Enhanced Strengthening and Toughening of T6-Treated 7046 Aluminum Alloy through Severe Plastic Deformation

Abstract

The 7046 aluminum alloy possesses a favorable fatigue property, corrosion resistance and weldability, but its moderate strength and plasticity limit its wider application and development. In the present study, severe plastic deformation (SPD) was applied prior to T6 treatment to significantly enhance the strength and toughness of the 7046 aluminum alloy. The results show that the alloy processed by four passes of equal channel angular pressing (ECAP) at 300 °C prior to T6 treatment exhibits an excellent mechanical performance, achieving an ultimate tensile strength (UTS) and elongation (EL) of 485 MPa and 19%, respectively, which are 18.6% and 375% higher than that of the T6 alloy. The mechanical properties of the alloy are further improved by an additional room temperature (RT) rolling process, resulting in a UTS of 508 MPa and EL of 23.4%, respectively. The increased presence of η′ and Al₆Mn phases in the 300°C4P-R80%-T6 and 300°C4P-T6 alloys contributes to a strengthening and toughening enhancement in the SPD-processed T6 alloy. The findings from this work may shed new insights into enhancing the 7046 aluminum alloy.

1. Introduction

7xxx series (Al-Zn-Mg-Cu) aluminum alloys are typical precipitation-hardened alloys characterized by high specific strength, good fracture toughness and excellent machinability [1,2,3]. These alloys are widely utilized in aerospace, advanced equipment manufacturing, rail transportation, 5G applications and new infrastructure projects [4,5]. However, the challenges posed by harsh and complex service environments, along with the high-quality connection requirements between different materials in practical applications, demand superior fatigue properties, corrosion resistance and weldability from 7xxx series aluminum alloys [6,7]. Consequently, new Cu-free 7xxx series aluminum alloys containing small amounts of Zr and Ti (such as 7046) have emerged as promising alternatives [8,9]. Research on the 7046 aluminum alloy mainly focuses on its mechanical properties [10,11,12], fatigue resistance [9,13], corrosion resistance [11,14,15,16] and weldability [17,18]. However, relatively little attention has been given to the mechanisms underlying its strengthening and toughening. Thus, it is crucial to further investigate effective processing techniques for producing more strengthened and toughened 7046 aluminum alloys and to deepen their associated mechanisms.

Solid solution with peak aging (T6 treatment) is a common method for enhancing the strength of 7xxx aluminum alloys [19,20]. After T6 treatment, the GP zones precipitated in large quantities in grains can be transformed into η′ (MgZn₂) phases. These coherent GP zones and semi-coherent η′ phases induce significant lattice distortion, thereby strengthening the alloy [21]. For instance, an approximate 115 MPa increment in ultimate tensile strength (UTS) can be achieved in a 7A99 aluminum alloy processed by T6 treatment [22]. Both ultrafine-grained (UFG) and coarse-grained (CG) 7075 alloys exhibit notable UTS increases after T6 treatment, with increments of about 143 MPa and 223 MPa, respectively [23]. Similar studies have also indicated that while T6 treatment significantly enhances the strength of these alloys, the improvement is limited and often comes with a strength–ductility trade off [24,25]. Therefore, it is imperative to develop new processing technologies to echo the call for developing super higher-performance (high-strength and high-ductility) alloys.

To address this issue, some researchers have proposed applying room temperature (RT) severe plastic deformation (SPD) during T6 treatment (before aging and after solid solution treatment) to improve the comprehensive mechanical properties of the 7xxx aluminum alloys [26,27,28]. Studies indicate that one pass of RT equal channel angular pressing (ECAP) after solid solution greatly shortens the peak-aging time of experimental 7055 aluminum alloy extrusion bars. This treatment improves the ultimate UTS and YS (yield strength) by 39 MPa and 33 MPa, respectively, while the elongation (EL) remains approximately 9% compared to that of T6 treatment [29]. The UTS of Al-6.6Zn-1.96Mg-2.35Cu alloy subjected to two passes of cross accumulative extrusion bonding (CAEB) after solid solution treatment (470 °C for 1 h) and before aging (120 °C for 24 h) is 562 MPa, which is 40% more than that of the T6-treated alloy. Meanwhile, the alloy maintains a good EL of 21.5% [30]. By embedding a RT SPD process in solid solution and aging (T6 treatment), the grains of the alloys become finer more easily because of the shear strain and the resultant recrystallization, which contributes to fine-grain strengthening. Additionally, the increased number of grain and subgrain boundaries formed during RT SPD provides more nucleation sites for the η′ phases, which enhances precipitation strengthening [31,32,33].

