Effect of microstructure evolution on mechanical properties of a TiZrAlB alloy rolled by different processes

Materials Science & Engineering A 766 (2019) 138348

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Effect of microstructure evolution on mechanical properties of a TiZrAlB alloy rolled by different processes

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S.G. Liua, B.H. Chena, H.L. Luob, S. Yangc, X.Y. Zhanga, M.Z. Maa, R.P. Liua,∗ a

State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, China Central Iron & Steel Research Institute, Beijing, 100081, China c Department of Materials Physics and Chemistry, University of Science and Technology Beijing, Beijing, 100083, China b

A R T I C LE I N FO

A B S T R A C T

Keywords: Hot-rolled TiZrAlB alloys Mechanical properties Rolling technologies Microstructure evolution

For the development of structural titanium alloys with excellent mechanical properties, the effect of rolling temperature and cooling rate on phase composition, microstructure evolution, and mechanical properties in a full-α TiZrAlB alloy were investigated. Optical microscopy, scanning electron microscopy, transmission electron microscopy, and statistical analysis, were employed to further understand and control the processing-structureproperty relationships. The experimental results of specimens rolled at a temperature range from the α to β phase region demonstrated a significant grain refinement trend as the rolling temperature increased, and the work hardening effect was stronger in specimens rolled at the α+β double phase region than those rolled in the single phase region. The high density of the original β grain boundaries also enhanced the mechanical properties when the specimens were rolled at temperatures around/above the β-transus temperature. The highest strength (σ0.2 = 1121 MPa and σb = 1387 MPa) with a failure elongation of 6.1% was obtained in 840 °C-rolling waterquenched specimens, however, the furnace cooling led to worsened mechanical properties. In sum, the variation in the mechanical properties of the series of rolled alloys was mainly ascribed to the microstructure evolution under the different rolling processes.

1. Introduction Titanium (Ti) and its alloys exhibit an excellent combination of mechanical and physical properties in key applications in the aerospace, chemical, and medical industries [1–4]. Both the performance and service capabilities of structural Ti alloys are important criteria for their practical application. It is worth noting that zirconium (Zr) has similar physicochemical properties with Ti, since they belong to the same IVB group, and they can form infinite α phase solid-solutions (hexagonal close-packed structure) at low temperature or β phase (body-centered cubic structure) at high temperature. In addition, Zr can improve the microstructure and enhance the strength of Ti alloys [5]. Therefore, a series of Ti–Zr based alloys with excellent-performance, such as Ti-Zr [6,7], Ti-Zr-Al [8,9], Ti-Zr-Nb [10], Ti–Zr–Y [11], Ti-ZrMo-Sn [12], and our novel Ti–Zr–Al–B [13,14], have been investigated in recent years. These Ti–Zr based alloys have not only improved and reinforced the traditional Ti alloys, but have also broadened the application area. However, the development of such alloys usually stops at the research stage, due to the high processing costs accompanied by inefficient processing of some new techniques for Ti alloys. The rolling

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process, as one of the simplest and efficient manufacturing technologies, is suitable for industrialization, and has tremendous application potential for Ti alloys. In previous studies, mechanical processing has been a frequent way used to fabricate high-performance structural materials [15–18], in which microstructure characteristics, such as grain size, shape, distribution, and orientation can be well controlled. For example, Zhang et al. [15] reported that β-metastable TiZr alloys rolled at ambient temperature with different accumulative strains demonstrated high elastic admissible strain, which could be attributed to the nanoscale low-elastic modulus α″ lamellae with high density phase boundaries produced by cold-rolling. In addition, surface mechanical attrition treatment was used to process the grain-size gradient-structured interstitial free steel, where the gradient structures converted the applied uniaxial stress into multiaxial stresses, due to the evolution of incompatible deformation along the gradient depth, leading to the high ductility in interstitial-free steel [19]. In particular, thermomechanical processing, such as hot rolling, hot extrusion, and hot forging, can further affect the phase transformation, recovery, and recrystallization of materials [19,20]. It has been demonstrated that precipitated TiB

Corresponding author. E-mail address: [email protected] (R.P. Liu).

https://doi.org/10.1016/j.msea.2019.138348 Received 11 July 2019; Received in revised form 26 August 2019; Accepted 27 August 2019 Available online 28 August 2019 0921-5093/ © 2019 Elsevier B.V. All rights reserved.

