- Email: [email protected]
Microstructures and mechanical properties of Zr–Al binary alloys processed by hot-rolling
Materials Science & Engineering A xxx (xxxx) xxx
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructures and mechanical properties of Zr–Al binary alloys processed by hot-rolling X. Zhang a, B. Zhang a, S.G. Liu a, C.Q. Xia b, X.Y. Zhang a, M.Z. Ma a, R.P. Liu a, * a b
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, China School of Materials Science and Engineering, Hebei University of Technology, Tianjin, 300130, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Zr–Al alloys Microstructure Mechanical properties Solid solution strengthening Dynamic precipitation
The microstructures and mechanical properties of Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) that underwent the hot-rolling at 1143 K and the subsequent water-quenching were investigated. Microstructural analyses show that the pure Zr, Zr–3Al and Zr–6Al alloys mainly consisted of the α and α0 phases and exhibited bimodal micro structures. The Zr–9Al and Zr–12Al alloys consisted of the α, β, Zr2Al and Zr3Al phases and exhibited equiaxed microstructures. Tensile tests show that, with the increase in Al content, the yield and ultimate strength of the alloys increased from 324.9 MPa and 443.8 MPa (pure Zr) to 665.2 MPa and 787.6 MPa (Zr–12Al), respectively. Meanwhile, the elongation to failure initially decreased from 17.3% (pure Zr) to 10.3% (Zr–6Al) and subse quently increased to 19.5% (Zr–12Al). Analyses show that the significant enhancement of yield strength resulted from the solid solution strengthening by Al. In addition, dynamic precipitation occurred in the Zr–9Al and Zr–12Al alloys. The precipitated Zr3Al particles resulted in the obvious strain hardening behavior of the Zr–9Al and Zr–12Al alloys. The high strain hardening rate and elongation at the uniform plastic deformation stage of the Zr–12Al alloy resulted in the simultaneous enhancement of the ultimate strength and ductility.
1. Introduction Given their low thermal neutron absorption cross section, excellent corrosion resistance, good biocompatibility and low magnetic suscep tibility, zirconium (Zr) and its alloys have been widely used in nuclear, chemical and medical industries [1–3]. In addition, Zr and its alloys also show many other excellent properties such as low density, low expan sion coefficient and high melting point [4]. Therefore, Zr has great po tential application as structural materials in aerospace industry. However, the mechanical properties of Zr are insufficient to meet the requirements of structural materials. The current studies of Zr alloys mainly focus on the properties such as thermal neutron absorption cross section, corrosion resistance, hydrogen embrittlement and high-temperature creep for nuclear use [5–8]. Few studies have focused on the improvement of mechanical properties of Zr alloys. Therefore, it is necessary to develop novel Zr alloys to achieve improved mechanical properties to broaden their applications. Titanium (Ti) and Zr are transition metals belonging to Group ⅣB of the periodic table of elements. The similarity in the dispositions of the outer electrons is responsible for the similarities in some of the chemical and physical properties of Zr and Ti [9]. Zr exists as the body-centered
cubic (BCC) β phase at the high temperature above 1139 K–2128 K and atmospheric pressure. Under slow cooling, the high-temperature β phase will transform into the close-packed hexagonal (HCP) α phase by diffusive transformation. Under rapid cooling, the high-temperature β phase will transform into the HCP-α0 phase by martensitic trans formation. Aluminium (Al) has a good solid solution strengthening ef fect in Ti and is a widely used alloying element in Ti alloys [10]. In addition, Al has considerable solid solubility in the high-temperature α-Zr and the solubility can reach 10.5 at.% at 1213 K [11]. The study by Jiang et al. shows that the ultimate strength of Ti–25Zr at.% alloy increased by 41% after the addition of 15 at.% Al [12]. The study by Feng et al. shows that the addition of Al obviously caused the lattice distortion of α-Zr and had an obvious strengthening effect in Zr–Be al loys [13]. In addition, the ordered Zr3Al phase with the face-centered cubic (FCC) structure of L12 type can be precipitated to strength the Zr–Al alloys (Al content is less than 25 at.%) [11]. Therefore, Al is considered to be an ideal strengthening element in Zr. Kematick et al. analyzed the heats of formation of the Zr–Al intermetallic compounds and the solid solution of Al in BCC-Zr using the Knudsen cell mass spectrometric technique [14]. Mukhopadhyay et al. investigated the formation of a D019 phase in Zr–Al martensites and the precipitation of
* Corresponding author. E-mail address: [email protected] (R.P. Liu). https://doi.org/10.1016/j.msea.2019.138723 Received 16 September 2019; Received in revised form 18 November 2019; Accepted 21 November 2019 Available online 21 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: X. Zhang, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138723
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 1. OM micrographs of the as-cast (a) pure Zr, (b) Zr–3Al, (c) Zr–6Al, (d) Zr–9Al and (e) Zr–12Al alloys.
Zr3Al in a Zr–4.6 wt% Al martensite [15,16]. Schulson investigated the tensile and corrosion behavior of ordered Zr3Al-based alloys [17]. Howe et al. conducted the transmission electron microscopy investigations of ordered Zr3Al [18]. Alatalo et al. studied the stability of Zr–Al com pounds [19]. However, the previous studies on Zr–Al system have focused on Zr–Al intermetallic compounds. The study of microstructures and mechanical properties of Zr–Al alloys in the hypostoichiometric composition range (α þ Zr3Al phase region) has not yet been system atically conducted. Hot-rolling is a widely used method for forming technique and microstructure adjustment of metallic materials. In addition, the grain boundary α (αGB) and Widmanst€ atten α (αW) formed in the as-cast Zr and Ti alloys have negative effects on the mechanical properties such as strength and ductility [20]. Hot-rolling process is often used to eliminate the αGB and αW and obtain equiaxed or duplex microstructure to enhance the mechanical properties of alloys. The study by Chao et al. shows that
an equiaxed ultrafine-grained microstructure was successfully produced in a Ti–6Al–4V alloy through optimum hot-rolling process of a martensitic starting microstructure and the strength and plasticity of the alloy were simultaneously enhanced [21]. In this work, the microstructures and mechanical properties of the hot-rolled Zr–Al binary alloys in the hypostoichiometric composition range (α þ Zr3Al phase region) were investigated. The strengthening mechanisms, as well as the microstructure evolution were discussed. The findings of this work may provide some new insights to develop novel Zr alloys with enhanced mechanical properties. 2. Experimental procedure Industrially pure Zr (Zr þ Hf � 99.5 wt%) and industrially pure Al (99.9 wt%) were used to prepare the needed alloys in this study. Zr–xAl alloys containing different Al contents (x ¼ 0, 3, 6, 9 and 12 at.%) were 2
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Table 1 XRD diffraction angles of the α/α0 -{1011} lattice planes of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Alloy no.
Pure Zr
Zr–3Al
Zr–6Al
Zr–9Al
Zr–12Al
Angle (degree)
30.401
31.941
32.061
32.199
32.108
Fig. 2. XRD patterns of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
prepared to investigate the effects of Al contents on the microstructures and mechanical properties. The ingots (approximately 85 g) of the investigated alloys were melted via a non-consumable tungsten-elec trode arc furnace on a water-cooled copper crucible in a protective argon atmosphere. Pure Ti was used as a getter. The ingots were turned over and re-melted five times to achieve homogeneity. The ingots were heated to 1143 K and held for 40 min, subjected to multiple rolling passes and water-cooled to room temperature. The reduction of each pass was 2.16 mm. The samples were held in the furnace for 5 min be tween each pass. The total deformation reached 60%. All of the ingots were rolled under the same conditions. X-ray diffraction (XRD) was used to determine the phases of the al loys using an X-ray diffractometer (Rigaku D/MAX 2500/PC) with Cu Kα radiation. The microstructures of the alloys were examined using optical microscopy (OM, Zeiss Ax overt 200MAT), scanning electron micro scopy (SEM, Hitachi S-3400 N) and transmission electron microscopy (TEM, FEI Talos F200X). Specimens for SEM observation were me chanically polished and chemically etched. The etchant consisted of distilled water, nitric acid and hydrofluoric acid (85:10:5 by volume). TEM foils were prepared through twin-jet electrochemical polishing in a solution containing perchloric acid and methanol (10:90 by volume) at 238 K and 15 V. The tensile properties of the alloys were tested using an Instron 5982 testing machine at an initial strain rate of 5 � 10 4 s 1. An extensometer with a 12.5 mm gauge length was mounted on the speci mens to measure the strain. The specimens for the tensile tests were bone-shaped plates with an original gauge length of 20 mm and crosssectional dimensions of 3 mm � 2 mm. At least three specimens were tested for repeatability, and the average values were reported.
