ω-Assisted refinement of α phase and its effect on the tensile properties of a near β titanium alloy

Journal Pre-proof ␻-assisted refinement of ␣ phase and its effect on the tensile properties of a near ␤ titanium alloy Ruifeng Dong, Jinshan Li, Hongchao Kou, Jiangkun Fan, Yuhong Zhao, Hua Hou, Li Wu

PII:

S1005-0302(20)30055-4

DOI:

https://doi.org/10.1016/j.jmst.2019.10.031

Reference:

JMST 1922

To appear in:

Journal of Materials Science & Technology

Received Date:

30 July 2019

Revised Date:

18 September 2019

Accepted Date:

14 October 2019

Please cite this article as: Dong R, Li J, Kou H, Fan J, Zhao Y, Hou H, Wu L, ␻-assisted refinement of ␣ phase and its effect on the tensile properties of a near ␤ titanium alloy, Journal of Materials Science and amp; Technology (2020), doi: https://doi.org/10.1016/j.jmst.2019.10.031

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Research Article

ω-assisted refinement of α phase and its effect on the tensile properties of a near β titanium alloy Ruifeng Dong1,2, Jinshan Li2, Hongchao Kou2,* , Jiangkun Fan2, Yuhong Zhao1,*, Hua Hou1, Li Wu1

College of Materials Science and Engineering, North University of China, Taiyuan

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1

030051, China 2

State Key Laboratory of Solidification Processing, Northwestern Polytechnical

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University, Xi’an 710072, China

Corresponding authors.

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E-mail addresses: [email protected] (Hongchao Kou); [email protected]

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(Yuhong Zhao).

[Received 30 July 2019; Received in revised form 18 September 2019; Accepted 14

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October 2019]

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Abstract

In this work, the phase transformation sequence during the continuous heating process (3 oC/min) was investigated in a near β titanium alloy. The results show that the staring formation of ω phase is about 267 oC, and the ending precipitation temperature about 386 oC during the heating process. When the heating temperature is greater than 485 oC, there are no ω phase detected within the β matrix. Combined with the microstructural characterization, it is found that ω phase facilitates the nucleation of α phase nearby the 1

ω/β interface and has a great effect on the refinement for α phase. As compared with the specimens directly aged, the specimens with ω-assisted refinement of α phase possess high tensile strength, but there is no yield stage detected on their stress-strain curve. Combined with the analyses of the fracture morphology, the specimens with ω-assisted refinement of α phase present a brittle fracture. This is mainly ascribed to its relatively lager width of grain boundaries and the absence of widmanstätten α precipitates.

nucleation; α precipitation; Tensile properties

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1. Introduction

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Key words: Near β titanium alloy; Continuous heating; Phase transformation; ω-assisted

Near β titanium alloys are widely used in the aircraft and aerospace industries due

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to their high specific strength, low density, good fatigue and crack resistance for manufacturing land gears and fasteners[1-4]. As compared with traditional (α+β)

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titanium alloys (e.g. Ti-6Al-4V, wt%), near β titanium alloys have great capability to obtain the high strength and balanced by a considerable ductility. Their high strengths are mainly ascribe to the precipitation of α phase, including its size, volume faction,

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morphology and distribution[4-7]. Thus the adjustment for α precipitation is always a hot issue for exploring the optimum balance of the mechanical properties for near β

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titanium alloys[8-10]. The thermal mechanical process and the heat treatment are usually applied to realize the α phase control. Chen et al.[11] have studied the effect of

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deformation reduction on the mechanical properties of Ti-15Mo-3Al-2.7Nb-0.2Si (wt%) alloy and found that by 80% rolling reduction the alloy could obtained a higher tensile strength companied by a better ductility as compared with 20% reduction. By related to microstructural characteristics, they ascribe this good combination of mechanical properties to the refined grain size and the fragment of the α precipitation on grain boundaries. Different heat treatments have a great influence on the morphology and size of α precipitation. Dong et al.[4] have studied the correlation between the volume 2

fraction and the size of α precipitates and the mechanical properties in a Ti-7Mo-3Al-3Cr-3Nb (wt%) alloy, and indicated that the high volume fraction and fine size contribute to improve the strength of the alloy. In conclusion, refining α precipitates by thermal mechanical process and heat treatment is an effective way to improve the mechanical properties of the alloys. It is well known that the precipitation of α phase traditionally originates from parent β phase by β→α transformation. Thus all of factors that can affect the phase transformation are the breakthrough points to refine the α precipitates. Owing to the lower stability of β phase in near β titanium alloys, some

