梯度热处理快速优化Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn合金获得高强高韧性能的微观组织

1 Introduction

Metastable β titanium alloy has great potential application in aerospace and biomedical industries due to their high specific strength, good combination of strength and ductility [1-2]. The development of aerospace industry requires higher performance of titanium alloy, facilitating the design and development of new β-Ti alloys [3-6], Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy is one of them. As a high strength metastable β-Ti alloy, the strength and ductility is very sensitive to microstructure mainly depending on solution and ageing treatment. It is a key way to tailoring mechanical properties through establishing the relationship between heat treatment and microstructures.

Previous studies found that although the ductility of metastable β titanium alloy significantly depends on the size of β-grain and primary α phase, strength considerably depends on the morphology, size and volume fraction of αs phase [7-9]. Therefore, controlling the size, morphology, and distribution of α phase through heat treatment is one of important ways for tailoring mechanical properties of metastable β-Ti alloys [4, 10-12]. Establishing the relationship between heat treatment parameters and microstructure of metastable β-Ti alloy is a relatively complex process, including annealing, solution and ageing. After different solution and ageing treatments, the microstructure of the alloy could be tailored to equiaxed, lamellar, bimodal and even more hierarchical features [13-14]. Besides, the morphology and volume fraction of the primary α phase, size and spacing of the secondary phase are also very sensitive to solution and ageing. In Ti-5Al-4Zr-8Mo-7V alloy, αs with    51 nm in width and 85 nm in spacing resulted in ultimate strength of 1390 MPa with elongation of 10.3% after the solution treatment at 800 ℃ and ageing at 570 ℃ for 8 h [9]. REN et al [15] achieved good combination of strength and ductility in Ti5231 alloy, ultimate strength of 1238 MPa and elongation of 20%, owing to the microstructure consisted of 13 vol% of αp and αs of 187 nm in spacing after solution treatment at 830 ℃ and ageing at 620 ℃ for 6 h. Through changing the solution and ageing temperature, the size and fraction of both αp and αs phases could be tailored for achieving good combination of strength and ductility, such as Ti7333 [16-17], Ti1023 [18] and Ti55531 [11, 19-20] alloys.

In recent years, some efforts have been made to take advantage of high-throughput technologies to acquire amounts of microstructure features rapidly in order to tailor and optimize microstructure and mechanical properties of titanium alloy. AFONSO et al [21] obtained different cooling rates for           Ti-20Nb alloy by Jorminy quenching test to study the relationship among cooling rates, different microstructures and mechanical properties. XU et al [22] accurately determined the pseudo-spinodal decomposition temperature of Ti5553 alloy through gradient heat treatment and obtained a high volume fraction of small size α phase by pseudo-spinodal decomposition, resulting in a very high strengthening effect. The continuous component gradient can be achieved by diffusion multiple experiment, and the effect of components on performance can be determined conveniently and accurately. By this method, WU et al [23] studied the effect of Mo element and V content on the microstructure of Ti-Mo-V alloy in gradient composition by high-throughput multiple sample, and designed Ti-6Mo-3V alloy with ultrafine α phase, which has yield strength of 1411 MPa and elongation of 6.5%. ZHANG et al [24] rapidly established the relastionship of “composition-microstructure-elastic modulus” of Ti-Nb-Zr system and bulid the elastic modulus and hardness database.

In this study, a convenient high-throughput heat treatment approach was developed, which could create temperature gradient for solution and ageing treatment in only one sample. The ageing hardening behavior and microstructural evolution of Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy were rapidly studied for optimizing microstructure for improving strength and ductility.

2 Experimental

2.1 Materials

The as-received alloy was forged rod supplied by the Baotai Group Co., Ltd. The chemical composition of the alloy is listed in Table 1. The initial microstructure shown in Figure 1 consists of fine bimodal α+β microstructure with approximately 14 vol% equiaxed αp phase and fine dispersed lamellar αs phase, and the average size of αp phase is about 2-3 μm.

Table 1  The chemical composition of Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy( wt% )

Al V Mo Cr Nb Fe Zr Sn O Ti

3.89 1.22 6.78 2.78 2.01 0.048 1.05 1.00 0.107 Bal.