The above-mentioned method (embedding RT SPD in T6 treatment) is effective for the more easily deformable 7xxx series aluminum alloys. However, for 7xxx series aluminum alloys with low plasticity, such as the 7046 alloy, a high-temperature (HT) SPD treatment is required [34,35]. It should be pointed out that dynamic precipitation would be significantly accelerated by the HT SPD [36], which means the GP zones and η′ strengthening phases would easily be transformed into η phases [36,37]. The η phases have a weaker strengthening effect, so this simple HT SPD process often does not play a role in the strengthening and toughening of the alloys. Consequently, some studies have explored performing a T6 treatment after HT SPD. For example, 7055 aluminum alloys processed by a two-stage solution treatment (470 °C for 16 h and 475 °C for 8 h) and double-step hot rolling (250 °C for 20% reduction and 430 °C for 75% reduction), combined with a T6 treatment, shows excellent comprehensive mechanical properties. It exhibits a UTS, YS and EL of 610 MPa, 542 MPa and 21.6%, respectively [38]. This successful approach is inspirational for the simultaneous strengthening and toughening of hard-to-deform 7xxx series aluminum alloys. But so far, few studies have paid attention to why the SPD process can enhance the strengthening and toughening effect of T6 treatment, let alone in 7046 alloys. Considering that ECAP is the most common and probably the most cost-effective SPD process for the production of hard-to-deform 7xxx series aluminum alloys, we have chosen ECAP combined with RT rolling as the SPD process in the present study. Simultaneously, high strength and toughness are realized in 7046 aluminum alloys by SPD processing prior to T6 heat treatment, and the corresponding strengthening and toughening mechanism is systematically discussed.

2. Experimental Material and Procedures

The studied as-cast 7046 ingot (185 mm × 60 mm × 25 mm) was obtained by remelting commercial 7046 aluminum alloys in an electrical resistance furnace (SG2-5-10, Shanghai Yunyue Instrument Equipment Limited Company, Shanghai, China). In order to ensure that the composition of the alloy does not change significantly after remelting, a refinement agent was not used. The actual chemical composition was measured using a direct reading spectrometer (ARL-3460, Thermo Fisher Scientific, Waltham, MA, USA), as shown in

Table 1. Actual chemical composition of the 7046 alloy (wt.%).

Zn	Mg	Mn	Cu	Cr	Ti	Zr	Fe	Si	Al
7.26	1.7	0.26	0.22	0.01	0.02	0.13	0.14	0.09	Bal.

.

The as-cast alloy was cut into samples with a size of 19.5 mm × 19.5 mm × 45 mm by wire electrical discharge machining (DK7735, Aier Cnc Machine tools, Taizhou, China). The samples were solid-solutionized at 470 °C for 1 h, and then quenched in water. Since 7046 alloys can hardly be ECAPed at RT, the ECAP die was heated to 300 °C in a resistance furnace and maintained at this temperature for 1 h before inserting a sample into the entrance channel [39]. To prevent crack formation, all the samples were ECAPed in 4 passes quickly and continuously. ECAP was conducted via route Bc in a 90° die with an average press speed of 5 mm/min [40,41]. Sheets with a thickness of 6 mm were cut from 4P-ECAP samples along the extrusion direction. They were rolled on a rolling mill (BWD-87-Y4-ZP, Shanghai Zhouyi Machinery Equipment Limited Company, Shanghai, China) with a 5% average reduction per pass at RT until a total of 80% reduction. Finally, all these samples were heat treated by T6 treatment (solid solution at 470 °C for 1 h and further aging at 120 °C for 24 h). The samples were named as 300°C4P-T6 and 300°C4P-R80%-T6, respectively. Additionally, as-cast, T6 and R80%-T6 samples were prepared for performance comparison.

