Investigation of Microstructures and Mechanical Properties of Ultra-High Strength Al-Zn-Mg-Cu Alloy Prepared by Rapid Solidification and Hot Extrusion

Abstract

Al-Zn-Mg-Cu aluminum alloys have the advantages of high specific strength, easy processing, and high toughness, showing great potential application in the aerospace field. However, ultra-high strength aluminum alloys usually contain coarse microstructures, micro-segregation, and casting defects that seriously deteriorate mechanical properties. Here, we report a high-strength aluminum alloy (Al-10.5Zn-2.0Mg-1.2Cu-0.12Zr-0.1Er) prepared by rapid solidification and hot extrusion to explore the microstructure modification of the alloy based on this strategy. The results show that: rapid-solidification technology can significantly refine alloy grains, alloy ribbons were composed of α (Al) equiaxed fine grains, and the average grain size was less than 6 μm. After extrusion, the alloy had partially recrystallized, existing coarse second-phase (T-phase) and needle-shaped precipitates were MgZn₂ (η-phase), and the tensile strength and elongation of the extruded bar were 466.4 MPa and 12.9%, respectively. After T6 heat treatment, the tensile strength of the alloy reached 635.8 MPa, while elongation decreased to 10.5%. According to microstructure analysis and considering the contributions of grain boundary, dislocation, and precipitation-strengthening to the improvement of the mechanical properties, it was found that precipitation-strengthening is the main strengthening mechanism. Our research shows that rapid-solidification and hot-extrusion technology have great potential for improving the microstructures and mechanical properties of aluminum alloys.

1. Introduction

Al-Zn-Mg-Cu aluminum alloys, as important lightweight structural materials for the aerospace and aircraft industries, provide low-density, high-specific strength, low-melting point, and high-hardness materials that are highly valued by the majority of materials researchers [1,2,3]. In recent years, research on the ultra-high strength Al-Zn-Mg-Cu alloy mainly has focused on optimized alloy-composition design, exploring the preparation technology of ultra-fine grain microstructure and adopting a new type of billet-making method, the development of new molding processing, and precision heat treatment, etc. [4].

As for high-Zn aluminum alloys, in a certain range, the mechanical properties of the alloy can be improved by increasing the Zn content, but the yield of traditional casting-metallurgical methods is low. These alloys consist of coarse, α (Al) dendrites and network-eutectoid structures due to the slow cooling rate. Generally, the traditional casting process cannot maintain the high strength and plasticity of the alloy after the Zn content exceeds 8% [5]. Meanwhile, hot cracks, micro-segregation, and solidification porosities of the casting defects usually form in these alloys. Therefore, more attempts are needed to further enhance the mechanical properties of Al-Zn-Mg-Cu aluminum alloys. Until recently, mainly three methods have been used to increase the mechanical properties of the alloy: (Ι) Alloying technique, by adding some rare-earth elements (such as Sc or Er) to enhance grain boundary strengthening and precipitation strengthening. The elements Sc and Al can form Al₃Sc particles which act as heterogeneous nucleation sites for the Al grains, and the Al₃Sc particles effectively pin the grain boundaries, resulting in a significant refinement of grain size of the alloy [6]. (ΙΙ) In a rapid-solidification technique, a high cooling rate is beneficial for increasing the rate of nucleation and grain refinement. (ΙΙΙ) Plastic deformation (such as extrusion or rolling) can eliminate the cast defects and increase the dislocation density to enhance the strength of the alloy [7].

The rapid-solidification technique has become one of the most common means to develop the potential of new materials. During this process, the cooling rate can reach 10⁶ K/s, and it is beneficial to microstructure refinement, increasing the solid solubility of alloying elements and reductions of composition segregation [8,9]. The materials prepared by rapid solidification are mainly divided into powder (gas atomization), ribbons (melt spinning), and non-compact bulk material (spray deposition) [10]. Moreover, melt spinning is considered an effective method and, due to the advantage of a faster solidification rate, it is widely used in aluminum alloys [11,12,13,14]. Kapinos et al. [15] used melt spinning and extrusion to prepare an Al-9Zn-2.5Mg-1.8Cu alloy in which the tensile strength of the hot-extruded sample was 405 MPa and the elongation was 17%. Meng et al. [16] also used melt spinning and extrusion to obtain an Al-27Zn-1.5Mg-1.2Cu-0.08Zr alloy in which the tensile strength of the extruded alloy was 485 MPa and the elongation was 5.2%. Nowadays, research on the melt-spinning method is mostly focused on the microstructures and properties of the ribbons. However, relevant research reports on the preparation of ultra-high strength aluminum alloys by melt spinning and hot extrusion are rare, and there is a lack of in-depth investigation.