Fig. 3. α/α0 -Zr lattice parameters of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
results with the selected area electron diffraction (SAED) results, the phases of the alloys were identified. The pure Zr and Zr–3Al alloy con sisted of the α and α0 phases. With the increase in Al content, Zr–Al intermetallic compounds formed in the Zr–6Al, Zr–9Al and Zr–12Al al loys. The Zr–9Al and Zr–12Al alloys consisted of the α, β, Zr2Al (tI12) and Zr3Al (cP4) phases. According to the relative intensity of the diffraction peaks, the phase fraction of Zr2Al of the Zr–9Al and Zr–12Al alloys were very low. The phase fraction of Zr3Al of the Zr–12Al alloy increased substantially compared with that of the Zr–9Al alloy. The simultaneous precipitation of the Zr2Al and Zr3Al phases in the Zr–9Al and Zr–12Al alloys may be related to their similar formation enthalpy [19]. The Zr–6Al alloy consisted of the α, α0 and a small amount of the compound phases. However, the SAED results show that the measured diffraction d-spacing of the compound of the Zr–6Al alloy did not match any of the Zr–Al compounds in the International Centre for Diffraction Data (ICDD) pdf card library. The unknown compound phase may be an unreported allotrope of the Zr3Al or Zr2Al phase. Table 1 shows the XRD diffraction angles of the α/α0 -{1011} lattice planes of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Obviously, the addition of Al led to the shift of the α/α0 -{1011} diffraction peak toward high angles. The atomic radius of Al (1.82 Å) is smaller than that of Zr (2.16 Å). Al atoms replaced the positions of Zr atoms in the α/α0 -Zr lattice, resulting in the decrease in the α/α0 -Zr lat tice parameters. The α/α0 -Zr lattice parameters of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) are shown in Fig. 3. Corresponding to the change of the α/α0 -Zr lattice parameters, with the increase in Al content from 0 at.% to 9 at.%, the shift of the α/α0 -(1011) diffraction peak toward high angles increased. However, the α/α0 -(1011) diffraction peak of the Zr–12Al alloy shifted toward an angle lower than that of the Zr–9Al alloy. The exsolution of the α phase of the Zr–12Al alloy was considered to result in the decrease in Al content and the increase in the lattice parameters. Figs. 4a, b, 4c, 4d and 4e show the SEM micrographs of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Al stabilizes α-Zr thermody namically [11]. With the increase in Al content, the α→β transformation temperatures of the Zr–xAl alloys increased. It can be seen from the SEM micrographs that the pure Zr, Zr–3Al and Zr–6Al alloys consisted of the thick primary α (αp) and fine α0 grains. The Zr–9Al and Zr–12Al alloys
3. Results and discussion 3.1. Phase and microstructure Fig. 1 shows the OM micrographs of the as-cast Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). It can be seen that all the as-cast Zr–Al alloys consisted of the basketweave microstructure. The thick αGB lay along the primary β grain boundary. In addition, Widmanst€ atten structures were also observed in the as-cast Zr–Al alloys. Fig. 2 shows the XRD patterns of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Identifying the phases of the alloys using only XRD was difficult because of the variety of Zr–Al compounds and their crystal structures. This situation was aggravated by the low intensity and overlap of some diffraction peaks. From the combination of the XRD 3
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 4. SEM micrographs of the hot-rolled (a) pure Zr, (b) Zr–3Al, (c) Zr–6Al, (d) Zr–9Al and (e) Zr–12Al alloys; (f) Average widths of the αp grains of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
Fig. 5. (a) TEM micrograph of the hot-rolled Zr–6Al alloy; (b) STEM micrograph of the hot-rolled Zr–9Al alloy.
consisted of the equiaxed αp grains. The αp grains in the pure Zr and Zr–6Al alloy did not get dynamic globularization fully and exhibited a high length–width aspect ratio. The αGB grains also remained after hot-rolling. Fig. 4f shows the average widths of the αp grains of the hot-rolled Zr–xAl alloys. According to Fig. 4f, the average width of the αp grains of the Zr–12Al alloy was about 1.7 μm and was obviously smaller than those of the pure Zr, Zr–3Al, Zr–6Al and Zr–9Al alloys (2.7 μm–4.2 μm). Fig. 5a shows the TEM micrograph of the Zr-6Al alloy. It can be seen that the unknow compound phase precipitated from β phase and was surrounded by α0 phase, which was transformed from β phase. Fig. 5b shows the scanning transmission electron microscopy (STEM) micrograph of the Zr–9Al alloy. The diffusion paths provided by the grain boundaries can be observed. Meanwhile, the obvious contrast between the boundary area and the α grain interior indicated that the formation of the compound phase led to the exsolution of the α phase along the grain boundaries. Fig. 6a and 6b shows the SEM and STEM micrographs of the Zr–12Al alloy. Figs. 6c and 6d shows the element
mapping maps of Al and Zr corresponding to Fig. 6b. A large number of Zr3Al particles distributed at the grain boundaries can be observed. The high contrast around the compound particles and the α grain boundaries indicates that the Al content in these areas was relatively lower than that inside the α grains. Fig. 6e shows the TEM micrograph of the Zr–12Al alloy. The Zr3Al and β phases located at the α grain boundaries were identified by SAED in Fig. 6f and 6g, respectively. It can be seen from Fig. 6g that there were some diffuse intensity distributions in the diffraction patterns. It has been well documented that the occurrence of diffuse intensity indicates the tendency for the ω phase transformation [9]. Fig. 6h shows the energy disperse spectroscopy (EDS) of the β phase in Fig. 6e. It indicates that the Al content of the β phase was 7.5 at.%. Nucleation of the Zr3Al particles was observed to occur exclusively at the α grain boundaries, combined with the exsolution phenomenon of the α phase, indicating that the dynamic precipitation process occurred in the Zr–9Al and Zr–12Al alloys during the hot-rolling. Dynamic pre cipitation behavior has been reported in steel, Al and magnesium (Mg) 4
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 6. (a) SEM and (b) STEM micrographs of the hot-rolled Zr–12Al alloy; Element mapping maps of (c) Al and (d) Zr corresponding to (b); (e) TEM micrograph of the hot-rolled Zr–12Al alloy; SAED patterns of the (f) Zr3Al and (g) β phases in (e); (h) EDS of the β phase in (e).
alloys but seldom in Ti and Zr alloys [22–25]. Mukhopadhyay et al. investigated the static precipitation behavior of Zr3Al in a Zr–4.6 wt% Al martensite and concluded that the precipitation of the equilibrium Zr3Al (cP4) phase in a supersaturated, martensitic α-Zr(Al) matrix occurred in a discontinuous manner and the equilibrium-transformed structure comprised duplex cells, each containing a lamellar distribution of the α-Zr(Al) and Zr3Al phases [26]. In this case, however, the morphology of Zr3Al was quite different from cell growth of static precipitation. It is well known that dynamic precipitation is accelerated by deformation, compared with the precipitation rate measured during isothermal transformations in undeformed structures at the same temperature [27–29]. This acceleration of precipitation is believed to be due to the introduction of favourable nucleation sites by the deformation process [22]. In addition, the dynamic globularization process led to the refinement of α grains and provided a large number of nucleation sites for the precipitation of the Zr3Al particles. Meanwhile, the grain boundaries that contained a large number of vacancies and dislocations
provided diffusion paths for atoms [30,31]. The Zr3Al particles distrib uted on the α grain boundaries pinned the grain boundaries, thus inhibiting the dynamic coarsening process of the grains and resulting in the further refinement of the α grains. Therefore, with the increase in the phase fraction of Zr3Al, the α grains of the Zr–12Al alloy obviously exhibited a finer size than those of the Zr–9Al alloy. The dynamic precipitation resulted in the exsolution of the α phase near the Zr3Al particles and grain boundaries. The decrease in Al content of the local α phase resulted in the decrease in its α→β transformation temperature. Thus, the α phase near the grain boundaries and Zr3Al particles transformed to β phase at 1143 K. However, it is well known that Al is an α stabilizer and broadens α phase region [11]. It seems that the athermal product of the high-temperature Al-containing β phase should be martensite α0 . Banerjee et al. have studied the transformation sequence in the Zr–27 at.% Al alloy during rapid quenching from the liquid state [32]. This abnormal β phase was also observed in the liquid-quenched Zr–27 at.% Al alloy. In addition, the study by Jiang
Fig. 7. (a) Engineering stress–strain curves and (b) strain hardening rate of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). 5
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
microstructure with a high strain hardening capacity inducing a desir able increase in both strength and ductility [34]. The pure Zr, Zr–3Al and Zr–6Al alloys mainly consisted of the αp and α0 phases. α0 phase has the same crystal structure (close-packed hexag onal structure) with α phase which is the equilibrium transformation product of high-temperature β phase. In general, the thick αp is considered to be softer than the fine α0 . According to the deformation theory of multiphase alloy mentioned above, it is considered that the plastic deformation started in the αp . Therefore, the relative deviation of the yield strengths ðΔσx Þ of the pure Zr and the Zr–xAl (x ¼ 3, 6, 9 and 12) alloys are mainly affected by two factors. One factorðΔσ1 Þ is the strength change caused by the change in Al content. Another factor ðΔσ 2 Þ is the strength change caused by the size change of the αp grains. Thus, Δσ x can be calculated as following:
Table 2 Tensile properties of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Alloy no.