Tensile tests were conducted to evaluate the mechanical properties of all samples. The type and dimensions of the samples are shown in Figure 1. Tensile tests were carried out at RT using a universal testing machine (Suns-UTM4294X, Jinan Hensgrand Instrument Limited Company, Jinan, China) at a loading rate of 0.5 mm/min. Three parallel tensile specimens were tested for each state and the average values were used. The microstructures of the alloys were characterized using optical microscopy (OM, BX6OM, Olympus Corporation, Tokyo, Japan), transmission electron microscopy (TEM, FEI Tecnai G2 F20 S-TWIN, Thermo Fisher Scientific, Hillsboro, MA, USA) and electron backscatter diffraction (EBSD, Hitachi S-3400N, Hitachi, Tokyo, Japan). For OM observations, samples were ground with SiC abrasive papers, polished with 3.5 µm diamond spray and finally etched with Keller’s reagent for 75 s. EBSD samples were electropolished in an electrolytic polishing solution (10% HClO₄ + 90% C₂H₅OH) at approximately −20 °C, with a constant voltage of 32 V, for 60 s. For TEM observations, samples were first ground to about 60 µm by SiC abrasive papers and then thinned in an ion-thinning apparatus (GATAN-691, Gatan Incorporated, Santa Barbara, CA, USA).

3. Results

3.1. Microstructure Observation

Figure 2 shows the optical microstructures of the as-cast, T6, 300°C4P-T6 and 300°C4P-R80%-T6 7046 alloys. As can be seen, the as-cast alloy primarily consists of an α-Al matrix and a small amount of second phases distributed along the grain boundaries (Figure 2a). These second phases mainly comprise primary MgZn₂ phases and Al₇Cu₂Fe phases [42]. After T6 treatment, most of the MgZn₂ phases dissolve into the Al matrix, while only a small number of larger MgZn₂ phases still remain, along with insoluble Al₇Cu₂Fe phases, as illustrated in Figure 2b. Following ECAP, ECAP and RT rolling, before T6 treatment, the grain size of the alloy is obviously refined, and the originally coarse second phases become fine and dispersed, as shown in Figure 2c,d.

Figure 3 exhibits the EBSD maps of the as-cast, T6, 300°C4P-T6 and 300°C4P-R80%-T6 alloys, respectively. The grains of the as-cast alloy are nearly equiaxed, but the grain size is heterogeneous, ranging from 40 to 1100 µm, as shown in Figure 3a. After T6 treatment, the grains grow slightly (ranging from 50 to 1500 µm), but remain heterogeneous, as depicted in Figure 3b. Most of the grains exhibit elongation when subjected to four passes of ECAP at 300 °C combined with T6 treatment, with some of the grains becoming fine, recrystallized grains (as shown in Figure 3c). The average grain size (AGS) of the 300°C4P-T6 alloy decreases to about 40 µm (Figure 3e). With the additional RT rolling process after ECAP and before T6 treatment, the grain size of the alloy is further refined, the elongated state of the grains is weakened and almost all of the grains become fine recrystallized grains (Figure 3d). The AGS of 300°C4P-R80%-T6 alloy is further reduced to 19 µm (Figure 3f).

Figure 4 visualizes the polar figures of T6, 300°C4P-T6 and 300°C4P-R80%-T6 alloy, respectively. As shown in Figure 4a, the peak texture intensity of the T6 alloy is 5.41. After performing ECAP extrusion prior to T6 treatment, the peak texture intensity of the alloy decreases slightly to 5.31 (Figure 4b). Following ECAP and RT rolling prior to T6 treatment, the peak texture intensity of the alloy is further weakened to 3.02 (Figure 4c) with a more uniform weave distribution.

Since all the samples after different processes were subjected to T6 treatment, the dislocation intensity in the samples is rather low, so the precipitates are paid more attention. Figure 5 illustrates three typical precipitates distributed along the grain boundaries and within the grains of T6 alloy observed by TEM. Firstly, as shown in Figure 5a,b, the η phases, with a large dimensional difference, are observed along the grain boundaries in the T6 alloy, with length dimensions ranging from approximately 30 nm to 450 nm and length–diameter ratios ranging from 1.6 to 5.1. A precipitate-free zone (PFZ) with a width between 120 nm and 180 nm is observed at the grain boundaries of the alloys. Secondly, a lot of biggish short rod-like Al₆Mn phases (its average length is about 102 nm, and its length–diameter ratio is about 2.6) are distributed inside the grains (Figure 5c). Finally, as shown in Figure 5c,d, a large number of small oval η′ phases (sizes around 5 nm) are diffusely distributed in the grains.