Materials Science & Engineering A 766 (2019) 138348

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particles, induced by hot-rolling, can enhance the strength of B-containing Ti alloys [20]. Moreover, dynamic recrystallization behavior has been advantageously realized through thermomechanical treatment to tailor the microstructure and mechanical properties of Ti alloys [21]. In several previous studies [9,16,20,22], the thermomechanical processing temperature of Ti–Zr based alloys was chosen in single α or β and/or α+β phase regions, which revealed a relatively monotonous microstructure evolution. In general, only few reports on the hot rolling temperature in all phase regions of Ti–Zr based alloys have been documented. Moreover, the type of cooling after hot rolling is another important factor affecting the grain growth, microstructure, and mechanical properties of Ti alloys. Feng et al. [23] discussed the relationship between cooling rate and wear resistance in Ti–6Al–4V, which exemplified a positive correlation between cooling rate and surface hardness and a negative correlation between cooling rate and wear resistance. This was attributed to that the wear mechanism changed from plastic deformation to a more brittle-fracture behavior of the surface. Therefore, the effect of cooling rate on the microstructure and mechanical properties of studied alloys has been simultaneously taken into consideration. In order to obtain excellent mechanical properties with optimized microstructure in TiZrAlB alloys, in this study, the hot-rolling deformation under various rolling temperatures with water quenching was employed, and different cooling processes after rolling at 840 °C were adopted. The microstructural evolution and mechanical properties of the rolled alloys were investigated in depth.

Fig. 2. Process flow diagram.

processes was designed according to the DSC curves. Thus, the strip samples were heated to 720, 760, 800, 840, and 880 °C for 30 min (heating rate: 10 °C/min). At the final 65% deformation, the strip samples were rolled into 7 mm thick plates via multi-pass rolling (2 mm per pass and strain rate of 0.6 s−1). The samples rolled at 720, 760, 800, and 880 °C were immediately quenched in water, and the other samples rolled at 840 °C were cooled through water quenching (WQ), air cooling (AC), and furnace cooling (FC) to room temperature. The series hotrolled (HR) experimental samples were denoted as HR720, HR760, HR800, HR840, HR880, HR840AC, and HR840FC as defined in Table 1. Finally, 1-mm-thick portions (thicker than the oxide skin) were removed from both sides of the rolled surface, and all the test specimens were taken along the rolling direction (RD) for the following tests. The phase composition and crystal structure of the series hot-rolled T40ZAB alloy specimens were confirmed by conventional X-ray diffraction (XRD) with Cu Kα radiation (D/max-2500/PC), and the diffraction angle ranged from 20° to 100° (step size: 0.02°). An image of the bone-shaped tensile specimen dimensions is shown in the top-right corner of Fig. 6. The uniaxial tensile tests were conducted at room temperature at a strain rate of 5 × 10-4 s−1 using an Instron 5892 Universal Material Testing Machine. The microhardness of polished specimens was analyzed by a Vickers microhardness tester at a load of 150 g for 10 s. Metallography and fractography of the rolled specimens were carried out using optical microscopy (OM) and scanning electron microscopy (SEM), and the high-magnification microstructure was observed using transmission electron microscopy (TEM). The metallographic specimens were mechanically polished to a final level of 0.5 μm diamond plaster and subsequently etched in a solution of 5% hydrofluoric acid, 15% nitric acid, and 80% deionized water. The TEM specimens were prepared by the twin-jet electrochemical polishing technique within a solution containing 10% perchloric acid and 90% methanol at 13 V and at −35 °C.