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transient phases could firstly generate under some special conditions. For example, as for near β titanium alloys, there are ω phase firstly precipitated at lower aging

temperature or at the early stage of the aging process[12-14]. Previous research has

indicated that ω phase formed in the early stage has a great effect on the precipitation of

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α phase. Although previous work have confirmed that ω phase as an brittle phase deteriorates the mechanical properties, the effect of ω-assisted refinement of α

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precipitates on the mechanical properties still remains limited.

In this work, a near β titanium alloy, Ti-7333 (Ti-7Mo-3Nb-3Cr-3Al, wt%), was

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investigated. Firstly, thermal dilatometer was conducted on the β-quenched Ti-7333 alloy with a slow heating rate to investigate the phase transformation sequence during

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the heating process. According to the results of thermal expansion test, phase transformation sequence during the heating process was investigated, including the formation of ω phase, ω-assisted precipitation of α phase and the refinement of α phase.

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The mechanical properties of the alloy with the ω-assisted refinement of α phase were investigated, and for the sake of contrast the mechanical properties of the alloys directly

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aged were investigated as well. Eventually, based on the microstructure characteristics and the fracture morphology, their fracture behaviors are discussed. 2. Experimental procedure The alloy studied in this work was the hot-rolled Ti-7333 bar with the nominal composition of Ti-7Mo-3Nb-3Cr-3Al, wt%. The Ti-7333 ingot was firstly prepared by vacuum self-consumable arc melting, and forged in β and (α+β) fields to the bars with 3

the diameter of 27 mm. Then the cross-rolling was conducted on the forged bars. Before the rolling, the forged bars were heated to 820 oC and held for 60 min. The bars were rotated by 90o around the rolling direction (RD) after every pass rolling over. Eventually, the bars with the diameter of 18 mm were obtained by six passes followed by air cooling (AC). The β transus temperature of the alloy is about 850 oC, determined by the metallographic method. The rolled-bar was solution-treated at 900 oC for 15 min followed by water quenching (WQ) to obtain the microstructure with the single β phase, and the X-ray

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diffraction (XRD) was carried out to confirm it. The specimens with 5 mm in diameter and 25 mm in height were cut out of the β-quenched alloy. Thermal dilatometer

(Netzsch® DIL-402C) was firstly conducted on the β-quenched specimens at a heating rate of 3 oC/min to investigate the phase transformation sequence during the heating

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process from room temperature to the different heating temperatures. Correspondingly, the specimens with the size of 8 mm × 8 mm × 8 mm were cut from the β-quenched

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alloy as well. On the basis of the dilatometry curve obtained, the specimens were continuously heated (3 oC/min) to different temperatures (300 oC, 350 oC, 450 oC, 600 o

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C) followed by WQ to retain the transformed microstructures. Then the specimens

directly aged at 600 oC for 192 min were prepared as well, which corresponds to the

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specimens continuously heated to 600 oC. For microstructural characterization, scanning electron microscopy (SEM, MIRA3 XMU) and transmission electron microscopy (TEM, FEI Tecnai G2 F30) were used. As for SEM observations, specimens were firstly ground

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and polished according to the standard metallographic methods, then followed by electropolishing with a solution of 5% perchloric acid in alcohol using a voltage of 35 V

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for 60 s at lower than 20 oC. Kroll’s etching solution (10% HF + 10% HNO3 + 80% H2O, vol.%) was used to reveal the microstructure. With regard to TEM observation, thin foils with about 60 μm were mechanically prepared, then followed by twin-jet electropolishing in a solution of 5% perchloric acid in 95% ethanol at -25 oC. TEM was operated at 300 kV. The size of β grain and the morphological parameters of α precipitates were statistically analyzed by using Photoshop and Image-Pro Plus software. 4

