengthening distributions of both Ω and θ'.

These precipitated phases are easily distinguished in transmission electron microscope images because the well-known θ'' phase forms preferentially as large plates on {100} planes. Wrought aluminum alloys containing about 6% Cu, 0.5% Mg and 0.5% Ag can have an excellent range of mechanical and electrical properties as a result of the presence of these phases.

However, while the structure of the Ω phase is now well established, its composition is only poorly characterized. It has been generally assumed that these precipitates must contain both silver and magnesium, and Taylor et al. have suggested that they nucleate as Mg3Ag Guinier Preston (GP) zones and grow by the collection of copper and aluminium atoms to attain an overall stoichiometry of approximately Al2Cu.

More recently, Cousland and Tate identified only GP zones of nominal composition MgAg in the early stages of decomposition of AI-Mg-Ag alloys. In addition, these authors show that silver may segregate to the interface between the n phase and the matrix, and that the 0' precipitates contain neither silver nor magnesium.

Precipitation of Ω increases the spall strength and decreases localized shear through its multiple cutting interactions with dislocations at the matrix interface. Dispersed particles also increase the strength of the alloy in high strain-rate applications by resisting localized shear.

Aluminum can be alloyed with small amounts of elements such as Cu, Mg, Mn, Li, Zn, or Si to increase the strength of the naturally soft material. The relatively low density of aluminum makes it beneficial in comparison with other metals for numerous civilian and military applications. Its high strength makes it suitable for shock loading, and its toughness at different temperatures makes it beneficial for aerospace applications.

Quite generally, the additional impurities, which have to be accepted in industrial manufacture of the alloys, should be kept as low as possible and should not exceed a total value of 0.25% by weight for all elements taken together. The silicon content should be kept as low as possible in order to avoid the formation of low-melting eutectics in the grain boundaries.

Moreover, intermetallic compounds with magnesium, which would represent a loss of the latter metal for its advantageous effect in conjunction with silver, should be avoided. For this reason, the silicon content should remain below 0.10% by weight. The transition metals manganese, zirconium and vanadium are intended for grain refinement and for the formation of intermetallic phases which, in a finely divided form, effect dispersion-hardening and above all contribute to an increase in high-temperature strength.

Further additions of iron, nickel and chromium, having similar effects, to the claimed alloy compositions are feasible. However, these elements have the disadvantage that they form additional intermetallic compounds with copper, so that the content of this later element available for the precipitation hardening and the strength of the matrix is reduced. In any case, caution is advisable in the use of iron and/or nickel, which can at most be added in contents from 0.1 to 1.5% by weight as a maximum.

Aluminum-copper-magnesium-silver (Al-Cu-Mg-Ag) alloys that were developed for thermal stability also offer attractive ambient temperature strength-toughness combinations, and therefore, can be considered for a broad range of airframe structural applications.

High-strength, low-density Al-Cu-Mg-Ag alloys were initially developed to replace conventional 2000 (Al-Cu-Mg) and 7000 (Al-Zn-Cu-Mg) series aluminum alloys for aircraft structural applications. During the High Speed Civil Transport (HSCT) program, improvements in thermal stability were demonstrated for candidate aircraft wing and fuselage skin materials through the addition of silver to Al-Cu-Mg alloys based on Al 2519 chemistry.

Thermal stability of the resulting Al-Cu-Mg-Ag alloys, C415-T8 and C416-T8, was due to co-precipitation of the thermally stable Ω (AlCu) and θ' (Al2Cu) strengthening phases. The strength and toughness behavior was investigated for these alloys produced as 0.090-inch thick rolled sheet in the T8 condition and after various thermal exposures. The mechanical properties were shown to be competitive with conventional aircraft alloys, 2519-T8 and 2618-T8.

During the Integral Airframe Structure (IAS) program, advanced aluminum alloys were examined for use in an integrally stiffened airframe structure where the skin and stiffeners would be machined from plate and extruded frames would be mechanically attached. The advantages of integrally stiffened structure include reduced part count, and reduced assembly times compared to conventional built-up airframe structure. The near-surface properties of a thick plate are of significance for a machined integrally stiffened airframe structure since this represents the skin location. Properties measured at the mid-plane of the plate are more representative of the stiffener web.

The use of aluminum alloys for commercial and military applications has increased substantially due to the alloys’ low areal density, toughness, and processability. It has recently been shown that the aluminum alloy 2139 with copper, magnesium, and silver can be significantly toughened and strengthened by combinations of θ’ and Ω precipitates and dispersed manganese particles.

April, 2010
Resuelve tus desafíos de materiales.
Descubre cómo podemos ayudarte.