2. Experimental procedure A 5-kg ingot was prepared using an electromagnetic induction furnace with a water-cooled copper crucible. The raw materials were sponge Ti (99.9 wt%) mixed with sponge Zr (Zr + Hf ≥ 99.7 wt%), commercial-purity Al (99.9 wt%), and B powder (99.9 wt%). The nominal composition of the experimental alloy ingot was Ti–40Zr–4Al0.005B (abbreviated as T40ZAB) in wt.%. For compositional homogeneity, the experimental ingot was flipped and remelted thrice, and then homogenized at 1000 °C (above the β-transus temperature) for 12 h (furnace cooling). The strip samples (20 mm × 20 mm × 150 mm) were cut from the as-cast ingot by electro-discharge machining (EDM) for thermo-mechanical processing. Before hot-rolling, differential scanning calorimetry (DSC) was utilized to detect the phase transition temperature of the experimental alloy at a heating rate of 10 °C/min in an argon-protected environment using a corundum crucible. The result is shown in Fig. 1. As it can be seen in Fig. 2, the flow chart of the hot rolling

Table 1 Different rolling technologies and corresponding abbreviations as used throughout. Rolling technology 720 °C, 760 °C, 800 °C, 840 °C, 880 °C, 840 °C, 840 °C,

Fig. 1. The DSC curves of the T40ZAB alloy. 2

30 min, 30 min, 30 min, 30 min, 30 min, 30 min, 30 min,

rolling, rolling, rolling, rolling, rolling, rolling, rolling,

Corresponding nomenclature water quenching water quenching water quenching water quenching water quenching air cooling furnace cooling

HR720 HR760 HR800 HR840 HR880 HR840AC HR840FC

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Fig. 3. The referenced diffraction peak positions of each common phase in Ti alloys, and XRD patterns of the HR-T40ZAB alloys.

under rolling around the α transus temperature at 760 °C. For HR760, during the heat preservation process, a part of primary α phase transformed into β phase, while in the subsequent cooling process, the β phase entirely decomposed into α′ martensite phase. Afterwards, as the rolling temperature increased, the microstructure evolution exhibited similar characteristics, where the metallographic morphology was nearly filled with interlaced acicular α′ martensite phase (basket-weave microstructure), and only the density of the original β grain boundaries demonstrated a decreasing trend (Fig. 4(c–e)). The TEM bright field micrographs in Fig. 5 show the microstructural details of the developed alloys, and selected area electron diffraction patterns (SADP) of HR880 are shown in Fig. 5(e). The coarse primary α lamellae carried the entire deformation and produced a large number of dislocations (Fig. 5(a)). When rolling at 760 °C, the microstructure (Fig. 5(b)) consisted of break-up primary α lamellae interminglint with secondary acicular α′ martensite. As the rolling temperature further increased (at 800, 840, and 880 °C), a monotonous microstructure with acicular α′ lamellae and nearly dislocation-free, as well as obvious grain refinement could be observed (Fig. 5(c–e)). In addition, the SADP shown in Fig. 5(e), jointly with the XRD patterns, confirmed the phase composition. Fig. 6(a) and (b) show the typical engineering stress-strain curves until fracture and the true stress-strain curves until necking of the HRT40ZAB alloys under different conditions. Three specimens were tested in each condition in order to ensure repeatability of the material properties, however only one typical curve is presented. Meanwhile, Fig. 6(c) visualizes the variation trend of mechanical properties of the developed alloys and Table 2 lists the mechanical property data. As the rolling temperature increased until it reached 880 °C, the engineering ultimate tensile strength (UTS) (σb), true UTS (σm), and microhardness (HV) were all enhanced with concomitant failure elongation (εf) loss. However, when the rolling temperature increased from 720 to 760 °C, the yield strength (σ0.2) initially decreased, and then it increased until the temperature reached 880 °C. Among the current alloys, HR720 exhibited the highest εf of 11.3% at the expense of the lowest σb, σm, and microhardness. When rolling at 840 °C, HR840 exhibited the highest σ0.2 of 1121 MPa, σb of 1387 MPa, σm of 1443 MPa, microhardness of 448.1 HV, and still retained an εf of 6.1%. On the whole, there was no significant difference between the respective mechanical properties of HR840 and HR880. As it is clearly manifested in Table 2 and Fig. 6(c), the mechanical properties of series HR-T40ZAB alloys exhibited a nonuniform trend caused by the various microstructures, which would be separately discussed. The strain-hardening rate (Θ = dσ/dε) versus the logarithmic strain of the examined alloys is shown in Fig. 7, and the fitted curves are simultaneously presented. In specimen HR720, a drastic decrease in work hardening was observed at a logarithmic strain between 1 and 3%. When the rolling temperature increased, the strain-hardening rate mainly tended to increase before necking, although that of HR840 was slightly higher than that of HR880. A comparison between the