The tensile specimens with the gague part of 6 mm in width, 3 mm in thickness and 25 mm in length were machined along the RD after they were subjected to the above heat treatments. The tensile tests were carried out on the INSTRON-5581 with a strain rate of 10-3 s-1 at room temperature on the basis of GB/T 228-2002[15]. Two tensile specimens were prepared for each condition, and their corresponding stress-strain curves were compared. 3. Results and discussion

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3.1. Phase transformation sequence during the heating process Fig. 1(a) gives the microstructure of the as-solution treated specimen. It consists of equiaxed β grains with an average grain size of about 190 μm. Its corresponding XRD

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pattern, as shown in Fig. 1(b), indicates that only β phase is detected after the β-solution treatment. Differentiated from other near β titanium alloys, such as

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Ti-5Mo-5Al-5V-3Cr[13, 16] and Ti-10V-2Fe-3Al[17], there is no athermal ω phase formed within the β matrix after WQ from the single β phase field owing to its

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relatively high β phase stability with the molybdenum equivalent of about 9.64. The dialtometry curve of the β-quenched Ti-7333 is shown in Fig. 2(a). The heating rate was 3 oC/min. It can be seen that there are three temperature turning points

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on the curve. Fig. 2(b) is the fraction of ω phase as a function of heating temperature, which corresponds to the result of Fig. 2(a). Thus it can be determined that temperature

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A (TA, 267 oC) stands for the point at which isothermal ω (ωiso) phase begins to form and TB (386 oC) indicates that α phase precipitates, while TC (485 oC) suggests that ω

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phase is about to consume completely. Thus the phase transformation sequence under this condition can be concluded. When the heating temperature is less than TA, only single β matrix is retained. When greater than TA and less than TB, ωiso phase gradually generates by the phase transformation of β → β + ωiso. When greater than TB and less than TC, α precipitates gradually form by the β + ωiso → β + α transformation. With the increase of heating temperature, ωiso phase is completely consumed until the temperature is greater than Tc owing to its thermal stability[13,14]. Based on the results 5

of Fig. 2, the corresponding phase transformation and microstructural evolution during the heating process would be specifically investigated as following by interrupted tests. 3.2. ω precipitation and ω-assisted nucleation for α phase The β-quenched specimens are continuously heated to 300 oC and 350 oC, respectively, at 3 oC/min from room temperature followed by WQ. According to the results of Fig. 2(a), this quenched temperature is located in the range between TA an TB in which the transformation of β → β + ωiso occurs, i.e. the precipitation of ωiso phase.

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Fig. 3(a) shows the bright-field image of the specimens water-quenched at 300 oC, and there are no obvious precipitates detected within the matrix. Observed from the

corresponding SAED pattern given in inset of Fig. 3(a), one can detect very weak

diffraction spots located at the positions of 1/3 {112}β and 2/3 {112}β, as marked by the

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dotted circle in inset. These spots originate from one of ω variants (as marked by dotted

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circle in inset), and the crystallographic orientation relationship between ω phase and β matrix can be determined: {0001}//{111}β, <11-20>//<110>β. Fig. 3(b) is its dark-filed image and ω phase presents a low density distribution which corresponds to its weak

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diffraction intensity. In contrast, when heating temperature reaches up to 350 oC, the strong diffraction intensity and relatively higher density distribution of ω phase can be

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observed, as shown in Fig. 3(c) and (d). Meanwhile, judged from the SAED pattern shown in inset of Fig. 3(c), one can determine that at least two of ω variants generate

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within the matrix. It should be mentioned here that four ω variants could be formed in one parent β grain without the variant selection, and only two ω variants can be

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distinguished in one of <110>β zone axis[18]. These nano-sized ω precipitates present an ellipsoidal morphology, and such morphology feature is typical for the near β titanium alloys with low ω/β misfits, such as in Ti-6.8Mo-4.5Fe-1.5Al[18], Ti-5Mo-6Cr-5V-4Al[19], Ti-5Mo-5Al-5V-3Al[20]. Fig. 4 gives the TEM observations of specimen after quenching at 400 oC. This quenched temperature situates in the temperature range between TB and TC in which the transformation of β + ωiso → β + α occurs, i.e. the precipitation of α phase. The SAED 6

patterns shown in Fig. 4(a) and (b) indicate that, besides the spots of the ω phase and the β matrix, there are additional spots located at the position of 1/2{112}β. These spots belong to α phase. Fig. 4(c) is the dark-field micrograph obtained by the reflection of ω phase and α phase. It can be confirmed that the needle-like precipitates are α phase (as indicated by the green arrows) and the ellipsoids are the existed ω phase (as indicates by cyan circles) according to their morphology characteristics[12, 18, 19, 21]. Note that the needle-like α phase develops nearby and is distanced from the ω/β interface. Fig. 4(d) is the HRTEM micrograph. According to its corresponding FFT diffractogram shown in

