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Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy with icosahedral phase
Author’s Accepted Manuscript Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg6.52Zn-0.95Y alloy with icosahedral phase K.N. Li, Y.B. Zhang, Q. Zeng, G.H. Huang, B. Ji, D.D. Yin www.elsevier.com/locate/msea
PII: DOI: Reference:
S0921-5093(19)30255-2 https://doi.org/10.1016/j.msea.2019.02.083 MSA37602
To appear in: Materials Science & Engineering A Received date: 23 October 2018 Revised date: 22 February 2019 Accepted date: 23 February 2019 Cite this article as: K.N. Li, Y.B. Zhang, Q. Zeng, G.H. Huang, B. Ji and D.D. Yin, Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy with icosahedral phase, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.02.083 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 galley proof before it is published in its final citable 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.
Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy with icosahedral phase K.N. Li, Y.B. Zhang, Q. Zeng, G.H. Huang, B. Ji, D.D. Yin* Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Sichuan, Chengdu, 610031, PR China
Abstract: The microstructure evolutions and mechanical properties of Mg-6.52Zn-0.95Y (wt.%) alloy with icosahedral phase (I-phase) by semisolid treatment and equal channel angular processing (ECAP) were systematically investigated. After semisolid treatment, lamellar ultra-fine eutectic phases with spacing of 100 ± 19 nm formed at the grain boundaries. After subsequent 4 passes ECAP, fine grains (2.9 ± 0.4 μm) were obtained and the lamellar eutectic phases were fragmented into 50-500 nm granular particles. The best comprehensive mechanical properties with ultimate tensile strength (UTS) of 385 ± 18 MPa, tensile yield strength (TYS) of 290 ± 3 MPa and elongation to failure (EL) of 22.0 ± 1.6% were obtained by the combination of semisolid treatment and subsequent 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C, compared with that by the initial alloy. The improvement of mechanical properties resultes from combined effects of the grain refinement, dispersion of nanoscale I-phase particles and weakened basal texture. In addition, the I-phase particles contributed to the grain refinement during dynamic recrystallization by particle stimulated nucleation (PSN), which contributed to both strength and ductility. The average Schmid factor (SF) of basal slip increased from 0.2 to 0.4 after semisolid treatment and ECAP processing, which benefit the ductility. Keywords: Mg-Zn-Y alloys; icosahedral quasicrystal; semisolid treatment; ECAP; mechanical properties; Corresponding authors at: Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Sichuan, Chengdu, 610031, PR China. Tel.: +86-28-87634673; E-mail address: [email protected] (Y.B. Zhang), [email protected] (D.D. Yin). 1
1. Introduction
As the lightest structural metal, the ductility of magnesium (Mg) and its alloys at room temperature is limited due to lack of slip system [1]. Many efforts have been done to develop high-performance Mg alloys, and the Mg-Zn-Y-(Zr) alloy containing in situ I-phase during solidification exhibits great potential mechanical properties [2-6]. The I-phase particles are believed to be the keystrengthening factor due to theirhigh hardness, thermal stability and coherent interface between the α-Mg matrix, which have strong pinning effects on dislocations and high resistance to grain coarsening at high temperatures [7]. Generally, after extruding, rolling or forging, I-phase exists in the α-Mg matrix as a micron-scale particle. According to Zener equation [8], the pinning force (Pz) is proportional to the volume fraction of the second phase ( ) and inversely proportional to particle diameter ( ). Barlow et al. [9] observed that the introduction of nanoscaledispersoids into an aluminum matrix could accelerate the grain refinement because of enhanced dislocation generation and the reduction of the slip distances. Therefore, the refined size of the I-phase from micron scale to nano scale is likely to further improve the mechanical properties of Mg alloys containing quasicrystal I-phase.
Severe plastic deformation (SPD), such as high-pressure torsion (HPT) [10], cyclic-extrusion compression (CEC) [11] and equal channel angular pressing (ECAP) [12, 13], has been proved as an effective method for grain refinement and second phase fragmentation. Zheng et al. [14] reported that ECAP could refine the grain size effectively, break and disperse the I-phase in Mg-Zn-Y alloy, resulting in an in-situ I-phase/magnesium composites Tang et al. [15] reported the superplastic behavior of Mg-5.8Zn-1.0Y-0.48Zr alloy after ECAP with a large elongation to failure of about 800%. Garcés et al. [16] found hot forging could refine the grain size of as-cast Mg-Zn-Y alloy, but it was not effective to obtain a 2
homogeneous I-phase distribution within the magnesium matrix. Kwak et al. [17] tried to refine both α-Mg matrix and I-phase through high-ratio differential speed rolling (HRDSR), which showed that the refinement effect on the grain size was significant. However, it was found that the most of the I-phase particles were still micron scale and the uniform distribution in the matrix appeared to be poor after HRDSR. The formation of micro scale I-phase particles may be related to the micro scale lamellar spacing of the eutectic phase before SPD. Our previous study [18] observed an (α-Mg + I-phase) nano-eutectic phase of Mg-Zn-Y alloys cooling in water from semisolid state, and the fine (α-Mg + I-phase) quasi-eutectic phase may provide a foundation for the further fragment of I-phase.
In this study, an Mg-6.52Zn-0.95Y alloy reinforced by I-phase was fabricated byconventional extrusion. The combined effect of semisolid treatment and ECAP on the I-phase, grain size, texture and mechanical properties was systematically investigated.The comprehensive effect of fine-grain strengthening, I-phase precipitation strengthening and texture weakening on mechanical properties were emphatically discussed.The present work can provide a further development on the high-performance Mg alloys.
2. Experimental procedures
Pure Mg (99.99 wt. %), pure Zn (99.99 wt. %) and master alloy (Mg-30Y, wt. %) were melted in a crucible furnace under a mixed gas atmosphere of SF6 (1 vol%) and CO2 (99 vol%) and cast into an iron model of 95 mm diameter. The as-cast billets were homogenized at 420 °C for 12 h by air cooling and then the billets were machined into Ø90 mm for extrusion. Afterwards, the billets were extruded into rods with a diameter of 30 mm (extrusion ratio 9:1) at 350 °C, and the extrusion velocity of rods was 0.4 m/min.
The semisolid process was to heat the extruded samples with iron sulfide in an electric resistance furnace at 500 °C for 1 h and then cooled in water at 60 °C. The rectangular samples prepared for ECAP, 3
with the geometry of 20 × 20 × 70 mm3, were machined from the center of the semisolid bar. ECAP was carried out on the semisolid material through a die with an internal angle (φ) of 90° and a curvature angle (ψ) of 37° (the equivalent stain of each pass was ~1 [19]). A high-temperature lubricant (MoS2) was used to smear the samples and the channels. The rectangular sample was pressed through the Bc route (the sample rotated by 90° in the same direction between consecutive passes). After each pass of ECAP, the sample wasquenched in water to maintain the final microstructure and for further ECAP, it was held at the experimental temperature for 15 min, which could guarantee the experiment conducted at the stated temperature. Research [20] suggested that 4-pass ECAP could make the deformation more uniform and improve the comprehensive properties of the samples. However, with the continue increase of ECAP pass, the weakening of textures is enhanced which make basal slip in soft orientation for extruded direction. As a result, the yield strength in the extruded direction may be reduced. Two ECAP experimental methods were used, the first one was conducted at 300 °C for four passes, and the other one was conducted at 300 °C for the first two passes and 200 °C for the last two passes. For repeatability and accuracy, each test in the same condition was repeated at least three times. The sketch map of the experiment and details of ECAP were shown in Fig. 1 (a) and (b), respectively.
The microstructural observation was investigated by optical microscopy (OM, Zeiss Axio Lab A1), scanning electron microscopy (SEM, JSM-7001F) and electron backscattered diffraction (EBSD, Oxford Instrument Nordlys Nano EBSD detector and HKL channel 5 data acquisition software). Specimens for OM and SEM were mechanically polished and then etched in a solution containing 1 ml nitric acid, 1 ml acetic acid, 1 g oxalic acid, and 150 ml distilled water. For EBSD, specimens were taken through the same procedure to OM and further electropolished in a solution containing 10 vol.% perchloric acid and 90 vol.%
4
ethyl alcohol at about -40 °C to -20 °C and a voltage of 4 V for ~20 s. Constituent phases were identified by X-ray diffraction (XRD, Rigaku D/MAX 2500PC) using monochromatic Cu Kα radiation.
