Effect of Al-Sr master alloy on the formation of long period stacking ordered phase and mechanical properties of Mg-Gd-Zn alloy

Effect of Al-Sr master alloy on the formation of long period stacking ordered phase and mechanical properties of Mg-Gd-Zn alloy

Materials Science & Engineering A 738 (2018) 125–134 Contents lists available at ScienceDirect Materials Science & Engineering A journal homepage: w...

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Materials Science & Engineering A 738 (2018) 125–134

Contents lists available at ScienceDirect

Materials Science & Engineering A journal homepage: www.elsevier.com/locate/msea

Effect of Al-Sr master alloy on the formation of long period stacking ordered phase and mechanical properties of Mg-Gd-Zn alloy

T

Liping Biana,b, , Yuanliang Zhaoa,b, Yichen Zhouc, Tao Wangd, Lipeng Wanga,b, Wei Lianga,b ⁎

a

College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, Shanxi 030024, PR China Shanxi Key Laboratory of Advanced Magnesium-based Materials, Taiyuan, Shanxi 030024, PR China c School of Mechanical Engineering, Xi’an Jiaotong University, Xi’an, Shaanxi 710049, PR China d College of Mechanical Engineering of Taiyuan University of Technology, Taiyuan, Shanxi 030024, PR China b

ARTICLE INFO

ABSTRACT

Keywords: Mg-Gd-Zn alloy LPSO Al-Sr master alloy Microstructure Mechanical properties

The rapid formation of profuse fine Long Period Stacking Ordered Structure (LPSO) lamellae is especially desired for Mg-Gd-Zn alloys as type II wherein LPSO phases are difficult to form in as-cast state and prolonged heat treatments are necessary. The aim of this work is to investigate the effect of various Al-10 Sr master alloy additions (0, 0.6, 1.8, in wt. %) on LPSO formation in as-cast Mg-14.02Gd-2.33Zn alloy and the microstructural evolution during subsequent heat treatment and ECAP processing as well as their mechanical behavior variations. Minor addition of 0.6 wt. % Al-Sr master alloy introduces abundant 18R and 14H-LPSO phases in as-cast Mg-Gd-Zn alloy, and accelerates the microstructural evolution into high density of well-aligned long fiber-like 14H nanolamellae only through short-time solid solution and 1-pass ECAP processing as well as low-temperature aging. While a higher level of 1.8 wt. % (Al-Sr) addition produces short rod-like 14H particulates and considerable hard Al2Gd, network (Mg, Al)3Gd intermetallic compounds deteriorating the mechanical property. Consequently, processed Mg-Gd-Zn-0.6(Al-Sr) alloy exhibits a superior mechanical property due to synergistic effects of long fiber-like LPSO phase strengthening and grain refinement strengthening and so on. This paper provides an effective and economical method to fabricate high-performance Mg-Gd-Zn alloys reinforced with LPSO phases.

1. Introduction Magnesium and its alloys have widespread applications as light alloys to substitute some conventional structural materials for weight reduction in vehicles such as cars, trucks, trains and aircrafts and so on, owing to their low density, excellent machinability and other advantages [1,2]. Due to the unique crystallographic features and excellent strengthening and toughening effects on Mg alloys, long period stacking ordering (LPSO) structure have received considerable interest in recent years and their strengthening mechanisms were studied [3,4]. As a type II alloy, however, LPSO phases were seldom observed in the as-cast Mg-Gd-Zn alloy which were generally formed by solid solution treatment and aging heat treatment for prolonged time [5–9], although there were subsequently a few reports about LPSO structures in as-cast Mg-Gd-Zn system alloys [10–12]. For example, as-cast Mg96.5Zn1Gd2.5 alloy consists of α-Mg and Mg3Gd (D03, a = 0.700 nm) phases. Heat treatment produces the coherent precipitation of a 14H-type LPSO structure. Warm-extrusion of Mg96.5Zn1Gd2.5 (at. %) alloys with LPSO structure results in the high dispersion of the bent LPSO structure and the



