Effect of precipitates evolution on mechanical properties of Al 7050 alloy during secondary aging

The microstructure and mechanical properties were investigated in Al 7050 alloy after aging treatments. Results showed that refining the η′ precipitates and increasing the proportion of the η′ precipitates significantly improved its mechanical properties. High density dislocations created by cold deformation promoted the nucleation of the main strengthening η′ phase during aging. High-density fine precipitates was created, which increases the yield strength. However, the strength of the deformed sample decreased by ∼31.6% after the secondary aging at 157 °C, with more degradation than other samples. Microstructure study presented that some precipitates were composed of two distinctive areas. An in situ transformation from a simple triclinic Mg2Zn3 to the η phase (MgZn2) was observed during the secondary aging. The loss in strength after secondary aging at 157 °C proved that the η′ precipitates transformed into the η precipitates by absorbing solute atoms at a temperature lower than equilibrium precipitation temperature of η precipitates. The residual high-density dislocations in the cold-deformed samples promoted the diffusion of solute atoms, which accelerated the η′ → η transformation during the secondary aging. Significant reduction in the volume of η′ precipitations led to the rapid deterioration of mechanical properties.


Introduction
The 7000 series aluminum alloys (Al-Zn-Mg-Cu) are widely used as structural materials in aerospace and transportation fields due to their high strengths, low densities, excellent toughness, and good corrosion resistance [1][2][3][4][5]. The series of aluminum alloys are aging strengthened aluminum alloys, and their strengths primarily depend on the types and sizes of precipitates [4,6,7]. The precipitation sequence of the precipitates during solution and aging treatment was summarized as follows [8]-supersaturated solid solution (SSS) → GP zones → metastable η′ precipitates → equilibrium η precipitates. The GP zones are formed in the early staging of artificial aging and transformed into η′ precipitates when it grows to the critical size [5]. The η′ precipitates generally have hexagonal structures, chemical compositions of Al 2 Zn 5-x Mg 2+x (x = 2-4), are semi-coherent with the matrix, and are the major strengthening phases in 7000 series aluminum alloys [8][9][10][11][12]. The η precipitates (MgZn2) are non-coherent equilibrium phases with hexagonal structures and have low strengthening effects. High density η′ phases confer high strengths to this series of aluminum alloys during peak aging treatment [5,7,13,14]. After solution treatment and cold deformation, the tensile strength of Al 7050 alloy increased by more than 20% after peak aging [15]. Significant improvement of mechanical properties is obtained at a larger strain or lower aging temperature [16], because high-density dislocations promoted the nucleation of η′ precipitates during peak aging; higher-density η′ precipitates are thus formed.
The peak aging of Al 7050 alloy could render the resultant alloys with high strength, however the stress corrosion resistance are poor. In order to improve the stress corrosion resistance, secondary aging was carried out after pre-aging treatment. Grain boundary precipitates with discontinuous distribution are obtained by coarsening the precipitates and reducing their densities during secondary aging, which improves the stress corrosion resistance due to anodic corrosion channel inhibition [17]. The GP zones formed during peak aging promote η′ and η precipitates formations during secondary aging [18]. Increased aging temperature and times resulted in coarsening of the precipitates and loss of strength [19,20]. However, other researches showed that the transformation of a large number of η′ precipitates into η precipitates also resulted in performance degradation [3,14]. Usually, the transition temperature from η′ to η precipitates occurs from 237°C-247°C [21]. In related research, long-term preservation at temperatures far below the phase equilibrium transition temperature (longterm heat preservation at 160-180°C) could result in the transformation of a higher volume fraction of η′ to η precipitates [22,23], which led to a decrease in strength. On the transition processes from η′ to η precipitates, study suggested that when the temperature rises to about 165°C, the η′ precipitates were the most and the material was the hardest. If the temperature increased continuously, η′ precipitates dissolved and η precipitates appeared [24]. Another study suggested the enrichment of solute atoms Zn to η′ precipitate during heat preservation caused the transformation of η′ into η precipitates [25]. Therefore, there is no consensus in understanding the precipitates evolution.
In 7xxx series aluminum alloys, cold deformation will produce high-density dislocations in large aluminum components. A certain density of dislocations remains in the matrix, which affects the subsequent aging of the alloys. Therefore, it is necessary to study the effect of residual dislocations introduced by cold deformation on the microstructure and mechanical property during secondary aging. In the present study on Al 7050 alloy, cold deformation improved the strength of the alloy after aging, but huge performance degradation appeared after secondary aging at 157°C. Precipitation behavior is investigated to provide an understanding of microstructure evolution with/without a pre-deformation and its effect on mechanical properties for industrial application of high-performance Al 7050 alloy.

