Quench Rate Effects on the Natural Aging Behavior of 7xxx Al-mg-zn-cu Aluminum Alloys

The effect of quench rate on the natural aging behavior was examined using microhardness, conductivity and the differential scanning calorimeter. The data indicated that GPZ formed during extended natural aging of 7075, while extended aging of 7050 resulted in the precipitation of η’. Introduction The aluminum alloy system Al-Zn-Mg-Cu is the mainstay of aluminum alloys used in the aerospace industry. This alloy system is being used in fatigue and fracture critical applications in military and commercial aircraft. With increasing applications, and shrinking budgets, it is important that the most performance be obtained from any alloy system. Strength, as well as affordability, must be achieved for modern aircraft applications. An understanding of precipitation during quenching and aging can be understood by nucleation theory applied to diffusion controlled solid-state reactions [1]. The kinetics of precipitation during quenching are dependent on the degree of solute supersaturation and the diffusion rate. As an alloy is quenched, there is greater supersaturation (assuming no solute precipitates) and the diffusion rate decreases with decreasing temperature. When either the supersaturation or the diffusion rate is low, the precipitation rate is low. At intermediate temperatures, the amount of supersaturation is relatively high, as is the diffusion rate. Therefore the precipitation rate is the greatest at intermediate temperatures. The amount of time spent in this critical temperature range is governed by the quench rate and quench path. Quenching, and the cooling effect of quenchants have been extensively studied [2] [3] [4] [5]. The first systematic attempt to correlate properties to the quench rate in Al-Zn-Mg-Cu alloys was performed by Fink and Wiley [6] for thin (0.064") sheet. A Time-Temperature-Tensile Property curve was created and was probably the first instance of a TTT diagram for aluminum. It was determined that the critical temperature range for 75S is 400°C to 290°C. This is similar to the critical temperature range found for Al-Zn-Mg-Cu alloys [7]. At quench rates exceeding 450°C/sec., it was determined that the maximum strength and corrosion resistance were obtained. At intermediate quench rates of 450 to 100°C/sec., the strength obtained is lowered (using the same age treatment), but the corrosion resistance was unaffected. Between 100°/sec and 20°C/sec, the strength decreased rapidly, and the corrosion resistance is at a minimum. At quench rates below 20°C/sec, the strength decreases rapidly, but the corrosion resistance improved. After quenching, parts are generally artificially aged to produce either peak strength, or overaged to produce improved corrosion properties. This is a complex process, and requires an understanding of vacancies, and the interaction of precipitation and metastable phases. In general, the sequence of precipitation occurs by clustering of vacancies, formation of GP Zones; nucleation of η ́; precipitation of η, and finally the coarsening of the precipitates. The effect of quench rate on the formation of Frank and prismatic dislocation loops has been studied [8] [9] [10]. The number of dislocation loops increased when the specimens were water quenched. The formation of loops is promoted by a small amount of solute atoms, and this effect depends on the kind of solute atoms and the difference of atomic radii between the trace element and aluminum. As the difference in solute radii increases, the formation of loops is promoted. As the Mg content in aluminum is increased, there is a decreasing dependence of the quench rate [11]. The density of dislocation loops is higher with increasing Mg content at the same aging temperature as the pure aluminum. This is due to the difference between the aluminum and magnesium atomic radii. It was concluded that Frank loops are formed from vacancy clusters [12] of approximately 10 vacancies [13]. As the density of dislocation loops increases, there is a finer and more uniform distribution of nucleation sites. These vacancy rich clusters, present immediately after quenching, form ordered GP zones [14] [15]. In addition, the formation rate of these GP zones is enhanced, because of the shorter distance solute atoms have to diffuse to a vacancy rich cluster or Frank Loop. The kinetics of vacancy-rich cluster formations are changed radically by the presence of Mg and this is interpreted as diffusion of Zn atoms along Mg-vacancy couples [16]. These couples readily move throughout the matrix at room temperature, and there is a strong binding energy between vacancies and Mg atoms. Polmear [17] investigated the aging characteristics of ternary Al-Zn-Mg alloys and found that the aging process was greatly modified by the addition of Mg to the binary Al-Zn alloys. He attributed this to the clustering of zinc and magnesium atoms as suggested