refinement of LPSO structures and α-Mg grains and thus the alloy exhibits high strength (345 MPa) and large elongation (6.9%) [7]. In terms of the formation of LPSO phase in as-cast Mg-Gd-Zn alloys, Wu et al. [10] reported that as-cast Mg96.82Gd2Zn1Zr0.18 alloy produced by conventional ingot metallurgy mainly comprise α-Mg solid solution, fine-lamellae 14H-type LPSO and (Mg, Zn)3Gd eutectic phase. Du et al. [11] reported that as-cast Mg96Gd3Zn1 alloy produced by induction melting process is composed of the α-Mg solid solution, lamellar 14H-type LPSO structure within α-Mg matrix, lamellar 18R-type LPSO structure at grain boundaries and WMg3Zn3Gd2 phase. However, there are still considerable amounts of coarse eutectic compounds remained in these as-cast Mg-Gd-Zn alloys and subsequent solid solution and aging heat treatment is still necessary to complete the transformation of the remaining eutectic compounds into LPSO structures. In our previous work, AlCa2 can promote the almost entire transformation of eutectic compound into 18R-LPSO in as-cast Mg-Gd-Zn alloy [12]. These results indicate that whether LPSO structure is formed in as-cast MgGd-Zn alloys or not is highly dependent on the alloy composition. Al element is often used in Mg-RE-Zn alloys for grain refinement,

Corresponding author at: College of Materials Science and Engineering, Taiyuan University of Technology, Taiyuan, Shanxi 030024, PR China. E-mail address: [email protected] (L. Bian).

https://doi.org/10.1016/j.msea.2018.09.072 Received 28 July 2018; Received in revised form 19 September 2018; Accepted 20 September 2018 Available online 21 September 2018 0921-5093/ © 2018 Elsevier B.V. All rights reserved.

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particulate strengthening [13,14], weakening the texture of Mg alloys [15], improving corrosion resistance [16] and so on. Similarly, Sr as a surface-active element can exert the effects analogous to REs when alloyed into Mg, such as refining the Mg matrix and second phases [17,18], improving the corrosion and oxidation resistance of Mg-RE-Zn alloys. Moreover, both Al and Sr can decrease the stack fault energy (SFE) of Mg, even more significant for Sr [19]. While Sr inhibited the formation of LPSO phases in LPSO-contained Mg-6Gd-3Y-0.1Zn-xSr (x = 0, 0.2, 0.6) alloys [20]. But pure Al (0.3 at. %) and Sr (0.1 at. %) additions can refine and increase the volume fraction of LPSO phases in as-cast Mg-Gd-Ni alloy containing LPSO [21]. The above effects of Al and/or Sr on the formation of LPSO in Mg alloys, however, are involved in as-cast Mg/LPSO alloys, there is no report about the influence of AlSr coaddition on Mg-Gd-Zn alloy without LPSO in as-cast state. In the present work, we are focused on the effect of Al-Sr master alloy on the formation of LPSO in as-cast Mg-Gd-Zn alloy, and further investigate the microstructural evolution during the subsequent solid solution heat treatment and severe plastic deformation induced by ECAP and their mechanical responses. Mg96.5Gd2.5Zn1 alloy is typical in Mg-Gd-Zn alloys containing LPSO. Also, Mg/LPSO alloys are generally processed by solid solution treatment, plastic deformation and aging for profuse fine lamellar LPSO formation and structural modification to achieve excellent mechanical properties. Hence here, Mg-14.02Gd2.33Zn-x(Al-10 Sr) (x = 0,0.6,1.8, in wt. %) alloys with Gd/Zn atomic ratio of 2.5:1 are designed. A short solid solution duration of 10 h at 500 °C, and only 1-pass ECAP processing as well as low-temperature aging at 200 °C were adopted to investigate the influence of Al-Sr additions on the subsequent microstructural evolution rate. Furthermore, excessive plastic deformation will lead to serious dynamic recrystallization (DRX) in Mg/LPSO alloys and soften the alloys too.