σ0.2 (MPa)
σb (MPa)
Е (%)
Pure Zr Zr–3Al Zr–6Al Zr–9Al Zr–12Al
323.9 432.3 537.6 632.5 665.2
443.8 520.5 610.6 694.1 787.6
17.3 11.9 10.3 16.1 19.5
et al. shows that, with the increase in Al content, the athermal product of the β-quenched Ti–40Zr–xAl (at.%) alloys changed sequentially as fol lows: α0 →α’0 [33]. When the Al content reached 15 at.%, the β phase was retained to room temperature completely. Although the detailed for mation mechanism of this abnormal β phase remains to be studied, Al as an α stablizer does stabilize β phase under certain conditions.
Δσ x ¼ Δσ 1 þ Δσ 2
(1)
The solid solution strengthening of alloys is known to arise from the elastic interaction between the strain field of a solute and that of a dislocation, and the strength is related to the concentration of solute atoms [38,39]. Fleischer [39] reported that the relationship between the concentration of solute atoms and the strength in binary metallic ma terials could be described as following:
3.2. Mechanical properties Fig. 7a shows the engineering stress–strain curves of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) and related parameters are shown in Table 2. With the increase in Al content, the yield (σ0.2) and ultimate (σ b) strength increased from 324.9 MPa (pure Zr) and 443.8 MPa (pure Zr) to 665.2 MPa (Zr–12Al) and 787.6 MPa (Zr–12Al), respectively. Meanwhile, the elongation to failure (E) initially decreased from 17.3% (pure Zr) to 10.3% (Zr–6Al) and subsequently increased to 19.5% (Zr–12Al). Compared with those of the pure Zr, the yield and ultimate strength of the Zr–12Al alloy increased about 105% and 177%, respectively. Meanwhile, the elongation to failure of the Zr–12Al alloy even exceeded that of the pure Zr. In addition, the yield strength of the Zr–12Al alloy only increased about 33 MPa compared with that of the Zr–9Al alloy, while the ultimate strength increased about 120 MPa. It can be inferred that a large number of precipitated Zr3Al particles on the grain boundaries had an obvious effect on the ultimate strength rather than the yield strength of the Zr–12Al alloy. Fig. 7b shows the strain hardening rate (dσT/dET) at the uniform plastic deformation stage of the Zr–xAl alloys. It can be seen that the strain hardening behavior of the Zr–9Al and Zr–12Al alloys was more obvious than that of the pure Zr, Zr–3Al and Zr–6Al alloys. The elongations of the Zr–9Al and Zr–12Al alloys at the uniform plastic deformation stage were significantly greater than those of the other three alloys. The mechanical properties of alloys are related to their compositions and microstructures. Recent studies show that the presence of a me chanical contrast between the two phases in the heterogeneous mate rials, such as dual-phase steels will cause obvious strain hardening behavior [34,35]. When deforming such a composite microstructure, the softer phase is the first to deform plastically, while the harder remain elastic. After a certain amount of deformation, the initial apparent strain hardening is enhanced when compared to the strain hardening of the softer phase. The difference in behavior between soft and hard compo nents of the microstructure will sustain this effect as the soft phases will deform more than hard phases. A dual-phase microstructure thus naturally induces an additional strain hardening when compared to the individual constituting phases. In this case, the obvious strain hardening behavior of the Zr–9Al and Zr–12Al alloys was caused by the Zr3Al particles which acted as the hard phase and obstructed dislocation motion during deformation. This simultaneous enhancement of strength and ductility induced by high strain hardening has been reported in several studies. Lu et al. reported that high strain hardening which was caused by the presence of ordered interstitial complex resulted in the simultaneous enhancement of strength and ductility of the oxygen-doped TiZrHfNb high-entropy alloy [36]. Han et al. reported that high strength and ductility of the Zr–2.5Nb alloy with a high strain hardening was achieved by randomly oriented α/β phase boundary strengthening [37]. Formanoir et al. reported that a sub-transus thermal treatment followed by water quenching generated a dual-phase α þ α0
σ ∝C2=3
(2)
where σ is the yield strength and C is the concentration of the solute atom. However, the expression was modified by She [40] and is described as follows: Δσ ∝C2=3
(3)
where Δσ is the increase in the yield strengths caused by solute atoms. In general, lattice parameters will change monotonically as solute is added in solution and the empirical expression is described as following [41, 42]: ax ¼ a0 þ k1 C
(4)
where ax is the lattice parameter of the α phase in the Zr–xAl alloys (x ¼ 3, 6, 9 and 12 at.%); a0 is the lattice parameter of the α phase in the pure Zr; k1 is a constant. The expression can be revised as follows: C ¼ Δa=k1
(5)
where Δα is the relative deviation of the lattice parameter a of the α in the pure Zr and the Zr–xAl (x ¼ 3, 6, 9 and 12) alloys. In combination with the current work, the relational expression (2) can be revised as follows: Δσ 1 ∝Δα2=3
(6)
According to the Hall–Petch relation, the yield strength is related to the grain size by the equation [43]:
σ ¼ σ0 þ kDGB1=2
(7)
where σ0 and k are constants; DGB is the grain size. The Hall–Petch slope k for slip in Zr is 8 MPa mm1/2 [44]. Therefore, Δσ2 can be calculated as following: � � Δσ 2 ¼ k DxAl1=2 DPZ1=2 (8) where D xAl is the grain size of the αp in the Zr–xAl alloy (x ¼ 3, 6, 9 and 12); D PZ is the grain size of the αp in the pure Zr. Therefore, Δσ 1 can be calculated as following: � � Δσ 1 ¼ Δσx k DxAl1=2 DPZ1=2 (9) In this study, Δσ 1 values are a function of the relative deviation of lattice parameter of α in the Zr–xAl alloy (x ¼ 3, 6, 9 and 12) and pure Zr 6
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
provided the research topic and funds. Declaration of competing interest We declare that we do not have any commercial or associative in terest that represents a conflict of interest in connection with the work submitted. Acknowledgments This work was supported by the National Natural Science Foundation of China (51531005/51671166/51571174). References [1] T.R.G. Kutty, K. Ravi, C. Ganguly, Studies on hot hardness of Zr and its alloys for nuclear reactors, J. Nucl. Mater. 265 (1) (1999) 91–99. [2] W. He, X. Chen, N. Liu, B. Luan, G. Yuan, Q. Liu, Cryo-rolling enhanced inhomogeneous deformation and recrystallization grain growth of a zirconium alloy, J. Alloy. Comp. 699 (2017) 160–169. [3] R. Kondo, N. Nomura, Suyalatu, Y. Tsutsumi, H. Doi, T. Hanawa, Microstructure and mechanical properties of as-cast Zr–Nb alloys, Acta Biomater. 7 (12) (2011) 4278–4284. [4] R.B. Russell, Coefficients of thermal expansion for zirconium, JOM 6 (9) (1954) 1045–1052. [5] M. Grosse, J.R. Santisteban, J. Bertsch, B. Schillinger, A. Kaestner, M.R. Daymond, N. Kardjilov, Investigations of the hydrogen diffusion and distribution in zirconium by means of neutron imaging, Kerntechnik 83 (6) (2018) 495–501. [6] J. Zhang, Y. Hu, J. Huang, L. Tu, M. Yao, B. Zhou, The corrosion resistance of Zr–0.7Sn–1Nb–0.2Fe–xCu–xGe alloys in 360� C lithiated water, Corros. Sci. 111 (2016) 132–138. [7] G.P. Grabovetskaya, E.N. Stepanova, A.S. Dubrovskaya, Effect of hydrogen on the creep of the ultrafine-grained zirconium Zr–1Nb alloy at 673 K, Int. J. Hydrogen Energy 42 (35) (2017) 22633–22640. [8] S. Suman, M.K. Khan, M. Pathak, R.N. Singh, J.K. Chakravartty, Hydrogen in Zircaloy: mechanism and its impacts, Int. J. Hydrogen Energy 40 (17) (2015) 5976–5994. [9] P. Mukhopadhyay, S. Banerjee, Phase Transformations: Examples from Titanium and Zirconium Alloys, Elsevier, 2010. [10] H.W. Rosenberg, W.D. Nix, Solid solution strengthening in Ti–Al alloys, Metall. Trans. 