Experimental procedure
Ingots of Al 7050 alloy were cut into 101 × 52 × 51.5 mm 3 rectangular blocks by wire-electrode cutting. Its composition is listed in table 1.
The solid solution and aging treatment of the samples were carried out in an electric furnace. The as-cast Al 7050 alloy was homogenized at 480°C for 3 h, water-cooled, and followed by hot rolling. Rolling experiments were conducted on a two-roll reversible rolling mill with roll diameters of ø350 × 300 mm. The roll speed was 300 r/min, the rolling temperature was 430°C. The samples were held for 30 min at 430°C, rolled with a thickness reduction of 2 mm per pass, and returned to the furnace for 10 min after each pass. The final deformation was 69%. Following hot rolling, solution treatment at 470°C for 3 h and 480°C for 13 h was conducted immediately, followed by water cooling and aging. The samples were divided into three groups. The first group was cold rolled with 26% deformation before aging, then aged at 130°C for 24 h and marked as 26#. During cold rolling, each pass was pressed by 0.5 mm. The roll speed was 300 r min −1 . The second group of samples was not rolled but instead subjected directly to first aging (FA) at 130°C for 24 h and marked as N#. The third group of samples was not rolled and directly subjected to first aging at 140°C for 24 h and marked as C#. After the first aging, all samples were subjected to secondary aging (SA) at 157 for 20h. The technical processes were listed in table 2.
The tensile test was performed on the MTS-810 testing machine with a tensile speed of 1 mm min −1 . The tensile specimens were axial along the rolling direction (RD) with at least three samples from each group tested. Rolling schematic and tensile sample size were shown in figure 1. The fracture morphology of the samples was observed using a FEINova450 scanning electron microscope (SEM). These samples were polished using an argon ion beam, then a field emission scanning electron microscope with EBSD (SUPRA 40) was used to study the grain structure. The scanning step was 0.6 μm. The data from EBSD were post-processed using TSL OIM Analysis 7 software (EDAX, USA) to analyze the kernel average misorientation (KAM) and to count the sample grain sizes. After the FA and SA treatments, DSC testing was performed, using a Mettler DSC1 differential  scanning calorimeter (heating rate of 5°C min −1 ), on N#, 26#, and C# samples. TEM observations were carried out using a FEI Tecnai F20 field-emission-gun scanning transmission electron microscope (TEM), the TEM samples prepared by cutting discs from the selected samples along their loading direction and thinning the discs mechanically to 0.07 mm before twin-jet electropolishing in a mixture of 33% nitric acid and 67% methanol at -20°C and a working voltage of 10 V. Statistics of the precipitates and grain sizes were conducted using imaging-pro plus 6.0 software. To ensure statistical validity, the same magnification images were used in the statistics of precipitate sizes. The average grain size was determined by averaging the values of the long and short axes. More than 200 precipitates taken into account for each sample; the average size of the precipitates was the average size of the long axis direction.

Tensile property
Based on the tensile test, the strength was the highest in the 26# sample aged after cold deformation, but the plasticity dropped significantly. The strength of the C# sample was the lowest. The N# sample showed higher strength and plasticity (see the solid line curve in figure 2). After the SA, the strength of the samples decreased significantly (see