Figure 6 presents the three typical precipitates in the 300°C4P-T6 alloy. Compared to Figure 6a and Figure 5a,b, the size distribution of the η phases at the grain boundaries of the 300°C4P-T6 alloy is more uniform and finer than that of T6 alloy, with an average length of about 38 nm and a length–diameter ratio of about 2.7. The width of the PFZ at the grain boundaries is significantly reduced to 44 nm (Figure 6a). As shown in Figure 6b,c, a large number of biggish Al₆Mn phases (its average length is about 50 nm, and its length–diameter ratio is about 1.5) and small oval η′ phases (its size is about 5 nm) are diffusely distributed in the grains. Compared to T6 alloy, the size and shape of η′ phases in 300°C4P-T6 alloy after ECAP extrusion did not change significantly, but the Al₆Mn phases became finer in size and transformed from a short rod-like shape to an elliptical shape (its length–diameter ratio decreases). A small number of Al₆Mn phases also appear at the grain boundaries of the 300°C4P-T6 alloy (Figure 6a).

Figure 7 shows the three typical precipitates in the 300°C4P-R80%-T6 alloy. Compared to Figure 7a and Figure 6a, the continuous distributed η phases at the grain boundaries of the 300°C4P-R80%-T6 alloy become more elongated compared to those of the 300°C4P-T6 alloy, with the average length increasing to 55 nm and the average width decreasing to 18 nm (a length–diameter ratio of approximately 3). The width of the PFZ at the grain boundaries of the 300°C4P-R80%-T6 alloy (the average width is about 40 nm) does not change too much compared to that of the 300°C4P-T6 alloy. The biggish Al₆Mn phases and small η′ phases are abundantly distributed within the grains, as shown in Figure 7b, c. The average length of the Al₆Mn phase is about 48 nm with a length–diameter ratio of 1.4, which is slightly smaller than that of the Al₆Mn phases in 300°C4P-T6 alloy. The average size of the η′ phases remains around 5 nm.

To better exhibit the distribution of the typical Al₆Mn precipitates of a 7046 alloy in a broader range, Figure 8 is presented. As shown in Figure 8a, the T6 alloy contains numerous short rod-like Al₆Mn phases with a large size (its average length is about 102 nm, and its length–diameter ratio is about 2.6). After ECAP extrusion, the Al₆Mn phases in the 300°C4P-T6 alloy become finer and the shape changes from short and rod-like to elliptical (its average length is about 50 nm, and its length–diameter ratio decreases to 1.5). The quantity of Al₆Mn phases in the 300°C4P-T6 alloy does not change significantly, but the distribution of them become more uniform (Figure 8b). In the 300°C4P-R80%-T6 alloy, the content of Al₆Mn phases increases notably, while their size and shape remain similar to those in the 300°C4P-T6 alloy (Figure 8c).

3.2. Mechanical Properties

The typical engineering and true stress–strain curves of the as-cast, T6, R80%-T6, 300°C4P-T6 and 300°C4P-R80%-T6 alloys are shown in Figure 9a,b, respectively. The corresponding mechanical properties obtained from Figure 9a, b are summarized in

Table 2. Summary of the mechanical properties of the as-cast, T6, R80%-T6, 300°C4P-T6 and 300°C4P-R80%-T6 alloys.