3. Results and discussion 3.1. Rolling at different temperatures and water quenching The general phase-transition temperature range of the T40ZAB alloy was acquired from the DSC curves (Fig. 1). Owing to the discrepancy of onset/end transition temperature in the heating and cooling processes, which resulted from the overheating of the α→β phase transition, the undercooling of the β→α phase transition, and the hysteresis effect in the DSC test [24], the α phase in T40ZAB was stable below 738 °C and started to transform into β phase when the temperature ranged between 738 and 781 °C. At high temperatures over 856 °C, only β phase exists. The phase composition information of HR-T40ZAB alloys is presented in Fig. 3. Fig. 3(a) lists the peak positions in the XRD patterns of various common phases in pure Ti according to previous studies [22,25]. The XRD patterns (Fig. 3(b)) of the current alloys, under different rolling temperatures, reveal that the diffraction peak positions only matched with α/α′ phases (hexagonal close-packed structure) without other phases, such as β phase or intermetallic phase, which are corresponding to the standard diffraction peaks of α-Ti in Fig. 3(a). Because Ti and Zr are neutral elements to each other, Al is an α-stabilizing element for Ti and no β-stabilizing element exists. The same result about the phase transformation of single-phase α-Ti alloys has been thoroughly reported in our previous work [14] and the study of Jiang et al. [8]. Therefore, the phase transition of specimens HR760 and HR800 simply underwent the α→(α+β)→(α+α′) process, and analogously, the α→(α+β)→β→ α′ phase transition process occurred in specimens HR840 and HR880. That is, for the studied alloys, no β phase, ω phase, and α′′ phase were retained after quenching from temperatures above the α transus temperature, but only α phase or/and α′ martensitic phase. Fig. 3(c) shows Fig. 3(b) in detail with 2θ from 37° ~ 40° and 71° ~ 77°. As the rolling temperature increased, the peak positions showed no shift, but a broad trend appeared on the diffraction peaks. These results indicate that the grains were increasingly refined in the condition of identical composition, which was caused by the martensitic transformation of HRT40ZAB alloys during water quenching, and concomitantly with the formation of closely spaced interfaces between martensitic lamellae [26]. The microstructure evolution of series HR-T40ZAB alloys is illustrated in Figs. 4 and 5. As it can be observed in both Figures, the microstructure was gradually refined as the rolling temperature increased, which was also confirmed by the XRD results. After rolling at 720 °C, the lamella α phase was severely distorted or broken (Fig. 4(a)). As it has been previously reported by Weiss et al. [27], this can be attributed to dual mechanisms. One is that both low- and high-angle boundaries form across the α lamella, and another is that localized shear and rotation of the α lamella occur with misorientation. In Fig. 4(b), it can be observed that the duplex microstructure of coarse primary α phase coexisting with fine α′ martensite phase appeared