(TEM, JEOL 2010) and inductively coupled plasma atomic emission spectroscope (ICP-AES). Phase constituent analysis was performed with Y-2000 × -ray diffraction (XRD) using monochromatic Cu-Kα radiation. The samples for OM and SEM observation were machined from the longitudinal or transversal sections of the billets and etched with a solution of 5 ml nitric acid +95 ml ethyl alcohol before observation. Microhardness testing was measured by HVS-1000A Vickers hardness testing machine with load of 300 gf and loading time of 15 s. The dogbone-shaped tensile samples have the gauge length of 15 mm, the width of 3 mm and the thickness of 2 mm. The tensile tests were carried out using CMT-5105 electronic universal testing machine at room temperature with an initial strain rate of 0.4 mm/min. The grain size and phase fraction were measured and determined by linear intercept method and image analysis technique, respectively. Thin foils for TEM observation were prepared by cutting the bulk sample into slices, grinding to the thickness of about 50 µm, followed by twin-jet electropolishing with 5% perchloric acid and 95% ethanol solution and low-angle ion milling using the ion beam milling system Leica EM RES102 finally. 3. Results and discussion 3.1. Microstructures of as-cast alloys Fig. 1 and Fig. 2 show the microstructural features of the as-cast alloys with different Al-Sr contents. Typical constituent phases in the studied alloys, as indicated by white arrows A~F in Fig. 1 and Fig. 2, are identified further by EDS analysis (Table 2) and XRD patterns (Fig. 3). It is shown that phase constituents, the morphology and distribution of component phases in the alloys are pronouncedly influenced by the Al-Sr additions. As shown in Fig. 1(a) (b) combined with Table 2 and Fig. 3, (Al-Sr)free alloy I mainly consists of α-Mg solid solution (pointed by arrow A), coarse integrated W-Mg3Zn3Gd2 network eutectic structure (pointed by arrow B) [12,22] at α-Mg interdendritic boundaries and a little fine lamellae 14H (pointed by white arrow in Fig. 1(b)) [23] around the eutectic structure. While alloyed by Al-Sr master alloy, the intergranular phase in alloy II (Fig. 1(c) (d)) is substituted for relatively isolated gray short 18RLPSO plates (indicated by arrow C, identified as Mg10Gd1Zn1 by EDS and the report [11]) within which small amounts of residual WMg3Zn3Gd2 eutectic are laminated. Further, abundant short 14H lamellae are formed around the α-Mg grains which are extending from grain boundaries to interior grains to form a halo structure. In addition, new bright stamens-like phase and polygonal particle phase are present in the α-Mg interior grains, which are identified as (Mg, Al)3Gd and Al2Gd by EDS of arrows D and E as well as the work [24], respectively. They are prone to form in Mg alloys with small content of Al and large amount of Gd additions due to the difference of the Pauling electronegativity (Mg:1.31; Zn: 1.65; Gd:1.20; Al: 1.61). Meanwhile, the dendritic α-Mg grains tend to be finer and spherical with an average grain size of 57 µm, indicating much more dispersed 18 R particles formed at grain boundaries. Fig. 2(a) presents a block 18 R phase in alloy II at higher magnification wherein some straight light gray needlelike phases indicated by arrow F are observed, which is identified as Mg10GdAlZn by EDS analysis. Alloy III shown in Fig. 1(e)(f) presents the similar microstructure to that of alloy II in constituent phases but with the more profuse polygonal Al2Gd particulates and developed stamen-like (Mg, Al)3Gd eutectic structure which takes on flake shape at the end of the branches [also in Fig. 2(b)]. The stamens-like (Mg, Al)3Gd phases are regarded as the aggregation of discontinuous fine strip-like phases. Polygonal Al2Gd particles in as-cast Mg-Gd-Al-(Zn) alloys have been researched by Chen [25] and Pourbahari [26]. Al2Gd has high melting point (1525 ℃), which has grain refining effect on Mg alloys by inhibiting the migration of grain boundaries. However, the preferential formation of (Mg,

2. Experimental procedures The nominal compositions of the present alloys were Mg-14.02Gd2.33Zn, Mg-14.02Gd-2.33Zn-0.6(Al-Sr), Mg-14.02Gd-2.33Zn-1.8(AlSr), named as alloy I, II and III, respectively, as listed in Table 1. The alloys were melted in a magnetic levitation induction melting furnace under the protection of argon atmosphere with high-purity Mg (99.98 wt. %, Zn (99.98 wt. %), Mg-30Gd and Al-10 wt. % Sr master alloys as raw materials. The alloy components were mixed fully and then poured into carbon crucible to prepare cast ingots with dimension of φ30 × 65 mm. The as-cast ingots were machined into 12.2 × 12.2 × 65 mm3 for solid solution treatment (T4) by wrapping with aluminum foil and being embedded in fine graphite powder in an air furnace at 773 K for 10 h, immediately water quenching. Then the solid solution samples were mechanically abrased to form the dimensions of 12 × 12 × 65 mm3 for ECAP extrusion. ECAP extrusion was performed by using a die with φ = 90° (inner arc of curvature) and ψ = 20° (outer arc of curvature) for 1 pass at 633 K. The average pressing speed is 4 mm/min. The equivalent strain per pass is approximately 1. ECAPed billets were subsequently aged at 473 K for 10 h, 20 h, 30 h, 40 h, respectively, and water quenched. Microstructures and chemical compositions were measured with optical microscope (OM; Leica DM2500), scanning electron microscopy (SEM; TescanMIRA3 LMH) equipped with an energy dispersive spectrometer (EDS; OXFORD X-MAX 20), transmission electron microscopy Table 1 Chemical compositions of the experimental alloys. Alloys