Alloy	UTS (MPa)	YS (MPa)	EL (%)	Peak True Stress (MPa)	True Fracture Strain (%)
as-cast	272 ± 2.9	203 ± 3.1	5.0 ± 0.5	283 ± 4.3	4.9 ± 0.4
T6	409 ± 2.7	357 ± 1.8	4.0 ± 0.3	421 ± 2.8	3.9 ± 0.2
R80%-T6	462 ± 2.6	365 ± 3.8	15.6 ± 0.6	526 ± 3.0	14.5 ± 0.5
300°C4P-T6	485 ± 1.8	366 ± 2.7	19.0 ± 0.3	557 ± 3.0	17.4 ± 0.2
300°C4P-R80%-T6	508 ± 1.4	393 ± 2.3	23.4 ± 2.6	607 ± 1.7	21.0 ± 2.1

. The as-cast alloy exhibits a UTS of 273 MPa, a YS of 203 MPa and an EL of 5.0%. After T6 treatment, the alloy shows an increase in the UTS and YS, accompanied by a slight decrease in EL. It presents a UTS of 409 MPa, a YS of 357 MPa and an EL of 4.0%. By pre-rolling treatment, the R80%-T6 alloy shows a slight increase in YS and a significant increase in UTS and EL, with a UTS, YS and EL of 462 MPa, 365 MPa and 15.6%, respectively. The strengthening and toughening effect of the pre-ECAP treatment is more obvious. The UTS, YS and EL of the 300°C4P-T6 alloy are 485 MPa, 366 MPa and 19.0%, respectively. The mechanical properties of the alloy are further enhanced by the additional RT rolling process. The 300°C4P-R80%-T6 alloy shows a UTS, YS and EL of 508 MPa, 393 MPa and 23.4%, representing an increase of 24.2%, 10.1% and 485%, respectively, compared to that of the T6 alloy.

3.3. Tensile Fracture Surfaces

Figure 10 shows the tensile fracture surfaces of the 7046 alloy after different processes. The fracture surfaces of the as-cast alloys exhibit some dimples, but these dimples are large and there are some cleavage steps locally, as can be seen in Figure 10a. A tract of cleavage steps is observed in the T6 alloys, and the dimples are barely observed (Figure 10b), making it an obvious brittle fracture. Performing ECAP extrusion prior to T6 treatment, a large number of small-sized dimples are nested between large, icing sugar-like patterns (Figure 10c), which correspond to an increase in the ductility of the 300°C4P-T6 alloy. Performing ECAP and RT rolling prior to T6 treatment, the dimples become deeper, the size of them is smaller and the amount of them increases, transforming it into a ductile fracture (Figure 10d). These fracture characteristics are consistent with the high EL% (23.4%) of the 300°C4P-R80%-T6 alloy.

4. Discussion

As demonstrated in the above section, the 300°C4P-R80%-T6 alloy developed in this work exhibits excellent mechanical properties, indicating that performing ECAP and RT rolling prior to T6 treatment can greatly improve the mechanical properties of the alloy. In this section, the strengthening and toughening mechanisms of the 300°C4P-T6 and 300°C4P-R80%-T6 alloys are discussed.

4.1. Strengthening Mechanism

T6 treatment slightly decreases ductility but greatly enhances strength. To be specific, the UTS of the T6 alloy increases to 409 MPa, which is about 50% higher than that of the as-cast alloy (272 MPa). It can be reasonably concluded that the UTS increment (409 − 272 = 137 MPa) of the T6 alloy is primarily due to η′ precipitates (i.e., the second phase strengthening). When it comes to the 300°C4P-T6 and T6 alloys, the UTS increases by about 76 MPa. The main strategies for improving mechanical properties include (i) solid-solution strengthening, (ⅱ) fine-grain strengthening, (ⅲ) second phase strengthening and (ⅳ) dislocation strengthening [43]. In this study, all the samples after different processes are subjected to T6 treatment, which eliminates internal stresses, and a large number of strengthening phases precipitate. Dislocation entanglements can hardly be observed, and the textures are weak (Figure 4). Consequently, the contribution of dislocation strengthening and solid-solution strengthening can be neglected. Fine-grain strengthening can be calculated by the Hall–Petch relation [23,44], as shown in Equation (1).

$${\sigma }_{HP}={\sigma }_0+k{d}^{-1/2}$$

(1)

where ${\sigma }_0$ denotes the lattice friction stress, k is the Hall–Petch slope and d is the average grain size. As can be seen from Equation (1), fine-grain strengthening depends on the grain’s refinement. The grain size of the T6 alloy ranges from 50 to 1500 µm. Performing 300°C4P-ECAP before T6 treatment, the AGS of the 300°C4P-T6 alloy rapidly drops to 40 µm (Figure 3e), which implies that fine-grain strengthening does contribute to a partial increase in strength.

