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Influence of phase composition and microstructure on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys
Accepted Manuscript Influence of phase composition and microstructure on mechanical properties of hotrolled Ti-χZr-4Al-0.005B alloys S.G. Liu, C.Q. Xia, Z.H. Feng, X. Zhang, B.H. Chen, C.L. Xie, Y.K. Zhou, X.Y. Zhang, M.Z. Ma, R.P. Liu PII:
S0925-8388(18)31329-X
DOI:
10.1016/j.jallcom.2018.04.043
Reference:
JALCOM 45681
To appear in:
Journal of Alloys and Compounds
Received Date: 20 December 2017 Revised Date:
30 March 2018
Accepted Date: 3 April 2018
Please cite this article as: S.G. Liu, C.Q. Xia, Z.H. Feng, X. Zhang, B.H. Chen, C.L. Xie, Y.K. Zhou, X.Y. Zhang, M.Z. Ma, R.P. Liu, Influence of phase composition and microstructure on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys, Journal of Alloys and Compounds (2018), doi: 10.1016/ j.jallcom.2018.04.043. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT Influence of phase composition and microstructure on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys S.G. Liua, C.Q. Xiab, Z.H. Fengc , X. Zhanga, B.H. Chena, C.L. Xiea, Y.K. Zhoud, X.Y. Zhanga, M.Z.
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University,
Qinhuangdao 066004, China b
School of Materials Science and Engineering, Research Institute for Energy Equipment Materials,
c
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Hebei University of Technology, Tianjin 300130, China
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Maa, R.P. Liua,*
School of Materials Science and Engineering, Hebei Key Laboratory of Material Near-net Forming
Technology, Hebei University of Science and Technology, Shijiiazhuang 050018, China d
School of Mechanical Engineering, Yanshan University, Qinhuangdao 066004, China
ABSTRACT
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The effects of phase composition and microstructure evolution on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys (abbreviated as, TχZAB with χ=0, 10, 20, 30, and 40 wt.%) were investigated by means of theoretical calculations, optical microscopy, scanning electron microscopy,
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and transmission electron microscopy. The as-prepared TχZAB alloys showed an excellent combination of the tensile strength and ductility features. For example, the hot-rolled T40ZAB alloy
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displayed the ultrahigh tensile strength of 1535 MPa at 6.06% elongation. As Zr content increased, the fracture morphologies revealed typical transitional processes from dimple to quasi-cleavage fracture surfaces. The theoretical calculations agreed well with results from X-ray diffraction, where examined alloys showed only α/α′ crystal phases. Under the same hot-rolled condition, the microstructure evolved from distorted (or broken) lamellar α phase to crisscross acicular α′ martensite phase as the Zr content increased. Finally, in search for ideal structural titanium alloys, phase composition and microstructure and relationship with mechanical properties were examined and discussed. Keywords: Hot-rolled TiZrAlB alloys; Mechanical properties; Theoretical calculation; Phase composition; Microstructure
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ACCEPTED MANUSCRIPT *Corresponding author. address: State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, China e-mail: [email protected] (R.P. Liu). Tel: 0086-335-8074723
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Fax: 0086-335-8074545
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ACCEPTED MANUSCRIPT 1. Introduction Nowadays, structural titanium (Ti) alloys are increasingly applied in high-tech fields owing to their excellent mechanical properties combined with low density [1-4]. Because the requirements for these applications are becoming more and more strict, improving both the performance and
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affordability of Ti alloys become motivated and imperative. Hot-working processing methods, such as hot rolling and hot forging are effective for improving the mechanical properties of materials by eliminating defects in as-cast structures and solute segregations [5-9]. Furthermore, hot-working processing could not only influence phase composition and precipitate but also tailor the microstructure
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[8,9]. These two features of phase composition and microstructure could, in turn, notably modulate the
using thermal processing techniques.
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mechanical properties [10]. Therefore, it is practical to develop and prepare high-performance Ti alloys
At equilibrium conditions, Ti is either hexagonal close-packed (hcp) structural α phase or body-centered cubic (bcc) structural β phase. On the other hand, while Ti alloys are rapidly cooled from α+β dual-phase or single β phase regions, two possible types of martensite phases (hexagonal close-packed α′ phase and orthorhombic α′′ phase) and/or one possible non-equilibrium phase (simple
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hexagonal ω phase) could form. [11,12]. The stability of β phase strongly depends on the variation in alloy composition owing to the electron structure parameters bond energy and the number of shared electron pairs in the covalent bond [11]. It is well established that the martensitic structure of water
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quenched Ti and its alloys derivative at temperatures above α transition temperature occur following the order: α′, then transformed to α′′, then transformed to ω, then transformed to β. This is due to the
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increase in the content of β stabilizer [13]. Moreover, each phase makes a particular contribution to the mechanical properties. For instance, in quenched Ti-V and Ti-V-Sn alloys [14], the β and ω phases of binary Ti-V alloys would form as V content increased. Subsequently, Sn addition to Ti-V alloys enhances the martensitic transformation by suppressing athermal ω-phase formation. By comparing the tensile properties of β and α′ dominated Ti-V alloys, it could be noticed that α′ has higher strength but lower elongation [14]. Some studies also reported that the presence of α′′ martensites reduces hardness and strength of Ti alloys [15]. As mentioned above, several investigations have exampled phase transformations in Ti alloys, including the theoretical calculations based on valence electron theory and തതതത law. Here, the valence electron theory provides a relationship between the final crystal തതതത -Md Bo 3
ACCEPTED MANUSCRIPT തതതത law, തതതത-Md structure of quenched Ti alloys and its average number of valence electrons e/a [16]. In Bo തതതത is the average d-orbital energy level (eV) തതതത represents the average bond order between atoms and Md Bo തതതത diagram reveals the effect of alloying തതതത-Md of the elements present in the examined alloy. Also, the Bo തതതത diagram has been used for the elements on the stability of each phase of Ti alloys [17]. The തതത Boത-Md
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design of Ti-based alloys, which is useful for treating the phase stability problem. Therefore, theoretical calculations forecasting phase compositions of quenched Ti alloys would play a guidance role in alloy composition design and subsequent heat treatment.
Apart from tailoring phase composition, the control of microstructural is another effective way to
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adjust and improve the mechanical properties of Ti alloys [15,18,19]. The microstructural evolution strongly depends on cold/hot deformation process and subsequent heat-treatment processes. The
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microstructures of Ti alloys can typically be classified into basketweave, equiaxed, duplex, widmanstatten, among others. These various microstructures as well as various phases have their peculiar impacts on the mechanical properties. For instance, Shi et al. [15] reported that equiaxed microstructure of TiZr-based alloys has high-ductility but low tensile strength. After cold rolling, the size of equiaxed grains slightly decreased to form internal needle-like laths with abundant dislocations,
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leading to an increase in strength at the expense of drastic ductility loss. These results indicate that the equiaxed microstructure brings about favorable ductility and high fatigue strength. In addition, tailoring one or multiple microstructural variables (e.g. lamellar spacing, dislocation density, grain size) of
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similar microstructure of Ti alloys could influence the mechanical properties. The study of Ti-25Nb-10Ta-1Zr-0.2Fe alloy [20] showed that hardness and strength markedly enhance as cold
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reduction increase due to grain refinement caused by cold deformation. In sum, investigations associated with control of microstructure and effects of microstructure evolution on the mechanical properties surely contribute to develop advanced structural Ti alloy applications. In previous studies, novel as-cast TiZrAlB alloys were developed [21]. Meanwhile, the
mechanical properties were restricted due to the casting defects. For potential applications as structural materials, the mechanical properties of TiZrAlB alloys should be improved by further processing and heat treatment. Traditionally, Zr was used as a neutral element in Ti alloys but its introduction decreased phase-transition temperature according to binary Ti-Zr phase diagram, indicating the effect of Zr on β-phase stability. Therefore, the effect of Zr on phase stability of Ti alloys would still controversial and further studies should bring clarifications [13,16]. In this paper, a series of hot-rolled 4
ACCEPTED MANUSCRIPT TiZrAlB alloys with excellent mechanical properties were reported. The martensitic transformation mechanism in quenched Ti-Zr-Al-B alloys with high Zr contents was investigated. Also, the effects of phase composition and microstructure evolution on the mechanical properties of hot rolled TiZrAlB
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alloys were also examined and the results were discussed.