Tensile tests were performed at room temperature using a CMT-5105 universal test machine with a strain rate of 1 × 10-3 s-1. The specimens have a gauge length of 8 mm and cross-sectional areas of 2 × 3 mm2. The tensile axis was parallel to the extrusion direction (ED).
3. Results 3.1 Microstructure 3.1.1 Initial Microstructure
Fig. 2 shows the representative microstructure of the as-extruded Mg-6.52Zn-0.95Y alloy, and the inset is the high magnification SEM image. The as-extruded Mg-6.52Zn-0.95Y alloy exhibited a bimodal grain structure consisting of ultrafine grains with several microns and coarse unrecrystallized grains, which was also observed by other researchers [21]. Meanwhile, it was notable that fine grains were generally around the second phase. The inhomogeneous grains were usually attributed to the local occurrence of dynamic recrystallization (DRX), and it was commonly considered that second phase particles can inhibit the motion of dislocations as well as impede to the growth of newly formed recrystallized grains [22, 23].
The average grain size of the as-extruded alloy, measured by linear intercept method, is 19.9 ± 3.8 μm. The second phase distributed along the extrusion direction (ED). XRD analysis (Fig. 3(a)) indicated that the second phase was Mg3Zn6Y1 (I-phase). The size of the I-phase along ED was in the range of 1.0 - 5.0 μm and the volume fraction of I-phases was 4.8 ± 1.0%.
3.1.2 Microstructure evolution during semisolid and ECAP
5
Fig. 4a and 4d are representative OM and SEM images of Mg-6.52Zn-0.95Y alloy after semisolid treatment, and the insets in Fig. 4d are partially enlarged details. The XRD pattern of Mg-6.52Zn-0.95Y alloy after semisolid treatment shown in Fig. 3b proved that there was no change in phase constituent. As was shown in Fig. 4a and 4d, obvious eutectic structure could be found at the grain boundary of α-Mg, and the unfused α-Mg was not spherical but polygonal. The average grain size of the Mg-6.52Zn-0.95Y alloy after semisolid was 91.0 ± 6.5 μm. Under water-cooling condition at 60 °C, an ultra-fine eutectic phase formed at the grain boundaries with the lamellar spacing of 100 ± 19 nm. The volume fraction of the eutectic structure measured by image analyzer software was 4.0%, indicating that the volume fraction of I-phase was 2.0% (assume that the volume fraction of the I-phase and α-Mg in eutectic structure were equal [24]).
The typical microstructures of the semisolid Mg-6.52Zn-0.95Y alloy after 4 passes of ECAP at 300 °C are shown in Fig. 4b and 4e. The OM microstructure showed that grains were refined significantly to a level of 7.2 ± 0.8 μm (average grain size). However, there are still coarse unrecrystallized grains, and the volume fraction of unrecrystallized coarse grains was 19.7 ± 0.3%. Similar inhomogeneous microstructures were also discovered in AZ31 [25], ZK60 [26], Mg-Al-Y-Zn [27] and Mg-Gd-Nd-Zn-Zr[28] alloys after ECAP. The formation of bimodal grain structure was related to the processing temperature, the number of repetitive processing and non-uniform strain caused by the outer rounded corner [29]. The black particles in Fig. 4b were I-phase particles, and its XRD pattern was shown in Fig. 3d. It was clear that the lamellar eutectic phases were fragmented into 50-500 nm granular particles. Moreover, the distribution of the broken particles was in bands along the extrusion direction of ECAP.
Fig. 4c and 4f are OM and SEM images of Mg-6.52Zn-0.95Y alloy after 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. Comparing with the former process, the α-Mg matrix was further refined, with the average grain size of 2.9 ± 0.4 μm, and banded distributed broken I-phase particles were 6
also 50-500 nm in size. Whereas, it was noteworthy to see that there were much finer and more uniform nanoscale I-phase particles distributed throughout the α-Mg matrix. Zeng et al [30] observed the precipitation behavior of Mg-6Zn-1.5Y-0.5Zr alloy and it suggested that the aging process at 180 °C had an effect on the hardness of extruded Mg-Zn-Y-Zr alloys because of the precipitates. Meanwhile, the same precipitation behavior of I phase was also discovered in Mg-5Zn-0.92Y-0.16Zr alloy during ECAP at 200 °C [21]. It was reasonable to hypothesize that the present uniformly distributed I-phase particles was attributed to the occurrences of dynamic aging at lower temperature (200 °C) ECAP.
Overall, these results indicated that shear stress induced by ECAP promoted the refinement and dispersion of the I-phase particles. Meanwhile, the ECAP temperature had an important effect on grain refinement and precipitation of I phase. Comparing the relative intensities of the three most prominent diffraction peaks of (10 ̅ 0), (0002) and (10 ̅ 1) of Mg-6.52Zn-0.95Y alloy after different processes in Fig. 3, it could be easily found that after 4 passes ECAP at 300 °C, there was a sharp increase in the intensity of (0002) peak, while the intensity of the other two peaks fell after 4 passes ECAP at 300 °C. However, the specimen showed a resultant reversal in the relative intensities after 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. The change of diffraction peaks was correlated to textures formed at different temperatures [25, 31], and the texture evolution will be discussed in the next section.
3.1.3 Texture The inverse pole figure maps (IPF map), misorientation angle distribution histogram, {0001} pole figure (PF) and inverse pole figure for the ED (ED IPF) of Mg-6.52Y-0.95Y alloy in different states in ED-TD plane are shown in Fig. 5. The black zones in Fig.5a, 5e, 5i, and 5m were unidentified I-phase 7
particles. Grain boundaries were color-coded in IPF maps according to their misorientation angle, and the white lines and black lines correspond to low-angle grain boundaries (LAGBs, 2°≤θ<15°) and high-angle grain boundaries (HAGBs, θ≥15°). The misorientation angle quantitative distribution of as-extruded Mg-6.52Y-0.95Y alloy was shown in Fig. 5b, it could be seen that there are two distribution peaks at 20° and 86°, respectively. LAGBs account for 18.3% for the as-extruded state, which suggested the incomplete of DRX. The PF of as-extruded Mg-6.52Y-0.95Y alloy, as shown in Fig. 5c, exhibited a typical magnesium extrusion texture (ED // <10 ̅ 0> fiber texture), with the basal plane rotating along transverse direction (TD) so that c-axis of grains are perpendicular to the ED, which was consistent with the generally reported texture characteristics of extruded magnesium bars [32-34]. Analyzing with the ED IPF (Fig. 5d), the as-extruded Mg-6.52Y-0.95Y alloy exhibited a main peak around < ̅ 2 ̅ 0> indicating that the < ̅ 2 ̅ 0> direction of most grains was nearly parallel to the ED. The IPF map of Fig. 5e indicated that the semisolid bar exhibited a relatively random grain orientation distribution, comparing with the extruded bar. With the increase of grain size, the LAGBs frequency was decreased to 9.6%. The PF, as shown in Fig. 5g, was similar to the extruded bar. However, the basal texture at semisolid state was weaker than former and consisted of several main components. Our previous study [35] also observed similar basal texture weakening in Mg-Y alloy after annealing process, and it was ascribed to preferential growth of non-basal grains.
After 4 passes ECAP at 300 °C, the LAGBs frequency was increased to 19.7%. When the ECAP temperature was decreased to 200 °C, the LAGBs frequency further was increased to 29.1%. Mabuchi et al. [36] also discovered that the frequency of LAGBs was larger for ECAP than those for the conventional extrusion. ECAP was a type of SPD, which would introduce sever shear deformation and cause a lot of dislocations and sub-grains. A small number of repetitive processing at a low temperature would decrease 8
the volume fraction of DRX and simultaneously increase the frequency of LAGBs. The PF of Mg-6.52Y-0.95Y alloy after ECAP at 300 °C exhibited a {0002} basal plane texture with a maximum intensity of around 23.1 MRD (multiples of random distribution) and had an angle of about 82° with ED. Corresponding ED IPF (Fig. 5l) exhibited a main peak around < ̅ 2 ̅ 0>. However, the PF of Mg-6.52Y-0.95Y alloy after ECAP at 200 °C in Fig. 5o exhibited that the basal plane texture, with a maximum intensity of around 13.4 MRD, was about 44° with ED. Meanwhile, corresponding ED IPF (Fig. 5p) showed two obvious poles around <11 ̅ 2> and <10 ̅ 1>. Yu Yoshida [25, 31] also discovered the same texture development of AZ31 magnesium alloy during ECAP. He ascribed the differences between textures to plastic deformation mechanism of magnesium alloy at different temperatures. At low temperature (200 °C), the formation of texture was mainly dominated by tensile twinning and basal slip; While, at high temperature (300 °C), < c + a > slip predominated at the initial stage of the deformation. Finally, after 8 passes ECAP, a basal plane texture, about 45° or 90° with ED, was generated. Considering the same textures in this research, it is reasonable to hypothesize that it is caused by different plastic deformation mechanisms at different ECAP temperatures.