I II III

Nominal/ measured composition (wt%) Gd

Zn

Al

Sr

Mg

14.02/13.794 14.02/14.002 14.02/13.948

2.33/2.271 2.33/2.308 2.33/2.294

0.00 0.54/0.493 1.62/1.597

0.00 0.06/0.056 0.18/0.168

Bal. Bal. Bal.

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(b)

(a)

Fine lamellae

A:α-Mg

(c)

B:Eutectic structure

(d) C:Block phase

Al2Gd

(e)

(f) D:(Mg,Gd)3Al

E:Al2Gd

Fig. 1. SEM back scattered electron images of the as-cast alloys: (a-b) Alloy I, (c-d) Alloy II and ( e-f) Alloy III.

Al)3Gd and Al2Gd intermetallic compounds with high content of Al addition consumes a certain amount of Gd in the melt and hence decreases the volume fraction of Gd-rich long period stacking ordered

(a)

(LPSO) phase and W-Mg3Zn3Gd2 eutectic phase in alloy III. XRD patterns (Fig. 3) show that the diffraction peaks corresponding to 18R-LPSO phase (Mg10GdZn), (Mg, Al)3Gd phase, Al2Gd and Al4Sr

(b)

F Rod-like phase Flake-like phase

Fig. 2. SEM back scattered electron images of the as-cast alloys at high magnifications: (a) Alloy II, ( b) Alloy III. 127

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and then reduces slightly with the increase of Al-Sr additions owing to coarsening and the growth of the intermetallic compound phases.

Table 2 EDS analysis results of A~F constituents indicated by white arrows in Fig. 2 and Fig. 3 (at. %). position

Mg

Gd

Zn

Al

Sr

Phase

A B C D E F

98.0 68.6 88.6 51.0 5.6 84.4

1.3 17.0 4.6 19.7 30.9 7.7

0.7 14.4 6.4 2.2 2.0 6.6

0.0 0.0 0.4 27.1 61.5 1.3

0.0 0.0 0.0 0.0 0.0 0.0

α-Mg Mg3Zn3Gd2 Mg10GdZn (Mg, Al)3Gd Al2Gd Mg10GdAlZn

3.2. Microstructures of solid-solution alloys Fig. 4 shows the microstructural evolution of the solid-solution alloys. As shown in Fig. 4(a)(b), solid-solution alloy I consists of α-Mg solid solution, integrated network 18R-LPSO intergranular phases with some white particles embedded, and sparse short 14H lamellae along with 18R. Fig. 4(b) further identifies the constituent phases as follows. Block phase marked as 1 is 18R-Mg10Gd1Zn1 phase containing Mg (87.1 at. %), Zn (6.1 at. %) and Gd (6.9 at. %) with the ratio of Zn to Gd of 0.88, close to Zn/Gd ratio 1 for the 18R-Mg10Gd1Zn1 phase. The gray fine-lamellar phase marked as 2 is 14H containing Mg (94.9 at. %), Zn (2.2 at. %), Gd (2.9 at. %). The white strip-like phase marked as 3 is WMg3Zn3Gd2 with EDS analysis of Mg (62.1 at. %), Zn (18.7 at. %) and Gd (19.2 at. %). The white irregular particle marked as 4 is estimated to be Gd-rich particles with EDS analysis of Gd (99.5 at. %), minor Zn (0.5 at. %). The microstructure analysis indicates that most of W-Mg3Zn3Gd2 eutectic phases in as-cast alloy I have transformed into 18R LPSO during 500 °C × 10 h solid solution treatment, but only a little fraction of 18R start to transform to lamellar 14H phase. Amazingly, solid-solution alloy II as shown in Fig. 4(c)(d) comprises predominantly fine 14H lamellae penetrating through α-Mg interior grains to form ~100 µm cellular structures and small amounts of discontinuous 18R short plates located at 2H-Mg+14H cell boundaries. Residual W-Mg3Zn3Gd2 phase is substantially reduced in size and quantity. It is implied that the transformations of 18R→14H and W→ 14H have completed for solid-solution alloy II during short-time solid solution, that is, 14H formation process was greatly accelerated in solidsolution Mg-Gd-Zn alloy with minor Al-Sr addition. Fig. 4(e)(f) present the microstructure of the solid-solution alloy III at the various magnifications. Unlike long fiber-like 14H lamellae in solid-solution alloy II, densely discontinuous short rods/particles with the length of 4–32 µm and width of ~0.85 μm are formed in the α-Mg interior grains of solid-solution alloy III and aligned along Mg bases, which are identified as Mg93.3Al2.0Gd2.7Zn1.9Sr0.1 phase according to EDS analysis. Stamen-like (Mg,Al)3Gd eutectic has partially dissolved to long particles. The area fraction variations of constituent phases in the various alloys before and after solid solution treatment are shown in Table 3. The most significant increase in LPSO number by solid solution is alloy II raised from 21.3% to 74.5%, followed by alloy I from little to 16.2% and alloy III from 19.2% to 22.8%. The decreasing amplitudes of W phase in three alloys are in the order of alloy I, alloy II and alloy III. While the number of (Mg, Al)3Gd and Al2Gd phases keeps almost constant due to their high melting points with only slightly partial dissolution of (Mg, Al)3Gd and insolubilization of Al2Gd under the existing solid-solution process. The numbers of total second phases in three solid-solution alloys show increasing trends, but alloy II is the most significant with the most profuse 14H-LPSO lamellae forming.