4 (5) (1973) 1333–1338. [11] J. Murray, A. Peruzzi, J.P. Abriata, The Al–Zr (aluminum–zirconium) system, J. Phase Equilibria 13 (3) (1992) 277–291. [12] X.J. Jiang, R. Jing, C.Y. Liu, M.Z. Ma, R.P. Liu, Structure and mechanical properties of TiZr binary alloy after Al addition, Mater. Sci. Eng. A 586 (2013) 301–305. [13] Z.H. Feng, Z.G. Zhang, C.Q. Xia, R. Jing, S.X. Liang, J.Z. Zhao, X.Y. Zhang, M.Z. Ma, R.P. Liu, Study of microstructure evolution and strengthening mechanisms in novel ZrBeAl alloys, J. Alloy. Comp. 689 (2016) 48–55. [14] R.J. Kematick, H.F. Franzen, Thermodynamic study of the zirconium-aluminum system, J. Solid State Chem. 54 (1984) 226–234. [15] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Precipitation of Zr3Al in a Zr–4.6 wt.% Al martensite, J. Nucl. Mater. 82 (1979) 227–238. [16] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Formation of a D019 phase in zirconium aluminium martensite, J. Nucl. Mater. 82 (1979) 227–238. [17] E.M. Schulson, The tensile and corrosion behaviour of ordered Zr3Al-based alloys, J. Nucl. Mater. 50 (2) (1974) 127–138. [18] L.M. Howe, M. Rainville, E.M. Schulson, Transmission electron microscopy investigations of ordered Zr3Al, J. Nucl. Mater. 50 (2) (1974) 139–154. [19] M. Alatalo, M. Weinert, R.E. Watson, Stability of Zr Al alloys, Phys. Rev. B 57 (4) (1998) 2009–2012. [20] C. Sauer, G. Lütjering, Influence of α layers at β grain boundaries on mechanical properties of Ti-alloys, Mater. Sci. Eng. A 319–321 (2001) 393–397. [21] Q. Chao, P. Cizek, J. Wang, P.D. Hodgson, H. Beladi, Enhanced mechanical response of an ultrafine grained Ti–6Al–4V alloy produced through warm symmetric and asymmetric rolling, Mater. Sci. Eng. A 650 (2016) 404–413. [22] D.N. Crowther, Z. Mohamed, B. Mintz, The relative influence of dynamic and static precipitation on the hot ductility of microalloyed steels, Metall. Trans. A 18 (11) (1987) 1929–1939. [23] M.G. Akben, T. Chandra, P. Plassiard, J.J. Jonas, Dynamic precipitation and solute hardening in a titanium microalloyed steel containing three levels of manganese, Acta Metall. 32 (4) (1984) 591–601. [24] W.Z. Han, Y. Chen, A. Vinogradov, C.R. Hutchinson, Dynamic precipitation during cyclic deformation of an underaged Al–Cu alloy, Mater. Sci. Eng. A 528 (24) (2011) 7410–7416. [25] X. Hou, Z. Cao, X. Sun, L. Wang, L. Wang, Twinning and dynamic precipitation upon hot compression of a Mg Gd Y Nd Zr alloy, J. Alloy. Comp. 525 (2012) 103–109. [26] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Precipitation of Zr3Al in a Zr 4.6 wt% Al martensite, J. Nucl. Mater. 82 (2) (1979) 227–238. [27] M.G. Akben, I. Weiss, J.J. Jonas, Dynamic precipitation and solute hardening in A V microalloyed steel and two Nb steels containing high levels of Mn, Acta Metall. 29 (1) (1981) 111–121.
Fig. 8. △σ 1 values of the hot-rolled Zr–xAl alloys (x ¼ 3, 6, 9 and 12 at.%) as a function of the deviation of the lattice parameters a.
and are presented in Fig. 8. The fitting data of Δσ 1 and Δa2=3 of the Zr–3Al, Zr–6Al, Zr–9Al and Zr–12Al alloys have a well linear relation ship. In addition, according to Fig. 8, the increases in yield strengths of the Zr–xAl alloys are mainly contributed by the solid solution of Al. Meanwhile, the microstructure refinement of the Zr–12Al alloy contributed about 40 MPa to the increase in yield strength. Thus, the main strengthening mechanism of the Zr–xAl alloys is solid solution strengthening. 4. Conclusions The microstructures and mechanical properties of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) were investigated. The ef fects of solid solution and dynamic precipitation on the mechanical properties were discussed. The results are summarized as follows: 1. The Zr–12Al alloy processed by hot-rolling at 1143 K and subsequent water-quenching showed an ultimate strength of 787.6 MPa with a fracture elongation of 19.5%, however, those of the pure Zr after the same treatment were 443.8 MPa and 17.3%, respectively. The strength and ductility of the Zr–12Al alloy were simultaneously enhanced compared with those of the pure Zr. 2. The main strengthening mechanism of the Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) processed by hot-rolling at 1143 K and subsequently water-quenching was solid solution strengthening. 3. The precipitated Zr3Al particles increased the strain hardening rate of the Zr–9Al and Zr–12Al alloys. The high strain hardening rate and elongations at the uniform plastic deformation stage of the Zr–9Al and Zr–12Al alloys resulted in the simultaneous enhancement of strength and ductility. 4. Dynamic precipitation of Zr3Al particles occurred in the Zr 9Al and Zr 12Al alloys during hot-rolling at 1143 K, leading to the exsolu tion of the α phase near the compounds and the grain boundaries. The α phase near the compounds and grain boundaries transformed into the β phase, and retained to room temperature after waterquenching. Author contributions section X. Zhang wrote the manuscript and prepared the needed alloys. B. Zhang conducted the XRD, SEM and TEM characterizations. S.G. Liu conducted the mechanical tests. C.Q. Xia analyzed the experimental results. X.Y. Zhang and M.Z. Ma designed the experiments. R.P. Liu 7
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
[28] I. Weiss, J.J. Jonas, Interaction between recrystallization and precipitation during the high temperature deformation of HSLA steels, Metall. Trans. A 10 (7) (1979) 831–840. [29] A. le Bon, J. Rofes-Vernis, C. Rossard, Recrystallization and precipitation during hot working of a Nb-bearing HSLA steel, Met. Sci. 9 (1) (1975) 36–40. [30] R.W. Balluffi, Grain boundary diffusion mechanisms in metals, Metall. Trans. B 13 (4) (1982) 527–553. [31] H.S. Levine, C.J. MacCallum, Grain boundary and lattice diffusion in polycrystalline bodies, J. Appl. Phys. 31 (3) (1960) 595–599. [32] S. Banerjee, R.W. Cahn, An ordered ω-phase in the rapidly solidified Zr–27 at.% Al alloy, Acta Metall. 31 (10) (1983) 1721–1735. [33] X.J. Jiang, G. Yu, Z.H. Feng, C.Q. Xia, C.L. Tan, X.Y. Zhang, M.Z. Ma, R.P. Liu, Abnormal β-phase stability in TiZrAl alloys, J. Alloy. Comp. 699 (2017) 256–261. [34] C. de Formanoir, G. Martin, F. Prima, S.Y.P. Allain, T. Dessolier, F. Sun, S. Viv�es, B. Hary, Y. Br� echet, S. Godet, Micromechanical behavior and thermal stability of a dual-phase αþα0 titanium alloy produced by additive manufacturing, Acta Mater. 162 (2019) 149–162. [35] X. Wu, M. Yang, F. Yuan, G. Wu, Y. Wei, X. Huang, Y. Zhu, Heterogeneous lamella structure unites ultrafine-grain strength with coarse-grain ductility, Proc. Natl. Acad. Sci. 112 (47) (2015) 14501.
[36] Z. Lei, X. Liu, Y. Wu, H. Wang, S. Jiang, S. Wang, X. Hui, Y. Wu, B. Gault, P. Kontis, D. Raabe, L. Gu, Q. Zhang, H. Chen, H. Wang, J. Liu, K. An, Q. Zeng, T.-G. Nieh, Z. Lu, Enhanced strength and ductility in a high-entropy alloy via ordered oxygen complexes, Nature 563 (7732) (2018) 546–550. [37] J.W. Zhang, I.J. Beyerlein, W.Z. Han, Hierarchical 3D nanolayered duplex-phase Zr with high strength, strain hardening, and ductility, Phys. Rev. Lett. 122 (25) (2019) 255501. [38] R.L. Fleisgher, Solution hardening, Acta Metall. 9 (11) (1961) 996–1000. [39] R.L. Fleischer, Substitutional solution hardening, Acta Metall. 11 (3) (1963) 203–209. [40] J. She, Y. Zhan, C. Li, Novel in situ synthesized zirconium matrix composites reinforced with ZrC particles, Mater. Sci. Eng. A 527 (23) (2010) 6454–6458. [41] V.A. Lubarda, On the effective lattice parameter of binary alloys, Mech. Mater. 35 (1) (2003) 53–68. [42] R.L. Fleischer, Substitutional solutes in AlRu—I. Effects of solute on moduli, lattice parameters and vacancy production, Acta Metall. 41 (3) (1993) 863–869. [43] N. Hansen, Hall–Petch relation and boundary strengthening, Scr. Mater. 51 (8) (2004) 801–806. [44] R. Kapoor, A. Sarkar, J. Singh, I. Samajdar, D. Raabe, Effect of strain rate on twinning in a Zr alloy, Scr. Mater. 74 (2014) 72–75.