In this work, an Al-10.5Zn-2.0Mg-1.2Cu-0.12Zr-0.1Er alloy was designed and fabricated by rapid-solidification melt spinning and hot extrusion. The objective of this investigation is to reveal the evolution of microstructures, the mechanical properties of the new alloy, and the strengthening mechanisms of the alloy. Furthermore, the fracture mechanism of the alloy during tensile process is also clarified.

2. Materials and Methods

In this study, an ultra-high strength aluminum alloy, Al-10.5Zn-2.0Mg-1.2Cu-0.12Zr-0.1Er, was selected for investigation. Metal mold-casting technology was used to fabricate the aluminum alloy (crucible-resistance furnace). The raw materials of the alloys included high-purity Al (99.99%), high-purity Zn (99.99%), high-purity Mg (99.9%), and master alloys such as Al-50Cu, Al-4Zr, and Al-6Er. Subsequently, the rapid-solidification single-roll melt-spinning method was used to prepare alloy ribbons that were 3–5 mm in width and 70–100 μm in thickness under a vacuum. The prepared alloy ribbons were pulverized by a high-speed universal pulverizer (its speed is 26,000 rpm), which has a mesh size of 40 to 200, and then cold pressed into billets with a diameter of 40 mm under 18 MPa for 2 min at room temperature (100t four-column hydraulic press). Subsequently, cylindrical alloy billets were extruded into 10 mm rods with an extrusion ratio of 16 at 400 °C (100t vertical extruder), and then the rods were cooled to room temperature (RT) in air. The chemical compositions of the alloy were confirmed by inductively coupled plasma luminescence spectra (ICP) and the testing results were listed in

Table 1. Chemical composition of the experimental samples.

Elements	Zn	Mg	Cu	Zr	Er	Al
Wt (%) At (%)	10.45 4.70	1.99 2.21	1.24 0.54	0.09 0.03	0.08 0.01	Bal. 92.51

. Finally, in order to improve the overall performance of the alloys, T6-heat treatment was carried out for the hot extruded rods; the process parameters were as follows: a solid solution temperature of 435 °C, a solid solution time of 2 h, after water-quenching to room temperature, and 4 h of artificial aging at 120 °C.

The experimental alloys for the longitudinal section and cross-section microstructure observations were cut parallel to and perpendicular to the hot-extrusion direction. The microstructures of the alloys were tested using Quanta FEG 650 scanning electron microscopy (SEM, C528FEI, Holland). X-ray diffraction (XRD, D8 ADVACEX, Bruker, Billerica, Massachusetts, Germany) was used to analyze the phase constituents of the alloy. The radiation of XRD analysis was CuKa with scanning angles of 10°–90°and a scanning speed of 2°/min. The grain size and grain orientation were characterized using SEM equipped with an electron backscatter diffraction (EBSD) detector. JEOL-JEM2100F transmission electron microscopy (TEM, JEOL, Tokyo, Japan) was used to describe the microstructure characteristics. The samples used for TEM testing were fabricated by a two-jet electro-chemical polishing process after mechanically grinding them to 80 μm, and the current and temperature used were 90 mA and −25 °C, respectively. These samples were preserved after TEM observations and then used for the EBSD observations. Otherwise, only mechanical polishing was applied for the SEM observations under the backscatter electron (BSE) mode. A CMT 5205 GL mechanical testing machine (Sansi Taijie Electric Equipment Co., Ltd., Zhuhai, Guangdong Province, China)was used for room-temperature tensile tests at an initial strain rate of 6 × 10⁻⁴·s⁻¹. The tensile sample was processed according to the national standard (GB/T228.1-2010), and the specimen dimensions are shown in Figure 1. The fracture surfaces of the alloys were observed by SEM under secondary electron mode.