3.2 Mechanical properties
Fig. 6 shows representative tensile stress-strain curves of the Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP when testing along ED at room temperature. The corresponding mechanical properties of tensile yield strength (TYS), ultimate tensile strength (UTS) and elongation to failure (EL) are summarized in Table 1. As Table 1 shows, after semisolid treatment, the mechanical properties of Mg-6.52Zn-0.95Y alloy were decreased because of the obvious grain growth. After subsequent ECAP, the comprehensive mechanical properties of Mg-6.52Zn-0.95Y alloy was increased with ECAP pass at 300 °C, when the ECAP temperature was decreased to 200 °C, the UTS, TYS, and EL was increased more 9
significantly. The alloy exhibited best tensile properties with UTS of 385 MPa, TYS of 290 MPa and EL of 22.0%. There are 24.2%, 35.5%, and 56.0% increment of UTS, TYS, and EL, respectively, comparing with the as-extruded alloy.
Fig. 7 illustrates the SEM images and corresponding backscattered electron images of fracture morphologies of the tensile specimens, and the bright particles in Fig. 7d, 7e and 7f are I-phase particles. As can be seen in the Fig. 7, the fracture features were similar for all alloys, they were mainly composed of tear ridges and dimples, which suggested that the specimens underwent a substantial plastic deformation prior to fracture. However, more highly developed tearing ridges and tiny as well as deeper dimples were observed after ECAP, which was consistent with the significant improvement of elongation, as aforementioned in Table 1. Turning to I phase, it could be obviously seen that the I-phase particles after ECAP were finer andevenly distributed at the bottom of the dimples, resulting in the aggrandizement of strength and plasticity.
4. Discussions
In recent years, the interests in Mg-Zn-Y-(Zr) alloys with I-phase have been growing, mainly because of their desirable yield and ultimate tensile strength both at room and elevated temperature after thermomechanical processing, such as hot extrusion, forging or rolling. Fig. 8 is a comparison drawing in the mechanical property of current work and I-phase enhanced Mg-Zn-Y-(Zr) alloys in recent years [14, 16, 17, 24, 37-39]. Kwak et al. [24] used different speed rolling to refine the cast microstructure of Mg-9.25Zn-1.66Y alloy with I-phase, and the alloy showed high UTS approach 380 MPa but low EL of 4.0%, as shown in Fig. 8, the brown dotted box. The low EL was ascribed to the non-fully dispersed I-phase particle agglomerates. Zheng et al. [38]revealed the effect of ECAP on as-cast and as-extruded ZWK510 10
(Mg-5.0Zn-0.9Y-0.2Zr, wt. %) alloy containing I-phase. The result showed that both strength and ductility of the as-cast ZWK510 alloy were improved after ECAP, however, with the increase of EL of the as-extruded ZWK510 alloy after ECAP, its strength was decreased significantly. The as-ECAPed ZWK510 alloy exhibited best tensile properties with UTS of 330 MPa and EL of 13.0%, as shown in Fig. 8 indicated by the red points. Texture modification was considered to be the main reason for the decrease of UTS.
Comparing with the previous studies, it could be seen that the status of material before SPD affected their processed performance, and the as-semisolided Mg-6.52Zn-0.95Y alloy after ECAP showed simultaneous increase in strength and ductility, with UTS of 385 ± 18 MPa, TYS of 290 ± 3 MPa and EL of 22.0 ± 1.6%. As shown in Fig. 4d, the formation of ultra-fine eutectic phase provided the foundation for the generation of nanoscale I-phase particles. With the subsequent high shear strain during ECAP, the ultra-fine eutectic phase was broken into nanoscale I-phase particlesand dispersed into α-Mg matrix. The refinement and random distribution of I-phase particles, which can effectively reduce stress concentration, could guarantee the improvement of ductility.
Fig. 9 shows the SEM image of the semisolid Mg-6.52Zn-0.95Y alloy after 1 pass ECAP at 300 °C. It could be clearly seen that the eutectic phases were broken into particles initially and a continuous cluster of recrystallized α-Mg grains with the size about 1 ~ 3 μm preferentially generated around the I-phase particles. Generally, the main role of I-phase particles was to stimulate the recrystallization through the particle stimulated nucleation (PSN [40, 41]), in which nucleation of recrystallized α-Mg grains occurred preferentially over I-phase particles at an early stage. The presence of I-phase particle promoted the grain refinement and improved the strength of as-ECAPed Mg-6.52Zn-0.95Y alloy. Alok Singh [42] and Kwak[24] also discovered that α-Mg grain preferentially nucleated on I-phase surfaces after high-pressure torsion (HPT) and the uniform dispersion of the I-phase partlclewasimportant to the grain refinement. 11
Besides the I-phase precipitation strengthening, grain size refinement was considered to be a major strengthening mechanism. Generally, Hall-Petch relation was often used to explain the strength of a material. Fig. 10 is Hall-Petch relation for Mg-6.52Zn-0.95Y alloy after semisolid and ECAP treatment, and Hall-Petch relation of I-phase enhanced Mg-13Zn-1.55Y alloy [17], Mg-8.32Zn-1.47Y alloy [43] and Mg-7.58Zn-1.72Y alloy [44] are also plotted. As was shown in Fig.10, it is different from the as-extruded or HRDSRed Mg-Zn-Y alloys,the relation between the TYS and the average grain size could not be explained by Hall-Petch relation. Similar abnormal relation was also observed in AZ31 [31] and AZ61 [34, 45] alloys. Texture modification during ECAP was thought to be responsible for the above results.
In general, the textures of hcp metals after ECAP had a significant effect on mechanical properties, especially on yield stress and elongation. Main slip systems in magnesium and its alloys are [46]: basal slip ({0001} <11 ̅ 0>), prismatic slip ({10 ̅ 0} <11 ̅ 0>) and second order pyramidal slip ({11 ̅ 2} <11 ̅ 3>). Beside these slip systems, mechanical twinning played an important role in plastic deformation of magnesium alloys, including extension twins, contraction twins and double twins [32, 47]. It is generally accepted that the dominating deformation mechanism at room temperature are extension twinning and basal slip, for the critical resolved shear stress (CRSS) at room temperature follow the relationship: CRSS CRSS
extension twinning<
CRSS
prismatic<
CRSS
pyramidal[48,
basal<
49]. The Schmid factor (SF) evolution of basal slip
and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP treatment is shown in Fig. 11. After ECAP, the weighted average SF for extension twinning was decreased slightly whereas the SF for basal slipwas increased obviously; with the decrease of ECAP temperature, the SF for basal slip was increased from 0.19 to 0.40. It is believable that the lower TYS and higher EL of Mg-6.52Zn-0.95Y alloy after 3 ~ 4 passes ECAP may be ascribed to the highest average SF of the basal slip. However, another point to be noted in Table 1 was that the TYS of the as-ECAPed Mg-6.52Zn-0.95Y alloy at 200 °C was notably 12
higher than that of the as-extruded state, and the difference may be attributed to the obvious grain size refinement at low temperature.
As a conclusion, severe shear strain introduced by ECAP and PSN effect of I-phase particles led tothe grain refinement. Meanwhile, the formation of weakened basal texture and randomly distributed nanoscale I-phase particles prevented alloy fracturing at an early age. The combined effect of fine grain, texture and I-phase resulted in a simultaneously increase in the strength and plasticity of Mg-6.52Zn-0.95Y alloy.
5. Conclusions
The effects of semisolid and ECAP on the I-phase, microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy were systematically investigated. The main results of this study are outlined as follows:
(1) After the as-extruded Mg-6.52Zn-0.95Y alloy was semisolid treated at 500 °C for 1 h and then cooled in water at 60 °C, ultra-fine eutectic phases in the alloy were generated at the grain boundaries with the lamellar spacing of 100 ± 19 nm.
(2) Fine grain with average size of 2.9 ± 0.4 μm and high angle boundary fraction of 0.71 were obtained after semisolid treatment and 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. The as-ECAPed Mg-6.52Zn-0.95Y alloy, with a lower fraction of I-phase (2.0%) and high ductility, showed the UTS of 385 ± 18 MPa, TYS of 290 ± 3 MPa and EL of 22.0 ± 1.6%.