8
Contents lists available at ScienceDirect
Materials Science & Engineering A journal homepage: http://www.elsevier.com/locate/msea
Microstructures and mechanical properties of Zr–Al binary alloys processed by hot-rolling X. Zhang a, B. Zhang a, S.G. Liu a, C.Q. Xia b, X.Y. Zhang a, M.Z. Ma a, R.P. Liu a, * a b
State Key Laboratory of Metastable Materials Science and Technology, Yanshan University, Qinhuangdao, 066004, China School of Materials Science and Engineering, Hebei University of Technology, Tianjin, 300130, China
A R T I C L E I N F O
A B S T R A C T
Keywords: Zr–Al alloys Microstructure Mechanical properties Solid solution strengthening Dynamic precipitation
The microstructures and mechanical properties of Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) that underwent the hot-rolling at 1143 K and the subsequent water-quenching were investigated. Microstructural analyses show that the pure Zr, Zr–3Al and Zr–6Al alloys mainly consisted of the α and α0 phases and exhibited bimodal micro structures. The Zr–9Al and Zr–12Al alloys consisted of the α, β, Zr2Al and Zr3Al phases and exhibited equiaxed microstructures. Tensile tests show that, with the increase in Al content, the yield and ultimate strength of the alloys increased from 324.9 MPa and 443.8 MPa (pure Zr) to 665.2 MPa and 787.6 MPa (Zr–12Al), respectively. Meanwhile, the elongation to failure initially decreased from 17.3% (pure Zr) to 10.3% (Zr–6Al) and subse quently increased to 19.5% (Zr–12Al). Analyses show that the significant enhancement of yield strength resulted from the solid solution strengthening by Al. In addition, dynamic precipitation occurred in the Zr–9Al and Zr–12Al alloys. The precipitated Zr3Al particles resulted in the obvious strain hardening behavior of the Zr–9Al and Zr–12Al alloys. The high strain hardening rate and elongation at the uniform plastic deformation stage of the Zr–12Al alloy resulted in the simultaneous enhancement of the ultimate strength and ductility.
1. Introduction Given their low thermal neutron absorption cross section, excellent corrosion resistance, good biocompatibility and low magnetic suscep tibility, zirconium (Zr) and its alloys have been widely used in nuclear, chemical and medical industries [1–3]. In addition, Zr and its alloys also show many other excellent properties such as low density, low expan sion coefficient and high melting point [4]. Therefore, Zr has great po tential application as structural materials in aerospace industry. However, the mechanical properties of Zr are insufficient to meet the requirements of structural materials. The current studies of Zr alloys mainly focus on the properties such as thermal neutron absorption cross section, corrosion resistance, hydrogen embrittlement and high-temperature creep for nuclear use [5–8]. Few studies have focused on the improvement of mechanical properties of Zr alloys. Therefore, it is necessary to develop novel Zr alloys to achieve improved mechanical properties to broaden their applications. Titanium (Ti) and Zr are transition metals belonging to Group ⅣB of the periodic table of elements. The similarity in the dispositions of the outer electrons is responsible for the similarities in some of the chemical and physical properties of Zr and Ti [9]. Zr exists as the body-centered
cubic (BCC) β phase at the high temperature above 1139 K–2128 K and atmospheric pressure. Under slow cooling, the high-temperature β phase will transform into the close-packed hexagonal (HCP) α phase by diffusive transformation. Under rapid cooling, the high-temperature β phase will transform into the HCP-α0 phase by martensitic trans formation. Aluminium (Al) has a good solid solution strengthening ef fect in Ti and is a widely used alloying element in Ti alloys [10]. In addition, Al has considerable solid solubility in the high-temperature α-Zr and the solubility can reach 10.5 at.% at 1213 K [11]. The study by Jiang et al. shows that the ultimate strength of Ti–25Zr at.% alloy increased by 41% after the addition of 15 at.% Al [12]. The study by Feng et al. shows that the addition of Al obviously caused the lattice distortion of α-Zr and had an obvious strengthening effect in Zr–Be al loys [13]. In addition, the ordered Zr3Al phase with the face-centered cubic (FCC) structure of L12 type can be precipitated to strength the Zr–Al alloys (Al content is less than 25 at.%) [11]. Therefore, Al is considered to be an ideal strengthening element in Zr. Kematick et al. analyzed the heats of formation of the Zr–Al intermetallic compounds and the solid solution of Al in BCC-Zr using the Knudsen cell mass spectrometric technique [14]. Mukhopadhyay et al. investigated the formation of a D019 phase in Zr–Al martensites and the precipitation of
* Corresponding author. E-mail address: [email protected] (R.P. Liu). https://doi.org/10.1016/j.msea.2019.138723 Received 16 September 2019; Received in revised form 18 November 2019; Accepted 21 November 2019 Available online 21 November 2019 0921-5093/© 2019 Elsevier B.V. All rights reserved.
Please cite this article as: X. Zhang, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.138723
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 1. OM micrographs of the as-cast (a) pure Zr, (b) Zr–3Al, (c) Zr–6Al, (d) Zr–9Al and (e) Zr–12Al alloys.
Zr3Al in a Zr–4.6 wt% Al martensite [15,16]. Schulson investigated the tensile and corrosion behavior of ordered Zr3Al-based alloys [17]. Howe et al. conducted the transmission electron microscopy investigations of ordered Zr3Al [18]. Alatalo et al. studied the stability of Zr–Al com pounds [19]. However, the previous studies on Zr–Al system have focused on Zr–Al intermetallic compounds. The study of microstructures and mechanical properties of Zr–Al alloys in the hypostoichiometric composition range (α þ Zr3Al phase region) has not yet been system atically conducted. Hot-rolling is a widely used method for forming technique and microstructure adjustment of metallic materials. In addition, the grain boundary α (αGB) and Widmanst€ atten α (αW) formed in the as-cast Zr and Ti alloys have negative effects on the mechanical properties such as strength and ductility [20]. Hot-rolling process is often used to eliminate the αGB and αW and obtain equiaxed or duplex microstructure to enhance the mechanical properties of alloys. The study by Chao et al. shows that
an equiaxed ultrafine-grained microstructure was successfully produced in a Ti–6Al–4V alloy through optimum hot-rolling process of a martensitic starting microstructure and the strength and plasticity of the alloy were simultaneously enhanced [21]. In this work, the microstructures and mechanical properties of the hot-rolled Zr–Al binary alloys in the hypostoichiometric composition range (α þ Zr3Al phase region) were investigated. The strengthening mechanisms, as well as the microstructure evolution were discussed. The findings of this work may provide some new insights to develop novel Zr alloys with enhanced mechanical properties. 2. Experimental procedure Industrially pure Zr (Zr þ Hf � 99.5 wt%) and industrially pure Al (99.9 wt%) were used to prepare the needed alloys in this study. Zr–xAl alloys containing different Al contents (x ¼ 0, 3, 6, 9 and 12 at.%) were 2
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Table 1 XRD diffraction angles of the α/α0 -{1011} lattice planes of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Alloy no.
Pure Zr
Zr–3Al
Zr–6Al
Zr–9Al
Zr–12Al
Angle (degree)
30.401
31.941
32.061
32.199
32.108
Fig. 2. XRD patterns of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
prepared to investigate the effects of Al contents on the microstructures and mechanical properties. The ingots (approximately 85 g) of the investigated alloys were melted via a non-consumable tungsten-elec trode arc furnace on a water-cooled copper crucible in a protective argon atmosphere. Pure Ti was used as a getter. The ingots were turned over and re-melted five times to achieve homogeneity. The ingots were heated to 1143 K and held for 40 min, subjected to multiple rolling passes and water-cooled to room temperature. The reduction of each pass was 2.16 mm. The samples were held in the furnace for 5 min be tween each pass. The total deformation reached 60%. All of the ingots were rolled under the same conditions. X-ray diffraction (XRD) was used to determine the phases of the al loys using an X-ray diffractometer (Rigaku D/MAX 2500/PC) with Cu Kα radiation. The microstructures of the alloys were examined using optical microscopy (OM, Zeiss Ax overt 200MAT), scanning electron micro scopy (SEM, Hitachi S-3400 N) and transmission electron microscopy (TEM, FEI Talos F200X). Specimens for SEM observation were me chanically polished and chemically etched. The etchant consisted of distilled water, nitric acid and hydrofluoric acid (85:10:5 by volume). TEM foils were prepared through twin-jet electrochemical polishing in a solution containing perchloric acid and methanol (10:90 by volume) at 238 K and 15 V. The tensile properties of the alloys were tested using an Instron 5982 testing machine at an initial strain rate of 5 � 10 4 s 1. An extensometer with a 12.5 mm gauge length was mounted on the speci mens to measure the strain. The specimens for the tensile tests were bone-shaped plates with an original gauge length of 20 mm and crosssectional dimensions of 3 mm � 2 mm. At least three specimens were tested for repeatability, and the average values were reported.