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Fig. 4(e), it can be seen that ω phase and α phase co-exist within the matrix. The IFFT shown in Fig. 4(f) obviously indicates that α phase and ω phase are disconnected by a layer of the distorted β matrix. This suggests that α phase prefers to precipitate at the vicinities of the ω/β interface. One of the reasons is that the lattice strain induced by

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β→ω transformation could contribute to the nucleation of α phase (This work is going

on). The other one is closely related to the local enrichment of α-stabilizer. The work of

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Azimzadeh et al. have found an enrichment of α-stabilizer Al content at some distance from the ω/β interface[22]. This enrichment could arise from the interaction between the

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stress field associated with the ω/β interface and the solute elements diffusion and eventually facilitates α precipitate nearby the interface[22].

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Fig. 5 shows the microstructure of the specimen after continuously heated to 600 oC. It can be seen, from Fig. 5(a), that the obvious intragranular α phase and grain boundary α phase (αGB) are densely distributed within the β grain. Fig. 5(b) is the

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corresponding TEM micrograph of the intragranular α precipitates with the fine-scaled size and the acicular morphology. Fig. 5(c) displays the αGB precipitates along the grain

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boundaries. The statistical results indicate that the average width of the intragranular α precipitates is about 84 nm and the average length about 1.2 μm, and the average width of GBs about 142 nm, as listed in Table 1. Thus the phase transformation sequence and the microstructure evolution during this continuous heating process can be illustrated by Fig. 6.

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3.3. Tensile properties The stress-strain curves of the β-quenched specimens (in black) and the specimens continuously heated to 600 oC (in red) are given in Fig. 7. The yield strength ofβ-quenched specimen is about 760 MPa, the ultimate strength about 870 MPa and the elongation about 20.7%. Fig. 8(a) and (b) shows the corresponding facture morphologies. It can be seen that the cross section of the specimen presents a typical necking characteristic and there are numerous dimples densely distributed on the facture

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surface, as shown in Fig. 8(a1) and (b). This means that a significant ductile fracture occurs in the β-quenched specimens, which corresponds to its high elongation of about 20.7%.

Based on the results of Sections 3.1 and 3.2, the specimens continuously heated to

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600 oC at heating rate of 3 oC/min could obtain the fine-scale α precipitate within the

matrix. Its corresponding stress-strain curve is completely different. There is no yield

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stage detected on the curve, while its ultimate strength could be prominently improved to 1235 MPa by almost 360 MPa higher than that of the β-quenched specimen. Its

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corresponding fracture surface is shown in Fig. 7(e) and (f). The cross section of the specimens is very flat, as shown in Fig. 7(e1). Besides a few number of dimples, one can see that a large number of cleavages faces are mainly distributed on the fracture surface.

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Note that the sizes of the cleavage face are close to the size of β grains, and some chevron patterns (as evidenced by the blue arrows) could be observed on the faces, as

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shown in Fig. 7(d). This suggests that the intergranular fracture occurs in the specimens, and the same results have been reported by Mantri et al. in Ti-15Mo-3Al-3Nb-0.2Si

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alloys containing the fine-scaled α precipitates[23]. For comparison, the β-quenched specimens are directly aged at 600 oC for 192 min same with the heating time of the continuously heated specimens from room temperature to 600 oC at 3 oC/min. Its stress-strain curve (in blue) is shown in Fig. 6 as well. It can be seen that the curve presents an initial elastic stage followed by a non-linear plastic stage similar with that of the β-quenched specimens. Its yield strength is about 914 MPa, the ultimate strength about 968 MPa and the elongation about 13.9%. Thus the aged specimens could possess 8

a relatively high strength and a considerable ductility as compared with the continuously heated specimens. Fig. 8(c) and (d) gives its fracture surface micrographs. It is composed of a large quality of dimples and a few number of cleavage faces (as indicated by the red arrows), as shown in Fig. 8(d). This means that the specimen presents mixed ductile and brittle features during the tensile test, and the ductile fracture is dominant. In near β titanium alloys, the β phase with body-center cubic structure contributes to the ductility of the alloys. This is attributed to its large quality in slip systems as