(3) The excellent mechanical properties of Mg-6.52Zn-0.95Y alloy after semisolid treatment and ECAP, were attributed to grain refinement, the formation of weakened basal texture and the refinement of I-phase particles. 13
(4) The fragmentation of ultra-fine eutectic phase and the random distribution of I-phase particles reduced stress concentration and prevented alloy fracturing at an early age. Meanwhile, I-phase particles contributed the grain refinement during dynamic recrystallization by PSN.
Acknowledgements This work was supported by National Nature Science Foundation of China (Nos. 51201142 and 51401172) and Fundamental Research Funds for the Central Universities (No. 2682016CX073).
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[33] S. Godet, L. Jiang, A.A. Luo, J.J. Jonas, Use of Schmid factors to select extension twin variants in extruded magnesium alloy tubes, Scr. Mater. 55(2006) 1055-1058. [34] W.J. Kim, S.I. Hong, Y.S. Kim, S.H. Min, H.T. Jeong, J.D. Lee, Texture development and its effect on mechanical properties of an AZ61 Mg alloy fabricated by equal channel angular pressing, Acta Mater. 51(2003) 3293-3307. [35] G.H. Huang, D.D. Yin, J.W. Lu, H. Zhou, Y. Zeng, G.F. Quan, Q.D. Wang, Microstructure, texture and mechanical properties evolution of extruded fine-grained Mg-Y sheets during annealing, Mater. Sci. Eng. A (2018). [36] M. Mabuchi, Y. Chino, H. Iwasaki, T. Aizawa, K. Higashi, The Grain Size and Texture Dependence of Tensile Properties in Extruded Mg-9Al-lZn, Mater. Trans. 42(2001) 1182-1188. [37] W.N. Tang, R.S. Chen, J. Zhou, E.H. Han, Effects of ECAE temperature and billet orientation on the microstructure, texture evolution and mechanical properties of a Mg-Zn-Y-Zr alloy, Mater. Sci. Eng. A 499(2009) 404-410. [38] M.Y. Zheng, S.W. Xu, X.G. Qiao, W.M. Gan, K. Wu, S. Kamado, Y. Kojima, H.G. Brokmeier. Equal Channel Angular Pressing of Magnesium Alloy Containing Quasicrystal Phase. Mater. Sci. Forum 2006. [39] M.Y. Zheng, X.G. Qiao, S.W. Xu, K. Wu, S. Kamado, Y. Kojima, Effect of Equal Channel Angular Extrusion on the Microstructure and Mechanical Properties of Magnesium Alloy Containing Quasicrystallines, Mater. Sci. Forum 488-489(2005) 589-592. [40]J.D. Robson, D.T. Henry, B. Davis, Particle effects on recrystallization in magnesium-manganese alloys: Particle-stimulated nucleation, Acta Mater. 57(2009) 2739-2747. [41] T.Y. Kwak, H.K. Lim, W.J. Kim. Effect of the volume fraction of the icosahedral phase on the microstructures, hot compressive behaviors and processing maps of Mg-Zn-Y alloys. J. Alloy. Compd. 725(2017) 711-723. [42] A. Singh, D.A. Basha, H. Somekawa, K. Tsuchiya, Nucleation of recrystallized magnesium grains over quasicrystalline phase during severe plastic deformation of a Mg-Zn-Y alloy at room temperature, Scr. Mater. 134(2017) 80-84. [43] E. Mora, G. Garcés, E. Oñorbe, P. Pérez, P. Adeva, High-strength Mg-Zn-Y alloys produced by powder metallurgy, Scr. Mater. 60(2009) 776-779. [44] A. Singh, Y. Osawa, H. Somekawa, T. Mukai, Effect of microstructure on strength and ductility of high strength quasicrystal phase dispersed Mg-Zn-Y alloys, Mater. Sci. Eng. A 611(2014) 242-251. [45] W.J. Kim, H.G. Jeong, Mechanical Properties and Texture Evolution in ECAP Processed AZ61 Mg Alloys, Mater. Sci. Forum 419-422(2003) 201-206. [46] D.D. Yin, Q.D. Wang, C.J. Boehlert, Z. Chen, H.M. Li, R.K. Mishra, A. Chakkedath, In-Situ Study of the Tensile Deformation and Fracture Modes in Peak-Aged Cast Mg-11Y-5Gd-2Zn-0.5Zr (Weight Percent), Metall. Mater. Trans. A (2016) 6348-6452. [47] J.W. Lu, D.D. Yin, G.H. Huang, G.F. Quan, Y. Zeng, H. Zhou, Q.D. Wang, Plastic anisotropy and deformation behavior of extruded Mg-Y sheets at elevated temperatures, Mater. Sci. Eng. A 700(2017) 598-608. [48] L. Yao, Y. Wei, A polycrystal based numerical investigation on the temperature dependence of slip resistance and texture evolution in magnesium alloy AZ31B, Int. J. Plast. 55(2014) 80-93. [49] F. Kabirian, A.S. Khan, T. Gnäupel-Herlod, Visco-plastic modeling of mechanical responses and texture evolution in extruded AZ31 magnesium alloy for various loading conditions, Int. J. Plast. 68(2015) 1-20.
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Figure captions
Fig. 1 (a) The sketch maps of extrusion, semisolid treatment and ECAP processing; (b) experimental details of ECAP (ED: extrusion direction; TD: transverse direction; ND: normal direction). Fig. 2 The optical image of as-extruded Mg-6.52Zn-0.95Y bar at TD-ED plane, the top right picture is the zoomed-in SEM images. Fig. 3 XRD patterns of Mg-6.52Zn-0.95Y alloy after different processes. Fig. 4 Optical micrographs (a, b, c) and SEM images (d, e, f) of Mg-6.52Zn-0.95Y alloy after different processes
at
TD-ND
plane
(a,
d)
semisolid;
(b,
e)
semisolid+ECAP-300
°C-P4;
(c,
f)
semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 5 EBSD orientation maps (a, e, i, m) and corresponding misorientation angle distribution histograms (b, f, j, n), {0001} pole figures (c, g, k, o), inverse pole figures for the ED direction (d, h, l, p) of Mg-6.52Zn-0.95Y alloy after different processes; (a, b, c, d) as-extruded; (e, f, g, h) semisolid; (i, j, k, l) semisolid+ECAP-300 °C-P4; (m, n, o, p) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 6 Engineering stress-strain curves of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP processing. Fig. 7 Secondary electron images (a, b, c) and corresponding backscattered electron images (d, e, f) of fracture surface of Mg-6.52Zn-0.95Y alloy after different processes (a, d) as-extruded; (b, e) semisolid+ECAP-300 °C-P4; (c, f) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 8 Comparison of I-phase enhanced Mg-Zn-Y-(Zr) alloys (HRDSR: high-ratio differential speed rolling). Fig.9 The SEM image of the Mg-6.52Zn-0.95Y alloy after 300 °C 1 pass ECAP. Fig. 10 The Hall-Petch relation for the current alloy, I-phase enhanced Mg-13Zn-1.55Y alloy, Mg-8.32Zn-1.47Y alloy and Mg-7.58Zn-1.72Y alloy.
17
Fig. 11 Schmid factor (SF) evolution of basal slip and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP processing.
Fig. 1 (a) The sketch maps of extrusion, semisolid treatment and ECAP processing; (b) experimental details of ECAP (ED: extrusion direction; TD: transverse direction; ND: normal direction).
18
Fig. 2 The optical image of as-extruded Mg-6.52Zn-0.95Y bar at TD-ED plane, the top right picture is the zoomed-in SEM images.
19
Fig. 3 XRD patterns of Mg-6.52Zn-0.95Y alloy after different processes.
Fig. 4 Optical micrographs (a, b, c) and SEM images (d, e, f) of Mg-6.52Zn-0.95Y alloy after different processes at TD-ND plane (a, d) semisolid; (b, e) semisolid + ECAP-300 °C-P4; (c, f) semisolid + ECAP(300 °C-P2 + 200 °C-P2).
20
Fig. 5 EBSD orientation maps (a, e, i, m) and corresponding misorientation angle distribution histograms (b, f, j, n), {0001} pole figures (c, g, k, o), inverse pole figures for the ED direction (d, h, l, p) of Mg-6.52Zn-0.95Y alloy after different processes; (a, b, c, d) as-extruded; (e, f, g, h) semisolid; (i, j, k, l) semisolid+ECAP-300 °C-P4; (m, n, o, p) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). 21
Fig. 6 Engineering stress-strain curves of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP processing.