Fig. 3. α/α0 -Zr lattice parameters of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
results with the selected area electron diffraction (SAED) results, the phases of the alloys were identified. The pure Zr and Zr–3Al alloy con sisted of the α and α0 phases. With the increase in Al content, Zr–Al intermetallic compounds formed in the Zr–6Al, Zr–9Al and Zr–12Al al loys. The Zr–9Al and Zr–12Al alloys consisted of the α, β, Zr2Al (tI12) and Zr3Al (cP4) phases. According to the relative intensity of the diffraction peaks, the phase fraction of Zr2Al of the Zr–9Al and Zr–12Al alloys were very low. The phase fraction of Zr3Al of the Zr–12Al alloy increased substantially compared with that of the Zr–9Al alloy. The simultaneous precipitation of the Zr2Al and Zr3Al phases in the Zr–9Al and Zr–12Al alloys may be related to their similar formation enthalpy [19]. The Zr–6Al alloy consisted of the α, α0 and a small amount of the compound phases. However, the SAED results show that the measured diffraction d-spacing of the compound of the Zr–6Al alloy did not match any of the Zr–Al compounds in the International Centre for Diffraction Data (ICDD) pdf card library. The unknown compound phase may be an unreported allotrope of the Zr3Al or Zr2Al phase. Table 1 shows the XRD diffraction angles of the α/α0 -{1011} lattice planes of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Obviously, the addition of Al led to the shift of the α/α0 -{1011} diffraction peak toward high angles. The atomic radius of Al (1.82 Å) is smaller than that of Zr (2.16 Å). Al atoms replaced the positions of Zr atoms in the α/α0 -Zr lattice, resulting in the decrease in the α/α0 -Zr lat tice parameters. The α/α0 -Zr lattice parameters of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) are shown in Fig. 3. Corresponding to the change of the α/α0 -Zr lattice parameters, with the increase in Al content from 0 at.% to 9 at.%, the shift of the α/α0 -(1011) diffraction peak toward high angles increased. However, the α/α0 -(1011) diffraction peak of the Zr–12Al alloy shifted toward an angle lower than that of the Zr–9Al alloy. The exsolution of the α phase of the Zr–12Al alloy was considered to result in the decrease in Al content and the increase in the lattice parameters. Figs. 4a, b, 4c, 4d and 4e show the SEM micrographs of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Al stabilizes α-Zr thermody namically [11]. With the increase in Al content, the α→β transformation temperatures of the Zr–xAl alloys increased. It can be seen from the SEM micrographs that the pure Zr, Zr–3Al and Zr–6Al alloys consisted of the thick primary α (αp) and fine α0 grains. The Zr–9Al and Zr–12Al alloys
3. Results and discussion 3.1. Phase and microstructure Fig. 1 shows the OM micrographs of the as-cast Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). It can be seen that all the as-cast Zr–Al alloys consisted of the basketweave microstructure. The thick αGB lay along the primary β grain boundary. In addition, Widmanst€ atten structures were also observed in the as-cast Zr–Al alloys. Fig. 2 shows the XRD patterns of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Identifying the phases of the alloys using only XRD was difficult because of the variety of Zr–Al compounds and their crystal structures. This situation was aggravated by the low intensity and overlap of some diffraction peaks. From the combination of the XRD 3
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 4. SEM micrographs of the hot-rolled (a) pure Zr, (b) Zr–3Al, (c) Zr–6Al, (d) Zr–9Al and (e) Zr–12Al alloys; (f) Average widths of the αp grains of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%).
Fig. 5. (a) TEM micrograph of the hot-rolled Zr–6Al alloy; (b) STEM micrograph of the hot-rolled Zr–9Al alloy.
consisted of the equiaxed αp grains. The αp grains in the pure Zr and Zr–6Al alloy did not get dynamic globularization fully and exhibited a high length–width aspect ratio. The αGB grains also remained after hot-rolling. Fig. 4f shows the average widths of the αp grains of the hot-rolled Zr–xAl alloys. According to Fig. 4f, the average width of the αp grains of the Zr–12Al alloy was about 1.7 μm and was obviously smaller than those of the pure Zr, Zr–3Al, Zr–6Al and Zr–9Al alloys (2.7 μm–4.2 μm). Fig. 5a shows the TEM micrograph of the Zr-6Al alloy. It can be seen that the unknow compound phase precipitated from β phase and was surrounded by α0 phase, which was transformed from β phase. Fig. 5b shows the scanning transmission electron microscopy (STEM) micrograph of the Zr–9Al alloy. The diffusion paths provided by the grain boundaries can be observed. Meanwhile, the obvious contrast between the boundary area and the α grain interior indicated that the formation of the compound phase led to the exsolution of the α phase along the grain boundaries. Fig. 6a and 6b shows the SEM and STEM micrographs of the Zr–12Al alloy. Figs. 6c and 6d shows the element
mapping maps of Al and Zr corresponding to Fig. 6b. A large number of Zr3Al particles distributed at the grain boundaries can be observed. The high contrast around the compound particles and the α grain boundaries indicates that the Al content in these areas was relatively lower than that inside the α grains. Fig. 6e shows the TEM micrograph of the Zr–12Al alloy. The Zr3Al and β phases located at the α grain boundaries were identified by SAED in Fig. 6f and 6g, respectively. It can be seen from Fig. 6g that there were some diffuse intensity distributions in the diffraction patterns. It has been well documented that the occurrence of diffuse intensity indicates the tendency for the ω phase transformation [9]. Fig. 6h shows the energy disperse spectroscopy (EDS) of the β phase in Fig. 6e. It indicates that the Al content of the β phase was 7.5 at.%. Nucleation of the Zr3Al particles was observed to occur exclusively at the α grain boundaries, combined with the exsolution phenomenon of the α phase, indicating that the dynamic precipitation process occurred in the Zr–9Al and Zr–12Al alloys during the hot-rolling. Dynamic pre cipitation behavior has been reported in steel, Al and magnesium (Mg) 4
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
Fig. 6. (a) SEM and (b) STEM micrographs of the hot-rolled Zr–12Al alloy; Element mapping maps of (c) Al and (d) Zr corresponding to (b); (e) TEM micrograph of the hot-rolled Zr–12Al alloy; SAED patterns of the (f) Zr3Al and (g) β phases in (e); (h) EDS of the β phase in (e).
alloys but seldom in Ti and Zr alloys [22–25]. Mukhopadhyay et al. investigated the static precipitation behavior of Zr3Al in a Zr–4.6 wt% Al martensite and concluded that the precipitation of the equilibrium Zr3Al (cP4) phase in a supersaturated, martensitic α-Zr(Al) matrix occurred in a discontinuous manner and the equilibrium-transformed structure comprised duplex cells, each containing a lamellar distribution of the α-Zr(Al) and Zr3Al phases [26]. In this case, however, the morphology of Zr3Al was quite different from cell growth of static precipitation. It is well known that dynamic precipitation is accelerated by deformation, compared with the precipitation rate measured during isothermal transformations in undeformed structures at the same temperature [27–29]. This acceleration of precipitation is believed to be due to the introduction of favourable nucleation sites by the deformation process [22]. In addition, the dynamic globularization process led to the refinement of α grains and provided a large number of nucleation sites for the precipitation of the Zr3Al particles. Meanwhile, the grain boundaries that contained a large number of vacancies and dislocations
provided diffusion paths for atoms [30,31]. The Zr3Al particles distrib uted on the α grain boundaries pinned the grain boundaries, thus inhibiting the dynamic coarsening process of the grains and resulting in the further refinement of the α grains. Therefore, with the increase in the phase fraction of Zr3Al, the α grains of the Zr–12Al alloy obviously exhibited a finer size than those of the Zr–9Al alloy. The dynamic precipitation resulted in the exsolution of the α phase near the Zr3Al particles and grain boundaries. The decrease in Al content of the local α phase resulted in the decrease in its α→β transformation temperature. Thus, the α phase near the grain boundaries and Zr3Al particles transformed to β phase at 1143 K. However, it is well known that Al is an α stabilizer and broadens α phase region [11]. It seems that the athermal product of the high-temperature Al-containing β phase should be martensite α0 . Banerjee et al. have studied the transformation sequence in the Zr–27 at.% Al alloy during rapid quenching from the liquid state [32]. This abnormal β phase was also observed in the liquid-quenched Zr–27 at.% Al alloy. In addition, the study by Jiang
Fig. 7. (a) Engineering stress–strain curves and (b) strain hardening rate of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). 5
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
microstructure with a high strain hardening capacity inducing a desir able increase in both strength and ductility [34]. The pure Zr, Zr–3Al and Zr–6Al alloys mainly consisted of the αp and α0 phases. α0 phase has the same crystal structure (close-packed hexag onal structure) with α phase which is the equilibrium transformation product of high-temperature β phase. In general, the thick αp is considered to be softer than the fine α0 . According to the deformation theory of multiphase alloy mentioned above, it is considered that the plastic deformation started in the αp . Therefore, the relative deviation of the yield strengths ðΔσx Þ of the pure Zr and the Zr–xAl (x ¼ 3, 6, 9 and 12) alloys are mainly affected by two factors. One factorðΔσ1 Þ is the strength change caused by the change in Al content. Another factor ðΔσ 2 Þ is the strength change caused by the size change of the αp grains. Thus, Δσ x can be calculated as following:
Table 2 Tensile properties of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%). Alloy no.