Microstructure
The KAMs of the samples after the FA and SA are compared in figure 3. After the FA, the KAMs of the 26# sample were significantly higher than the N# and C# samples (see figures 3(a)-(c)) due to residual cold deformation. After the SA, the KAMs of the 26# sample remained higher. The statistical results of average grain  size and dislocation density were listed in table 4, and showed that the dislocation densities in all samples decreased by one order of magnitude after the SA. The grain size of the samples did not change significantly.
The bright field morphology of many precipitates of the samples after the FA were shown in figure 4. Based on the diffraction spots, the precipitates were mainly η and η′ phases. The sizes of the precipitates in 26 # sample are smaller and their densities are higher than that in N# sample. This is because the high-density dislocations introduced by cold deformation before aging promoted the heterogeneous nucleation of the precipitates and improved the nucleation rate [26]. The precipitates grew more quickly in the C# sample because of higher the FA temperature, which resulted in larger sizes ( figure 4(c)).
The precipitates with platelet-like, elliptical, and rod-like shapes were primarily observed in the matrix. The smaller particles or platelet-like shape of precipitates with a long axis direction between 5-15 nm were the η′ phases ( figures 5(a) and (b)). The η′ phases are semi-coherent relationship with the matrix lattice, they strongly  inhibits dislocation movements [4,27] ( figure 5(c)). The elliptical or rod-like precipitates with sizes greater than 15 nm or long axis directions of 15-55 nm are mainly η phases (figures 5(d) and (e)). The η phases are noncoherent with the matrix (figure 5(f)), which was consistent with the description of the morphologies and sizes of η′ and η precipitates in related studies [10,28,29]. Hence, morphologies and size were used to perform quantitative statistics. The precipitates significantly coarsened in the samples (figures 6(a)-(c)) after the SA. The coarsening rate in the 26# sample was faster than that in the N# and C# samples. The precipitates were distinguished by HRTEM and diffraction spots (figures 6(d)-(f)). The precipitates below 15 nm were still η′ phases. Other larger precipitates were mainly η phases. The sizes of the precipitates were larger due to the higher FA temperature in C# sample after the FA, which resulted in lager precipitates after SA.
On the basis of the corresponding TEM and HRTEM results, the precipitates types with related morphologies and sizes were distinguished. Precipitates and their corresponding amounts and sizes were counted in samples (listed in table 5). After the FA, the precipitates in the N# sample contained up to 83.2% of η′ precipitates and only 16.8% of η precipitates. However, the percentage of η′ precipitates decreased to 65% after the SA. The decreased tendency in 26# sample was more obvious than that in N# sample. The amount of η′  phases reached 85.7% and η precipitates were only 14.3% after the FA. But the proportion of η′ precipitates decreased to 52.6% after the SA.
Coarsening of η′ and η precipitates occurred in all samples after the SA. The precipitates showed the greatest size changes in the 26# sample; the η′ precipitates increased from 5.34 nm to 10.35 nm, the η′ phases increased from 17.65 nm to 26.51 nm. The size and the ratio of η′ and η precipitates changed during the SA.