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Fig. 4. Optical micrographs of the HR-T40ZAB alloys with various rolling temperatures: (a) 720 °C, (b) 760 °C, (c) 800 °C, (d) 840 °Cand (e) 880 °C.

the above results.

mechanical properties of the specimens rolled at 720 and 760 °C revealed that σ0.2 and both σb and σm had an opposite variation tendency. Δσ defined as the difference between σm and σ0.2 also reached 273 MPa in HR760, whereas it was only 109 MPa in HR720. Furthermore, according to Figs. 6(b) and Fig. 7, HR760 derived its higher UTS than that of HR720 from the high work-hardening rate, which was remarkable for the α+α′ phase microstructure.

3.3. Fracture surface morphologies The SEM fracture morphologies of HR-T40ZAB alloys with different processes are gathered in Fig. 9. The variation of fracture surfaces of post-rift tensile specimens depicted typical trends. The fracture morphology of HR720 alloy specimen consisted mainly of tiny and equiaxed dimples within large and deep dimples (Fig. 9(a)). In addition, the necking phenomenon is obvious. Therefore, HR720 exhibits excellent tensile ductility. As rolling temperature increased to 760 °C, the dimples shrunk, and the micro-pores generated by defects produced metastable dimples during the tensile loading (Fig. 9(b)). The increased rolling temperature reduced and shallowed the formed dimples with continuous rippled pattern. In regard to the fine-grained HR800, HR840 and HR880 alloys with approximate ductility, they showed the similar fracture morphology of shallows (Fig. 9(c–e)). Fig. 9(f) and (g) show the fracture surface morphologies of HR840AC and HR840FC after tensile failure. A typical quasi-cleavage fracture morphology is observed in Fig. 9(f), where evenly distributed facets, cracks, and cleavage planes with little necking appeared over the entire surface. These characteristics are more pronounced in the fracture morphology of HR840FC (Fig. 9(g)). Meanwhile, it is worth mentioning that, after FC, the tensile ductility deteriorated dramatically (Figs. 6(a) and Fig. 8(d)). This morphology can be caused by intergranular fractures, which are cracks

3.2. Different cooling methods after rolling at 840 °C The specimens rolled at 840 °C with the highest strength were selected in order to tailor their microstructure and mechanical properties by different cooling methods. Fig. 8 shows the phase composition, microstructure and variation trend of mechanical properties of specimens cooled through AC and FC after 840 °C rolling. The β→α phase transformation occurred via shear mode under AC and via full diffusional mode under FC. There were still only α/α′ phases existing (Fig. 8(a)) under the slower cooling rate. The lamellar microstructure became gradually coarse with the decreased cooling rate (Fig. 8(b) and (c)). Fig. 8(d) shows the relationship between mechanical properties and the examined alloys. In conclusion, the AC-induced lamella growth and microstructure promote tensile ductility with strength loss, while the FC-induced lamella growth and microstructure completely worsen the mechanical properties in the developed alloys. The following discussion will qualitatively investigate the mechanisms responsible for 4

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Fig. 5. TEM bright field microstructures of the HR-T40ZAB alloys with various rolling temperatures: (a) 720 °C, (b) 760 °C, (c) 800 °C, (d) 840 °Cand (e) 880 °C.

α transus temperature), a large number of dislocations were generated by multi-pass rolling, i. e. the significant deformation. As rolling proceeds, these dislocations tangle and form dislocation walls and networks (Fig. 5(a)), which can even form sub-structures [29]. Moreover, under tensile loading, the high-density original dislocation pile-ups lie at the lamella boundaries leading to restricted dislocation mobility, further improving the yield strength [30]. When the rolling temperature raises up to 760 °C, only a small number of dislocations are reserved in the primary α lamellae, while some dislocations disappear due to the recovery and recrystallization process (Fig. 5(b)). As aforementioned, HR760 demonstrated a duplex microstructure of coarse primary α phase + acicular α′ martensite phase. It is generally acknowledged that the α′ martensite phase is harder than other common