σ0.2 (MPa)
σb (MPa)
Е (%)
Pure Zr Zr–3Al Zr–6Al Zr–9Al Zr–12Al
323.9 432.3 537.6 632.5 665.2
443.8 520.5 610.6 694.1 787.6
17.3 11.9 10.3 16.1 19.5
et al. shows that, with the increase in Al content, the athermal product of the β-quenched Ti–40Zr–xAl (at.%) alloys changed sequentially as fol lows: α0 →α’0 [33]. When the Al content reached 15 at.%, the β phase was retained to room temperature completely. Although the detailed for mation mechanism of this abnormal β phase remains to be studied, Al as an α stablizer does stabilize β phase under certain conditions.
Δσ x ¼ Δσ 1 þ Δσ 2
(1)
The solid solution strengthening of alloys is known to arise from the elastic interaction between the strain field of a solute and that of a dislocation, and the strength is related to the concentration of solute atoms [38,39]. Fleischer [39] reported that the relationship between the concentration of solute atoms and the strength in binary metallic ma terials could be described as following:
3.2. Mechanical properties Fig. 7a shows the engineering stress–strain curves of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) and related parameters are shown in Table 2. With the increase in Al content, the yield (σ0.2) and ultimate (σ b) strength increased from 324.9 MPa (pure Zr) and 443.8 MPa (pure Zr) to 665.2 MPa (Zr–12Al) and 787.6 MPa (Zr–12Al), respectively. Meanwhile, the elongation to failure (E) initially decreased from 17.3% (pure Zr) to 10.3% (Zr–6Al) and subsequently increased to 19.5% (Zr–12Al). Compared with those of the pure Zr, the yield and ultimate strength of the Zr–12Al alloy increased about 105% and 177%, respectively. Meanwhile, the elongation to failure of the Zr–12Al alloy even exceeded that of the pure Zr. In addition, the yield strength of the Zr–12Al alloy only increased about 33 MPa compared with that of the Zr–9Al alloy, while the ultimate strength increased about 120 MPa. It can be inferred that a large number of precipitated Zr3Al particles on the grain boundaries had an obvious effect on the ultimate strength rather than the yield strength of the Zr–12Al alloy. Fig. 7b shows the strain hardening rate (dσT/dET) at the uniform plastic deformation stage of the Zr–xAl alloys. It can be seen that the strain hardening behavior of the Zr–9Al and Zr–12Al alloys was more obvious than that of the pure Zr, Zr–3Al and Zr–6Al alloys. The elongations of the Zr–9Al and Zr–12Al alloys at the uniform plastic deformation stage were significantly greater than those of the other three alloys. The mechanical properties of alloys are related to their compositions and microstructures. Recent studies show that the presence of a me chanical contrast between the two phases in the heterogeneous mate rials, such as dual-phase steels will cause obvious strain hardening behavior [34,35]. When deforming such a composite microstructure, the softer phase is the first to deform plastically, while the harder remain elastic. After a certain amount of deformation, the initial apparent strain hardening is enhanced when compared to the strain hardening of the softer phase. The difference in behavior between soft and hard compo nents of the microstructure will sustain this effect as the soft phases will deform more than hard phases. A dual-phase microstructure thus naturally induces an additional strain hardening when compared to the individual constituting phases. In this case, the obvious strain hardening behavior of the Zr–9Al and Zr–12Al alloys was caused by the Zr3Al particles which acted as the hard phase and obstructed dislocation motion during deformation. This simultaneous enhancement of strength and ductility induced by high strain hardening has been reported in several studies. Lu et al. reported that high strain hardening which was caused by the presence of ordered interstitial complex resulted in the simultaneous enhancement of strength and ductility of the oxygen-doped TiZrHfNb high-entropy alloy [36]. Han et al. reported that high strength and ductility of the Zr–2.5Nb alloy with a high strain hardening was achieved by randomly oriented α/β phase boundary strengthening [37]. Formanoir et al. reported that a sub-transus thermal treatment followed by water quenching generated a dual-phase α þ α0
σ ∝C2=3
(2)
where σ is the yield strength and C is the concentration of the solute atom. However, the expression was modified by She [40] and is described as follows: Δσ ∝C2=3
(3)
where Δσ is the increase in the yield strengths caused by solute atoms. In general, lattice parameters will change monotonically as solute is added in solution and the empirical expression is described as following [41, 42]: ax ¼ a0 þ k1 C
(4)
where ax is the lattice parameter of the α phase in the Zr–xAl alloys (x ¼ 3, 6, 9 and 12 at.%); a0 is the lattice parameter of the α phase in the pure Zr; k1 is a constant. The expression can be revised as follows: C ¼ Δa=k1
(5)
where Δα is the relative deviation of the lattice parameter a of the α in the pure Zr and the Zr–xAl (x ¼ 3, 6, 9 and 12) alloys. In combination with the current work, the relational expression (2) can be revised as follows: Δσ 1 ∝Δα2=3
(6)
According to the Hall–Petch relation, the yield strength is related to the grain size by the equation [43]:
σ ¼ σ0 þ kDGB1=2
(7)
where σ0 and k are constants; DGB is the grain size. The Hall–Petch slope k for slip in Zr is 8 MPa mm1/2 [44]. Therefore, Δσ2 can be calculated as following: � � Δσ 2 ¼ k DxAl1=2 DPZ1=2 (8) where D xAl is the grain size of the αp in the Zr–xAl alloy (x ¼ 3, 6, 9 and 12); D PZ is the grain size of the αp in the pure Zr. Therefore, Δσ 1 can be calculated as following: � � Δσ 1 ¼ Δσx k DxAl1=2 DPZ1=2 (9) In this study, Δσ 1 values are a function of the relative deviation of lattice parameter of α in the Zr–xAl alloy (x ¼ 3, 6, 9 and 12) and pure Zr 6
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
provided the research topic and funds. Declaration of competing interest We declare that we do not have any commercial or associative in terest that represents a conflict of interest in connection with the work submitted. Acknowledgments This work was supported by the National Natural Science Foundation of China (51531005/51671166/51571174). References [1] T.R.G. Kutty, K. Ravi, C. Ganguly, Studies on hot hardness of Zr and its alloys for nuclear reactors, J. Nucl. Mater. 265 (1) (1999) 91–99. [2] W. He, X. Chen, N. Liu, B. Luan, G. Yuan, Q. Liu, Cryo-rolling enhanced inhomogeneous deformation and recrystallization grain growth of a zirconium alloy, J. Alloy. Comp. 699 (2017) 160–169. [3] R. Kondo, N. Nomura, Suyalatu, Y. Tsutsumi, H. Doi, T. Hanawa, Microstructure and mechanical properties of as-cast Zr–Nb alloys, Acta Biomater. 7 (12) (2011) 4278–4284. [4] R.B. Russell, Coefficients of thermal expansion for zirconium, JOM 6 (9) (1954) 1045–1052. [5] M. Grosse, J.R. Santisteban, J. Bertsch, B. Schillinger, A. Kaestner, M.R. Daymond, N. Kardjilov, Investigations of the hydrogen diffusion and distribution in zirconium by means of neutron imaging, Kerntechnik 83 (6) (2018) 495–501. [6] J. Zhang, Y. Hu, J. Huang, L. Tu, M. Yao, B. Zhou, The corrosion resistance of Zr–0.7Sn–1Nb–0.2Fe–xCu–xGe alloys in 360� C lithiated water, Corros. Sci. 111 (2016) 132–138. [7] G.P. Grabovetskaya, E.N. Stepanova, A.S. Dubrovskaya, Effect of hydrogen on the creep of the ultrafine-grained zirconium Zr–1Nb alloy at 673 K, Int. J. Hydrogen Energy 42 (35) (2017) 22633–22640. [8] S. Suman, M.K. Khan, M. Pathak, R.N. Singh, J.K. Chakravartty, Hydrogen in Zircaloy: mechanism and its impacts, Int. J. Hydrogen Energy 40 (17) (2015) 5976–5994. [9] P. Mukhopadhyay, S. Banerjee, Phase Transformations: Examples from Titanium and Zirconium Alloys, Elsevier, 2010. [10] H.W. Rosenberg, W.D. Nix, Solid solution strengthening in Ti–Al alloys, Metall. Trans. 