 -Mg

 W-Mg Zn Gd





Intensity(a.u.)

 









Alloy III

LPSO phases (Mg, Al) Gd

 Al Gd





 

 









20





Alloy I







30

Al Sr

  



Alloy II





40

50 60 2θ(degree)



70

80

Fig. 3. XRD patterns of the as-cast Mg-Gd-Zn-Al-Sr series alloys.

phase are emerged in alloys II and III with Al-Sr additions. While Al4Sr particles have not been observed in Fig. 1 and Fig. 2 due to their small size and quantity. Furthermore, the intensity of the diffraction peaks corresponding to (Mg, Al)3Gd, Al2Gd and Al4Sr phase become stronger with increasing Al-Sr content while weaker for the 18R-LPSO phase and W-Mg3Zn3Gd2 eutectic phase, indicating the relative amount variations of the corresponding phases with Al-Si additions. Table 3 gives the area fractions of all phases in the as-cast alloys quantitatively. With increasing Al-Sr addition from 0.0 wt. % to 0.6 wt. %, the area fraction of 18R LPSO phases increases from 0% to 21.3%, W-Mg3Zn3Gd2 phase decreases obviously from 6.6 to 2.6. That indicates Al-Sr master alloy can significantly promote the formation of LPSO phase and inhibit the precipitation of eutectic W-Mg3Zn3Gd2 in as-cast Mg-Gd-Zn alloy. The reasons for the formation of abundant 18R LPSO in as-cast Mg-Gd-Zn alloy may be attributed to two points. Firstly, alloying elements in the Mg matrix can have a marked influence on SFEs. Besides Zn and Gd, Al and Sr can also considerably diminish the stacking fault energy (SFE) of Mg and thus produce more stacking faults so as to form more LPSO phases. Secondly, Al-Sr addition promotes the transformation of W-Mg3Zn3Gd2 eutectic into 18R in as-cast state with the evidence of residual W-Mg3Zn3Gd2 laminated within 18R plates. With increasing Al-Sr content up to 1.8 wt. %, the area fraction of 18R LPSO slightly diminishes to 19.2%, and the area fractions of (Mg, Al)3Gd and Al2Gd phase increase from 0% to 3.7% and 0.7%, respectively. The overall content of second phases first increases dramatically

3.3. Microstructures of the post-ECAP aged alloys Fig. 5 and Fig. 6 show the microstructures of various solid-solution alloys after 1-pass ECAP at 360 °C and aging at 200 °C for 30 h on longitudinal and cross-sectional planes, respectively. There is no obvious dynamic recrystallization (DRX) occurring in all three alloys except in a few local regions with particles agglomerated by particle stimulated nucleation (PSN). Fig. 5(a)(b) show that 18R intergranular phases in alloy I are progressively elongated along the shear stress direction. Only a small fraction of short 14H lamellae form around 18R shown in Fig. 6(a). Residual W inclusions in 18R (Fig. 6(b)) are still present. It is indicated that the transformations of 18R and WMg3Zn3Gd2 into 14H did not occur significantly in alloy I during severe plastic deformation.