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compared to the α phase with hexagonal close-packed structure. Besides the dislocation slip, the twining and the stress-induced martensite transformation could be activated in

the β phase as well[24-26]. All of those deformation modes have a positive effect on its improving ductility. Thus as for the β-quenched specimens with single β phase, the

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excellent ductility can be obtained. When the α phase precipitates within β matrix under aging condition or under continuous heating condition, twining and the stress-induced

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martensite transformation are hardly observed for β phase except for the dislocation slip[4, 27]. This results in obvious increasing strength companied by a significant

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decrease of the ductility due to the dispersion strengthening of α precipitates. Fig. 9 gives the micrograph of the specimen directly aged at 600 oC for 192 min.

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As shown in Fig. 8(a), except for the formation of the intragranular α precipitate and αGB, widmanstätten α precipitates (αWGB) are densely distributed nearby the GBs. The corresponding detailed TEM micrograph is shown in Fig. 8(b). The morphology

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parameters of α precipitates under the continuous heating condition are listed in Table.1 as well. It can be seen that under those two conditions the volume fractions of α

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precipitates are nearly equal, meaning that the precipitation of α phase is saturated under these two conditions. Under the former condition, the sizes of α precipitates are smaller than those of the directly aged specimens. On the one hand, the smaller sizes are attributed to the ω-assisted refinement for α precipitates during the continuous heating process. As evidenced in Figs. 2 and 3, the ω phase first forms before the precipitation of α phase during the heating process. With the increase of the heating temperature, α phase begins to precipitate and prefers to nucleate at the vicinities of the ω/β interface. 9

Thus as compared with the directly aged specimen, the densely-distributed ω phase could effectively improve the nucleation sits for α precipitation in the continuously heated specimens. Under the condition with the same volume fraction of α precipitates, their sizes must be smaller than those of the directly aged specimens. On the other hand, the driving force for the nucleation of α phase decreases and the driving force for the growth increases with the increasing aging temperature. All of those factors eventually result in the smaller size in the continuously heated specimens. Combined with the results shown in Fig. 7, the finer α precipitates bring in the significant increasing in

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strength. This result is consistent with the work by Fan et al. in which they found that the smaller acicular α precipitates bring in a higher strengthening effect[27]. This is due to the fact that the finer α precipitates bring more α/β interfaces into the matrix which effectively hinder the dislocation motion and eventually improve the high strength.

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Although the continuously heated specimens could possess a high strength, its ductility

is far from satisfactory and there is no yield stage during the tensile test, as shown in Fig.

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7. This corresponds to its brittle fracture. The same result has been found by He et al. that the nano-scaled α precipitates could lead to the brittle fracture in a near β titanium

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alloy of Ti-5Al-5Mo-5V-1Cr-1Fe alloy[28] and by Mantri et al. in Ti-15Mo-3Al-3Nb-0.2Si alloy[23, 29]. However, as shown in Fig. 8(f), the fracture

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surface presents an obvious transgranular feature, indicating that its fracture is also closely related to the grain boundary characteristics. As for the continuously heated specimens, the average width of GBs is larger than that of the directly aged specimens

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and there are no αWGB precipitated nearby GBs. The relative wide GBs can be attributed to the formation of αGB which is closely related with the ω-assisted nucleation of α

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phase. It has been evidenced that both of ω formation and the ω-assisted α precipitation involve the compositional partitioning. Thus the rejected Al, as the stabilizer for α phase and the destabiliser for ω phase, from the ω phase nearby GBs prefers to segregates to GBs and contributed to the formation of αGB, eventually resulting in relatively wide GBs and the absence of αWGB. Its high strength can be attributed to fine-scaled α precipitates and its brittle fracture is closely associated with the relative wide GBs and the absence of αWGB as compared with the directly aged specimens, and their 10

corresponding crack propagation path can refer to the schematic diagram shown in Fig. 10. 4. Conclusions The phase transformation sequence during continuous heating process and the effect of ω-assisted refinement of α phase on mechanical properties were investigated in a near β titanium alloy. The following conclusions can be drawn: (1) The ω phase is firstly formed in β-quenched Ti-7333 at the early stage of the continuously heating process. The temperature for the starting precipitation of ω phase