3. Results and Discussion

3.1. Phase and Microstructures Evolution

Figure 2 shows the experimental procedures of the rapidly solidified ribbons and hot extruded rods. As shown in Figure 2a, for the prepared alloy ribbons with 3–5 mm in width and 70–100 μm in thickness, it can clearly be seen that the alloy-ribbon surface has no holes and good continuity. Figure 2b shows that the alloy ribbons were pulverized by a high-speed universal pulverizer, which has a mesh size of 40 to 200. Figure 2d shows hot-extruded rods with 10 mm in diameter; the surface quality of hot extruded rods was very smooth, and no obvious defects such as distortion, pitting, and cracks were found. In a word, hot-extrusion deformation promotes the effective welding of rapidly solidified ultra-high strength aluminum alloy.

The XRD patterns of as-rapidly solidified and as-extruded alloys are shown in Figure 3. The rapidly solidified sample is mainly composed of an α (Al) phase, and no diffraction peaks of second-phase particles were observed. As to the extruded sample, it consists of an α (Al), η-phase with a closely packed hexagonal (hcp) structure. At the same time, the intensity of the diffraction peaks of the α (Al) phase and η-phase were notably increased when compared with the rapidly solidified sample. This indicates that the precipitates (η-phase) were formed during the extrusion process. In addition, the lattice parameters of α (Al) in the rapidly solidified and extruded samples were 0.4042 ± 0.022% nm and 0.4044 ± 0.025% nm, respectively, with both being slightly lower than that of pure Al (0.4050 nm, face-centered cubic structure, PDF#04-0787, Fm-3m(225)), which is mainly due to the addition of the Zn solution into the α (Al) matrix alloy since the Zn atom has a smaller atomic radius compared to the Al atom. From the XRD results of the alloy, it can be concluded that the extrusion process has great effect on the type of phases.

Figure 2. Schematic illustration of rapid-solidification and hot-extrusion experimental procedures for the alloy: (a) ribbons; (b) powders; (c) cylinder; (d) extruded rods.

Figure 2. Schematic illustration of rapid-solidification and hot-extrusion experimental procedures for the alloy: (a) ribbons; (b) powders; (c) cylinder; (d) extruded rods.

Figure 4 exhibits the SEM microstructures, grain-size statistics of the conventional casting, and rapidly solidified ribbons. As shown in Figure 4a, it can clearly be seen that there are coarse dendrites in the conventional casting of the α (Al) grains with a non-equilibrium eutectic phase with network distribution, serious dendrite segregation, and an average size was 88.63 μm. At this stage, the content of the solid solution precipitates of the alloy was 2.4%. In addition, as displayed by Figure 4c, the rapidly solidified ribbons’ microstructures are fine-equiaxed α (Al) grains with an average size of 5.98 μm; compared with the conventional casting, the grain size was refined by 93.3%. It was found that the rapid-solidification technology can effectively reduce segregation and second-phase particles since the content of the solid solution precipitates of the alloy was reduced to 0.5%. This is because alloying elements dissolve into an α (Al) matrix to form a supersaturated, solid solution structure and realize a large solid solubility of the matrix under rapid-solidification conditions. Therefore, the α (Al) grains of the rapidly solidified sample contained a higher solid solubility of alloying elements (Zn, Mg, and Cu), and the microstructure characteristics were notably different from the conventional casting ultra-high strength Al alloys.

In order to observe the microstructure of the as-extruded sample more intuitively, we investigated the cross-section and longitudinal section in two directions. The longitudinal section and cross-section microstructures of the hot-extruded samples are shown in Figure 5a and Figure 5b, respectively. As can be seen in Figure 5a, the longitudinal section microstructure belong to typical extrusion streamline structure. Figure 5c shows the magnified image of the rectangular area in Figure 5a in which many coarse second-phase particles are clearly seen. The EDS measurement (

Table 2. Chemical compositions of tested points with the experimental samples (at.%).