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Figure 1  SEM image of as-received alloy

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2.2 Preparation of gradient sample

A round-rod sample of 10 mm in diameter and 92 mm in length was wire-cut from the as-received alloy. Figure 2 shows the schematic of the gradient heat treatment. The tube furnace was used for gradient heat treatment with a precisely-programmable temperature-controlled zone in the middle part. The temperature was set at 950 ℃, and decreased gradually from 950 ℃ to 25 ℃ in the stokehole. The position of gradient temperature was tested using standard sample by thermocouple accurately, as shown in Figure 2. In order to achieve accurate gradient temperature, nine equally spaced holes were punched in the sample by electric spark drilling with 10-mm spacing, and then the thermocouple wires were placed in the holes. To prevent oxidation of the sample during gradient heat treatment, Cr2O3 powder was mixed with water and then evenly coated on the sample surface before gradient solution treatment. After water evaporated, the sample was put in a porcelain boat and placed to the gradient position in the furnace.

Figure 2  Schematic illustration of gradient heat treated sample preparation

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As can be seen in Table 2, the measured temperature from thermocouple were 746, 770, 795, 819, 844, 865, 885, 900 and 909 ℃, respectively. After gradient solution for 2 h, the sample was quenched in water immediately. The β-transus temperature of the alloy is about (845±5) ℃. The gradient solution in α+β and β phase regions could be obtained in only one sample. Then, the quenched gradient solution sample was cut into four identical sheets: one of them was not further aged, and the other three were aged for 8 h at 450, 550 and 600 ℃, respectively. Before ageing, Cr2O3 was also evenly coated on each sample as an anti-oxidation layer.

Table 2  Gradient temperature measured by thermocouple wires

Position Temperature/℃

TC1 746

TC2 770

TC3 795

TC4 819

TC5 844

TC6 865

TC7 885

TC8 900

TC9 909

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2.3 Microstructure observation

The microstructure of gradient solution treated and aged samples were observed using MIRA2 LMH scanning electron microscope (SEM) and Tecnai G2 transmission electron microscope (TEM) operated at 200 kV. For TEM observation, the thin foils were prepared by a twin-jet electro polishing technique using Kroll’s reagent, which composed of 5% perchlorate, 35% butyl alcohol and 60% methanol. The average grain size and volume fraction of αp phase, the thickness and length of αs phase were measured statistically by ImageJ software.

2.4 Mechanical property testing

Hardness testing was conducted on 200HBVS-30 Vickers hardness tester with 9.8 N load. Each group of data has seven hardness values and the average value was used. Tensile properties were tested according to GB/T228—2010 standard. Before heat treatment, samples were cut into the round-rod shape firstly, after heat treatment, the round-rod sample was processed into a standard 25-mm long tensile specimen 5 mm in diameter. The tensile specimen is shown in Figure 3. Tensile tests were conducted on MTS Landmark at room temperature with strain rate of 10-3 s-1, and a strain extensometer was adopted to ensure the accuracy of stress-strain data measurement.

Figure 3  The dimensions of tensile specimen (Unit: mm)

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3 Results

3.1 Microstructure of the solution treated alloy

Figure 4 shows the microstructure of gradient solution treated sample. The volume fraction of αp decreases with the increase of solution temperature, from 30% at 746 ℃ to 3% at 819 ℃. When the solution temperature was above the β-transus temperature, no αp phase was observed at 844 ℃, which means that the α phase transformed to            β between 819 and 844 ℃. When the sample was solution treated at 746 ℃ and 770 ℃, the primary α phase had two different morphologies: globular α phase (αp) with diameter of 1-5 µm and rod-shaped α phase (αr) with width of 0.1-0.4 µm and length of 0.5-4 µm. In Figures 4(c) and (d), only the globular α phase could be observed near to β-transus temperature.