2. Calculation and experimental methods
In the calculation part, the structures were optimized using the CASTEP code [22] with an implementation of the ultrasoft pseudopotential proposed by Vanderbilt [23] based on density
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functional theory. The previous study reported by M. Yoshino et al. could provide the detailed calculation method [24]. The Monkhorst-Pack k grid of 4×4×4 was employed. And the chosen
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plane-wave cutoff energy was 500 eV, ensuring the convergence of cohesive energies within 0.01 eV, as compared with the results with the cutoff energies up to 800 eV. Structural relaxation was stopped until the changes of the total energy, maximum displacement, maximum stress, and maximum force were less than 5.0×10-6 eV per atom, 5.0×10-4 Å, 0.02 GPa and 0.01 eVÅ-1, respectively. The calculations of cohesive energies (Ecoh) was the most basic for investigating the alloying
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effect of Zr and Al on the β-phase stability of Ti alloys. Based on the research by Abdel-Hady [25], supercells consisting of 15 Ti atoms and one M atom (M = Zr or Al) were constructed using the optimized lattice parameters. The positions of the first-nearest-neighbor Ti atoms from a M atom were relaxed, and the cohesive energies were calculated from the total energies of the supercells and the
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isolated atoms. For example, the cohesive energy of a Ti-Zr alloy was shown as follows: Eୡ୭୦ (TiZr) = (Eୟ୲୭୫ (Zr) + 15Eୟ୲୭୫ (Ti) − Eୱ୳୮ୣ୰ (Ti15Zr))/16
(1)
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where Eatom(Zr) and Eatom(Ti) were the energies of isolated Zr and Ti atoms, respectively. Esuper(Ti15Zr) was the total energy of the supercell. The cohesive energies of each supercell were normalized by the number of atoms which was set as 16 in total. For the Ti-Zr-Al alloy, the supercell consisted of 14 Ti atoms, one Zr atom and one Al atom. The Al atom was located at the first-nearest-neighbor site from the Zr atom, and the positions of the neighboring atoms around Al and Zr atoms were relaxed so as to minimize the total energy. Then the cohesive energy was calculated in a similar way as Eq. (1). For the experiments, the nominal compositions of the experimental alloys Ti-χZr-4Al-0.005B were set to χ=0, 10, 20, 30 and 40 wt.%. The raw materials were sponge Ti (99.9 wt.%) mixed with sponge Zr (Zr+Hf ≥ 99.7 wt.%), Al (99.9 wt.%) and B (99.9 wt.%), and melted through a vacuum 5
ACCEPTED MANUSCRIPT non-consumable electro-arc furnace in a water cooled copper crucible using Ti as the getter under an argon atmosphere. The temperature in contact zone could reach 3000 ℃ when the raw materials melting in vacuum non-consumable electro-arc furnace. And smelting time was about three minutes each time flipped the ingots over. For compositional homogeneity, the experimental ingots were
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flipped and remelted at least five times. Differential scanning calorimetry (DSC) was used to measure the phase transition temperature at a heating rate of 10 °C/min in an argon protective environment using a corundum crucible, and the specimens were columnar with diameters of 3 mm and heights of 3 mm. The as-cast ingots were heated
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to 930 °C (heating rate: 10 °C/min) then left for 30 min before rolled into 5 mm thick plated specimens at the final deformation of 65%. Multi-pass rolling was carried out, and interval distance of roller
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decreased 2 mm each pass. Next, the specimens were quenched in water. Henceforth, the hot-rolled experimental alloys were denoted as HR-TχZAB, with χ=0, 10, 20, 30, and 40 wt.%. After removal of 1 mm thick portions from both sides of rolled surface, the bone-shaped plate specimens with original gauge lengths of 20 mm and cross-sectional dimensions of 3.00×2.00 mm2 were prepared for room-temperature tensile testing. All the tensile samples were cut from the rolled plates in the rolling
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direction. The uniaxial tensile tests were performed at a strain rate of 5×10-4 s-1 using an Instron 5892 Universal Material Testing Machine. The microhardness of polished alloys were measured using a microhardness tester at loads of 150 g for 10 s. X-ray diffraction (XRD) patterns were acquired by Cu
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Kα radiation (D/max-2500/PC) to analyze the phase composition and crystal structure of the hot-rolled TχZAB alloy specimens. The diffraction angle ranged from 20° to 100°, step size was 0.02° and time in
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each step was 2 s. The macrographic microstructures of the obtained alloys were observed by optical microscopy (OM) and scanning electron microscopy (SEM). Specimens used for OM and SEM were mechanically polished and chemically etched using a solution of 5% hydrofluoric acid (HF), 15% nitric acid (HNO3), and 80% deionized water (H2O). SEM was also utilized to analyze tensile fractures of the tensile
specimens.
high-magnification
Transmission
microstructure
electron
microscopy
of
experimental
the
(TEM) alloys
was with
employed
to
specimens
obtain twin-jet
electrochemically polished within solution containing 10% perchloric acid and 90% methanol at 13 V and -35 ℃. To minimize errors, all tests were performed at least three times. The average values of experimental data were then taken to draw the charts and test results were used for comparison of only 6
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3. Results and discussion 3.1. Mechanical Properties To evaluate the mechanical properties of HR-TχZAB alloys, uniaxial tensile and microhardness
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tests were conducted. Fig. 1 shows the engineering stress-strain curves with dimensions of tensile specimens of the alloys, and the data are gathered in Table 1. The yield strength (σ0.2), ultimate strength (σb) and microhardness (HV) were all enhanced with concomitant failure elongation (εf) loss as the Zr content increased. The HR-TAB alloy exhibited the lowest yield strength of 682 MPa, ultimate strength
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of 775 MPa, microhardness of 286 HV, and highest failure elongation of 15.59%. As Zr content rose to 40 wt.%, σ0.2, σb and microhardness increased respectively to 1388 MPa, 1535 MPa and 442 HV at the
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expense of the decline in ductility by 9.53%. Young's moduli E showed a downward trend as the Zr content increased. Among the examined alloys, HR-T20ZAB did not only showed moderate ductility but also ultrahigh strength, indicating its comprehensive mechanical properties. On the other hand, HR-T40ZAB alloy revealed the highest ultimate strength. The comparison between HR-TχZAB alloys did not only exhibit nearly equal elongation-to-failure with as-cast TχZAB alloys [21] but also much
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higher strengths than the as-cast TχZAB alloys. Fig. 2(a) demonstrated the differences of mechanical performances between hot-rolled and as-cast alloys with identical compositions. The relationships between the mechanical properties and Zr contents are shown in Fig. 2(b). Compared to as-cast TχZAB
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alloys [21], σ0.2 and σb of (HR-TAB, HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB) alloys increased by (50, 123, 192, 403, and 395 MPa) and (34, 171, 309, 436, and 400 MPa),
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respectively. Therefore, the hot-rolling deformation of TχZAB alloys undoubtedly enhanced the strength with almost unvaried elongation in comparison to parallel as-cast TχZAB alloys. According to previous studies [2,6,18,26-30], tensile strength and elongation of some frequently-applied structural Ti alloys were compared to current alloys and the results are listed in Fig 3. It could be easily observed that the mechanical properties of the as-prepared alloys outperformed those of published alloys, indicating the good combination of tensile strength and ductility required by structural materials. 3.2. Fracture surface morphologies The low-magnification SEM fracture morphologies of hot-rolled TχZAB alloys with different Zr contents are gathered in Fig. 4 from a macro perspective. The variation of fracture surfaces of post-rift 7
ACCEPTED MANUSCRIPT tensile specimens depicted typical trends. The fracture morphology of HR-TAB alloy specimen consisted mainly of large, deep and equiaxed dimples filled with minor dimples (Fig. 4(a)). The addition of Zr reduced and shallowed the formed dimples with continuous rippled pattern. As Zr content increased to 20 wt.%, the dimples shrunk, and the micro-pores caused by defects formed
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metastable dimples under the tensile stress took the dominant in fracture surface (Fig. 4(c)). The quasi-cleavage fracture morphology was then observed in Fig. 4(d) and 4(e), where facets, cracks and cleavage planes with mini-sized dimples became apparent on most of the surface. The comparison between (d) with (e) in Fig. 4 revealed that the fracture morphology of quasi-cleavage steps and rubbly
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surface of HR-T40ZAB alloy was more distinct than HR-T30ZAB alloy. In addition, the necking phenomenon gradually changed from obvious to slight then vanished. The combined elongation of
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HR-TχZAB alloys with fracture mode transformed from ductile to brittle, evidencing a decrease in alloy ductility.