Discussion
The dislocation density and grain size changed little during the SA, so the strength change caused by the dislocation and the grain size change was insignificant. Even in extreme conditions, assuming all solute atoms are in a solid solution and the effects from the various solutes are additive, solid solution strengthening in this alloy accounts for at most an increase of ∼73 MPa [30]. Therefore, evolution of precipitates is main factor for the change of mechanical properties during the SA.
The DSC curves of the three groups were displayed in figure 7. After the FA, the obvious endothermic peaks from 137.8°C-204.1°C, 148.35°C-203.2°C, and 152.94°C-209.9°C (peak 1) were observed in 26#, N#, and C# samples, respectively. These endothermic phenomena were caused by dissolution of GP zones and η′ precipitates during heating [13,24]. Peak 2 was caused by the formation of η precipitates [13]. The DSC curves  reflect the evolution law of the precipitates in the Al 7050 alloy. The η′ precipitates dissolved during heating, and more stable η phases will precipitate at higher temperatures [13].
After the SA at 157°C, the thermal effect in the initial DSC curve stage was insignificant due to the substantial reduction of GP zones. The endothermic effect was due to dissolution of the remaining η′ precipitates. No obvious exothermic peak appeared above 200°C. The absence of the peak 2 showed that the transformation from η′ to η precipitates occurred during the SA below the equilibrium precipitation temperature of the η precipitates. So the thermal effect of the η precipitation peak in the DSC curve was not obvious.
The fine precipitates ( figure 8(a)) obtained after the FA were grown into large-size precipitates ( figure 8(b)) after the SA. The enrichment of Mg, Zn, and Cu in the larger precipitates increased. Though the SA temperature was lower than the equilibrium precipitation temperature of the η phase, the evolution of the precipitates was accelerated due to rapid diffusion and segregation of solute atoms along the crystal defects. They readily formed larger-size precipitates on the grain boundary via rapid enrichment of solute atoms. Crystal defects promoted diffusion and aggregation of elements and accelerated the formation of larger stable phases.
Although the SA temperature did not reach the temperature at which η precipitates form in equilibrium, their structures changed during long-term holding. Figure 9 shows the precipitate in 26# sample after the SA.  The different crystal structures were observed in regions 1 and 2 in this precipitate. Region 1 is the η phase with a hexagonal structure (As shown in figure 9(c)), whereas region 2 is the Mg 2 Zn 3 phase with a simple triclinic structure (As shown in figure 9(d)). The structural evolution differed in various parts of a single precipitate due to various elemental ratios. Higher Zn/Mg ratios facilitated formation of the stable η phase [25]; the metastable precipitates gradually transformed into a stable structure due to diffusion and aggregation of alloying elements [31,32].
The Zn and Mg solute atoms gradually diffused and segregated into the η′ phase during prolonged SA. As the composition eventually met the conditions for transformation, the in situ η′ → η transformation starts to occur. Therefore, the enrichment rate of solute elements to the η′ phase affects the transformation.
The changes of the size and proportion of the precipitates after aging were further counted (As shown in table 6). The significant differences in mechanical property were related to the number of precipitates in samples during the SA. The decrease of the percentage of η′ precipitates corresponds to the decrease of the strength. The difference in mechanical properties is mainly caused by differences in the ratio of η′ and η phases during the SA.
Among the three samples, the largest η′ phase decrease (33.1%) and fastest coarsening rate of η′ and η phases occurred in the 26# sample during the SA. Cold deformation created high-density dislocations and grain boundaries acts as quick diffusion routes for solute atoms to promote solute element enrichment and precipitates growth [33]. Figure 10(a) showed the precipitate sizes near the dislocation line increased (Marked by the white arrow) after the FA, The size is larger than that of the precipitates in the region far from the dislocations. Figures 10(b) and (c) showed elemental distribution maps correspond to figure 10(a). The enrichment of Zn and Mg elements was observed in the precipitates near the dislocation. The sizes of precipitate  on the grain boundary was also much larger than that in the matrix (Marked by the red arrow), and the Mg content was approximately 33% in the precipitate, which is consistent with other studies [34], Grain boundary defects accelerated formation of equilibrium phases. The dislocation and grain boundary both played a similar role to promote elemental diffusion, which significantly accelerated diffusion of Mg, Zn, and Cu and promoted the transition of η′ phase to the equilibrium η phase. As a result, the proportion of the η′ phase in the 26# sample decreased significantly during the SA. After the first aging, the precipitate densities in the 26# sample exceeded the N# sample that in the N# sample due to the large number of nucleation sites provided by dislocations. The precipitates are mainly η′ phases. The distortion field between the precipitates and the Al matrix severely hindered dislocation movement and increased the yield strength. So the yield strength of the 26# sample was higher. However, the residual highdensity dislocations in 26# sample led to a rapid transformation of the η′ precipitates to η precipitates during secondary aging. The proportion of η′ precipitates in the 26# sample decreased more than that in the N# sample, which resulted in a greater loss in strength (190MPa) in the sample than that (66MPa) of N# sample. The proportion of η′ precipitates in C# sample was lower than that in N# sample due to the higher first aging temperature, so the yield strength of C# sample was also lower than that of N# sample after the first aging. After secondary aging, the transformation amount of η′ to η precipitates in C# sample was more than that of N# sample. The loss in the yield strength of C# sample was also greater than that of N# sample.
After the SA, the η′ precipitates were transformed into η precipitates with poor strengthening effect. The dislocations bypassed η precipitates and moved during the deformations. The dislocation loops were left around the η precipitates that cause stress concentrations and microcracks, which led to a lot of dimples on the fracture [27]. The tensile fracture morphologies of 26# specimen following the FA and the SA were shown in figure 11. There were coarse second phases on the fracture of 26# after the FA (As marked by white arrow). The second phases are Al 2 CuMg by EDS analysis. The cleavage planes were also observed on the fracture surface (as shown in figure 11(b)). After the SA, the fracture types of 26# sample were mainly microvoid accumulation fractures, and Long-term holding below the equilibrium transition temperature of the η phase resulted in the continuous enrichment of alloying elements Cu, Zn, and Mg to the η′ phases, which led to growth of the main strengthening precipitate η′ phases and promoted their transformation to the equilibrium η phases and coarsening of the η phases. Cold deformation effectively increased the density of crystal defects in the matrix, accelerated the diffusion and enrichment of alloying elements, and increased the transformation rate of η′ to η and coarsening rate of the precipitated phase, all of which degraded the strength of the alloy after the SA.

Conclusions
The relationship between microstructure evolution and mechanical properties of Al 7050 alloy during aging was studied, and the following conclusions were drawn: (1)High-density dislocations can promote the nucleation of the main strengthening precipitation η' phase under low-temperature aging in 26# sample. High-density fine precipitates was created, which increases the yield strength. After solution treatment and aging at 130°C, the yield strength of 26# sample reached 601 MPa, much greater than 573 MPa for N# sample without cold deformation. The strength of C# sample aged at 140°C was only 534 MPa and caused by a decline in the amount of the η′ phase and a coarsening of the precipitates.
(2)Although the subsequent secondary aging temperature was lower than the equilibrium precipitation temperature of the η phase, the enrichment of alloying elements to the η′ phase during holding promoted its transformation to the η phase, which resulted in lower tensile properties after the secondary aging. Crystal defects accelerated elemental diffusion, coarsening of the η′ phase, and the transition to the η phase. The residual high-density dislocations in 26# sample led to a rapid transformation of the η′ to η phase during secondary aging, which decreased the yield strength to 411 MPa, below both N# (507 MPa) and C# (415 MPa) after the secondary aging.