4 (5) (1973) 1333–1338. [11] J. Murray, A. Peruzzi, J.P. Abriata, The Al–Zr (aluminum–zirconium) system, J. Phase Equilibria 13 (3) (1992) 277–291. [12] X.J. Jiang, R. Jing, C.Y. Liu, M.Z. Ma, R.P. Liu, Structure and mechanical properties of TiZr binary alloy after Al addition, Mater. Sci. Eng. A 586 (2013) 301–305. [13] Z.H. Feng, Z.G. Zhang, C.Q. Xia, R. Jing, S.X. Liang, J.Z. Zhao, X.Y. Zhang, M.Z. Ma, R.P. Liu, Study of microstructure evolution and strengthening mechanisms in novel ZrBeAl alloys, J. Alloy. Comp. 689 (2016) 48–55. [14] R.J. Kematick, H.F. Franzen, Thermodynamic study of the zirconium-aluminum system, J. Solid State Chem. 54 (1984) 226–234. [15] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Precipitation of Zr3Al in a Zr–4.6 wt.% Al martensite, J. Nucl. Mater. 82 (1979) 227–238. [16] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Formation of a D019 phase in zirconium aluminium martensite, J. Nucl. Mater. 82 (1979) 227–238. [17] E.M. Schulson, The tensile and corrosion behaviour of ordered Zr3Al-based alloys, J. Nucl. Mater. 50 (2) (1974) 127–138. [18] L.M. Howe, M. Rainville, E.M. Schulson, Transmission electron microscopy investigations of ordered Zr3Al, J. Nucl. Mater. 50 (2) (1974) 139–154. [19] M. Alatalo, M. Weinert, R.E. Watson, Stability of Zr Al alloys, Phys. Rev. B 57 (4) (1998) 2009–2012. [20] C. Sauer, G. Lütjering, Influence of α layers at β grain boundaries on mechanical properties of Ti-alloys, Mater. Sci. Eng. A 319–321 (2001) 393–397. [21] Q. Chao, P. Cizek, J. Wang, P.D. Hodgson, H. Beladi, Enhanced mechanical response of an ultrafine grained Ti–6Al–4V alloy produced through warm symmetric and asymmetric rolling, Mater. Sci. Eng. A 650 (2016) 404–413. [22] D.N. Crowther, Z. Mohamed, B. Mintz, The relative influence of dynamic and static precipitation on the hot ductility of microalloyed steels, Metall. Trans. A 18 (11) (1987) 1929–1939. [23] M.G. Akben, T. Chandra, P. Plassiard, J.J. Jonas, Dynamic precipitation and solute hardening in a titanium microalloyed steel containing three levels of manganese, Acta Metall. 32 (4) (1984) 591–601. [24] W.Z. Han, Y. Chen, A. Vinogradov, C.R. Hutchinson, Dynamic precipitation during cyclic deformation of an underaged Al–Cu alloy, Mater. Sci. Eng. A 528 (24) (2011) 7410–7416. [25] X. Hou, Z. Cao, X. Sun, L. Wang, L. Wang, Twinning and dynamic precipitation upon hot compression of a Mg Gd Y Nd Zr alloy, J. Alloy. Comp. 525 (2012) 103–109. [26] P. Mukhopadhyay, V. Raman, S. Banerjee, R. Krishnan, Precipitation of Zr3Al in a Zr 4.6 wt% Al martensite, J. Nucl. Mater. 82 (2) (1979) 227–238. [27] M.G. Akben, I. Weiss, J.J. Jonas, Dynamic precipitation and solute hardening in A V microalloyed steel and two Nb steels containing high levels of Mn, Acta Metall. 29 (1) (1981) 111–121.
Fig. 8. △σ 1 values of the hot-rolled Zr–xAl alloys (x ¼ 3, 6, 9 and 12 at.%) as a function of the deviation of the lattice parameters a.
and are presented in Fig. 8. The fitting data of Δσ 1 and Δa2=3 of the Zr–3Al, Zr–6Al, Zr–9Al and Zr–12Al alloys have a well linear relation ship. In addition, according to Fig. 8, the increases in yield strengths of the Zr–xAl alloys are mainly contributed by the solid solution of Al. Meanwhile, the microstructure refinement of the Zr–12Al alloy contributed about 40 MPa to the increase in yield strength. Thus, the main strengthening mechanism of the Zr–xAl alloys is solid solution strengthening. 4. Conclusions The microstructures and mechanical properties of the hot-rolled Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) were investigated. The ef fects of solid solution and dynamic precipitation on the mechanical properties were discussed. The results are summarized as follows: 1. The Zr–12Al alloy processed by hot-rolling at 1143 K and subsequent water-quenching showed an ultimate strength of 787.6 MPa with a fracture elongation of 19.5%, however, those of the pure Zr after the same treatment were 443.8 MPa and 17.3%, respectively. The strength and ductility of the Zr–12Al alloy were simultaneously enhanced compared with those of the pure Zr. 2. The main strengthening mechanism of the Zr–xAl alloys (x ¼ 0, 3, 6, 9 and 12 at.%) processed by hot-rolling at 1143 K and subsequently water-quenching was solid solution strengthening. 3. The precipitated Zr3Al particles increased the strain hardening rate of the Zr–9Al and Zr–12Al alloys. The high strain hardening rate and elongations at the uniform plastic deformation stage of the Zr–9Al and Zr–12Al alloys resulted in the simultaneous enhancement of strength and ductility. 4. Dynamic precipitation of Zr3Al particles occurred in the Zr 9Al and Zr 12Al alloys during hot-rolling at 1143 K, leading to the exsolu tion of the α phase near the compounds and the grain boundaries. The α phase near the compounds and grain boundaries transformed into the β phase, and retained to room temperature after waterquenching. Author contributions section X. Zhang wrote the manuscript and prepared the needed alloys. B. Zhang conducted the XRD, SEM and TEM characterizations. S.G. Liu conducted the mechanical tests. C.Q. Xia analyzed the experimental results. X.Y. Zhang and M.Z. Ma designed the experiments. R.P. Liu 7
X. Zhang et al.
Materials Science & Engineering A xxx (xxxx) xxx
[28] I. Weiss, J.J. Jonas, Interaction between recrystallization and precipitation during the high temperature deformation of HSLA steels, Metall. Trans. A 10 (7) (1979) 831–840. [29] A. le Bon, J. Rofes-Vernis, C. Rossard, Recrystallization and precipitation during hot working of a Nb-bearing HSLA steel, Met. Sci. 9 (1) (1975) 36–40. [30] R.W. Balluffi, Grain boundary diffusion mechanisms in metals, Metall. Trans. B 13 (4) (1982) 527–553. [31] H.S. Levine, C.J. MacCallum, Grain boundary and lattice diffusion in polycrystalline bodies, J. Appl. Phys. 31 (3) (1960) 595–599. [32] S. Banerjee, R.W. Cahn, An ordered ω-phase in the rapidly solidified Zr–27 at.% Al alloy, Acta Metall. 31 (10) (1983) 1721–1735. [33] X.J. Jiang, G. Yu, Z.H. Feng, C.Q. Xia, C.L. Tan, X.Y. Zhang, M.Z. Ma, R.P. Liu, Abnormal β-phase stability in TiZrAl alloys, J. Alloy. Comp. 699 (2017) 256–261. [34] C. de Formanoir, G. Martin, F. Prima, S.Y.P. Allain, T. Dessolier, F. Sun, S. Viv�es, B. Hary, Y. Br� echet, S. Godet, Micromechanical behavior and thermal stability of a dual-phase αþα0 titanium alloy produced by additive manufacturing, Acta Mater. 162 (2019) 149–162. [35] X. Wu, M. Yang, F. Yuan, G. Wu, Y. Wei, X. Huang, Y. Zhu, Heterogeneous lamella structure unites ultrafine-grain strength with coarse-grain ductility, Proc. Natl. Acad. Sci. 112 (47) (2015) 14501.
[36] Z. Lei, X. Liu, Y. Wu, H. Wang, S. Jiang, S. Wang, X. Hui, Y. Wu, B. Gault, P. Kontis, D. Raabe, L. Gu, Q. Zhang, H. Chen, H. Wang, J. Liu, K. An, Q. Zeng, T.-G. Nieh, Z. Lu, Enhanced strength and ductility in a high-entropy alloy via ordered oxygen complexes, Nature 563 (7732) (2018) 546–550. [37] J.W. Zhang, I.J. Beyerlein, W.Z. Han, Hierarchical 3D nanolayered duplex-phase Zr with high strength, strain hardening, and ductility, Phys. Rev. Lett. 122 (25) (2019) 255501. [38] R.L. Fleisgher, Solution hardening, Acta Metall. 9 (11) (1961) 996–1000. [39] R.L. Fleischer, Substitutional solution hardening, Acta Metall. 11 (3) (1963) 203–209. [40] J. She, Y. Zhan, C. Li, Novel in situ synthesized zirconium matrix composites reinforced with ZrC particles, Mater. Sci. Eng. A 527 (23) (2010) 6454–6458. [41] V.A. Lubarda, On the effective lattice parameter of binary alloys, Mech. Mater. 35 (1) (2003) 53–68. [42] R.L. Fleischer, Substitutional solutes in AlRu—I. Effects of solute on moduli, lattice parameters and vacancy production, Acta Metall. 41 (3) (1993) 863–869. [43] N. Hansen, Hall–Petch relation and boundary strengthening, Scr. Mater. 51 (8) (2004) 801–806. [44] R. Kapoor, A. Sarkar, J. Singh, I. Samajdar, D. Raabe, Effect of strain rate on twinning in a Zr alloy, Scr. Mater. 74 (2014) 72–75.
8