3.3. The effects of phases on microstructural evolution and mechanical properties The DSC curves of the examined alloys are shown in Fig. 5 as a reference to phase-transition temperature. During heating, phase-transition temperature could be obtained by measuring the
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exothermic and endothermic variation of original TχZAB alloy specimens. Overall, only α phase existed until temperature reached onset-transition temperature. Between onset-transition and end-transition temperatures, both α and β phases were present. This indicated that once the
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onset-transition temperature was reached, α phase should start to transform into β phase. The phase transformation was complete once the temperature reached the end-transition temperature, after which
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only β phase existed. Obviously, phase-transition temperatures of the alloys decreased as Zr content rose from 0 wt.% to 40 wt.%. The onset temperature of transition from α→β phase declined from 954 ℃ to 776 ℃, and the end-transition temperature decreased from 1009 ℃ to 860 ℃. In Fig. 5, the dotted line represented 930 ℃ intersected with five DSC curves, where five intersection points located at: single α phase region, α→(α+β) dual-phase region, (α+β)→β dual-phase region, and two single β phase region of TAB, T10ZAB, T20ZAB, T30ZAB and T40ZAB, respectively. The microstructure could be tailored by varying the phase composition due to phase transformation [5]. Hence, 930 ℃ was selected as the rolling temperature to further study the effects of phase transformation and microstructure evolution on the mechanical properties of the alloys with various Zr contents. Increasing β-stabilizing alloy element contents would govern the martensitic structure of water 8
ACCEPTED MANUSCRIPT quenched Ti alloys containing β-phase, and the solid-state transformation which was diffusionless, i.e. martensitic, changes as the order: α′, then transformed to α′′, then transformed to ω, then transformed to β. [31-33]. Though phase transition temperature of Ti alloys decreased as Zr content increased according to binary equilibrium phase diagram of Ti-Zr and Fig. 5, the effect of Zr on β-phase stability
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of TχZAB alloy was estimated by theoretical calculations. The value of e/a could be used as reference for the quenched phase of Ti alloys according to valence electron theory. A higher value of e/a was acquired by calculation, and stronger stability was found for β phase [34]. Previous studies [16] reported that the martensitic phase of Ti alloys with e/a value below 4.07 should be α′ phase. The
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calculated e/a ratios of TχZAB alloys are listed in Table 2, and the results were all below 4.07. Besides തതതത law would be considered as another criterion for water തതതത-Md the valence electron theory, the Bo
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തതതത diagram of Ti alloys quenched phases of HR-TχZAB alloys [17]. Fig. 6 shows the extended തതത Boത-Md presenting α′, α″, ω and β four phase regions, and Table 2 lists the values of തതത Boത and തതതത Md. T10ZAB, T20ZAB, T30ZAB and T40ZAB alloys quenched from temperatures above the α transus temperature തതതത diagram. Here, the valence തതതത-Md were uniformly located in martensite α′ region (shadow area) of Bo തതതത law were appropriate for low-alloy Ti alloys with alloying element തതതത -Md electron theory and Bo
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contents of less than 20 at.% [16]. However, T40ZAB alloy appeared outside the valid interval. To yield reasonable and consummate theoretical calculations, the cohesive energies were calculated on the basis of Eq. (1). The cohesive energies of body-centered cubic (β phase) Ti-Al, Ti-Zr
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and Ti-Zr-Al supercells were 6.298, 6.441 and 6.350 eV atom-1, respectively. It is clear to contrast that the cohesive energy of Ti-Zr alloy was higher than those of Ti-Al and Ti-Zr-Al alloys, indicating that
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the average chemical bond between atoms in β-phase was stronger in Ti-Zr alloy than in Ti-Al or Ti-Zr-Al alloys. Here, the mass ratio of the alloying element Al remained invariant whereas the atomic ratio increased. Thus, the effect of Al on stability of β-phase in TχZAB alloys should also be taken into account. The ∆Ecoh referring to the difference between the cohesive energies of two body-centered cubic supercells could be calculated as follows: ∆Eୡ୭୦ (a) = Eୡ୭୦ (TiZrAl) − Eୡ୭୦ (TiAl)
(2)
∆Eୡ୭୦ (b) = Eୡ୭୦ (TiZrAl) − Eୡ୭୦ (TiZr) where ∆Ecoh(a) is 0.052 and ∆Ecoh(b) is -0.092. |∆Eୡ୭୦ (a)| represents the degree of Zr on the stability of β-phase, and |∆Eୡ୭୦ (b)| represents the degree of Al in reducing the β-phase stability. It would be easy to deduce that |∆Eୡ୭୦ (b)| − |∆Eୡ୭୦ (a)| > 0, meaning that the degree of Al on reducing β-phase 9
ACCEPTED MANUSCRIPT stability was much stronger than the degree of Zr on β-phase stability. All trends obtained by the above theoretical calculations were confirmed experimentally and explained below. Fig. 7 shows the XRD patterns of HR-TχZAB alloys for their actual phase compositions. The diffraction peaks only matched peaks of α or α′ phases since the crystal structures of both α and α′
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consisted of hexagonal close-packed (hcp). The XRD patterns indicated that the phases of HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys quenched from α+β dual-phase region, and the single β region were merely formed of primary α or martensite α′ phase with no residual β phase or other intermetallic phases. In other words, even at rolling and quenching temperatures above
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β transus temperature, only the β→α′ phase transition occurred during the quenching process. This was in accordance with the estimated theoretical calculations. When combined with DSC curves, the
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analysis revealed that the volume fraction of α′ phase in HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys gradually increased after hot rolling and quenching. According to previous studies [9], α′ martensite is a hard phase compared to other common phases in Ti and Zr alloys, such as α, β, and ω phases. Therefore, the formation of α′ martensite would increase both the yield strength and tensile strength but reduce ductility. Moreover, the martensitic transformation was found quite effective
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in strengthening Ti- and Zr-base alloys. Unlike the Fe-C system, the hcp martensites in Ti- and Zr-based alloys were not supersaturated with interstitial elements. In general, β-phase region quenching would produce martensitic α′ structure and slow cooling from the β-phase region to form
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Widmanstatten α structure. This strengthening effect owed to their substructures. For isomorphous martensitic α′ and Widmanstatten α, the martensitic α′ structure contained more closely spaced
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interfaces separating neighboring lamellae or twins with high density of dislocation compared to Widmanstatten α structure with the same composition [35]. Here, the microhardness of HR-TχZAB alloys increased as volume fraction of α′ phase rose. Z. Tarzimoghadam et al. [33] reported that phase transition in transformation kinetics and kinematics would enable the design of complex microstructures with features spanning across broad range of length scaling. In Fig. 7(right), the width of the diffraction peaks broadened after hot rolling owing to the synergistic effect of increased micro residual stress and grain refinement caused by large deformation and reduced phase-transition temperature. Previous study [21] reported that solute element Zr-induced lattice distortion changed the lattice parameter of TχZAB alloys from XRD results. As Zr content increased, the variation of axial ratio “c/a” 10
ACCEPTED MANUSCRIPT influenced the slip mechanism of TχZAB alloys, which further caused the decrease of ductility. Meanwhile, addition of alloying element Zr led to increasing degrees of saturation, which also changed the distance between the atoms and resulted in a change of the Young's modulus. Young's modulus is an intrinsic property of materials and is determined by the bonding force between atoms [36]. This
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bonding force is related to the addition of alloying elements, which can be affected by the crystal texture and any change in the distances between the atoms. In current study, the Young's modulus decreased as Zr increased. For the TχZAB alloys with the same phase, the lattice parameter “a” increased following the increased Zr content owing to the larger atomic radius of Zr than Ti [21]. The
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aforementioned result was confirmed by the XRD peaks in turn shifting towards low angle orientation. Therefore, Young's modulus mainly depended on chemical composition, i.e., Zr content.