Second phase strengthening is generally considered to be controlled by the Orowan mechanism [45]. It can be calculated by

$${\sigma }_s={\displaystyle \frac{0.13Gb}{\lambda }}ln{\displaystyle \frac{d_s}{2b}}$$

(2)

where G denotes the shear modulus of Al, b is the Burgers vector for the basal slip of Al, λ is the average interparticle spacing and $d_s$ is the average particle size. As indicated by Equation (2), the strengthening effect of the second phase depends on the size and distribution of the precipitates. There are three typical precipitates (η phase, Al₆Mn phase and η′ phase) along the grain boundaries and inside the grains of the T6 and 300°C4P-T6 alloys (Figure 5 and Figure 6). First, the η phases, with a large dimensional difference, are observed at the grain boundaries of the T6 alloy, with length dimensions ranging from approximately 30 nm to 450 nm and length–diameter ratios ranging from 1.6 to 5.1, as shown in Figure 5a,b. In contrast, the η phases at the grain boundaries of the 300°C4P-T6 alloy are more uniform and finer than those of T6 alloy, with an average length size of about 38 nm and a length–diameter ratio of about 2.7 (Figure 6a). Second, the Al₆Mn phases precipitated inside the grains of the T6 alloy have an average length about 102 nm, and a length–diameter ratio about 2.6 (Figure 5c). When it comes to the 300°C4P-T6 alloy, the average length of the Al₆Mn phases decreases to 50 nm, and their morphology changes from short and rod-like to elliptical (Figure 6b). The average interparticle spacing of the η phases and Al₆Mn phases in both the T6 and 300°C4P-T6 alloys is much larger than their average particle sizes. According to Equation (2), the strengthening contribution of the η phases and Al₆Mn phases is also limited. Third, the η′ phases are the main strengthening phases in the present studied alloys [5,46], which have a similar size in both the T6 and 300°C4P-T6 alloys (Figure 5d and Figure 6c). Therefore, the strengthening contribution from the η′ phases depends on their quantity. While it is difficult to directly compare the quantity of η′ phases, it can be indirectly assessed by comparing the width of their PFZs. As the solute’s depletion is one of the main reasons for the occurrence of a PFZ [47,48], η phases nucleate first at grain boundaries by drawing the solutes from the nearby matrix, leading to the occurrence of PFZs close to the grain boundaries [47]. So, the more or coarser the η phases appear along the grain boundaries, the wider the PFZ becomes. Meanwhile, it is known that the η phases and η′ phases are composed of the same elements (both are MgZn₂) and that their chemical composition is identical in the present studied alloy. Consequently, it can be roughly estimated that the wider the PFZ there is, the more η phases there are or the coarser they are, and the fewer η′ phases there are. The average width of the PFZ in the T6 alloy is about 150 nm (Figure 5a,b), while that in the 300°C4P-T6 alloy significantly decreases to 44 nm (Figure 6a). Therefore, it can be deduced that the number of η′ phases is much higher in the 300°C4P-T6 alloy than in the T6 alloy.

As the composition of precipitates is different within the matrix, the solute atoms must both find vacancies and have enough energy to move into the vacant site to complete the diffuse distribution of the η′ phases in grains [49]. Meanwhile, the precipitate nucleation rate reaches a maximum at HT, and the nucleation and growth of precipitates are always competitive with each other, which causes some η′ phases to transform into η phases [32] and makes the primary strengthening η′ phases hopeless to reach the maximum density. The HT ECAP prior to T6 treatment actually increases the concentration of vacancies and provides additional strain energy [50]. During ECAP extrusion, the growth of small precipitates will be inhibited, and basically no small platelets will have the ability to grow into larger precipitates. The growth of large precipitates occurs mainly through coalescence rather than through the dissolution of small precipitates. The higher concentration of vacancies and the more strain energy, the further the precipitation of small η′ phases, and this will prevent them from transforming into large η phases. That is why there are many more η′ phases in the 300°C4P-T6 alloy than in the T6 alloy.

Above all of these, the higher strength of the 300°C4P-T6 alloy can be attributed to the combination of fine-grain strengthening and second phase strengthening by more η′ phases.