Fig. 7 Secondary electron images (a, b, c) and corresponding backscattered electron images (d, e, f) of fracture surface of Mg-6.52Zn-0.95Y alloy after different processes (a, d) as-extruded; (b, e) semisolid+ECAP-300 °C-P4; (c, f) semisolid+ECAP-(300 °C-P2 + 200 °C-P2).
22
Fig. 8 Comparison of I-phase enhanced Mg-Zn-Y-(Zr) alloys (HRDSR: high-ratio differential speed rolling).
Fig.9 The SEM image of the Mg-6.52Zn-0.95Y alloy after 300 °C 1 pass ECAP.
23
Fig. 10 The Hall-Petch relation for the current alloy, I-phase enhanced Mg-13Zn-1.55Y alloy, Mg-8.32Zn-1.47Y alloy and Mg-7.58Zn-1.72Y alloy.
Fig. 11 Schmid factor (SF) evolution of basal slip and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP processing.
24
Table 1 UTS, TYS (σ0.2), and EL of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid treatment and ECAP processing Table 1 UTS, TYS (σ0.2) and EL of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP State of Mg-6.52Zn-0.95Y alloy As-extruded Semisolid Semisolid+ECAP-300℃-P1 Semisolid+ECAP-300℃-P2 Semisolid+ECAP-300℃-P3 Semisolid+ECAP-300℃-P4 Semisolid+ECAP-(300℃-P2+200℃ -P1) Semisolid+ECAP-(300℃-P2+200℃ -P2)
25
UTS (MPa) 310 ± 8 240 ± 5 290 ± 4 310 ± 17 330 ± 20 335 ± 11 355 ± 13 385 ± 18
YS (MPa) 214 ± 13 179 ± 14 224 ±5 236 ±6 206 ±6 209 ±4 220 ±7 290 ±3
EL (%) 14.1 ± 2.0 7.3 ± 1.3 12.5 ± 2.7 15.5 ± 3.0 21.3 ± 0.6 20.1 ± 1.3 21.1 ± 2.7 22.0 ± 1.6
PII: DOI: Reference:
S0921-5093(19)30255-2 https://doi.org/10.1016/j.msea.2019.02.083 MSA37602
To appear in: Materials Science & Engineering A Received date: 23 October 2018 Revised date: 22 February 2019 Accepted date: 23 February 2019 Cite this article as: K.N. Li, Y.B. Zhang, Q. Zeng, G.H. Huang, B. Ji and D.D. Yin, Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy with icosahedral phase, Materials Science & Engineering A, https://doi.org/10.1016/j.msea.2019.02.083 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 galley proof before it is published in its final citable 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.
Effects of semisolid treatment and ECAP on the microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy with icosahedral phase K.N. Li, Y.B. Zhang, Q. Zeng, G.H. Huang, B. Ji, D.D. Yin* Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Sichuan, Chengdu, 610031, PR China
Abstract: The microstructure evolutions and mechanical properties of Mg-6.52Zn-0.95Y (wt.%) alloy with icosahedral phase (I-phase) by semisolid treatment and equal channel angular processing (ECAP) were systematically investigated. After semisolid treatment, lamellar ultra-fine eutectic phases with spacing of 100 ± 19 nm formed at the grain boundaries. After subsequent 4 passes ECAP, fine grains (2.9 ± 0.4 μm) were obtained and the lamellar eutectic phases were fragmented into 50-500 nm granular particles. The best comprehensive mechanical properties with ultimate tensile strength (UTS) of 385 ± 18 MPa, tensile yield strength (TYS) of 290 ± 3 MPa and elongation to failure (EL) of 22.0 ± 1.6% were obtained by the combination of semisolid treatment and subsequent 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C, compared with that by the initial alloy. The improvement of mechanical properties resultes from combined effects of the grain refinement, dispersion of nanoscale I-phase particles and weakened basal texture. In addition, the I-phase particles contributed to the grain refinement during dynamic recrystallization by particle stimulated nucleation (PSN), which contributed to both strength and ductility. The average Schmid factor (SF) of basal slip increased from 0.2 to 0.4 after semisolid treatment and ECAP processing, which benefit the ductility. Keywords: Mg-Zn-Y alloys; icosahedral quasicrystal; semisolid treatment; ECAP; mechanical properties; Corresponding authors at: Key Laboratory of Advanced Technologies of Materials, Ministry of Education, School of Materials Science and Engineering, Southwest Jiaotong University, Sichuan, Chengdu, 610031, PR China. Tel.: +86-28-87634673; E-mail address: [email protected] (Y.B. Zhang), [email protected] (D.D. Yin). 1
1. Introduction
As the lightest structural metal, the ductility of magnesium (Mg) and its alloys at room temperature is limited due to lack of slip system [1]. Many efforts have been done to develop high-performance Mg alloys, and the Mg-Zn-Y-(Zr) alloy containing in situ I-phase during solidification exhibits great potential mechanical properties [2-6]. The I-phase particles are believed to be the keystrengthening factor due to theirhigh hardness, thermal stability and coherent interface between the α-Mg matrix, which have strong pinning effects on dislocations and high resistance to grain coarsening at high temperatures [7]. Generally, after extruding, rolling or forging, I-phase exists in the α-Mg matrix as a micron-scale particle. According to Zener equation [8], the pinning force (Pz) is proportional to the volume fraction of the second phase ( ) and inversely proportional to particle diameter ( ). Barlow et al. [9] observed that the introduction of nanoscaledispersoids into an aluminum matrix could accelerate the grain refinement because of enhanced dislocation generation and the reduction of the slip distances. Therefore, the refined size of the I-phase from micron scale to nano scale is likely to further improve the mechanical properties of Mg alloys containing quasicrystal I-phase.
Severe plastic deformation (SPD), such as high-pressure torsion (HPT) [10], cyclic-extrusion compression (CEC) [11] and equal channel angular pressing (ECAP) [12, 13], has been proved as an effective method for grain refinement and second phase fragmentation. Zheng et al. [14] reported that ECAP could refine the grain size effectively, break and disperse the I-phase in Mg-Zn-Y alloy, resulting in an in-situ I-phase/magnesium composites Tang et al. [15] reported the superplastic behavior of Mg-5.8Zn-1.0Y-0.48Zr alloy after ECAP with a large elongation to failure of about 800%. Garcés et al. [16] found hot forging could refine the grain size of as-cast Mg-Zn-Y alloy, but it was not effective to obtain a 2
homogeneous I-phase distribution within the magnesium matrix. Kwak et al. [17] tried to refine both α-Mg matrix and I-phase through high-ratio differential speed rolling (HRDSR), which showed that the refinement effect on the grain size was significant. However, it was found that the most of the I-phase particles were still micron scale and the uniform distribution in the matrix appeared to be poor after HRDSR. The formation of micro scale I-phase particles may be related to the micro scale lamellar spacing of the eutectic phase before SPD. Our previous study [18] observed an (α-Mg + I-phase) nano-eutectic phase of Mg-Zn-Y alloys cooling in water from semisolid state, and the fine (α-Mg + I-phase) quasi-eutectic phase may provide a foundation for the further fragment of I-phase.
In this study, an Mg-6.52Zn-0.95Y alloy reinforced by I-phase was fabricated byconventional extrusion. The combined effect of semisolid treatment and ECAP on the I-phase, grain size, texture and mechanical properties was systematically investigated.The comprehensive effect of fine-grain strengthening, I-phase precipitation strengthening and texture weakening on mechanical properties were emphatically discussed.The present work can provide a further development on the high-performance Mg alloys.
2. Experimental procedures
Pure Mg (99.99 wt. %), pure Zn (99.99 wt. %) and master alloy (Mg-30Y, wt. %) were melted in a crucible furnace under a mixed gas atmosphere of SF6 (1 vol%) and CO2 (99 vol%) and cast into an iron model of 95 mm diameter. The as-cast billets were homogenized at 420 °C for 12 h by air cooling and then the billets were machined into Ø90 mm for extrusion. Afterwards, the billets were extruded into rods with a diameter of 30 mm (extrusion ratio 9:1) at 350 °C, and the extrusion velocity of rods was 0.4 m/min.