that take place along the grain boundary when uniform deformation proceeds to a high extent during tensile loading. In particular, the coarse α lamellae produced by slow cooling from the β phase region are favorable to slip-length increase and crack propagation [28]. Consequently, the morphology of the fracture surface coincides with the deformation characteristic of the tensile test results of HR-T40ZAB alloys. 3.4. The effects of microstructures on the mechanical properties According to the above microstructure and tensile properties, the significantly higher σ0.2 and lower UTS of HR720 compared to those of HR760 can be explained as follows. During rolling at 720 °C (below the 5

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Fig. 6. Mechanical properties of the HR-T40ZAB alloys: (a) engineering stress-strain curves under different rolling processes and tensile specimen dimensions, (b) true stress-strain curves before necking, (c) the variation trend of mechanical properties of series HR-T40ZAB alloys.

phases in Ti alloys, such as the α, β, and ω phases [14,22]. Moreover, from a chemical perspective, alloying element Zr partitioning occurred during the heating process, when the rolling temperature was above the α transus temperature. As a high temperature phase, the β phase contains a high Zr content compared to the primary α phase in the α+β phase region, so, the secondary α′ martensite phase transformed from β phase is enriched in Zr after quenching. This leads to the strengthening effect of Zr-induced lattice distortion in the α′ martensite phase, which is slightly stronger than that in the primary α phase. For another, the α′ phase structure contains closer lamellar spacing than primary α phase, providing a high level of hardness [26]. Therefore, the soft primary α lamellae with a small number of dislocations in HR760 would first initiate plastic deformation. The prior-yielding α lamellae were constrained by surrounding hard α′ martensite lamellae, thus dislocations were piled-up and blocked at the lamellae interfaces [19,31]. When the adjacent hard α′ martensite phase initiated yielding at a higher strain, the dislocations could be able to progressively spread into the α′ martensite phase. Wu et al. [19] reported that the heterogeneous lamella structure in Ti produced by asymmetric rolling was characterized by soft micro-grained lamellae embedded in a hard ultrafine-grained lamella matrix, which exhibited higher strain hardening than coarsegrained metals. Similar results were also discussed by Feng et al. [16] concerning Zr–V alloy systems, where the heterogeneous structure including coarse initial α phase, ultrafine β phase, and acicular α′ phase significantly improved the UTS. Overall, in comparison with HR720, the duplex microstructure of HR760 was the main reason for its higher UTS (and higher work hardening rate) and the lower σ0.2. In HR760, HR800, and HR840, which had similar type of microstructure, the changes in mechanical properties were mainly attributed to the grain refinement. The statistical microstructural parameters of the HR-T40ZAB alloys with various rolling temperatures are shown in Fig. 10(a) and (b). The statistical lamellar thickness of the α′ grains showed that major thickness values were about 2, 0.11, and 0.09 μm for specimens rolled at 760, 800, and 840 °C, respectively. More specifically, the α′ grain size decreased as the rolling temperature increased. It can be roughly estimated that the main formation mechanism of α′ martensitic lamellae is analogous when specimens are quenched from

Table 2 T40ZAB alloys under different rolling temperatures and resulting yield strength σ0.2, engineering ultimate tensile strength (UTS) σb, true UTS σm, Δσ = σb - σ0.2, failure elongation εf, logarithmic uniform strain εu and microhardness HV. Alloys

σ0.2 (MPa)

σb (MPa)

σm (MPa)

Δσ (MPa)

εf (%)

εu (%)

HV

HR720 HR760 HR800 HR840 HR880

1040 993 995 1121 1094

1104 1216 1246 1387 1340

1149 1266 1295 1443 1390

109 273 300 323 295

11.3 8.8 6.6 6.1 6.2

3.9 4.0 3.8 4.0 3.7

413.6 427.0 428.0 448.1 443.4

Fig. 7. Strain hardening rate (Θ = dσ/dε) versus logarithmic strain of series HR-T40ZAB alloys.