Points	Al	Zn	Mg	Cu	Zr	Er
1	85.84	7.89	3.60	2.36	0.05	0.25
2	88.75	7.18	3.11	0.95	-	0.02
3	88.47	6.98	3.71	0.81	0.03	-
4	91.52	5.29	2.53	0.64	0.03	-
5	94.17	3.67	1.74	0.39	0.04	-
6	90.54	5.54	2.95	0.90	-	0.06
7	78.99	12.12	6.80	2.09	-	-
8	81.91	10.30	6.20	1.57	-	0.02
9	77.47	11.54	9.15	1.81	0.03	-
10	78.09	11.67	8.34	1.88	-	-

) reveals that the bright-white particles (points 1, 2, 3, 4, 5, and 6) within the α (Al) grains are rich in Zn, Mg, Cu, and Al, with a Mg/Zn ratio of about 1:2. It can be considered that the bright-white phase particles distributed along the extrusion streamline are mainly T-phase (AlZnMgCu), and such needle-shaped η-phases were formed during the extrusion process. This is the reason why the density of the η-phase of the extruded sample is higher than that of the rapidly solidified sample. The cross-section image (Figure 5b) reveals that the bright precipitates were embedded into the α (Al) grains, and the distribution of η-phase particles becomes more even compared to the rapidly solidified sample. Figure 5d shows the magnified image of the rectangular area in Figure 5b. We analyzed the EDS measurement of the second-phase particles (points 7, 8, 9, and 10), as shown in Table 2, and in combination with related literature [15], and performed an analysis of the EDS results, in addition to the presence of the η-phase, a small amount of T-phase (AlZnMgCu), and Al₃(Er.Zr) particles. The T-phase was dynamically precipitated and continuously became coarser during hot extrusion.

The typical SEM microstructure of the aged sample is shown in Figure 6. Compared with the extrusion sample (Figure 5), the density of the coarse phase particles of the aged sample was significantly reduced. In addition, compared with the experimental results of related literature [17,18], after aging treatment, a large number of fine particles appeared in the aged microstructure of the alloy, and the size of the remaining large MgZn₂ phase was also significantly reduced. In summary, the SEM microstructure results in Figure 6 will be further discussion later.

In order to quantitatively analyze the grain size and misorientation angle distribution (MAD) of the alloy, EBSD was carried out. Figure 7 displays the EBSD microstructures of the rapidly solidified sample. Pixels with the same colour indicate the same crystallographic orientation in inverse pole-figure (IPF) maps. From the results of the inverse pole-figure (IPF) map in Figure 7a, the grain orientation of the rapidly solidified samples were rather random. The average grain size of the rapidly solidified sample in Figure 7b was about 5.64 μm, consistent with the SEM microstructure results of the rapidly solidified ribbons. The fraction of low-angle grain boundaries (LAGBs, θ ≤ 15°) of the rapidly solidified sample in Figwas about 4.9%, and its high-angle grain boundaries ratio (HAGBs, θ > 15°) was 95.1%. It was also proved that the rapidly solidified ribbons are composed of supersaturated solid solution α (Al) equiaxed fine grains. This is because rapid solidification is a non-equilibrium solidification process where a large initial nucleation undercooling is formed in the alloy melt, correspondingly achieving a high-solidification rate in the process and finally obtaining a fine and excellent solidified microstructure. Moreover, the grain size of the rapidly solidified Al alloy is related to the large under-cooling, increasing nucleation frequency and leading to a significant refinement of the α (Al) grains. At the same time, the rapid cooling rate can significantly reduce the diffusion coefficients of solute atoms in the α (Al) matrix, resulting in the formation of a supersaturated solid solution.

Figure 8 is the EBSD images of the extrusion and aged sample, which shows the inverse pole-figure (IPF) maps and distribution-frequency maps of the misorientation angles. As displayed by Figure 8a, during the process of hot extrusion, some grain orientations are mainly concentrated in the <101> direction, similarly to the research results of Meng et al. [16]. This shows that the grain alignment of the alloy along the extrusion direction tends to the <101> direction during hot extrusion. The hot extruded sample consists of equiaxed-fine grains and elongated grains. These elongated grains exhibit an arrangement along a specific direction parallel to the extrusion direction which is characteristic of the Al-alloy extrusion [19]. At the same time, this phenomenon indicates that the recrystallization has occurred. In this work, the hot-extruded sample was cooled in air, indicating a relatively slow cooling rate. Hence, the recrystallization behavior may take place during the hot-extrusion deformation and cooling stage. For the aged sample, similar characteristics were also discovered in Figure 8b, e.g., a high degree of recrystallization and larger grain size, which was further demonstrated by the reduction of LAGBs (from 43.2% to 16.5%) and the increase of HAGBs (from 56.8% to 83.5%). Therefore, the corresponding distribution-frequency maps of misorientation angles in Figure 8c,d are the better explanation.