The semisolid process was to heat the extruded samples with iron sulfide in an electric resistance furnace at 500 °C for 1 h and then cooled in water at 60 °C. The rectangular samples prepared for ECAP, 3
with the geometry of 20 × 20 × 70 mm3, were machined from the center of the semisolid bar. ECAP was carried out on the semisolid material through a die with an internal angle (φ) of 90° and a curvature angle (ψ) of 37° (the equivalent stain of each pass was ~1 [19]). A high-temperature lubricant (MoS2) was used to smear the samples and the channels. The rectangular sample was pressed through the Bc route (the sample rotated by 90° in the same direction between consecutive passes). After each pass of ECAP, the sample wasquenched in water to maintain the final microstructure and for further ECAP, it was held at the experimental temperature for 15 min, which could guarantee the experiment conducted at the stated temperature. Research [20] suggested that 4-pass ECAP could make the deformation more uniform and improve the comprehensive properties of the samples. However, with the continue increase of ECAP pass, the weakening of textures is enhanced which make basal slip in soft orientation for extruded direction. As a result, the yield strength in the extruded direction may be reduced. Two ECAP experimental methods were used, the first one was conducted at 300 °C for four passes, and the other one was conducted at 300 °C for the first two passes and 200 °C for the last two passes. For repeatability and accuracy, each test in the same condition was repeated at least three times. The sketch map of the experiment and details of ECAP were shown in Fig. 1 (a) and (b), respectively.
The microstructural observation was investigated by optical microscopy (OM, Zeiss Axio Lab A1), scanning electron microscopy (SEM, JSM-7001F) and electron backscattered diffraction (EBSD, Oxford Instrument Nordlys Nano EBSD detector and HKL channel 5 data acquisition software). Specimens for OM and SEM were mechanically polished and then etched in a solution containing 1 ml nitric acid, 1 ml acetic acid, 1 g oxalic acid, and 150 ml distilled water. For EBSD, specimens were taken through the same procedure to OM and further electropolished in a solution containing 10 vol.% perchloric acid and 90 vol.%
4
ethyl alcohol at about -40 °C to -20 °C and a voltage of 4 V for ~20 s. Constituent phases were identified by X-ray diffraction (XRD, Rigaku D/MAX 2500PC) using monochromatic Cu Kα radiation.
Tensile tests were performed at room temperature using a CMT-5105 universal test machine with a strain rate of 1 × 10-3 s-1. The specimens have a gauge length of 8 mm and cross-sectional areas of 2 × 3 mm2. The tensile axis was parallel to the extrusion direction (ED).
3. Results 3.1 Microstructure 3.1.1 Initial Microstructure
Fig. 2 shows the representative microstructure of the as-extruded Mg-6.52Zn-0.95Y alloy, and the inset is the high magnification SEM image. The as-extruded Mg-6.52Zn-0.95Y alloy exhibited a bimodal grain structure consisting of ultrafine grains with several microns and coarse unrecrystallized grains, which was also observed by other researchers [21]. Meanwhile, it was notable that fine grains were generally around the second phase. The inhomogeneous grains were usually attributed to the local occurrence of dynamic recrystallization (DRX), and it was commonly considered that second phase particles can inhibit the motion of dislocations as well as impede to the growth of newly formed recrystallized grains [22, 23].
The average grain size of the as-extruded alloy, measured by linear intercept method, is 19.9 ± 3.8 μm. The second phase distributed along the extrusion direction (ED). XRD analysis (Fig. 3(a)) indicated that the second phase was Mg3Zn6Y1 (I-phase). The size of the I-phase along ED was in the range of 1.0 - 5.0 μm and the volume fraction of I-phases was 4.8 ± 1.0%.
3.1.2 Microstructure evolution during semisolid and ECAP
5
Fig. 4a and 4d are representative OM and SEM images of Mg-6.52Zn-0.95Y alloy after semisolid treatment, and the insets in Fig. 4d are partially enlarged details. The XRD pattern of Mg-6.52Zn-0.95Y alloy after semisolid treatment shown in Fig. 3b proved that there was no change in phase constituent. As was shown in Fig. 4a and 4d, obvious eutectic structure could be found at the grain boundary of α-Mg, and the unfused α-Mg was not spherical but polygonal. The average grain size of the Mg-6.52Zn-0.95Y alloy after semisolid was 91.0 ± 6.5 μm. Under water-cooling condition at 60 °C, an ultra-fine eutectic phase formed at the grain boundaries with the lamellar spacing of 100 ± 19 nm. The volume fraction of the eutectic structure measured by image analyzer software was 4.0%, indicating that the volume fraction of I-phase was 2.0% (assume that the volume fraction of the I-phase and α-Mg in eutectic structure were equal [24]).
The typical microstructures of the semisolid Mg-6.52Zn-0.95Y alloy after 4 passes of ECAP at 300 °C are shown in Fig. 4b and 4e. The OM microstructure showed that grains were refined significantly to a level of 7.2 ± 0.8 μm (average grain size). However, there are still coarse unrecrystallized grains, and the volume fraction of unrecrystallized coarse grains was 19.7 ± 0.3%. Similar inhomogeneous microstructures were also discovered in AZ31 [25], ZK60 [26], Mg-Al-Y-Zn [27] and Mg-Gd-Nd-Zn-Zr[28] alloys after ECAP. The formation of bimodal grain structure was related to the processing temperature, the number of repetitive processing and non-uniform strain caused by the outer rounded corner [29]. The black particles in Fig. 4b were I-phase particles, and its XRD pattern was shown in Fig. 3d. It was clear that the lamellar eutectic phases were fragmented into 50-500 nm granular particles. Moreover, the distribution of the broken particles was in bands along the extrusion direction of ECAP.
Fig. 4c and 4f are OM and SEM images of Mg-6.52Zn-0.95Y alloy after 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. Comparing with the former process, the α-Mg matrix was further refined, with the average grain size of 2.9 ± 0.4 μm, and banded distributed broken I-phase particles were 6
also 50-500 nm in size. Whereas, it was noteworthy to see that there were much finer and more uniform nanoscale I-phase particles distributed throughout the α-Mg matrix. Zeng et al [30] observed the precipitation behavior of Mg-6Zn-1.5Y-0.5Zr alloy and it suggested that the aging process at 180 °C had an effect on the hardness of extruded Mg-Zn-Y-Zr alloys because of the precipitates. Meanwhile, the same precipitation behavior of I phase was also discovered in Mg-5Zn-0.92Y-0.16Zr alloy during ECAP at 200 °C [21]. It was reasonable to hypothesize that the present uniformly distributed I-phase particles was attributed to the occurrences of dynamic aging at lower temperature (200 °C) ECAP.
Overall, these results indicated that shear stress induced by ECAP promoted the refinement and dispersion of the I-phase particles. Meanwhile, the ECAP temperature had an important effect on grain refinement and precipitation of I phase. Comparing the relative intensities of the three most prominent diffraction peaks of (10 ̅ 0), (0002) and (10 ̅ 1) of Mg-6.52Zn-0.95Y alloy after different processes in Fig. 3, it could be easily found that after 4 passes ECAP at 300 °C, there was a sharp increase in the intensity of (0002) peak, while the intensity of the other two peaks fell after 4 passes ECAP at 300 °C. However, the specimen showed a resultant reversal in the relative intensities after 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. The change of diffraction peaks was correlated to textures formed at different temperatures [25, 31], and the texture evolution will be discussed in the next section.
3.1.3 Texture The inverse pole figure maps (IPF map), misorientation angle distribution histogram, {0001} pole figure (PF) and inverse pole figure for the ED (ED IPF) of Mg-6.52Y-0.95Y alloy in different states in ED-TD plane are shown in Fig. 5. The black zones in Fig.5a, 5e, 5i, and 5m were unidentified I-phase 7
particles. Grain boundaries were color-coded in IPF maps according to their misorientation angle, and the white lines and black lines correspond to low-angle grain boundaries (LAGBs, 2°≤θ<15°) and high-angle grain boundaries (HAGBs, θ≥15°). The misorientation angle quantitative distribution of as-extruded Mg-6.52Y-0.95Y alloy was shown in Fig. 5b, it could be seen that there are two distribution peaks at 20° and 86°, respectively. LAGBs account for 18.3% for the as-extruded state, which suggested the incomplete of DRX. The PF of as-extruded Mg-6.52Y-0.95Y alloy, as shown in Fig. 5c, exhibited a typical magnesium extrusion texture (ED // <10 ̅ 0> fiber texture), with the basal plane rotating along transverse direction (TD) so that c-axis of grains are perpendicular to the ED, which was consistent with the generally reported texture characteristics of extruded magnesium bars [32-34]. Analyzing with the ED IPF (Fig. 5d), the as-extruded Mg-6.52Y-0.95Y alloy exhibited a main peak around < ̅ 2 ̅ 0> indicating that the < ̅ 2 ̅ 0> direction of most grains was nearly parallel to the ED. The IPF map of Fig. 5e indicated that the semisolid bar exhibited a relatively random grain orientation distribution, comparing with the extruded bar. With the increase of grain size, the LAGBs frequency was decreased to 9.6%. The PF, as shown in Fig. 5g, was similar to the extruded bar. However, the basal texture at semisolid state was weaker than former and consisted of several main components. Our previous study [35] also observed similar basal texture weakening in Mg-Y alloy after annealing process, and it was ascribed to preferential growth of non-basal grains.