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3.4. The effects of microstructures on the mechanical properties
The microstructures of the as-prepared TχZAB alloys visibly changed after hot-rolling when compared to those of as-cast TχZAB alloys reported in our previous study [21]. At the same hot-rolled conditions, the microstructure is dependent on the β transition temperature [37]. Meanwhile, at constant composition conditions, the microstructure would be a critical factor influencing the mechanical
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properties. For example, texture is a factor that affects mechanical properties. As we know, this influence is pronounced for hcp metals such as α-Ti or α-Zr alloys due to the high volume fraction of the inherently anisotropic hexagonal α-phase [38]. Fig. 8 depicts the metallographs of HR-TχZAB
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alloys. Obviously, the lamellar α-phase of TAB and T10ZAB were distorted or broken after the hot-rolling process (Fig 8(a) and 8(b)). The microstructure of the as-cast TAB and T10ZAB alloys
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were completely filled with crisscross slender α lamella. After hot-rolling, the typical basketweave morphologies underwent severe changes. Jiang et al. [7] suggested mechanisms behind the distortion and break-up of α-phase lamella during hot deformation process. The localized shear and rotation of α lamella would occur with misorientation across the shear band to reach approximately 20°. Simultaneously, the low and high angle boundaries both formed across α lamella, where misorientation angles varied from a few degrees to up to 30°. For our proposed alloys, these mechanisms would also be appropriate. The distorted or fractured lamellae would increase the grain boundary area hindered the dislocation motion, which caused the grain boundary strengthening. The microstructure evolution of HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys appeared quite similar (Fig 8(c), 8(d) and 8(e), respectively). The metallographs revealed that the microstructures was characterized by fine α′ 11
ACCEPTED MANUSCRIPT martensite crystals within the original β grain boundaries. The aforementioned morphology with acicular plate α′-phase would be typical of martensite in titanium alloys [33]. For HR-T20ZAB alloy, considering the rolling temperature nearby the β transition temperature, distorted original α phases were hardly detected. In addition, the α′-lamella structure became gradually finer in HR-T20ZAB,
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HR-T30ZAB and HR-T40ZAB alloys. The microstructural details of HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys were revealed by TEM bright field micrographs (Fig 9). A large number of dislocations could clearly be observed in the deformed lamellae (Fig 9(a)), suggesting that HR-T10ZAB alloy suffered from
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significant deformation during the hot-rolling process. The dislocation density was so high with severe dislocation pile-up in the grain boundaries that dislocation mobility was restricted inducing an increase
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in strength. In α-Ti alloys (hcp), the principal deformation mode corresponded to a-type dislocations (a/3<112ത0>), mainly gliding in the prismatic planes. The deformation was controlled by the motion of screw segments, in which atomistic simulations showed lattice friction due to three-dimensional core structure spread in prismatic as well as in basal and pyramidal planes [1,3]. Moreover, the high lattice friction caused by core structure of the dislocations had a major impact on the strengthening in lamellar
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microstructures [1]. Hence, the mechanism of dislocation strengthening was responsible for improving the mechanical properties of the proposed hot-rolled alloys in contrast to as-cast alloys. Fig 9(b-d) illustrated that the width of α′ lamellae decreased dramatically with few apparent dislocations. The
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selected-area electron diffraction pattern in Fig 9(d) further confirmed the phase composition. According to Fig 5, β transition temperature decreased with the increased Zr content. During quenching,
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β phase was rapidly converted into metastable α′ martensite phase. Furthermore, under the same quenching temperature and cooling conditions, α′ martensite phase of the alloy with higher Zr content had insufficient time to grow thicker, resulting in thinner lamellar thickness and spacing. As Zr content increased, the microstructure of the hot-rolled alloys increasingly refined. Admittedly, the strength of polycrystalline materials like TχZAB alloys followed the Hall-Petch empirical relationship depending on the fineness of grain: σ=σ0+kd-0.5, where σ represents the yield strength of the TχZAB alloys. Previous investigations [3,14,39] reported that full acicular α′ martensite microstructures could effectively have high strength in Ti alloys (e.g. in Ti-V, Ti-V-Sn and Ti-V-Al alloys) via controlling the size and distribution of α′ lamellae. Here, Hall-Petch effect was controlled by thickness of the α/α′ lamellae (i.e. the d). Fig 10(a) shows the average thickness of α/α′ lamellae and the thickness 12
ACCEPTED MANUSCRIPT distribution. Average grain size was measured according to the semiautomatic image analysis (ASTM: E1382-97(2010)). In order to obtain accurate data, thirty lamellae were measured from each sample in this study. By using Hall-Petch equation, it could be found that the decreased grain dimensions gave rise to enhanced strength (Fig 10(b)). Fine crystal strengthening mainly depended on hindering of the
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dislocation slip by the grain boundaries [40]. On the one hand, reducing the grain size with increasing density of the grain boundaries in the same cross-sectional area increased the number of dislocations piling up. The severe stress concentration occurred on grain boundary. On the other hand, the different grain orientations in both sides of the grain boundary led to increased sliding resistance near the grain
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boundary during the movement of the dislocation obstacle in the polycrystal. Thus, the propagative slip band could not directly enter the neighbouring grains resulting in accumulated dislocation at the grain
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boundary. Dislocation pile-up substantially improved the strength of alloy. Tarzimoghadam et al. [33] also reported similar results in Ti-4Mo alloy. Considering a mass of dislocations and dislocation cells formed in HR-TAB, HR-T10ZAB and HR-T20ZAB alloys, the increase in strength was more significant than HR-T30ZAB and HR-T40ZAB alloys. The microstructures of TχZAB alloys under the hot-rolled and water-quenched condition would be stable at service temperature. Therefore, the
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microstructure evolution would be a key to improve the mechanical properties of TχZAB alloys, and the mechanisms could most likely be generalized to other Ti alloys.
4. Conclusions
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For further studying and developing the potential structural Ti alloys, the mechanical properties, fracture characteristics, phase composition and microstructure evolution of HR-TχZAB alloys were
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investigated and the effects of phase composition and microstructure on mechanical properties were discussed. The following conclusions could be drawn: (1) The ternary yield strength, ultimate strength and microhardness obviously increased as Zr content increased. The HR-T40ZAB alloy revealed the highest tensile strength of 1535 MPa at 6.06% elongation. The comparison between HR-TχZAB alloys with previous as-cast TχZAB alloys with identical compositions suggested that hot-rolled alloys did not only exhibit nearly equal elongation with the as-cast alloys but also much higher strength than as-cast alloys. (2) The fracture morphologies revealed typical transitional processes from ductile fracture surface to quasi-cleavage fracture surface as the Zr content increased. Also, the necking phenomenon 13
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(3) The theoretical calculations of phase composition were consistent with the experimental data, where only α/α′ phases were obtained in HR-TχZAB alloys. And the strength increased as the
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volume fraction of α′ phase rose. (4) The lamellae of HR-TAB and HR-T10ZAB alloys were distorted or broken after hot-rolling, forming large numbers of intragranular dislocations. The microstructure of HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys was completely filled with acicular α′ martensite, which
increased as the thickness of lamellae decreased.
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Acknowledgments
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gradually became finer as Zr content increased. Additionally, the strength of HR-TχZAB alloys
This work was supported by the NSFC (Grant no. 51531005/51671166/51434008) and Postdoctoral Science Foundation of Hebei Province of China (Grant no. B2017003008).
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ACCEPTED MANUSCRIPT Figure captions Fig. 1. Stress-strain curves and tensile specimen dimensions of HR-TχZAB alloys. Fig. 2. Changes of mechanical properties: (a) the difference values of mechanical performances between hot-rolled and as-cast alloys with identical composition, (b) the variation trend of
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mechanical properties of series HR-TχZAB alloys. Fig. 3. The mechanical properties of the current HR-TχZAB alloys and other available Ti alloys [2,6,18,26-30].
χ=30 and (e) χ=40. Fig. 5. DSC curves of TχZAB alloys.
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തതതത diagram of Ti alloys [15]. Fig. 6. The expanded തതത Boത-Md
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Fig. 4. SEM fracture morphologies of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d)
Fig. 7. XRD patterns of the HR-TχZAB alloys.
Fig. 8. Optical micrographs of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
Fig. 9. TEM bright field microstructures of the HR-TχZAB alloys (wt.%): (a) χ=10, (b) χ=20, (c) χ=30
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and (d) χ=40.
Fig. 10. Grain size analysis: (a) variation of the average thickness of α/α′ lamellae and thickness
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distribution, (b) variation of σ0.2 with the thickness of α/α′ lamellae Hall-Petch equation.
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Fig. 1. Stress-strain curves and tensile specimen dimensions of HR-TχZAB alloys.
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Fig. 2.Changes of mechanical properties: (a) the difference values of mechanical performances between
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hot-rolled and as-cast alloys with identical composition, (b) the variation trend of mechanical properties of series HR-TχZAB alloys.
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Fig. 3. The tensile properties of the current HR-TχZAB alloys and other available Ti alloys
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[2,6,18,26-30].
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Fig. 4. SEM fracture morphologies of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
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Fig. 5. DSC curves of TχZAB alloys.
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തതതതത diagram of Ti alloys [17]. തതതത-݀ܯ Fig. 6. The expanded ܤ
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Fig. 7. XRD patterns of the HR-TχZAB alloys.
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Fig. 8. Optical micrographs of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
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Fig. 9. TEM bright field microstructures of the HR-TχZAB alloys (wt.%): (a) χ=10, (b) χ=20, (c) χ=30
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Fig. 10. Grain size analysis: (a) variation of the average thickness of α/α′ lamellae and thickness distribution, (b) variation of σ0.2 with the thickness of α/α′ lamellae Hall-Petch equation.
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ACCEPTED MANUSCRIPT Table Captions Table 1 Mechanical properties of the HR-TχZAB alloys.
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തതതത and e/a values of TχZAB alloys. തതതത, Md Table 2 Nominal compositions (wt.% and at.%) and Bo
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ACCEPTED MANUSCRIPT Table 1 Mechanical properties of the HR-TχZAB alloys. E (GPa)
σ0.2 (MPa)
σb (MPa)
εf (%)
HV
HR-TAB
161
682
775
15.59
286
HR-T10ZAB
143
860
1028
10.61
318
HR-T20ZAB
122
1062
1290
8.43
391
HR-T30ZAB
120
1294
1459
6.99
406
HR-T40ZAB
120
1388
1535
6.06
442
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Alloys
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ACCEPTED MANUSCRIPT തതതതത and e/a values of TχZAB alloys. തതതത, ݀ܯ Table 2 Nominal compositions (wt.% and at.%) and ܤ wt.%
Alloys
at.%
തതതത ܤ
തതതതത ݀ܯ
e/a
0.02
2.764
2.429
3.93
7.22
0.02
2.779
2.455
3.93
11.21
7.58
0.02
2.795
2.482
3.92
74.27
17.72
7.99
0.02
2.813
2.513
3.92
65.57
24.96
8.44
0.03
2.804
2.523
3.87
Zr
Al
B
Ti
Zr
Al
B
HR-TAB
95.995
0
4
0.005
93.09
0
6.89
HR-T10ZAB
85.995
10
4
0.005
87.43
5.34
HR-T20ZAB
75.995
20
4
0.005
81.18
HR-T30ZAB
65.995
30
4
0.005
HR-T40ZAB
55.995
40
4
0.005
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Ti
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ACCEPTED MANUSCRIPT (1) The mechanical properties of TiZrAlB alloys are improved by hot rolling process. (2) The HR-T40ZAB alloy shows tensile strength of 1535 MPa at 6.06% elongation. (3) Theoretical calculations are confirmed by experimental results.