Figure 4  SEM images of Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy solution treated at different temperatures: (a) 746 ℃; (b) 770 ℃; (c) 795 ℃; (d) 819 ℃; (e) 844 ℃ (The number in the top right corner shows the volume fraction of αp in the β  matrix)

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Figure 5 shows the TEM images of 900 ℃ and 746 ℃ solution treated sample in order to observe whether the athermal ω phase (ωath) precipitated during quenching. Some periodic striations were observed in the bright-field TEM images. Inserted selected area electron diffraction patterns show that, other than the faint diffuse scattering, there are no reflections at the 1/3 and 2/3 (112)β positions, which could indicate the existence of ωath phase. This type of periodic striation was also observed in the other metastable titanium alloys [25-27], showing the spinodal decomposition feature. When the titanium alloy has sufficient quantities of β-stability elements, the strong driving force of phase separation during quenching leads to the differentiation of β phase into β-lean and β-rich region, and eventually to obvious lattice distortion. The lattice distortion results in localised atomic or nano-scale structural modulation in the β-lean region, i.e., embryonic ω, which has an intermediate structure between β and ω phase [28].

Figure 5  Bright field and selection electron diffraction pattern of (a) 746 ℃ ST sample; (b) 900 ℃ ST sample (ST stands for solution treatment)

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3.2 Gradient microstructures during solution and ageing

3.2.1 α+β solution following ageing

The SEM images of the alloy after α+β solution treatment at 746, 770, 795 and 819 ℃ for   2 h, and ageing at 450, 550 and 600 ℃ for 8 h are shown in Figure 6. The equiaxed αp phase appears at prior β grain boundary, which indicates that αp phase could restrain the growth of β grain, and improve the ductility [29]. The increase of solution temperature reduces the volume fraction of αp and affects the morphology and size of αs [6, 30]. As shown in Tables 3-5. The width and phase spacing of αs decrease with the increase of solution temperature. The statistical results show that when the ageing temperature is 600 ℃, the width and spacing of αs decrease from 57 and 142 nm to 47 and 65 nm as the solution temperature increases from 746 ℃ to 819 ℃, respectively, and the same regular was observed when the ageing temperature is 450 ℃ and 550 ℃.

Figure 6  SEM images of Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy after gradient heat treatment: (a1) 746 ℃, (b1) 770 ℃, (c1) 795 ℃ and (d1) 819 ℃ solution treatment for 2 h; (a2-d2) 450 ℃, (a3-d3) 550 ℃ and (a4-d4) 600 ℃ for 8 h

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It can be seen from Table 3, the αp phase slightly coarsens with the increase of ageing temperature, which indicates lower stability during ageing. The size of αs and the spacing between αs are very sensitive to ageing temperature. The comparison of (a2)-(a4), (b2)-(b4), (c2)-(c4), and (d2)-(d4) in Figure 6 reveals that, at the same solution temperature, αs coarsens with the increase of ageing temperature. Statistical results on spacing and width of αs after solution treated at 746 ℃, 795 ℃, 819 ℃ followed by ageing temperatures are shown in Tables 4 and 5. It can be seen that the width and phase spacing of αs increase by 3-4 times when the ageing temperature increases from 450 ℃ to 600 ℃ after solution treated in the temperature range of 746-819 ℃. Solution treated at 746 ℃, the width and phase spacing of αs increased from 17 nm and 35 nm ageing at 450 ℃ to 57 nm and 142 nm ageing at 600 ℃, respectively.

Table 3  Statistical results of diameter of αp (dp), width of αs (ws) and spacing of αs (λ) in 746 ℃ solution treated sample followed by ageing

Ageing temperature/℃ dp/μm ws/nm λ/nm

450 2.76 17 35

550 3.05 39 78

600 3.17 57 142

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Table 4  Statistical results of width of αs in the solution treated samples at 746 ℃, 795 ℃ and 819 ℃ followed by ageing

Ageing temperature/℃ ws/nm

746 ℃ 795 ℃ 819 ℃

450 17 15 14

550 39 34 29

600 57 51 47

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Table 5  Statistical results of spacing length of αs (λ) in the solution treated samples at 746 ℃, 795 ℃ and 819 ℃ followed by ageing