After 4 passes ECAP at 300 °C, the LAGBs frequency was increased to 19.7%. When the ECAP temperature was decreased to 200 °C, the LAGBs frequency further was increased to 29.1%. Mabuchi et al. [36] also discovered that the frequency of LAGBs was larger for ECAP than those for the conventional extrusion. ECAP was a type of SPD, which would introduce sever shear deformation and cause a lot of dislocations and sub-grains. A small number of repetitive processing at a low temperature would decrease 8
the volume fraction of DRX and simultaneously increase the frequency of LAGBs. The PF of Mg-6.52Y-0.95Y alloy after ECAP at 300 °C exhibited a {0002} basal plane texture with a maximum intensity of around 23.1 MRD (multiples of random distribution) and had an angle of about 82° with ED. Corresponding ED IPF (Fig. 5l) exhibited a main peak around < ̅ 2 ̅ 0>. However, the PF of Mg-6.52Y-0.95Y alloy after ECAP at 200 °C in Fig. 5o exhibited that the basal plane texture, with a maximum intensity of around 13.4 MRD, was about 44° with ED. Meanwhile, corresponding ED IPF (Fig. 5p) showed two obvious poles around <11 ̅ 2> and <10 ̅ 1>. Yu Yoshida [25, 31] also discovered the same texture development of AZ31 magnesium alloy during ECAP. He ascribed the differences between textures to plastic deformation mechanism of magnesium alloy at different temperatures. At low temperature (200 °C), the formation of texture was mainly dominated by tensile twinning and basal slip; While, at high temperature (300 °C), < c + a > slip predominated at the initial stage of the deformation. Finally, after 8 passes ECAP, a basal plane texture, about 45° or 90° with ED, was generated. Considering the same textures in this research, it is reasonable to hypothesize that it is caused by different plastic deformation mechanisms at different ECAP temperatures.
3.2 Mechanical properties
Fig. 6 shows representative tensile stress-strain curves of the Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP when testing along ED at room temperature. The corresponding mechanical properties of tensile yield strength (TYS), ultimate tensile strength (UTS) and elongation to failure (EL) are summarized in Table 1. As Table 1 shows, after semisolid treatment, the mechanical properties of Mg-6.52Zn-0.95Y alloy were decreased because of the obvious grain growth. After subsequent ECAP, the comprehensive mechanical properties of Mg-6.52Zn-0.95Y alloy was increased with ECAP pass at 300 °C, when the ECAP temperature was decreased to 200 °C, the UTS, TYS, and EL was increased more 9
significantly. The alloy exhibited best tensile properties with UTS of 385 MPa, TYS of 290 MPa and EL of 22.0%. There are 24.2%, 35.5%, and 56.0% increment of UTS, TYS, and EL, respectively, comparing with the as-extruded alloy.
Fig. 7 illustrates the SEM images and corresponding backscattered electron images of fracture morphologies of the tensile specimens, and the bright particles in Fig. 7d, 7e and 7f are I-phase particles. As can be seen in the Fig. 7, the fracture features were similar for all alloys, they were mainly composed of tear ridges and dimples, which suggested that the specimens underwent a substantial plastic deformation prior to fracture. However, more highly developed tearing ridges and tiny as well as deeper dimples were observed after ECAP, which was consistent with the significant improvement of elongation, as aforementioned in Table 1. Turning to I phase, it could be obviously seen that the I-phase particles after ECAP were finer andevenly distributed at the bottom of the dimples, resulting in the aggrandizement of strength and plasticity.
4. Discussions
In recent years, the interests in Mg-Zn-Y-(Zr) alloys with I-phase have been growing, mainly because of their desirable yield and ultimate tensile strength both at room and elevated temperature after thermomechanical processing, such as hot extrusion, forging or rolling. Fig. 8 is a comparison drawing in the mechanical property of current work and I-phase enhanced Mg-Zn-Y-(Zr) alloys in recent years [14, 16, 17, 24, 37-39]. Kwak et al. [24] used different speed rolling to refine the cast microstructure of Mg-9.25Zn-1.66Y alloy with I-phase, and the alloy showed high UTS approach 380 MPa but low EL of 4.0%, as shown in Fig. 8, the brown dotted box. The low EL was ascribed to the non-fully dispersed I-phase particle agglomerates. Zheng et al. [38]revealed the effect of ECAP on as-cast and as-extruded ZWK510 10
(Mg-5.0Zn-0.9Y-0.2Zr, wt. %) alloy containing I-phase. The result showed that both strength and ductility of the as-cast ZWK510 alloy were improved after ECAP, however, with the increase of EL of the as-extruded ZWK510 alloy after ECAP, its strength was decreased significantly. The as-ECAPed ZWK510 alloy exhibited best tensile properties with UTS of 330 MPa and EL of 13.0%, as shown in Fig. 8 indicated by the red points. Texture modification was considered to be the main reason for the decrease of UTS.
Comparing with the previous studies, it could be seen that the status of material before SPD affected their processed performance, and the as-semisolided Mg-6.52Zn-0.95Y alloy after ECAP showed simultaneous increase in strength and ductility, with UTS of 385 ± 18 MPa, TYS of 290 ± 3 MPa and EL of 22.0 ± 1.6%. As shown in Fig. 4d, the formation of ultra-fine eutectic phase provided the foundation for the generation of nanoscale I-phase particles. With the subsequent high shear strain during ECAP, the ultra-fine eutectic phase was broken into nanoscale I-phase particlesand dispersed into α-Mg matrix. The refinement and random distribution of I-phase particles, which can effectively reduce stress concentration, could guarantee the improvement of ductility.
Fig. 9 shows the SEM image of the semisolid Mg-6.52Zn-0.95Y alloy after 1 pass ECAP at 300 °C. It could be clearly seen that the eutectic phases were broken into particles initially and a continuous cluster of recrystallized α-Mg grains with the size about 1 ~ 3 μm preferentially generated around the I-phase particles. Generally, the main role of I-phase particles was to stimulate the recrystallization through the particle stimulated nucleation (PSN [40, 41]), in which nucleation of recrystallized α-Mg grains occurred preferentially over I-phase particles at an early stage. The presence of I-phase particle promoted the grain refinement and improved the strength of as-ECAPed Mg-6.52Zn-0.95Y alloy. Alok Singh [42] and Kwak[24] also discovered that α-Mg grain preferentially nucleated on I-phase surfaces after high-pressure torsion (HPT) and the uniform dispersion of the I-phase partlclewasimportant to the grain refinement. 11
Besides the I-phase precipitation strengthening, grain size refinement was considered to be a major strengthening mechanism. Generally, Hall-Petch relation was often used to explain the strength of a material. Fig. 10 is Hall-Petch relation for Mg-6.52Zn-0.95Y alloy after semisolid and ECAP treatment, and Hall-Petch relation of I-phase enhanced Mg-13Zn-1.55Y alloy [17], Mg-8.32Zn-1.47Y alloy [43] and Mg-7.58Zn-1.72Y alloy [44] are also plotted. As was shown in Fig.10, it is different from the as-extruded or HRDSRed Mg-Zn-Y alloys,the relation between the TYS and the average grain size could not be explained by Hall-Petch relation. Similar abnormal relation was also observed in AZ31 [31] and AZ61 [34, 45] alloys. Texture modification during ECAP was thought to be responsible for the above results.