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(5) The microstructure is refined as the Zr content increased.
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(4) Only α/α′ phases are obtained in hot-rolled TiZrAlB alloys.
S0925-8388(18)31329-X
DOI:
10.1016/j.jallcom.2018.04.043
Reference:
JALCOM 45681
To appear in:
Journal of Alloys and Compounds
Received Date: 20 December 2017 Revised Date:
30 March 2018
Accepted Date: 3 April 2018
Please cite this article as: S.G. Liu, C.Q. Xia, Z.H. Feng, X. Zhang, B.H. Chen, C.L. Xie, Y.K. Zhou, X.Y. Zhang, M.Z. Ma, R.P. Liu, Influence of phase composition and microstructure on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys, Journal of Alloys and Compounds (2018), doi: 10.1016/ j.jallcom.2018.04.043. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
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ACCEPTED MANUSCRIPT Influence of phase composition and microstructure on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys S.G. Liua, C.Q. Xiab, Z.H. Fengc , X. Zhanga, B.H. Chena, C.L. Xiea, Y.K. Zhoud, X.Y. Zhanga, M.Z.
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University,
Qinhuangdao 066004, China b
School of Materials Science and Engineering, Research Institute for Energy Equipment Materials,
c
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Hebei University of Technology, Tianjin 300130, China
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a
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Maa, R.P. Liua,*
School of Materials Science and Engineering, Hebei Key Laboratory of Material Near-net Forming
Technology, Hebei University of Science and Technology, Shijiiazhuang 050018, China d
School of Mechanical Engineering, Yanshan University, Qinhuangdao 066004, China
ABSTRACT
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The effects of phase composition and microstructure evolution on mechanical properties of hot-rolled Ti-χZr-4Al-0.005B alloys (abbreviated as, TχZAB with χ=0, 10, 20, 30, and 40 wt.%) were investigated by means of theoretical calculations, optical microscopy, scanning electron microscopy,
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and transmission electron microscopy. The as-prepared TχZAB alloys showed an excellent combination of the tensile strength and ductility features. For example, the hot-rolled T40ZAB alloy
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displayed the ultrahigh tensile strength of 1535 MPa at 6.06% elongation. As Zr content increased, the fracture morphologies revealed typical transitional processes from dimple to quasi-cleavage fracture surfaces. The theoretical calculations agreed well with results from X-ray diffraction, where examined alloys showed only α/α′ crystal phases. Under the same hot-rolled condition, the microstructure evolved from distorted (or broken) lamellar α phase to crisscross acicular α′ martensite phase as the Zr content increased. Finally, in search for ideal structural titanium alloys, phase composition and microstructure and relationship with mechanical properties were examined and discussed. Keywords: Hot-rolled TiZrAlB alloys; Mechanical properties; Theoretical calculation; Phase composition; Microstructure
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ACCEPTED MANUSCRIPT *Corresponding author. address: State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao 066004, China e-mail: [email protected] (R.P. Liu). Tel: 0086-335-8074723
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Fax: 0086-335-8074545
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ACCEPTED MANUSCRIPT 1. Introduction Nowadays, structural titanium (Ti) alloys are increasingly applied in high-tech fields owing to their excellent mechanical properties combined with low density [1-4]. Because the requirements for these applications are becoming more and more strict, improving both the performance and
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affordability of Ti alloys become motivated and imperative. Hot-working processing methods, such as hot rolling and hot forging are effective for improving the mechanical properties of materials by eliminating defects in as-cast structures and solute segregations [5-9]. Furthermore, hot-working processing could not only influence phase composition and precipitate but also tailor the microstructure
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[8,9]. These two features of phase composition and microstructure could, in turn, notably modulate the
using thermal processing techniques.
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mechanical properties [10]. Therefore, it is practical to develop and prepare high-performance Ti alloys
At equilibrium conditions, Ti is either hexagonal close-packed (hcp) structural α phase or body-centered cubic (bcc) structural β phase. On the other hand, while Ti alloys are rapidly cooled from α+β dual-phase or single β phase regions, two possible types of martensite phases (hexagonal close-packed α′ phase and orthorhombic α′′ phase) and/or one possible non-equilibrium phase (simple
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hexagonal ω phase) could form. [11,12]. The stability of β phase strongly depends on the variation in alloy composition owing to the electron structure parameters bond energy and the number of shared electron pairs in the covalent bond [11]. It is well established that the martensitic structure of water
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quenched Ti and its alloys derivative at temperatures above α transition temperature occur following the order: α′, then transformed to α′′, then transformed to ω, then transformed to β. This is due to the
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increase in the content of β stabilizer [13]. Moreover, each phase makes a particular contribution to the mechanical properties. For instance, in quenched Ti-V and Ti-V-Sn alloys [14], the β and ω phases of binary Ti-V alloys would form as V content increased. Subsequently, Sn addition to Ti-V alloys enhances the martensitic transformation by suppressing athermal ω-phase formation. By comparing the tensile properties of β and α′ dominated Ti-V alloys, it could be noticed that α′ has higher strength but lower elongation [14]. Some studies also reported that the presence of α′′ martensites reduces hardness and strength of Ti alloys [15]. As mentioned above, several investigations have exampled phase transformations in Ti alloys, including the theoretical calculations based on valence electron theory and തതതത law. Here, the valence electron theory provides a relationship between the final crystal തതതത -Md Bo 3
ACCEPTED MANUSCRIPT തതതത law, തതതത-Md structure of quenched Ti alloys and its average number of valence electrons e/a [16]. In Bo തതതത is the average d-orbital energy level (eV) തതതത represents the average bond order between atoms and Md Bo തതതത diagram reveals the effect of alloying തതതത-Md of the elements present in the examined alloy. Also, the Bo തതതത diagram has been used for the elements on the stability of each phase of Ti alloys [17]. The തതത Boത-Md
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design of Ti-based alloys, which is useful for treating the phase stability problem. Therefore, theoretical calculations forecasting phase compositions of quenched Ti alloys would play a guidance role in alloy composition design and subsequent heat treatment.
Apart from tailoring phase composition, the control of microstructural is another effective way to
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adjust and improve the mechanical properties of Ti alloys [15,18,19]. The microstructural evolution strongly depends on cold/hot deformation process and subsequent heat-treatment processes. The
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microstructures of Ti alloys can typically be classified into basketweave, equiaxed, duplex, widmanstatten, among others. These various microstructures as well as various phases have their peculiar impacts on the mechanical properties. For instance, Shi et al. [15] reported that equiaxed microstructure of TiZr-based alloys has high-ductility but low tensile strength. After cold rolling, the size of equiaxed grains slightly decreased to form internal needle-like laths with abundant dislocations,
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leading to an increase in strength at the expense of drastic ductility loss. These results indicate that the equiaxed microstructure brings about favorable ductility and high fatigue strength. In addition, tailoring one or multiple microstructural variables (e.g. lamellar spacing, dislocation density, grain size) of
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similar microstructure of Ti alloys could influence the mechanical properties. The study of Ti-25Nb-10Ta-1Zr-0.2Fe alloy [20] showed that hardness and strength markedly enhance as cold
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reduction increase due to grain refinement caused by cold deformation. In sum, investigations associated with control of microstructure and effects of microstructure evolution on the mechanical properties surely contribute to develop advanced structural Ti alloy applications. In previous studies, novel as-cast TiZrAlB alloys were developed [21]. Meanwhile, the
mechanical properties were restricted due to the casting defects. For potential applications as structural materials, the mechanical properties of TiZrAlB alloys should be improved by further processing and heat treatment. Traditionally, Zr was used as a neutral element in Ti alloys but its introduction decreased phase-transition temperature according to binary Ti-Zr phase diagram, indicating the effect of Zr on β-phase stability. Therefore, the effect of Zr on phase stability of Ti alloys would still controversial and further studies should bring clarifications [13,16]. In this paper, a series of hot-rolled 4
ACCEPTED MANUSCRIPT TiZrAlB alloys with excellent mechanical properties were reported. The martensitic transformation mechanism in quenched Ti-Zr-Al-B alloys with high Zr contents was investigated. Also, the effects of phase composition and microstructure evolution on the mechanical properties of hot rolled TiZrAlB
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alloys were also examined and the results were discussed.