by atomic size and thermodynamics. At low temperatures, nucleation proceeds predominately by the migration of Zn atoms. The consensus is that during the growth of the GP zones, Mg + Zn or Mg + Zn + vacancy complexes play a role [18] [19] [20]. Comparing the characteristic properties of different trace elements on the precipitate dimensions and density after aging showed [21] that a combination of high solute/vacancy binding energy and diffusivity was required to significantly affect the length, thickness, and density of η'. The mechanism by which these trace elements affect the microstructure is to reduce the number of quenched-in dislocation loops for heterogeneous nucleation of precipitates. The decreased number of nucleation sites leads to a prolonged matrix super-saturation, which causes an increase in the length, thickness, and density of η'. After quenching, small diffuse clusters of solute atoms are formed [22] [23] prior to the formation of GP Zones. This was first demonstrated by Guinier [24] [25] and Preston [26] [27]. GP zones and the vacancy rich clusters (VRC) with a critical solute concentration are η' nucleation sites [28]. The precipitation mechanism is VRC → η' for water quenched specimens, and GPZ → η' for direct and step quenched specimens [29]. The vacancy rich clusters are most likely Frank Loops at aging temperatures. Nucleation of η' precipitates occurs in the matrix when no excess vacancies are available; and on quenched-in vacancy clusters. GP zones were not observed to be nuclei for the equilibrium precipitates (η) [30]. The critical temperature for η precipitation corresponds to the intersection of the temperature of the C-curves for η and the η ́ phases and does not correspond to the GP zone solvus temperature [31]. Staley [32] found that tensile strength and conductivity followed a sigmoidal curve in 7050 as a function of aging time that was subsequently followed by a linear portion. The sigmoidal curve was attributed to the formation of G.P. Zones. The linear portion of the curve was attributed to the growth of the G. P. zones. Two stages for GP zone formation have been found [33]. In the first stage, the kinetics can be described by Cottrell-Bilby [34] type kinetics. Duration of this stage is strongly influenced by the aging temperature and the composition of the alloy. It was thought that the formation of GP zones starts by the clustering of Zn atoms. Shortly thereafter, Mg atoms join the clusters. The process proceeds by the formation of nuclei of zones on these clusters. It is rate controlled by the diffusion of Mg-vacancy couples. In the second stage, the GP zones are growing. The growth of these GP zones is slowed by the formation of coherency strains around the GP zones, and the decrease in supersaturation. As discussed, GP Zones are thought to be precursors for the intermediate precipitate η'. Mukhopadhyay [35] found direct evidence of η' precipitating on pre-existing GP Zones. Using small angle scattering measurements, it was found that the size of the GP Zones was limited to approximately 35-40 Å [36] prior to the formation of the intermediate precipitate η'. This led to the conclusion that there was a critical GP Zone size for the precipitation of η' [37] [36]. Precipitation processes in Al-Mg-Zn-Cu alloys can be divided into three different types [38] [39] [17] defined by the temperatures utilized. The first type is for alloys quenched and aged above the GP zone solvus. Since no GP zones are formed, no easy nuclei for precipitation occurs. A very coarse dispersion of precipitates occurs with nucleation of η' occurring primarily on dislocations. The second type is for alloys quenched and aged below the GP zone solvus. GP zones form continuously and grow to a size where η' precipitation can occur. Nucleation and precipitation of η' is not effected by the quenching process except when vacancy-rich clusters form. The third type are alloys quenched below the GP solvus, and aged above the GP solvus. The final dispersion of precipitates and the PFZ width is controlled by the nucleation treatment below the GP zone solvus. This is the common situation for commercial heat treatments, and is typical of the T73 and T76 Tempers. When Al-Zn-Mg-Cu alloys are quenched slowly, solute depletion and the loss of vacancies occurs. This depresses the critical nucleation temperature. This indicates that the solute lost during slow quenching (and the associated loss of strength from the loss of solute) is not recoverable. This loss can be partially compensated by adjusting the first stage aging practice to a temperature that allows the GP zones to grow to a stable size. This can also be accomplished by slow heating rates to the first stage age temperature [32]. In 7XXX Al-Zn-Mg-Cu alloys, several phases have been identified that occur as a function of precipitation sequence. Four precipitation sequences have been identified. These are: η ⇒ α η ⇒ η′ ⇒ ⇒ ⇒ α ⇒ ′ ⇒ α ⇒ α