The strengthening of the alloy is further improved by the additional RT rolling process after ECAP and before T6 treatment. The grain size of the 300°C4P-R80%-T6 alloy is further refined with an AGS of 19 µm (Figure 3f). The average width of the PFZ in the 300°C4P-R80%-T6 alloy (40 nm) is similar to that in the 300°C4P-T6 alloy (Figure 6a and Figure 7a), which means that the η′ phases in the 300°C4P-R80%-T6 alloy are equivalent to those in the 300°C4P-T6 alloy. The second phase strengthening of them is almost the same. Therefore, the further increase in the strength of the 300°C4P-R80%-T6 alloy is mainly due to its fine-grain strengthening.

4.2. Toughening Mechanism

By pre-ECAP treatment, the 300°C4P-T6 alloy exhibits an EL of 23.7%, exhibiting an improvement of 485% compared to that of T6 alloy. The plasticity of the alloy is affected by a series of microstructure characteristics, including the crystal structure (grain size and orientation), the density of dislocations, coarse intermetallic phases, PFZ and precipitates in grains and at the grain boundaries [14,51,52].

Fine-grain strengthening is considered to improve the strength and plasticity of the alloy simultaneously [53]. When compared to that of T6 alloy (ranging from 50 to 1500 µm), the AGS of the 300°C4P-T6 alloy sharply decreases to 40 µm (Figure 3c,e). However, the enhancement in plasticity through micron-level grain refinement is relatively limited [54,55]. The density of dislocations also affects the plasticity of the alloy [56]. As all the samples under different processes were subjected to T6 treatment, the internal stress was eliminated. As shown in Figure 8a–c, dislocation entanglements are scarcely visible in the TEM images, indicating a very low dislocation density in T6, 300°C4P-T6 and 300°C4P-R80%-T6 alloys. Therefore, in the present study, the evolution of grain size and the density of dislocations have no significant impact on plasticity.

Coarse intermetallic phases have an adverse effect on the plasticity of the alloy, as they are usually the source of voids/cracks [57]. As shown in Figure 2b, the coarse intermetallic phases (mainly Al₇Cu₂Fe phases) in the T6 alloy are large in size and distributed unevenly. By performing four passes of ECAP at 300 °C prior to T6 treatment, the coarse intermetallic phases become finer and distributed diffusely due to fragmentation (Figure 2c), which is conducive to improving the plasticity of the alloy.

Furthermore, by comparing Figure 6a with Figure 5a,b, it can be seen that pre-ECAP results in a significant reduction in the width of the PFZ. The width of the PFZ at the grain boundaries of the 300°C4P-T6 alloy is significantly reduced to 44 nm (Figure 6a). PFZs distributed along the grain boundaries are much softer than the matrix. The preferential deformation of PFZs during deformation leads to the accumulation of dislocations and the concentration of stress. Gradually, microcracks nucleate at the interface of the PFZ and the matrix, which are detrimental to the plasticity of alloys [58,59]. With the PFZs refining, the negative influence on the plasticity will diminish.

As plastic deformation progresses, coarse grain boundary precipitates (GBPs) are usually the source of voids/cracks, leading to intergranular fractures and a significant reduction in plasticity [47]. The size of the GBP η phases of the 300°C4P-T6 alloy is more uniform and finer than that of the T6 alloy, with an average length of about 38 nm and a length–diameter ratio of about 2.7 (Figure 6a). This refinement of GBP η phases also contributes to the higher plasticity of the 300°C4P-T6 alloy.

In addition, there are a considerable number of nano-sized A1₆Mn precipitates in the grains, which are finer and elliptical in shape in the 300°C4P-T6 alloy (Figure 8a,b). During the deformation, when dislocations are obstructed by the A1₆Mn precipitates, they tend to change the slip system by means of cross-slip. This cross-slip mechanism allows material to maintain uniformly good plasticity throughout the deformation process [60,61].

In summary, the high plasticity of the 300°C4P-T6 alloy is mainly attributed to the optimized size and shape of its Al₆Mn phases. The fragmentation of the coarse intermetallic phases and the refinement of the GBP η phases and PFZs diminish their negative influence on plasticity.