2. Calculation and experimental methods
In the calculation part, the structures were optimized using the CASTEP code [22] with an implementation of the ultrasoft pseudopotential proposed by Vanderbilt [23] based on density
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functional theory. The previous study reported by M. Yoshino et al. could provide the detailed calculation method [24]. The Monkhorst-Pack k grid of 4×4×4 was employed. And the chosen
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plane-wave cutoff energy was 500 eV, ensuring the convergence of cohesive energies within 0.01 eV, as compared with the results with the cutoff energies up to 800 eV. Structural relaxation was stopped until the changes of the total energy, maximum displacement, maximum stress, and maximum force were less than 5.0×10-6 eV per atom, 5.0×10-4 Å, 0.02 GPa and 0.01 eVÅ-1, respectively. The calculations of cohesive energies (Ecoh) was the most basic for investigating the alloying
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effect of Zr and Al on the β-phase stability of Ti alloys. Based on the research by Abdel-Hady [25], supercells consisting of 15 Ti atoms and one M atom (M = Zr or Al) were constructed using the optimized lattice parameters. The positions of the first-nearest-neighbor Ti atoms from a M atom were relaxed, and the cohesive energies were calculated from the total energies of the supercells and the
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isolated atoms. For example, the cohesive energy of a Ti-Zr alloy was shown as follows: Eୡ୭୦ (TiZr) = (Eୟ୲୭୫ (Zr) + 15Eୟ୲୭୫ (Ti) − Eୱ୳୮ୣ୰ (Ti15Zr))/16
(1)
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where Eatom(Zr) and Eatom(Ti) were the energies of isolated Zr and Ti atoms, respectively. Esuper(Ti15Zr) was the total energy of the supercell. The cohesive energies of each supercell were normalized by the number of atoms which was set as 16 in total. For the Ti-Zr-Al alloy, the supercell consisted of 14 Ti atoms, one Zr atom and one Al atom. The Al atom was located at the first-nearest-neighbor site from the Zr atom, and the positions of the neighboring atoms around Al and Zr atoms were relaxed so as to minimize the total energy. Then the cohesive energy was calculated in a similar way as Eq. (1). For the experiments, the nominal compositions of the experimental alloys Ti-χZr-4Al-0.005B were set to χ=0, 10, 20, 30 and 40 wt.%. The raw materials were sponge Ti (99.9 wt.%) mixed with sponge Zr (Zr+Hf ≥ 99.7 wt.%), Al (99.9 wt.%) and B (99.9 wt.%), and melted through a vacuum 5
ACCEPTED MANUSCRIPT non-consumable electro-arc furnace in a water cooled copper crucible using Ti as the getter under an argon atmosphere. The temperature in contact zone could reach 3000 ℃ when the raw materials melting in vacuum non-consumable electro-arc furnace. And smelting time was about three minutes each time flipped the ingots over. For compositional homogeneity, the experimental ingots were
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flipped and remelted at least five times. Differential scanning calorimetry (DSC) was used to measure the phase transition temperature at a heating rate of 10 °C/min in an argon protective environment using a corundum crucible, and the specimens were columnar with diameters of 3 mm and heights of 3 mm. The as-cast ingots were heated
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to 930 °C (heating rate: 10 °C/min) then left for 30 min before rolled into 5 mm thick plated specimens at the final deformation of 65%. Multi-pass rolling was carried out, and interval distance of roller
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decreased 2 mm each pass. Next, the specimens were quenched in water. Henceforth, the hot-rolled experimental alloys were denoted as HR-TχZAB, with χ=0, 10, 20, 30, and 40 wt.%. After removal of 1 mm thick portions from both sides of rolled surface, the bone-shaped plate specimens with original gauge lengths of 20 mm and cross-sectional dimensions of 3.00×2.00 mm2 were prepared for room-temperature tensile testing. All the tensile samples were cut from the rolled plates in the rolling
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direction. The uniaxial tensile tests were performed at a strain rate of 5×10-4 s-1 using an Instron 5892 Universal Material Testing Machine. The microhardness of polished alloys were measured using a microhardness tester at loads of 150 g for 10 s. X-ray diffraction (XRD) patterns were acquired by Cu
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Kα radiation (D/max-2500/PC) to analyze the phase composition and crystal structure of the hot-rolled TχZAB alloy specimens. The diffraction angle ranged from 20° to 100°, step size was 0.02° and time in
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each step was 2 s. The macrographic microstructures of the obtained alloys were observed by optical microscopy (OM) and scanning electron microscopy (SEM). Specimens used for OM and SEM were mechanically polished and chemically etched using a solution of 5% hydrofluoric acid (HF), 15% nitric acid (HNO3), and 80% deionized water (H2O). SEM was also utilized to analyze tensile fractures of the tensile
specimens.
high-magnification
Transmission
microstructure
electron
microscopy
of
experimental
the
(TEM) alloys
was with
employed
to
specimens
obtain twin-jet
electrochemically polished within solution containing 10% perchloric acid and 90% methanol at 13 V and -35 ℃. To minimize errors, all tests were performed at least three times. The average values of experimental data were then taken to draw the charts and test results were used for comparison of only 6
ACCEPTED MANUSCRIPT experimental specimens.
3. Results and discussion 3.1. Mechanical Properties To evaluate the mechanical properties of HR-TχZAB alloys, uniaxial tensile and microhardness
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tests were conducted. Fig. 1 shows the engineering stress-strain curves with dimensions of tensile specimens of the alloys, and the data are gathered in Table 1. The yield strength (σ0.2), ultimate strength (σb) and microhardness (HV) were all enhanced with concomitant failure elongation (εf) loss as the Zr content increased. The HR-TAB alloy exhibited the lowest yield strength of 682 MPa, ultimate strength
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of 775 MPa, microhardness of 286 HV, and highest failure elongation of 15.59%. As Zr content rose to 40 wt.%, σ0.2, σb and microhardness increased respectively to 1388 MPa, 1535 MPa and 442 HV at the
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expense of the decline in ductility by 9.53%. Young's moduli E showed a downward trend as the Zr content increased. Among the examined alloys, HR-T20ZAB did not only showed moderate ductility but also ultrahigh strength, indicating its comprehensive mechanical properties. On the other hand, HR-T40ZAB alloy revealed the highest ultimate strength. The comparison between HR-TχZAB alloys did not only exhibit nearly equal elongation-to-failure with as-cast TχZAB alloys [21] but also much
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higher strengths than the as-cast TχZAB alloys. Fig. 2(a) demonstrated the differences of mechanical performances between hot-rolled and as-cast alloys with identical compositions. The relationships between the mechanical properties and Zr contents are shown in Fig. 2(b). Compared to as-cast TχZAB
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alloys [21], σ0.2 and σb of (HR-TAB, HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB) alloys increased by (50, 123, 192, 403, and 395 MPa) and (34, 171, 309, 436, and 400 MPa),
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respectively. Therefore, the hot-rolling deformation of TχZAB alloys undoubtedly enhanced the strength with almost unvaried elongation in comparison to parallel as-cast TχZAB alloys. According to previous studies [2,6,18,26-30], tensile strength and elongation of some frequently-applied structural Ti alloys were compared to current alloys and the results are listed in Fig 3. It could be easily observed that the mechanical properties of the as-prepared alloys outperformed those of published alloys, indicating the good combination of tensile strength and ductility required by structural materials. 3.2. Fracture surface morphologies The low-magnification SEM fracture morphologies of hot-rolled TχZAB alloys with different Zr contents are gathered in Fig. 4 from a macro perspective. The variation of fracture surfaces of post-rift 7
ACCEPTED MANUSCRIPT tensile specimens depicted typical trends. The fracture morphology of HR-TAB alloy specimen consisted mainly of large, deep and equiaxed dimples filled with minor dimples (Fig. 4(a)). The addition of Zr reduced and shallowed the formed dimples with continuous rippled pattern. As Zr content increased to 20 wt.%, the dimples shrunk, and the micro-pores caused by defects formed
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metastable dimples under the tensile stress took the dominant in fracture surface (Fig. 4(c)). The quasi-cleavage fracture morphology was then observed in Fig. 4(d) and 4(e), where facets, cracks and cleavage planes with mini-sized dimples became apparent on most of the surface. The comparison between (d) with (e) in Fig. 4 revealed that the fracture morphology of quasi-cleavage steps and rubbly
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surface of HR-T40ZAB alloy was more distinct than HR-T30ZAB alloy. In addition, the necking phenomenon gradually changed from obvious to slight then vanished. The combined elongation of
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HR-TχZAB alloys with fracture mode transformed from ductile to brittle, evidencing a decrease in alloy ductility.