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is determined to be about 267 oC, and the ending precipitation temperature about 386 oC. After the heating temperature is greater than 386 oC, α phase begins to precipitate and the ω phase gradually disappears when the temperature is greater than about 485 oC. (2) During the heating process, nano-sized ω phase with high density facilitates the

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nucleation for α phase nearby the ω/β interface and has a great influence on its refinement.

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(3) The specimens with ω-assisted refinement of α phase present high tensile strength,

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but hardly no ductility. Its fracture presents a typical intergranular fracture character, which is mainly attributed to the larger width of grain boundaries and the absence of the

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widmanstätten α precipitates.

Acknowledgments

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This work was supported financially by the National Natural Science Foundation of China (Nos. 51711530151, 51801156 and 51804279), the Major State Research

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Development Program of China (Nos. 2016YFB0701303 and 2016YFB0701305) and the Natural Science Basic Research Plan in Shaanxi Province of China (No. 2018JQ5035).

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Figure list:

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Fig. 1. (a) Micrograph of the specimen after the solution treatment at 900 oC for 15 min

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and (b) the corresponding XRD pattern.

Fig. 2. (a) Dilatometry curve of the β-quenched specimen continuously heated at the

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temperature.

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rate of 3 oC/min and (b) the corresponding fraction of ω phase as a function of heating

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Fig. 3. (a, c) Bright-field micrographs (the inset showing the selected-area electron diffraction (SAED) pattern in the [011]β zone axis) and (b, d) the corresponding

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dark-field micrographs of ω phase using the (0001)ω reflection indicated by the dotted circle in inset of specimens quenched at different heating temperatures: (a, b) 300 oC; (c,

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d) 350 oC.

Fig. 4. TEM observation of the specimen quenched at 400 oC: (a) the SAED pattern 14

with the [011]β zone axis; (b) the corresponding key diagram of Fig. 4(a); (c) the dark-field micrograph taken by using the reflection circled by green circle in Fig. 4(a); (d) the high-resolution TEM (HRTEM) showing the co-existence of ω phase and α phase within β matrix; the corresponding fast Fouriter transform (FFT) diffractogram (e)

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and Inverse FFT lattice image (f) of Fig. 4(d).

Fig. 5. Micrographs of specimens after quenching at 600 oC: (a) SEM micrograph

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showing the overall microstructure; (b, c) the detailed TEM micrographs showing the

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intragranular α phase and the grain boundary α phase.

Fig. 6. Schematic image illustrating the phase transformation sequence and the

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microstructure evolution during the continuous heating process.

Fig. 7. Tensile engineering stress-strain curves of the specimens under different thermal 15

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conditions at room temperature with the strain rate of 10-3 s-1.

Fig. 8. Fracture surface morphologies and the corresponding magnification micrographs of specimens under different thermal treatment conditions: (a, b) the as-solution treated

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specimens; (c, d) the directly aged specimens; (e, f) the continuously heated specimens.

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Fig. 9. TEM images of the specimen directly aged at 600 oC for 192 min showing the (a)

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intragranular α precipitates and (b) the grain boundary α precipitates.

Fig. 10. Schematic diagram showing the crack propagation path of (a) the continuously heated specimens and (b) the directly aged specimens.

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Table list Table 1 Morphological parameters of α precipitates and the average width of GBs of specimens under different thermal conditions. Heat treatment / Time

Average width of

Average length of

Volume fraction

Average width of

αIG (nm)

αIG (μm)

of α (%)

GBs (nm)

84 ± 8

1.2 ± 0.5

62.3 ±1.2

142.4 ± 5.6

282 ± 30

4.8 ± 0.3

61.5 ± 2.1

85.6 ± 7.8

Continuously heating / ~192 min Directly Aging

Jo

ur

na

lP

re

-p

ro of

/ 192 min

17