Table 3 Area fractions of all phases in the as-cast/solid-solution alloys. Alloy

I II III

Al-Sr

0 0.6 1.8

Area fraction (%) W-phase

LPSO phases

(Mg, Al)3Gd

Al2Gd

Total content

6.6/1.0 2.6/0.6 2.3/1.6

Little/16.2 21.3/74.5 19.2/22.8

0/0 3.5/3.5 3.7/3.7

0/0 Little/little 0.7/0.7

6.6/17.2 27.4/78.6 25.9/28.8

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(b)

(a)

1 2 3 4

Lamellar LPSO phases

(c)

(d)

14H-LPSO phases

(f)

(e)

Rod-like phases

Fig. 4. OM and SEM images of solid-solution alloys: (a-b) Alloy I, (c-d) Alloy II and ( e-f) Alloy III.

Whereas a homogeneous long-fiber 14H reinforced composite structure is formed in the post-ECAP aged alloy II as shown in Fig. 5(c) (d). The original 14H+2H-Mg cellular structure with random orientations in solid-solution alloy II were elongated and aligned along the shear direction to produce a mixture of the well-aligned 14H long fibers and Mg matrix. Meanwhile, 14H and 18R LPSO phases as shown in Fig. 5(d) are also kinked severely and bent to form multiple Zigzag kink bands and bent bands, respectively, or even broken into small particles. Some thick 18R LPSO plates were delaminated, elongated and redistributed along the shear flow stress to form isolated thin long strips. TEM image of Fig. 7(a) further demonstrates the occurrences of bending, kinking and refining of 18R-LPSO phases by delamination into thin strips and fragmentation into fine particles, or decomposing/ transforming into 14H-LPSO phases. Fig. 7(b) shows a cluster of fine needle-like phases near several particles, which was identified as 14H LPSO phase from the corresponding SAED pattern of Fig. 7(e) wherein fourteen consistent diffraction spots are stacked between the transmission and diffraction spots. It was inferred that new 14H-LPSO phases have precipitated from the α-Mg matrix and grown during thermal mechanical process. High density of dislocations shown in Fig. 7(a)(c)

are piled up inside 18R and 14H-LPSO phases and α-Mg slices, indicating that these phases can effectively hinder dislocation movement so as to significantly reinforce the alloy. Fig. 7(f) shows SAED pattern of a broken plate-shaped particle in Fig. 7(d) which is determined as (Mg, Al)3Gd phase, wherein the brightest spots should correspond to atomic columns rich in Gd element since Gd is much heavier than Mg and Al. It was manifested that hard (Mg, Al)3Gd phase was fragmented into small particles during ECAP processing. It was concluded that isolated short 18R plates in alloy II are more prone to refine, redistribute and transform into 14H-LPSO during ECAP processing in contrast to integrated network 18R configuration in alloy I. The longitudinal and cross-sectional microstructures of the postECAP aged alloy III are present in Fig. 5(e)(f) and Fig. 6(d)(f), respectively. Fig. 5(e) shows that the post-ECAP aged alloy III are predominated by short 14H-type Mg93.3Al2.0Gd2.7Zn1.9Sr0.1 rods/particles and α-Mg matrix grains which comprise the coarse elongated cellular structures. Small amounts of retained intergranular 18R phases are significantly bent and elongated along the shear direction, wherein residual W eutectic compounds are broken into fine particle agglomerates as shown in Fig. 5(f) and Fig. 6(f). Developed (Mg, Al)3Gd

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(a)

(c)

(b)

(d)

(c)

Zigzag kinking bands

Bent bands (e)

(f)

Delamination

Fig. 5. Microstructure of the post-ECAP aged alloy: observed on the longitudinal section along the extrusion direction: (a-b) Alloy I, (c-d) Alloy II and ( e-f) Alloy III.

eutectic structures with several hundreds of microns are pronouncedly fragmented into large clusters of fine particles. Polygonal Al2Gd particles are distributed more homogeneously along the flow stress.