When it comes to the 300°C4P-R80%-T6 and 300°C4P-T6 alloys, the EL further increases from 19% to 23.4%. The coarse intermetallic phases (Al₇Cu₂Fe phases) in 300°C4P-T6 and 300°C4P-R80%-T6 alloys are all fragmented into smaller sizes and diffusely distributed (Figure 2c,d). The width of the PFZs in the 300°C4P-R80%-T6 alloy is similar to that in the 300°C4P-T6 alloy (Figure 6a and Figure 7a). The distribution and the size of the η and η′ phases are almost the same, as can be seen in comparing Figure 6a,c with Figure 7a,c. However, the quantity of Al₆Mn precipitates in the 300°C4P-R80%-T6 alloy is significantly higher than that in 300°C4P-T6 alloy.

It is widely accepted that the transition temperature of Al₆Mn phases can be reduced by ECAP extrusion, which is closely related to the strain energy accumulated within the material as the number of passes increases [62]. The dislocations formed during severe plastic deformation increase the strain energy and this can be expressed in the following equation:

$$E_{stored}=G{b}^2{\displaystyle \frac{\rho }{4\pi \kappa }}In\left(b\sqrt{\rho }\right)$$

(3)

where G is the shear modulus of Al, b is the Burgers vector for the basal slip of Al, κ is arithmetic average of 1 and (1 − ν), ρ is the dislocation density and ν is Poisson’s ratio. Equation (3) is governed by ρ, where the higher the dislocation density ρ, the greater the strain energy. For the 300°C4P and 300°C4P-R80% alloys, the dislocation density increases with the accumulation of plastic deformation, leading to an increase in strain energy. This, in turn, lowers the precipitation temperature for the Al₆Mn phase, making it easier to precipitate. This is consistent with the observation that the 300°C4P-R80%-T6 alloy has a greater number of finer and elliptical Al₆Mn phases. The 300°C4P-R80%-T6 alloy has a significantly higher content of Al₆Mn phases, which are similar in size and shape to those in the 300°C4P-T6 alloy. Therefore, the further increase in plasticity of the 300°C4P-R80%-T6 alloy is mainly due to a larger number of fine Al₆Mn precipitates.

5. Conclusions

In this study, a 7046 aluminum alloy with high strength and excellent ductility was prepared using four passes of ECAP, RT rolling and T6 treatment. Based on the findings, we conclude the following:

(1)

The 7046 aluminum alloy processed by four passes of HT ECAP prior to T6 treatment realizes an obvious improvement in both strength and ductility, with a UTS of 485 MPa, YS of 366 MPa and EL of 19%, respectively. The mechanical properties of the alloy are further improved by an additional RT rolling process, with a UTS, YS and EL of 508 MPa, 393 MPa and 23.4%, exceeding those of the T6 alloy by 24.2%, 10.1% and 485%, respectively.

(2)

The strength enhancement in the 300°C4P-T6 and 300°C4P-R80%-T6 alloys is attributed to the combination of fine-grain strengthening and second phase strengthening by the presence of more η′ phases.

(3)

The superior plasticity of the 300°C4P-R80%-T6 and 300°C4P-T6 alloys compared to the T6 alloy can mainly be attributed to the optimization of the size and shape of their Al₆Mn phases. The fragmentation of the coarse intermetallic phases and the refinement of GBP η phases and PFZs diminish their negative influence on plasticity.

(4)

There are more precipitates of η′ and Al₆Mn phases in the 300°C4P-R80%-T6 and 300°C4P-T6 alloys, which is the main reason for the SPD process’s enhancement of the strengthening and toughening effect of T6-treated 7046 aluminum alloys.

6. Adding

All acronyms used in this paper have been summarized in

Table 3. Full name of acronyms used in this article.

Full Name	Acronyms
severe plastic deformation	SPD
equal channel angular pressing	ECAP
ultimate tensile strength	UTS
elongation	EL
room temperature	RT
ultrafine-grained	UFG
coarse-grained	CG
yield strength	YS
cross accumulative extrusion bonding	CAEB
high temperature	HT
wire electrical discharge machining	WEDM
optical microscopy	OM
transmission electron microscopy	TEM
electron backscatter diffraction	EBSD
average grain size	AGS

.