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Fig. 8. Phases, microstructures and mechanical properties of 840 °C hot-rolled HR-T40ZAB alloys under different cooling methods: (a) XRD patterns, (b) and (c) combined OM-SEM images of HR840AC and HR840FC alloys, (d) variation trend of mechanical properties.

Fig. 9. SEM fracture morphologies of the HR-T40ZAB alloys: (a) HR720, (b) HR760, (c) HR800, (d) HR840, (e) HR880, (f) HR840AC and (g) HR840FC.

the α+β phase region (i. e. 760, 800, and 840 °C). The formation mechanism of α′ lamellae changes from interface instability nucleation to sympathetic nucleation in sequence for near α Ti alloy, α+β Ti alloy, near β Ti alloy, and metastable β Ti alloy with the increase of the β stability parameter, which here refers to the rolling temperature. That is to say, for specimens HR760, HR800, and HR840, the volume fraction of the residual primary α phase and dislocations produced by rolling were decreased with the increase of the rolling temperature. As nucleation site, the residual primary α phase and the dislocations accelerate the β→α transformation, due to that, atoms can easily diffuse through the stress field of these defects during phase separation [32,33]. Under the same cooling rate (WQ), the higher nucleation rate in HR760 compared to HR800 and HR840, led the prior growth of α′ lamellae producing the coarser martensitic microstructure. Figs. 6 and 10 indicate that the σ0.2, UTS, and microhardness increased with the gradual refinement of the lamellae. These results follow a Hall-Petch empirical relationship (σ = σ0 + kd−0.5) with dependence on the lamella thickness (d). In addition, some previous reports [34,35] have verified this Hall-Petch behavior in Ti alloys with similar lamellar microstructure. As abovementioned, the Hall-Petch relationship is significantly effective on the hot-rolled alloys with different cooling methods (HR840AC and HR840FC). Due to the cooling rate, great differences in the lamellae size were observed (Fig. 10(c)), further

affecting the mechanical properties. The lamella thickness of HR840, HR840AC, and HR840FC were mainly distributed at about 0.05–0.15 μm, 1–1.5 μm, and 10–20 μm, respectively. Fig. 10(d) shows a good linear relationship between σ0.2 and the lamella thickness. According to the study by Jiang [36], fine-grain strengthening is mainly described from two aspects. On the one hand, the effect of grain size on the irreversible energy to overcome the frictional resistance which is proportional to the volume swept by the interface. Fine-grained structure contains large interfacial area. A decrease in the grain size could result in an increase in the internal frictional resistances and the yield strength. On the other hand, increasing density of the grain boundaries caused by reduction of grain size in the same cross-sectional area increased the number of dislocations piling up. Furthermore, the different grain orientations in both sides of the grain boundary also promote the dislocation piling up at the grain boundary. Thus, in this work, the grain refinement resulted in an increase in the strength. According to the above discussion, the changes in the nucleation mechanism are also applied to HR840 and HR880. Moreover, the other important nucleation and growth processes that occurred at the original β grain boundaries would further affect the mechanical properties. For β-quenched specimens (840 °C is near the β-transus temperature), the nucleation of grain boundaries α (αGB) preferentially occurs at the original β grain boundaries, with the growth of αGB, Widmanstätten α 7

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Fig. 10. Grain size analysis of series HR-T40ZAB alloys: (a) statistical thickness distribution of the specimens rolled at different temperatures, (b) average thickness of α/α′ lamellae and partial enlarged detail of microstructures of the specimens rolled at different temperatures, (c) statistical thickness distribution of the specimens with different cooling methods, (d) σ0.2 as function of the lamellar thickness of the specimens with different cooling methods.