In general, the textures of hcp metals after ECAP had a significant effect on mechanical properties, especially on yield stress and elongation. Main slip systems in magnesium and its alloys are [46]: basal slip ({0001} <11 ̅ 0>), prismatic slip ({10 ̅ 0} <11 ̅ 0>) and second order pyramidal slip ({11 ̅ 2} <11 ̅ 3>). Beside these slip systems, mechanical twinning played an important role in plastic deformation of magnesium alloys, including extension twins, contraction twins and double twins [32, 47]. It is generally accepted that the dominating deformation mechanism at room temperature are extension twinning and basal slip, for the critical resolved shear stress (CRSS) at room temperature follow the relationship: CRSS CRSS
extension twinning<
CRSS
prismatic<
CRSS
pyramidal[48,
basal<
49]. The Schmid factor (SF) evolution of basal slip
and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP treatment is shown in Fig. 11. After ECAP, the weighted average SF for extension twinning was decreased slightly whereas the SF for basal slipwas increased obviously; with the decrease of ECAP temperature, the SF for basal slip was increased from 0.19 to 0.40. It is believable that the lower TYS and higher EL of Mg-6.52Zn-0.95Y alloy after 3 ~ 4 passes ECAP may be ascribed to the highest average SF of the basal slip. However, another point to be noted in Table 1 was that the TYS of the as-ECAPed Mg-6.52Zn-0.95Y alloy at 200 °C was notably 12
higher than that of the as-extruded state, and the difference may be attributed to the obvious grain size refinement at low temperature.
As a conclusion, severe shear strain introduced by ECAP and PSN effect of I-phase particles led tothe grain refinement. Meanwhile, the formation of weakened basal texture and randomly distributed nanoscale I-phase particles prevented alloy fracturing at an early age. The combined effect of fine grain, texture and I-phase resulted in a simultaneously increase in the strength and plasticity of Mg-6.52Zn-0.95Y alloy.
5. Conclusions
The effects of semisolid and ECAP on the I-phase, microstructure and mechanical properties of Mg-6.52Zn-0.95Y alloy were systematically investigated. The main results of this study are outlined as follows:
(1) After the as-extruded Mg-6.52Zn-0.95Y alloy was semisolid treated at 500 °C for 1 h and then cooled in water at 60 °C, ultra-fine eutectic phases in the alloy were generated at the grain boundaries with the lamellar spacing of 100 ± 19 nm.
(2) Fine grain with average size of 2.9 ± 0.4 μm and high angle boundary fraction of 0.71 were obtained after semisolid treatment and 2 passes ECAP at 300 °C followed by 2 passes ECAP at 200 °C. The as-ECAPed Mg-6.52Zn-0.95Y alloy, with a lower fraction of I-phase (2.0%) and high ductility, showed the UTS of 385 ± 18 MPa, TYS of 290 ± 3 MPa and EL of 22.0 ± 1.6%.
(3) The excellent mechanical properties of Mg-6.52Zn-0.95Y alloy after semisolid treatment and ECAP, were attributed to grain refinement, the formation of weakened basal texture and the refinement of I-phase particles. 13
(4) The fragmentation of ultra-fine eutectic phase and the random distribution of I-phase particles reduced stress concentration and prevented alloy fracturing at an early age. Meanwhile, I-phase particles contributed the grain refinement during dynamic recrystallization by PSN.
Acknowledgements This work was supported by National Nature Science Foundation of China (Nos. 51201142 and 51401172) and Fundamental Research Funds for the Central Universities (No. 2682016CX073).
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Figure captions
Fig. 1 (a) The sketch maps of extrusion, semisolid treatment and ECAP processing; (b) experimental details of ECAP (ED: extrusion direction; TD: transverse direction; ND: normal direction). Fig. 2 The optical image of as-extruded Mg-6.52Zn-0.95Y bar at TD-ED plane, the top right picture is the zoomed-in SEM images. Fig. 3 XRD patterns of Mg-6.52Zn-0.95Y alloy after different processes. Fig. 4 Optical micrographs (a, b, c) and SEM images (d, e, f) of Mg-6.52Zn-0.95Y alloy after different processes
at
TD-ND
plane
(a,
d)
semisolid;
(b,
e)
semisolid+ECAP-300
°C-P4;
(c,
f)
semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 5 EBSD orientation maps (a, e, i, m) and corresponding misorientation angle distribution histograms (b, f, j, n), {0001} pole figures (c, g, k, o), inverse pole figures for the ED direction (d, h, l, p) of Mg-6.52Zn-0.95Y alloy after different processes; (a, b, c, d) as-extruded; (e, f, g, h) semisolid; (i, j, k, l) semisolid+ECAP-300 °C-P4; (m, n, o, p) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 6 Engineering stress-strain curves of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP processing. Fig. 7 Secondary electron images (a, b, c) and corresponding backscattered electron images (d, e, f) of fracture surface of Mg-6.52Zn-0.95Y alloy after different processes (a, d) as-extruded; (b, e) semisolid+ECAP-300 °C-P4; (c, f) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). Fig. 8 Comparison of I-phase enhanced Mg-Zn-Y-(Zr) alloys (HRDSR: high-ratio differential speed rolling). Fig.9 The SEM image of the Mg-6.52Zn-0.95Y alloy after 300 °C 1 pass ECAP. Fig. 10 The Hall-Petch relation for the current alloy, I-phase enhanced Mg-13Zn-1.55Y alloy, Mg-8.32Zn-1.47Y alloy and Mg-7.58Zn-1.72Y alloy.
17
Fig. 11 Schmid factor (SF) evolution of basal slip and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP processing.
Fig. 1 (a) The sketch maps of extrusion, semisolid treatment and ECAP processing; (b) experimental details of ECAP (ED: extrusion direction; TD: transverse direction; ND: normal direction).
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Fig. 2 The optical image of as-extruded Mg-6.52Zn-0.95Y bar at TD-ED plane, the top right picture is the zoomed-in SEM images.
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Fig. 3 XRD patterns of Mg-6.52Zn-0.95Y alloy after different processes.
Fig. 4 Optical micrographs (a, b, c) and SEM images (d, e, f) of Mg-6.52Zn-0.95Y alloy after different processes at TD-ND plane (a, d) semisolid; (b, e) semisolid + ECAP-300 °C-P4; (c, f) semisolid + ECAP(300 °C-P2 + 200 °C-P2).
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Fig. 5 EBSD orientation maps (a, e, i, m) and corresponding misorientation angle distribution histograms (b, f, j, n), {0001} pole figures (c, g, k, o), inverse pole figures for the ED direction (d, h, l, p) of Mg-6.52Zn-0.95Y alloy after different processes; (a, b, c, d) as-extruded; (e, f, g, h) semisolid; (i, j, k, l) semisolid+ECAP-300 °C-P4; (m, n, o, p) semisolid+ECAP-(300 °C-P2 + 200 °C-P2). 21
Fig. 6 Engineering stress-strain curves of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP processing.
Fig. 7 Secondary electron images (a, b, c) and corresponding backscattered electron images (d, e, f) of fracture surface of Mg-6.52Zn-0.95Y alloy after different processes (a, d) as-extruded; (b, e) semisolid+ECAP-300 °C-P4; (c, f) semisolid+ECAP-(300 °C-P2 + 200 °C-P2).
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Fig. 8 Comparison of I-phase enhanced Mg-Zn-Y-(Zr) alloys (HRDSR: high-ratio differential speed rolling).
Fig.9 The SEM image of the Mg-6.52Zn-0.95Y alloy after 300 °C 1 pass ECAP.
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Fig. 10 The Hall-Petch relation for the current alloy, I-phase enhanced Mg-13Zn-1.55Y alloy, Mg-8.32Zn-1.47Y alloy and Mg-7.58Zn-1.72Y alloy.
Fig. 11 Schmid factor (SF) evolution of basal slip and extension twinning of Mg-6.52Zn-0.95Y alloy after semisolid and ECAP processing.
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Table 1 UTS, TYS (σ0.2), and EL of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid treatment and ECAP processing Table 1 UTS, TYS (σ0.2) and EL of Mg-6.52Zn-0.95Y alloy after extrusion, semisolid and ECAP State of Mg-6.52Zn-0.95Y alloy As-extruded Semisolid Semisolid+ECAP-300℃-P1 Semisolid+ECAP-300℃-P2 Semisolid+ECAP-300℃-P3 Semisolid+ECAP-300℃-P4 Semisolid+ECAP-(300℃-P2+200℃ -P1) Semisolid+ECAP-(300℃-P2+200℃ -P2)
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UTS (MPa) 310 ± 8 240 ± 5 290 ± 4 310 ± 17 330 ± 20 335 ± 11 355 ± 13 385 ± 18
YS (MPa) 214 ± 13 179 ± 14 224 ±5 236 ±6 206 ±6 209 ±4 220 ±7 290 ±3
EL (%) 14.1 ± 2.0 7.3 ± 1.3 12.5 ± 2.7 15.5 ± 3.0 21.3 ± 0.6 20.1 ± 1.3 21.1 ± 2.7 22.0 ± 1.6