3.3. The effects of phases on microstructural evolution and mechanical properties The DSC curves of the examined alloys are shown in Fig. 5 as a reference to phase-transition temperature. During heating, phase-transition temperature could be obtained by measuring the
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exothermic and endothermic variation of original TχZAB alloy specimens. Overall, only α phase existed until temperature reached onset-transition temperature. Between onset-transition and end-transition temperatures, both α and β phases were present. This indicated that once the
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onset-transition temperature was reached, α phase should start to transform into β phase. The phase transformation was complete once the temperature reached the end-transition temperature, after which
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only β phase existed. Obviously, phase-transition temperatures of the alloys decreased as Zr content rose from 0 wt.% to 40 wt.%. The onset temperature of transition from α→β phase declined from 954 ℃ to 776 ℃, and the end-transition temperature decreased from 1009 ℃ to 860 ℃. In Fig. 5, the dotted line represented 930 ℃ intersected with five DSC curves, where five intersection points located at: single α phase region, α→(α+β) dual-phase region, (α+β)→β dual-phase region, and two single β phase region of TAB, T10ZAB, T20ZAB, T30ZAB and T40ZAB, respectively. The microstructure could be tailored by varying the phase composition due to phase transformation [5]. Hence, 930 ℃ was selected as the rolling temperature to further study the effects of phase transformation and microstructure evolution on the mechanical properties of the alloys with various Zr contents. Increasing β-stabilizing alloy element contents would govern the martensitic structure of water 8
ACCEPTED MANUSCRIPT quenched Ti alloys containing β-phase, and the solid-state transformation which was diffusionless, i.e. martensitic, changes as the order: α′, then transformed to α′′, then transformed to ω, then transformed to β. [31-33]. Though phase transition temperature of Ti alloys decreased as Zr content increased according to binary equilibrium phase diagram of Ti-Zr and Fig. 5, the effect of Zr on β-phase stability
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of TχZAB alloy was estimated by theoretical calculations. The value of e/a could be used as reference for the quenched phase of Ti alloys according to valence electron theory. A higher value of e/a was acquired by calculation, and stronger stability was found for β phase [34]. Previous studies [16] reported that the martensitic phase of Ti alloys with e/a value below 4.07 should be α′ phase. The
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calculated e/a ratios of TχZAB alloys are listed in Table 2, and the results were all below 4.07. Besides തതതത law would be considered as another criterion for water തതതത-Md the valence electron theory, the Bo
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തതതത diagram of Ti alloys quenched phases of HR-TχZAB alloys [17]. Fig. 6 shows the extended തതത Boത-Md presenting α′, α″, ω and β four phase regions, and Table 2 lists the values of തതത Boത and തതതത Md. T10ZAB, T20ZAB, T30ZAB and T40ZAB alloys quenched from temperatures above the α transus temperature തതതത diagram. Here, the valence തതതത-Md were uniformly located in martensite α′ region (shadow area) of Bo തതതത law were appropriate for low-alloy Ti alloys with alloying element തതതത -Md electron theory and Bo
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contents of less than 20 at.% [16]. However, T40ZAB alloy appeared outside the valid interval. To yield reasonable and consummate theoretical calculations, the cohesive energies were calculated on the basis of Eq. (1). The cohesive energies of body-centered cubic (β phase) Ti-Al, Ti-Zr
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and Ti-Zr-Al supercells were 6.298, 6.441 and 6.350 eV atom-1, respectively. It is clear to contrast that the cohesive energy of Ti-Zr alloy was higher than those of Ti-Al and Ti-Zr-Al alloys, indicating that
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the average chemical bond between atoms in β-phase was stronger in Ti-Zr alloy than in Ti-Al or Ti-Zr-Al alloys. Here, the mass ratio of the alloying element Al remained invariant whereas the atomic ratio increased. Thus, the effect of Al on stability of β-phase in TχZAB alloys should also be taken into account. The ∆Ecoh referring to the difference between the cohesive energies of two body-centered cubic supercells could be calculated as follows: ∆Eୡ୭୦ (a) = Eୡ୭୦ (TiZrAl) − Eୡ୭୦ (TiAl)
(2)
∆Eୡ୭୦ (b) = Eୡ୭୦ (TiZrAl) − Eୡ୭୦ (TiZr) where ∆Ecoh(a) is 0.052 and ∆Ecoh(b) is -0.092. |∆Eୡ୭୦ (a)| represents the degree of Zr on the stability of β-phase, and |∆Eୡ୭୦ (b)| represents the degree of Al in reducing the β-phase stability. It would be easy to deduce that |∆Eୡ୭୦ (b)| − |∆Eୡ୭୦ (a)| > 0, meaning that the degree of Al on reducing β-phase 9
ACCEPTED MANUSCRIPT stability was much stronger than the degree of Zr on β-phase stability. All trends obtained by the above theoretical calculations were confirmed experimentally and explained below. Fig. 7 shows the XRD patterns of HR-TχZAB alloys for their actual phase compositions. The diffraction peaks only matched peaks of α or α′ phases since the crystal structures of both α and α′
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consisted of hexagonal close-packed (hcp). The XRD patterns indicated that the phases of HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys quenched from α+β dual-phase region, and the single β region were merely formed of primary α or martensite α′ phase with no residual β phase or other intermetallic phases. In other words, even at rolling and quenching temperatures above
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β transus temperature, only the β→α′ phase transition occurred during the quenching process. This was in accordance with the estimated theoretical calculations. When combined with DSC curves, the
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analysis revealed that the volume fraction of α′ phase in HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys gradually increased after hot rolling and quenching. According to previous studies [9], α′ martensite is a hard phase compared to other common phases in Ti and Zr alloys, such as α, β, and ω phases. Therefore, the formation of α′ martensite would increase both the yield strength and tensile strength but reduce ductility. Moreover, the martensitic transformation was found quite effective
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in strengthening Ti- and Zr-base alloys. Unlike the Fe-C system, the hcp martensites in Ti- and Zr-based alloys were not supersaturated with interstitial elements. In general, β-phase region quenching would produce martensitic α′ structure and slow cooling from the β-phase region to form
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Widmanstatten α structure. This strengthening effect owed to their substructures. For isomorphous martensitic α′ and Widmanstatten α, the martensitic α′ structure contained more closely spaced
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interfaces separating neighboring lamellae or twins with high density of dislocation compared to Widmanstatten α structure with the same composition [35]. Here, the microhardness of HR-TχZAB alloys increased as volume fraction of α′ phase rose. Z. Tarzimoghadam et al. [33] reported that phase transition in transformation kinetics and kinematics would enable the design of complex microstructures with features spanning across broad range of length scaling. In Fig. 7(right), the width of the diffraction peaks broadened after hot rolling owing to the synergistic effect of increased micro residual stress and grain refinement caused by large deformation and reduced phase-transition temperature. Previous study [21] reported that solute element Zr-induced lattice distortion changed the lattice parameter of TχZAB alloys from XRD results. As Zr content increased, the variation of axial ratio “c/a” 10
ACCEPTED MANUSCRIPT influenced the slip mechanism of TχZAB alloys, which further caused the decrease of ductility. Meanwhile, addition of alloying element Zr led to increasing degrees of saturation, which also changed the distance between the atoms and resulted in a change of the Young's modulus. Young's modulus is an intrinsic property of materials and is determined by the bonding force between atoms [36]. This
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bonding force is related to the addition of alloying elements, which can be affected by the crystal texture and any change in the distances between the atoms. In current study, the Young's modulus decreased as Zr increased. For the TχZAB alloys with the same phase, the lattice parameter “a” increased following the increased Zr content owing to the larger atomic radius of Zr than Ti [21]. The
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aforementioned result was confirmed by the XRD peaks in turn shifting towards low angle orientation. Therefore, Young's modulus mainly depended on chemical composition, i.e., Zr content.