(αWGB) nucleate at protuberance of αGB or original β grain boundaries due to interface instability and then grow into original β grains, becoming lamellae [33,37]. Thus, contrary to HR880 (Fig. 4(e)), the specimens rolled at 840 °C with the high density of the original β grain boundaries (Fig. 4(d)) presented a high nucleation rate of α′ lamellae and further determined the slightly broader thickness of the α′ lamellae. However, the relationship between mechanical properties and grain size is abnormal. Here, the mechanical property changes of HR840 and HR880 were attributed to the reduced density of the original β grain boundaries (i. e. the original β grain size). Fig. 11 shows the inverse pole figure, the SEM morphology, and their overlay figure of HR880 at a representative region within an original β grain boundary. It can be clearly observed that the larger misorientation appeared between two adjacent original β grains (Fig. 11(a) and (c)), and the contiguous α lamellae nucleated from the original β grain boundary showed the same orientation (i. e. the αWGB). The misorientation was the transformation necessary to move from one crystal frame to another crystal frame, and the large misorientation between two sides of the original β grain boundary arisen from texture inheritance or variant selection occurring in the β→α transformation is common in Ti alloys. Bhattacharyya et al. [38] and Gey et al. [39] reported the relationship between crystallographic orientations and growth directions of grain boundary-allotriomorphic-α (αGB) and secondary Widmanstätten α lamellae growing from the αGB at grain boundaries separating β grains, revealing the microstructural characteristics at and adjacent to the grain boundaries. This specific interface (original β grain boundary) acts as barrier against dislocation motion [34], contributing to the stronger improvement of the mechanical properties than low-angle boundary between secondary α′ lamellae. Li et al. [11] also reported the positive effect of refined prior-β grain size on the compression properties in Ti–Zr–Y alloy systems. Therefore, in HR840, the enhancement of σ0.2, UTS, and microhardness with a small decrease in εf can be ascribed to the high density of the original β grain boundaries.

4. Conclusions The effect of rolling temperature and cooling method on microstructure evolution and mechanical properties of T40ZAB alloys has been investigated. The above-mentioned mechanisms of strengthening, nucleation and growth of α lamella were generalized to other Ti alloys. The following conclusions can be drawn from the investigations: (1) Only α/α′ phases were detected in series HR-T40ZAB alloys. (2) The lamellar structure gradually was refined resulted due to the increase in the nucleation rate with the evanescent dislocations as the rolling temperature increased, nevertheless, the prior-β grain size increased when the rolling temperature increased from 840 to 880 °C. Lamella coarsening occurred with the decreasing cooling rate (AC, FC). (3) The fracture surfaces revealed typical transitional processes from ductile to non-ductile as the rolling temperature increased. HR840AC and HR840FC showed quasi-cleavage fracture morphologies, and the characteristics of cleavage fracture morphology were more obvious as the cooling rate decreased. The transformation of fracture mode was in accordance with the tensile ductility of HRT40ZAB alloys series. (4) Both UTS and microhardness significantly increased as the rolling temperature increased from 720 to 840 °C. When the rolling temperature further increased to 880 °C, the strength slightly decreased with increased tensile ductility compared to the specimens rolled at 840 °C. In addition, the HR840 alloy revealed the highest σ0.2 (1121 MPa) and UTS (1387 MPa) at 6.1% elongation. (5) The work hardening of the specimens rolled at the α+β phase region induced by the heterogeneous lamella structure was stronger than that of specimens rolled at the single phase region. (6) The effect of the high-angle boundary (original β grain boundary) on hindering the dislocation motion was stronger than that of the 8

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Fig. 11. Inverse pole figure and its morphology overlay figure of 880 °C hot-rolled HR-T40ZAB alloy.

low-angle boundary between lamellae. (7) The εf of air-cooled specimens increased at the expense of strength loss compared to that of the water-quenched specimens, and the furnace-cooled method resulted in complete deterioration of the mechanical properties of the developed alloys.

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