Figure 9 shows the distribution characteristics of recrystallized microstructures of the extrusion and aged sample. The blue areas represent the recrystallized grains and yellow areas represent the sub-grains, in addition, areas colored by red represent the deformed grains of the alloy. There are more recrystallized grains distributed in the aged sample, and its volume fraction reached 86.1%, while the volume fraction was 67% in the extrusion sample. As for the sub-grains, the volume fraction was 13.9% in the extrusion sample, close to 13% in the aged sample. Meanwhile, the volume fraction of the deformed grains was 19.1% in the extrusion sample compared with 0.9% in the aged sample. The results further illustrate that the alloy has a higher degree of recrystallization during the aging process, which more intuitively proves the EBSD results of Figure 8.

Figure 8. EBSD results of (a) and (b) IPF maps of the extrusion and aged sample; (c) and (d) the corresponding distribution-frequency maps of misorientation angles of (a) and (b), respectively.

Figure 8. EBSD results of (a) and (b) IPF maps of the extrusion and aged sample; (c) and (d) the corresponding distribution-frequency maps of misorientation angles of (a) and (b), respectively.

Figure 10 shows the IPF maps of recrystallized microstructures acquired by an EBSD analysis of the extrusion and aged sample, and the corresponding distribution-frequency maps of grain sizes. As displayed in Figure 10a,b, the recrystallized grains of the hot-extruded and aged sample are arranged along a specific direction, which is parallel to the extrusion direction. The average grain size of the extrusion sample was about 8.28 μm (Figure 10c). It is obvious that the grain size of the extrusion sample was slight larger than that of the rapidly solidified sample since the recrystallization already occurred after extrusion, and the grains of the alloy had grown. In addition, the average grain size of the aged sample further increased from 8.28 μm to 22.78 μm (Figure 10d), and it could clearly be seen that heat treatment also had a certain effect on the grain size of the alloy. However, a slight increase in grain size had little effect on the strengthening contribution of the alloy, which is clarified in the following grain boundary-strengthening mechanism section.

Figure 11 is the TEM images of the hot-extruded and aged samples for which different parts of the thin region of the samples were observed. As displayed by Figure 11a, it can clearly be seen that there are some coarse T-phases with an average length of about 0.64 μm in the α (Al) grains, and such equilibrium needle-shaped η-phase of about 150 nm in length had an unfixed orientation relationship with the α (Al) matrix alloy [20]. In Figure 11b, there were some precipitated phases of about 60 nm in size at the grain boundaries (GBs); these precipitates may be η-phase. From an observation of Figure 11c, on one hand, high-density dislocations were discovered in the grains, with a great number of dislocation lines interweaving and interwining together, resulting in the formation of dislocation tangles. On the other hand, some Al₃(Er,Zr) particles were located near dislocations. This indicates that a large amount of deformation takes place during the hot-extrusion process. In Figure 11d, part of the dislocations also embraced the Al₃(Er,Zr) particles with the average size about 42 nm and a volume fraction of 1.1%. The Al₃(Er,Zr) particles can pin GBs and dislocations and improve the strength of the alloy, consistent with the research results of Wu et al. [21]. As displayed in Figure 11e, there were intermittently distributed precipitates at the GBs and a narrow precipitation-free zone (PFZ) near the GBs. As seen in Figure 11f, a high-magnification TEM image further shows high-density nano-sized precipitates of the alloy and the selected area’s electron diffraction pattern (SAED), indicating the existence of a GP-zone and η’-phase. The average size of one precipitate (GP-zone) was about 4 nm, and its volume fraction was 0.8%. Moreover, the average length of the other precipitate (η′-phase) was about 8 nm, and its volume fraction was 1.6%. As a result, the microstructure of the aged sample mainly consisted of GP-zones, η′-phases, and Al₃(Er,Zr) particles.

3.2. Mechanical Property Measurements

Figure 12 exhibits the tensile stress-strain curves of the extruded and aged samples. The corresponding values of the tensile-property parameters, ultimate tensile strength (UTS), yield strength (YS), and elongation (EL) are summarized in

Table 3. Tensile mechanical-property parameters of the experimental samples.