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3.4. The effects of microstructures on the mechanical properties
The microstructures of the as-prepared TχZAB alloys visibly changed after hot-rolling when compared to those of as-cast TχZAB alloys reported in our previous study [21]. At the same hot-rolled conditions, the microstructure is dependent on the β transition temperature [37]. Meanwhile, at constant composition conditions, the microstructure would be a critical factor influencing the mechanical
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properties. For example, texture is a factor that affects mechanical properties. As we know, this influence is pronounced for hcp metals such as α-Ti or α-Zr alloys due to the high volume fraction of the inherently anisotropic hexagonal α-phase [38]. Fig. 8 depicts the metallographs of HR-TχZAB
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alloys. Obviously, the lamellar α-phase of TAB and T10ZAB were distorted or broken after the hot-rolling process (Fig 8(a) and 8(b)). The microstructure of the as-cast TAB and T10ZAB alloys
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were completely filled with crisscross slender α lamella. After hot-rolling, the typical basketweave morphologies underwent severe changes. Jiang et al. [7] suggested mechanisms behind the distortion and break-up of α-phase lamella during hot deformation process. The localized shear and rotation of α lamella would occur with misorientation across the shear band to reach approximately 20°. Simultaneously, the low and high angle boundaries both formed across α lamella, where misorientation angles varied from a few degrees to up to 30°. For our proposed alloys, these mechanisms would also be appropriate. The distorted or fractured lamellae would increase the grain boundary area hindered the dislocation motion, which caused the grain boundary strengthening. The microstructure evolution of HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys appeared quite similar (Fig 8(c), 8(d) and 8(e), respectively). The metallographs revealed that the microstructures was characterized by fine α′ 11
ACCEPTED MANUSCRIPT martensite crystals within the original β grain boundaries. The aforementioned morphology with acicular plate α′-phase would be typical of martensite in titanium alloys [33]. For HR-T20ZAB alloy, considering the rolling temperature nearby the β transition temperature, distorted original α phases were hardly detected. In addition, the α′-lamella structure became gradually finer in HR-T20ZAB,
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HR-T30ZAB and HR-T40ZAB alloys. The microstructural details of HR-T10ZAB, HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys were revealed by TEM bright field micrographs (Fig 9). A large number of dislocations could clearly be observed in the deformed lamellae (Fig 9(a)), suggesting that HR-T10ZAB alloy suffered from
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significant deformation during the hot-rolling process. The dislocation density was so high with severe dislocation pile-up in the grain boundaries that dislocation mobility was restricted inducing an increase
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in strength. In α-Ti alloys (hcp), the principal deformation mode corresponded to a-type dislocations (a/3<112ത0>), mainly gliding in the prismatic planes. The deformation was controlled by the motion of screw segments, in which atomistic simulations showed lattice friction due to three-dimensional core structure spread in prismatic as well as in basal and pyramidal planes [1,3]. Moreover, the high lattice friction caused by core structure of the dislocations had a major impact on the strengthening in lamellar
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microstructures [1]. Hence, the mechanism of dislocation strengthening was responsible for improving the mechanical properties of the proposed hot-rolled alloys in contrast to as-cast alloys. Fig 9(b-d) illustrated that the width of α′ lamellae decreased dramatically with few apparent dislocations. The
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selected-area electron diffraction pattern in Fig 9(d) further confirmed the phase composition. According to Fig 5, β transition temperature decreased with the increased Zr content. During quenching,
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β phase was rapidly converted into metastable α′ martensite phase. Furthermore, under the same quenching temperature and cooling conditions, α′ martensite phase of the alloy with higher Zr content had insufficient time to grow thicker, resulting in thinner lamellar thickness and spacing. As Zr content increased, the microstructure of the hot-rolled alloys increasingly refined. Admittedly, the strength of polycrystalline materials like TχZAB alloys followed the Hall-Petch empirical relationship depending on the fineness of grain: σ=σ0+kd-0.5, where σ represents the yield strength of the TχZAB alloys. Previous investigations [3,14,39] reported that full acicular α′ martensite microstructures could effectively have high strength in Ti alloys (e.g. in Ti-V, Ti-V-Sn and Ti-V-Al alloys) via controlling the size and distribution of α′ lamellae. Here, Hall-Petch effect was controlled by thickness of the α/α′ lamellae (i.e. the d). Fig 10(a) shows the average thickness of α/α′ lamellae and the thickness 12
ACCEPTED MANUSCRIPT distribution. Average grain size was measured according to the semiautomatic image analysis (ASTM: E1382-97(2010)). In order to obtain accurate data, thirty lamellae were measured from each sample in this study. By using Hall-Petch equation, it could be found that the decreased grain dimensions gave rise to enhanced strength (Fig 10(b)). Fine crystal strengthening mainly depended on hindering of the
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dislocation slip by the grain boundaries [40]. On the one hand, reducing the grain size with increasing density of the grain boundaries in the same cross-sectional area increased the number of dislocations piling up. The severe stress concentration occurred on grain boundary. On the other hand, the different grain orientations in both sides of the grain boundary led to increased sliding resistance near the grain
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boundary during the movement of the dislocation obstacle in the polycrystal. Thus, the propagative slip band could not directly enter the neighbouring grains resulting in accumulated dislocation at the grain
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boundary. Dislocation pile-up substantially improved the strength of alloy. Tarzimoghadam et al. [33] also reported similar results in Ti-4Mo alloy. Considering a mass of dislocations and dislocation cells formed in HR-TAB, HR-T10ZAB and HR-T20ZAB alloys, the increase in strength was more significant than HR-T30ZAB and HR-T40ZAB alloys. The microstructures of TχZAB alloys under the hot-rolled and water-quenched condition would be stable at service temperature. Therefore, the
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microstructure evolution would be a key to improve the mechanical properties of TχZAB alloys, and the mechanisms could most likely be generalized to other Ti alloys.
4. Conclusions
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For further studying and developing the potential structural Ti alloys, the mechanical properties, fracture characteristics, phase composition and microstructure evolution of HR-TχZAB alloys were
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investigated and the effects of phase composition and microstructure on mechanical properties were discussed. The following conclusions could be drawn: (1) The ternary yield strength, ultimate strength and microhardness obviously increased as Zr content increased. The HR-T40ZAB alloy revealed the highest tensile strength of 1535 MPa at 6.06% elongation. The comparison between HR-TχZAB alloys with previous as-cast TχZAB alloys with identical compositions suggested that hot-rolled alloys did not only exhibit nearly equal elongation with the as-cast alloys but also much higher strength than as-cast alloys. (2) The fracture morphologies revealed typical transitional processes from ductile fracture surface to quasi-cleavage fracture surface as the Zr content increased. Also, the necking phenomenon 13
ACCEPTED MANUSCRIPT gradually weakened. The transformation of fracture mode was in accordance with the decrease in elongation of HR-TχZAB alloys series.
(3) The theoretical calculations of phase composition were consistent with the experimental data, where only α/α′ phases were obtained in HR-TχZAB alloys. And the strength increased as the
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volume fraction of α′ phase rose. (4) The lamellae of HR-TAB and HR-T10ZAB alloys were distorted or broken after hot-rolling, forming large numbers of intragranular dislocations. The microstructure of HR-T20ZAB, HR-T30ZAB and HR-T40ZAB alloys was completely filled with acicular α′ martensite, which
increased as the thickness of lamellae decreased.
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Acknowledgments
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gradually became finer as Zr content increased. Additionally, the strength of HR-TχZAB alloys
This work was supported by the NSFC (Grant no. 51531005/51671166/51434008) and Postdoctoral Science Foundation of Hebei Province of China (Grant no. B2017003008).
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ACCEPTED MANUSCRIPT Figure captions Fig. 1. Stress-strain curves and tensile specimen dimensions of HR-TχZAB alloys. Fig. 2. Changes of mechanical properties: (a) the difference values of mechanical performances between hot-rolled and as-cast alloys with identical composition, (b) the variation trend of
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mechanical properties of series HR-TχZAB alloys. Fig. 3. The mechanical properties of the current HR-TχZAB alloys and other available Ti alloys [2,6,18,26-30].
χ=30 and (e) χ=40. Fig. 5. DSC curves of TχZAB alloys.
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തതതത diagram of Ti alloys [15]. Fig. 6. The expanded തതത Boത-Md
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Fig. 4. SEM fracture morphologies of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d)
Fig. 7. XRD patterns of the HR-TχZAB alloys.
Fig. 8. Optical micrographs of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
Fig. 9. TEM bright field microstructures of the HR-TχZAB alloys (wt.%): (a) χ=10, (b) χ=20, (c) χ=30
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Fig. 10. Grain size analysis: (a) variation of the average thickness of α/α′ lamellae and thickness
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Fig. 1. Stress-strain curves and tensile specimen dimensions of HR-TχZAB alloys.
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Fig. 2.Changes of mechanical properties: (a) the difference values of mechanical performances between
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Fig. 3. The tensile properties of the current HR-TχZAB alloys and other available Ti alloys
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Fig. 4. SEM fracture morphologies of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
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Fig. 5. DSC curves of TχZAB alloys.
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Fig. 7. XRD patterns of the HR-TχZAB alloys.
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Fig. 8. Optical micrographs of the HR-TχZAB alloys (wt.%): (a) χ=0, (b) χ=10, (c) χ=20, (d) χ=30 and (e) χ=40.
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Fig. 9. TEM bright field microstructures of the HR-TχZAB alloys (wt.%): (a) χ=10, (b) χ=20, (c) χ=30
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Fig. 10. Grain size analysis: (a) variation of the average thickness of α/α′ lamellae and thickness distribution, (b) variation of σ0.2 with the thickness of α/α′ lamellae Hall-Petch equation.
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ACCEPTED MANUSCRIPT Table Captions Table 1 Mechanical properties of the HR-TχZAB alloys.
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തതതത and e/a values of TχZAB alloys. തതതത, Md Table 2 Nominal compositions (wt.% and at.%) and Bo
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ACCEPTED MANUSCRIPT Table 1 Mechanical properties of the HR-TχZAB alloys. E (GPa)
σ0.2 (MPa)
σb (MPa)
εf (%)
HV
HR-TAB
161
682
775
15.59
286
HR-T10ZAB
143
860
1028
10.61
318
HR-T20ZAB
122
1062
1290
8.43
391
HR-T30ZAB
120
1294
1459
6.99
406
HR-T40ZAB
120
1388
1535
6.06
442
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ACCEPTED MANUSCRIPT തതതതത and e/a values of TχZAB alloys. തതതത, ݀ܯ Table 2 Nominal compositions (wt.% and at.%) and ܤ wt.%
Alloys
at.%
തതതത ܤ
തതതതത ݀ܯ
e/a
0.02
2.764
2.429
3.93
7.22
0.02
2.779
2.455
3.93
11.21
7.58
0.02
2.795
2.482
3.92
74.27
17.72
7.99
0.02
2.813
2.513
3.92
65.57
24.96
8.44
0.03
2.804
2.523
3.87
Zr
Al
B
Ti
Zr
Al
B
HR-TAB
95.995
0
4
0.005
93.09
0
6.89
HR-T10ZAB
85.995
10
4
0.005
87.43
5.34
HR-T20ZAB
75.995
20
4
0.005
81.18
HR-T30ZAB
65.995
30
4
0.005
HR-T40ZAB
55.995
40
4
0.005
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(5) The microstructure is refined as the Zr content increased.
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(4) Only α/α′ phases are obtained in hot-rolled TiZrAlB alloys.