Ageing temperature/℃ λ/nm

746 ℃ 795 ℃ 819 ℃

450 35 25 22

550 78 58 32

600 142 105 65

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After solution at 746-770 ℃ for 2 h and following ageing at 550 ℃, 600 ℃ for 8 h, the hierarchical structure composed of micron scale αp phase, submicron scale αr phase and nanometre scale αs phase was created. Solution treated at 746 ℃ for 2 h following ageing at 600 ℃ for 8 h, the αs phase becomes thicker with the average thickness of 57 nm. The increase of thickness of αs phase makes the crack propagation path become more tortuous and require more energy to bypass αs, which in turn increase the ductility. The homogeneity of the strain gradient in the hierarchical α-structure is beneficial for the ductility enhancement of titanium alloy [13].

3.2.2 β solution following ageing

Figure 7 shows the gradient microstructure features of the alloy solution treated at 844-909 ℃ for 2 h following ageing at 450-600 ℃ for 8 h. When the aging temperature is 450 ℃, the αs phase did not appear in β matrix. When the ageing temperature is 550 ℃, fine αs phase precipitated in β matrix. When the ageing temperature is 600 ℃, the αs phase slightly coarsens. Moreover, when solution treated near β-transus temperature, αs is much        thicker. At the temperature range of 865-909 ℃, the solution treatment temperature has no obvious effect on the morphology and size of αs.

Figure 7  SEM images of Ti-6.8Mo-3.9Al-2.8Cr-2Nb-1.2V-1Zr-1Sn alloy after gradient heat treatment: (a1) 844 ℃, (b1) 865 ℃, (c1) 885 ℃, (d1) 900 ℃ and (e1) 909 ℃ solution treatment for 2 h; Represent ageing at (a2-e2) 450 ℃, (a3-e3) 550 ℃, and (a4-e4) 600 ℃ for 8 h

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3.3 Mechanical properties of the aged alloy

The age hardening behavior of the alloy after gradient solution treatment is shown in Figure 8. The hardness values of gradient solution treated sample were within HV 283-295. After ageing, the gradient solution treated sample shows obvious age hardening. Solution treated sample in β-phase region has more significant age hardening than that in solution treated sample in α+β region. Under the same ageing condition at 450 ℃, when the solution temperature increases from 746 ℃ to 819 ℃, the microhardness increases from HV 419.8 to             HV 482.8; but when the solution temperature increases from 844 ℃ to 909 ℃, the microhardness increases from HV 507.8 to HV 514.6, ageing at 550 ℃ and 600 ℃ the samples show similar results. This indicates that the age hardening of solution in α/β phase region is more sensitive to solution temperature than that in β phase region. In addition, at 450 ℃ the age hardening is stronger than that at 550 ℃ and 600 ℃.

Figure 8  Age hardening curves of gradient solution treated alloy aged at 450, 550 and 600 ℃ for 8 h

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XRD patterns in Figure 9 show that only α and β phase peaks appeared in the samples after solution treated at 746, 795, 819 and 900 ℃ following ageing at 550 ℃ for 8 h, indicating that there is only α phase in β matrix in those samples, which is consistent with the SEM microstructure shown in Figures 6 and 7, indicating that the hardening effect of the alloy resulted from αs phase. The increase of ageing temperature accelerated the decomposition of metastable phase and promoted the formation of equilibrium phase [31]. The increase of ageing temperature from 550 ℃ to 600 ℃ accelerated the precipitation of αs phase. Therefore, it can be inferred that there is also only α phase in β matrix ageing at 600 ℃ for 8 h.

Figure 9  XRD patterns of alloy solution treated at 746, 795, 819 and 900 ℃ for 2 h, and then aged at 550 ℃ for 8 h

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Based on the age hardening curve and microstructure characterization of gradient sample, the conditions possibly combined high strength and ductility were selected to evaluate the tensile properties, as shown in Figure 10. The alloy exhibits high yield strength (YS) of 1457 MPa but a relatively low elongation of 2.1% after solution treatment (ST) at 819 ℃ followed by ageing at 550 ℃. When aged at 600 ℃, the strength is decreased, but the ductility is largely improved. A superior combination of strength and ductility is achieved when ageing at 600 ℃ after ST at 746 ℃, the alloy obtained a good combination of elongation of 15% and yield strength of 1140 MPa. The strength decreases with ageing temperature increase, while the ductility shows an opposite trend. This is attributed to the decrease of αs phase length and increased thickness with the increase of ageing temperature.