Samples State	Ultimate Strength (MPa)	Yield Strength (MPa)	Elongation (%)
Extrusion	466.4	296.3	12.9
Aging	635.8	540.9	10.5

. As may be observed, the tensile strength of the hot-extruded rod after heat treatment was significantly improved, and elongation was reduced; after the flow stress reached its peak rapidly, a fracture soon occurred. The tensile strength of the hot-extruded bar after extrusion was 466.4 MPa, with an elongation reaching 12.9%. Combined with the literature [22,23] and the SEM microstructure of the aged sample, it was found that as the hot-extruded bar was subjected to T6 heat treatment, more solute atoms were precipitated from the solid solution, strengthened by the fine-precipitation phase, and the tensile strength increased from 466.4 MPa to 635.8 MPa, with an elongation of 10.5%. The grain size and precipitates distribution of the hot-extruded bar is one of the most important factors that affect the tensile strength and elongation of the alloy. Meng et al. [18] revealed that different types of strengthening mechanisms, which include grain-boundary strengthening, dislocation strengthening, and precipitate strengthening, play an important role in the strengthening of the alloy. In this work, the total contributions of the strengthening mechanisms are as follows [24,25,26,27]:

$${\sigma }_{YS}{=\sigma }_i{+\sigma }_g{+\sigma }_{\rho }{+\sigma }_{pct}$$

(1)

$$\Delta {\sigma }_g{=\sigma }_0{+K}_y{d}^{-1/2}$$

(2)

$$\Delta {\sigma }_{\rho }{=M\alpha \mathrm{Gb}\rho }^{1/2}$$

(3)

$$\Delta {\sigma }_{pct}=M\frac{0.4\mathrm{Gb}}{\mathsf{\pi }\sqrt{1-v}}·\frac{\ln (\frac{\sqrt{\frac{2}{3}}}{b}d)}{\sqrt{\frac{2}{3}}d(\sqrt{\frac{\mathsf{\pi }}{{4V}_p}}-1)}$$

(4)

Where the ${\sigma }_{YS}$ is the total yield strength of the alloy, ${\sigma }_i$ is the lattice-friction stress (35 MPa), $\Delta {\sigma }_g$ is the contribution of grain-boundary strengthening to yield strength, ${\sigma }_0$ is a constant (16 MPa), which is equivalent to the value of yield strength of a single crystal, $K_y$ is a constant (0.12 MPa m^1/2), which indicates the influence of grain boundaries on the strength of the alloy, and $d$ is the average grain size (22.78 μm). According to the Hall–Petch formula, with a decrease in grain size, the strength of the alloy increases. $\Delta {\sigma }_{\rho }$ is the contribution of dislocation strengthening to yield strength, the $\alpha$ is 0.33, the ${\rm M}$ is the Taylor factor (3.0), the $G$ is the shear modulus (26 GPa), the $b$ is the Burgers vector (0.286 nm), and the $\rho$ is the dislocation density (3.5 × 10¹³ m⁻²) of the alloy, which can be calculated based on the full-width at half-maxima (FWHM) of the XRD peaks by the modified Williamson–Hall analysis [28,29]. $\Delta {\sigma }_{pct}$ is the contribution of precipitation strengthening to yield strength, $\mathsf{\nu }$ is the poisson ratio (0.33),$d$ is the average size of the precipitates, and $V_P$ is the volume fraction of the precipitates [30]. For the precipitation strengthening, three kinds of the precipitates were estimated in the ultra-high strength Al alloy, and all of the precipitates were distributed within the grains. The contribution of the strengthening mechanism to the yield strength of the alloy was calculated based on Equations (1)–(4) and the calculation results are summarized in

Table 4. Contributions from different strengthening mechanisms in the alloy.

Calculated Data (MPa)					Experimental Data (MPa)
$\Delta {\sigma }_i$	$\Delta {\sigma }_g$	$\Delta {\sigma }_{\rho }$	$\Delta {\sigma }_{pct}$	${\sigma }_{YS}$	${\sigma }_{YS}$
35.0	41.2	102.7	343.8	522.7	540.9

. It can be found that the precipitation strengthening is the main strengthening mechanism of the alloy. In addition, grain boundary strengthening and dislocation strengthening also play an important role in strengthening of the alloy.