post-ECAP aged alloys are summarized in Fig. 9 and Table 4. The ultimate tensile strengths of the alloys tend to increase significantly first and then slightly diminish with increasing Al-Sr additions. And the percent elongations show the similar tendency. The post-ECAP aged alloy II exhibits the highest strength with TYS of 213 MPa and UTS of 303 MPa and the best ductility with EL of 6.7%, which are increased by 15.14%, 18.36%, 34.0% and 8.12%, 12.22%, 76.32% compared to alloys I and III, respectively. It was concluded that minor Al-10Sr master alloy can remarkedly improve the mechanical properties of processed Mg-Gd-Zn alloys by inducing profuse LPSO phase formation. The excellent combination of strength and ductility in alloy II is attributed to the dominant well-aligned 14H long-fiber ordered structure reinforcement as well as strengthening and toughening by kinking and bending of 14H and 18R [4]. Furthermore, grain refinement of Mg matrix and 18R-LPSO phase by delamination, fragmentation and homogeneous redistribution plays a significant role in strengthening and toughening the alloy. Also, a small number of dispersed secondary phase particles act as dispersion strengthening effect by inhibiting dislocation motion. Nevertheless, large amounts of the agglomerated, hard and brittle W, (Mg, Al)3Gd and polygonal Al2Gd particles in alloy III deteriorate seriously its mechanical

3.4. Mechanical property tests and analysis 3.4.1. Microhardness test-age hardening response Fig. 8 shows the microhardness variations of ECAP-post aged alloy II with various durations at 200 °C. The hardness of ECAP-post aged alloy II increases slowly at the early stage of aging, and reaches a peak of 113 HV at 30 h, followed by a sharp drop, which indicates the excellent thermal stability of the alloy II. The hardness increment is attributed to the formation of more and more long fiber-like 14H lamellae during aging process. The subsequent diminishment of hardness is caused by reduction in the length and number of 14H lamellae due to DRX occurrence in α-Mg matrix during the prolonged aging. 3.4.2. Tensile mechanical property and fracture analysis of post-ECAP aged alloys Ambient-temperature tensile mechanical properties values of various

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(a)

(b)

(c)

(d)

Gd-rich particle

(e)

(f) Delamination

Fig. 6. Microstructure of the post-ECAP aged alloys observed on the transverse section: (a-b) Alloy I, (c-d) Alloy II and ( e-f) Alloy III.

property. While the poor mechanical property in post-ECAP aged alloys I is ascribed to a small fraction of 14H LPSO and coarse 18R structure.

fracture surface with the cracks traversing both α-Mg matrix and LPSO phases. Large amounts of fine dispersed 18R plates are observed in Mg matrix from the fracture profile (Fig. 10(b) and (d)), which correspond to the fairly fine and homogeneous dimples on the cross-sectional fracture in Fig. 11(b). No interface debonding is observed between αMg/18R and α-Mg/14H indicating strong bond existing in α-Mg/18RLPSO and 2H-Mg/14H-LPSO coherent interfaces as well as the excellent intrinsic plasticity of 14H-LPSO phase. Further, cleavage fracture is not observed in processed alloy II due to significant refinement of 18R phases. As indicated by white arrow in Fig. 10(d), however, the segregated W particles embedding into 18R phases disrupt the strong bond between α-Mg/18R interface resulting in crack initiation and rapid propagation in these regions and eventually the macroscopic catastrophic failure. This case is further aggravated in the post-ECAP aged alloy III due to higher volume fractions of segregated hard and brittle eutectic particles (Mg, Al)3Gd, Al2Gd and residual W embedded in 18R plates (Fig. 11(c) and Fig. 6(f)). The post-ECAP aged alloy III exhibits a coarser fracture surface compared to alloy II as shown in Fig. 11(c). Small W phase particles with the chemical composition of Mg‐(12.7 at. %) Zn‐(15.2 at.

3.4.3. Fracture analysis of post-ECAP aged alloys Fig. 10 and Fig. 11 are SEM fractographs from the longitudinal section and cross section of various post-ECAP aged alloys to further verify the fracture mechanisms of the alloys. As shown in Fig. 10(a)(b), the post-ECAP aged alloys I and II display great differences in fracture profiles with regards to secondary phase structural features including LPSO phases and other particles. The processed alloy I presents a bumpy fracture surface with sparse 18R LPSO long strips and coarse residual W eutectic distributed at α-Mg grain boundaries which are detrimental to the mechanical property. It was confirmed further by Fig. 10(c) where the crack mainly propagates along the residual network W phases at grain boundaries, indicating the crack and fracture of the coarse network W eutectic brittle phases during ECAP deformation induces the premature failure of the processed alloy I. Moreover, cleavage fracture is observed in coarse 18R LPSO particles as shown in Fig. 11(a), implying 18R LPSO much harder than α-Mg matrix. While the post-ECAP aged alloy II exhibits a relatively smooth