Figure 10  The engineering stress-strain curves of solution treated and aged alloy samples (“ST” stand for solution treament; “UTS” stands for ultimate tensile strength; “EL” stands for elongation; “A” stands for ageing treatment)

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The solution temperature greatly influences strength and ductility at the same ageing temperature. Ageing at 600 ℃ for 8 h, when the solution temperature was increased from 746 ℃ to 819 ℃, YS was improved from 1140 to 1383 MPa, while the elongation decreased from 15.0% to 4.4%. Ageing at 550 ℃ for 8 h showed the same trend. This is due to the decrease of volume fraction of αp phase significantly, and the thickness and spacing of αs phase decrease as well.

3.4 Fractography of the aged alloy

Figure 11 shows the tensile fractographs of the specimen after solution treated at 746 ℃ for 2 h following ageing at 600 ℃ for 8 h. It can be seen that there are many dimples and some secondary cracks on the fracture surface. The fluctuation of fracture surface is significant, indicating that the crack path is tortuous. The fracture surface can be clearly divided into the shearing area and dimpled region, which has a typical cup cone shape with rough edges, indicating a considerable macroscopic plastic deformation before the final fracture. The fracture morphology shows a complete ductile fracture with a small secondary crack near the shear zone. The dimples in Figure 11(b) indicate the improved ductility.

Figure 11  Fractographs of the alloy ageing at 600 ℃ for 8 h with a solution treatment of 746 ℃ for 2 h: (a) Macro-fractography; (b) Dimple fracture region

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4 Discussion

4.1 Influence of solution treatment on microstructure

The solution treatment temperature has an important effect on the morphology, size and volume fraction of primary α phase, which affects the strength and ductility of the alloy. At lower solution temperature 746-770 ℃, there are two kinds of primary phases in the alloy: one is the globular αp phase; the other one is submicron αr phase. Nano-scale αs phase precipitated in β matrix during the subsequent higher temperature ageing. These three kinds of α phase constitute the feature of hierarchical structure. After solution at relatively higher temperature, αr phase disappeared, and globular αp phase existed. After ageing at higher temperature, the alloy exhibited bimodal microstructure. When the solution temperature exceeded β transition temperature, no αp phase was observed. In the subsequent higher temperature ageing, αs nano-precipitate dispersed in β matrix.

In this work, the hierarchical structure has a superior combination of strength and plasticity, mainly due to the following reasons: 1) The αs/β interface could block the movement of dislocation [31], which is the main contribution for high strength; 2) The hierarchical structure has high volume fractions of αp and αr phase, which produce strain hardening compatible with transformed β matrix to maintain uniform deformation; 3) The soft αp phase and fine αs phase endow the alloy high strength and good ductility. However, the microstructure composed of different sizes and morphologies are plastically non-homogeneous to some extent [32]. For the bimodal titanium alloy, plastic deformation is initiated in soft αp phase, resulting in higher plastic strain than global tensile strain [33]. During further deformation, strain incompatibility between αp and transformed β increases, which in turn decreases the plasticity. The hierarchal distribution of α phase causes more homogeneous strain partitioning and improves the plasticity [14].

The solution temperature could also affect the width and spacing of αs phase. The solution temperature affects the volume fraction of primary α phase, the distribution of elements and the stability of β phase [34-36]. Different elements tend to be concentrated in different phase, for example, β phase is rich with β-stabilising elements, such as V, Mo and Cr, while α phase is rich with α-stabiliser Al [35]. When the alloy is solution treated in α+β phase region, alloying elements are essential for phase formation and microstructure formation in β-Ti alloy, and their diffusion determines the composition and stability of α and β phase in the alloy [36-37]. This means that with the increase of solution temperature, the solute concentration of β stabilizers retained in matrix decreases, and so does the phase stability of the residual β phase as a result. The difference of β-stability exerts a remarkable influence on the precipitation of αs, which, in turn, results in an increment of driving force for α phase nucleation during ageing [31, 38-39]. This is the reason why at the same ageing temperature, αs phase becomes finer and the spacing becomes smaller with the increase of solution temperature.