Figure 13 is the comparison of the mechanical properties, ultimate tensile strength vs. the elongation of high-Zn aluminum alloys, from the related literature and the present work [15,16,31,32,33,34,35,36,37,38,39]. It was found that the Al-10.5Zn-2.0Mg-1.2Cu-0.12Zr-0.1Er alloy reveals a relative higher tensile strength, compared with that of the other high-Zn aluminum alloys prepared by die-casting Al-20Zn-2.5Cu alloy [31], gravity casted Al-Zn alloys [32,36], rolling [33], melt-spun and extruded-Al alloys [15,35], modified powder hot extrusion 7055 alloy [37], spray-formed and hot extrusion alloy [38], and melt-spun and extrusion ratio-39.1 alloy [39]. In addition, because of the higher Zn content, the tensile strength of two melt-spun and extruded Al alloys [16,34] were slightly higher than the alloy in this work. On one hand, the elongation of melt-spun and extruded Al-27Zn-1.5Mg-1.2Cu-0.08Zr alloy was relatively low. On the other hand, the melt-spun and extruded Al-10.7Zn-2.4Mg-0.9Cu alloy has higher extrusion ratio, and no rare earth elements were added to form nanoscale precipitates. Hence, the comprehensive mechanical properties of the alloy are very excellent. In conclusion, ultra-high strength Al-Zn-Mg-Cu alloy with high strength and reasonable ductility was also successfully prepared by rapid solidification and hot extrusion in this paper.

3.3. Fracture Characteristics

In order to study the fracture mechanisms of the extrusion and aged sample in-depth, the fracture-surface morphologies of the alloy were shown as in Figure 14. It can be seen that a lot of secondary cracks and tearing ridges were observed in Figure 14a. Further magnifying the SEM image of the rectangular area revealed that a large number of equiaxed dimples with less than 2 μm in diameter were distributed in the fracture surface of the extruded sample (Figure 14c), which is characteristic of ductile fracture, wherein the dimples are commonly considered as characteristic of ductile fracture. Figure 14b shows that for the fracture surface of the aged sample, the dimple sizes were not uniform and the depth was shallow, with partially seen, flat and smooth interfaces and cleavage facets, and the fracture mechanism becomes a ductile and brittle mixed fracture. The magnified image of the rectangular area of Figure 14b is shown in Figure 14d. The formation of dimples is centered on the second phase, due to the existence of a large number of second phases, and these second phase particles may become dimple centers. As a result, the size of the dimples in the fracture morphology is uneven. The tensile-fracture morphologies are consistent with the mechanical-property test results shown in Figure 12.

4. Conclusions

The Al-10.5Zn-2.0Mg-1.2Cu-0.12Zr-0.1Er alloy was fabricated by rapid solidification and hot extrusion in this study. The microstructure characteristics and mechanical property of the alloy were studied by XRD, SEM, EBSD, and TEM. The major conclusions are as follows:

(1)

The alloy ribbons prepared by the rapid-solidification single-roll melt-spinning method has a fine microstructure, consisting of fine-equiaxed grains with an average grain size of less than 6 μm; compared with the conventional casting, the grain size of a cast alloy was 88.63 μm, so the grain size was refined by 93.3%. The alloying elements were dissolved into the α (Al) matrix to form a supersaturated solid solution.

(2)

After hot extrusion, the alloy has partially recrystallized, and existing coarse T-phases and the needle-shaped precipitates were η-phases. Part of the coarse second phases redissolved after the solid-solution treatment and nano-sized fine precipitates (GP-zone, η´-phase and Al₃(Er,Zr)) were formed after the aging treatment.

(3)

The tensile strength of the rod prepared by rapid solidification and hot extrusion was 466.4 MPa, with the elongation reaching 12.9%. After T6-heat treatment, the aged alloy presented an ultimate tensile strength of 635.8 MPa and an elongation of 10.5%. The combination of grain-boundary strengthening, dislocation strengthening, and precipitation strengthening synergy improve the mechanical properties of the alloy. It can be concluded that precipitation strengthening is the main strengthening mechanism of the alloy.

(4)

The tensile fracture of the extruded alloy shows the characteristics of a ductile fracture. Moreover, the T6 state that tensile fractures show the characteristics of a ductile-brittle mixed fracture, and the elongation of the alloy has decreased.