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350

114

300

112

250

110 108

104

50

20 30 Aging time (h)

6 4

150 100

10

8

200

106

0

Yield Tensile stress Ultimate Tensile stress Elongation

0

40

Elongation/%

116

Stress /MPa

Microhardness(HV)

Fig. 7. TEM bright field images of (a) 18R, (b) (c) 14H LPSO and (d) (Mg, Al)3Gd phases in the post-ECAP aged alloy II, (e) (f) corresponding SAED patterns obtained from 14H LPSO and (Mg, Al)3Gd phase in (b)(d), respectively.

2

I

II

0

III

Fig. 9. Tensile properties of the post-ECAP aged alloys with different (Al-Sr) additions.

Fig. 8. Age hardening curves of post-ECAP alloy II aged at 200 ℃ for different time.

%) Gd by EDS analysis are observed to be segregated in the coarse bent 18 R strips, which is detrimental to the integrate of 18R and Mg matrix and further deteriorates the mechanical property of the alloy III.

Table 4 Ambient-temperature tensile properties of post-ECAP aged alloys.

4. Conclusion The effects of Al-Sr master alloy additions on the formation of the long period stacking ordered phase in the as-cast Mg-Gd-Zn alloy and 132

Alloys

UTS (MPa)

Elongation (%)

TYS (MPa)

I II III

256 303 270

5.0 6.7 3.8

185 213 197

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L. Bian et al.

Fig. 10. Longitudinal section near the tensile fracture surface of post-ECAP aged alloys: (a)(c) Alloy I, (b)(d) Alloy II.

(c)

(b)

(a)

Cleavage plane

Tearing ridge Dimples

W phase particles

Tearing ridge Fig. 11. Fracture surface morphology of the post-ECAP aged alloys (a) Alloy I, (b) Alloy II and (c) Alloy III.

the microstructural evolution during subsequent heat treatment and ECAP processing as well as the mechanical property variation were systematically investigated. The following conclusions have been drawn:

processing. High density of well-aligned long fiber-like 14H lamellae are formed in alloy II containing 0.6(Al-Sr) only through 1pass ECAP extrusion and low-temperature aging. While integrated network 18R intergranular phase in solid-solution alloy I are difficult to refine and only elongated, and small amounts of 14H-LPSO lamellae are produced in alloy I after ECAP processing and aging. Higher level of 1.8(Al-Sr) addition results in formation of short rodlike14H and too much hard and brittle (Mg, Al)3Gd and Al2Gd agglomerates which are detrimental to the mechanical property of alloy III. With the homogeneous long fiber-like 14H nanolamellae and dispersed fine 18R reinforced composite structure, the postECAP aged alloy II containing 0.6(Al-Sr) exhibits the highest mechanical property with YTS of 213 MPa, UTS of 303 MPa, and EL of 6.7%.

(1) Low level of Al-Sr master alloy can significantly promote the almost entirely transformation of network W-Mg3Zn3Gd2 eutectic phase into isolated 18R-LPSO plates and the formation of abundant 14H lamellae in as-cast Mg-Gd-Zn alloy. New stamen-like (Mg, Al)3Gd and polygonal Al2Gd phases are formed and increase in quantity and size with increasing Al-Sr additions. (2) Minor Al-Sr addition can greatly accelerate the microstructural evolution into lamellae 14H LPSO of Mg-Gd-Zn alloys during solid solution treatment. After short-time solid solution of 10 h at 500 °C, high density of 14H lamellae penetrating through α-Mg grains are formed by precipitation from α-Mg matrix and the transformations from 18R and residual W phase in alloy II containing 0.6 (Al-Sr). Contrastly, only the transformation from W eutectic to the network intergranular 18R are completed while the transformation from 18R to 14H is just initiated in alloy I without Al-Sr. (Mg, Al)3Gd have just partially dissovled into long particles while Al2Gd phases undissolved due to their high melting points. (3) Minor Al-Sr addition promotes the refinement and redistribution of isolated 18R-LPSO plates in solid-solution alloys by delamination, fragmentation and in situ transformation into 14H during ECAP

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