Since the αs/β interface strengthening is the primary strengthening mechanism in metastable β titanium alloy, the spacing of αs (λ) determines the distance that dislocation could slide freely, which in turn determines the dislocation accumulation at αs/β interface. This strengthening mechanism is similar to fine grain strengthening [12, 33]. Therefore, thinner αs phase and smaller αs phase spacing mean shorter distance at which the dislocation could slip freely, which increases strength but decreases ductility with the increase of solution temperature at the same ageing conditions.

4.2 Influence of ageing on microstructure

Ageing process could tailor the size and spacing of αs phase which affects strength and ductility of β-Ti alloy [30, 40]. The primary α phase with different morphologies and volume fractions was obtained by different solution treatments followed ageing at 450-600 ℃. The size of αs and the grain boundary become coarser with increasing ageing temperature, as presented in Figure 7.

Previous studies showed that the transition driving force of αs phase from β phase is insufficient at lower ageing temperature [41]. The isothermal ω phase (ωiso) could be possibly transition phase during ageing at lower temperature. The appearance of nanometre ωiso would significantly increase strength but decrease ductility [42]. Ageing at 450 ℃, the microhardness of alloy is considerably higher than that at other ageing conditions, which means that ωiso phase might precipitate in the alloy during lower temperature ageing. In addition, as shown in Figure 4, with the increase of solution temperature, the volume fraction of primary α phase decreases, which means that more metastable phases could be obtained after quenching, and more ωiso phase formed in the subsequent 450 ℃ ageing treatment. Therefore, the sample aged at 450 ℃ after gradient solution treated showed step-like increase in hardness.

When solution treated alloy was subjected to ageing at higher temperature 550 ℃ and 600 ℃, β phase could directly transform to αs phase due to the sufficient driving force of phase transformation [43]. Thus, samples exhibited larger-size αs phase with lower volume fraction after ageing at higher temperature, and with the increase of ageing temperature αs phase coarsened. This also resulted in the decrease of strength and improvement of ductility, especially solution treatment at 746 ℃ for 2 h following 600 ℃ ageing for 8 h. In this case, αs is much thicker and slightly shorter, making the superior combination of strength and ductility.

Primary α phase also coarsens slightly with increasing ageing temperature, which means that the primary α phase grows up gradually with atomic long-distance diffusion during ageing at higher temperature [9, 31]. The slightly coarsening of primary α phase almost has no obvious effect on the strength, but is beneficial to the improvement of ductility.

5 Conclusions

In order to rapidly optimize the microstructure of high strength and high ductility titanium alloy, the gradient heat treatment method was used. A good combination between strength and ductility was achieved, the conclusions are as followings.

1) The volume fraction and morphologies of the αp phase were sensitive to solution temperature, with the volume fraction decreasing from 30% at 746 ℃ to 3% at 819 ℃. Solution treated at 746 ℃ and 770 ℃, rod-shaped primary α phase with a

width of 0.1-0.6 µm and length of 0.3-4 µm appeared. At 795 ℃ and 819 ℃, only the global primary α phase was observed.

2) At the same ageing conditions, the width and spacing of αs phase decreased with the increase of solution temperature, resulting in the increase of strength and decrease of ductility. While the increase of the width and spacing of αs phase results in the improvement of ductility.

3) Samples solution treated at 746 ℃ for 2 h and aged at 600 ℃ for 8 h otained the hierarchical structure composed of the primary α phase (αp), sub-micro α-rods (αr), and αs phase, achieving good combination of 1140 MPa yield strength and 15% elongation. Improved ductility resulted from effective strain partition and compatibility inside the hierarchical α structure.