Journal of Materials Science & Technology, 2020, 54(0): 160-170 DOI: 10.1016/j.jmst.2020.04.031

Research Article

Achieving ultra-high strength in Mg-Gd-Ag-Zr wrought alloy via bimodal-grained structure and enhanced precipitation

Yu Zhang,a,b,*, Wei Rongb, Yujuan Wub, Liming Peng,b,*

a College of Materials Science and Engineering, Chongqing University, Chongqing 400044, China

b National Engineering Research Centre of Light Alloy Net Forming and State Key Laboratory of Metal Matrix Composite, School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China

Corresponding authors: *National Engineering Research Centre of Light AlloyNet Forming and State Key Laboratory of Metal Matrix Composite, School of Materi-als Science and Engineering, Shanghai Jiao Tong University, Shanghai 200240, China;College of Materials Science and Engineering, Chongqing University, Chongqing400044, China.E-mail addresses:yu.zhang@cqu.edu.cn(Y. Zhang),plm616@sjtu.edu.cn(L. Peng).

Received: 2020-03-14   Accepted: 2020-04-18   Online: 2020-10-1

Abstract

Mg-13.1Gd-1.6Ag-0.4 Zr (wt%) alloy was either iso-thermally extruded at 350 °C or differential-thermally extruded with respectively pre-heated billet at 500 °C and die at 350 °C. The iso-thermal extrusion leads to a near fully recrystallized structure and a [0001]//ED (extrusion direction) texture. In contrast, the differential-thermally extruded alloy develops a bimodal-grained structure composed of fine equiaxed recrystallized grains and coarse elongated unrecrystallized grains with a 011¯0//ED texture. The differential-thermally extruded alloy has a higher number density of precipitates after post-extrusion ageing than that of the iso-thermally extruded counterpart. Moreover, precipitation in the differential-thermally extruded alloy is further enhanced with cold rolling before ageing. Finally, the alloy obtains room temperature tensile yield strength of 421 MPa and ultimate tensile strength of 515 MPa via differential-thermal extrusion, cold rolling and ageing, mainly ascribed to the coupled strengthening from the bimodal-grained structure and enhanced precipitation. Strength of the alloy is noticeably higher than those of Mg-Gd(-Y)-Ag extruded alloys with similar compositions reported previously and is comparable to those of other high-strength Mg wrought alloys. The findings suggest that differential-thermal extrusion plus strain ageing is a suitable approach for achieving high strength in age-hardenable Mg alloys.

Keywords: Mg alloy ; Extrusion ; Bimodal-grained structure ; Precipitation ; Electron back scattering diffraction

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Cite this article

Yu Zhang, Wei Rong, Yujuan Wu, Liming Peng. Achieving ultra-high strength in Mg-Gd-Ag-Zr wrought alloy via bimodal-grained structure and enhanced precipitation. Journal of Materials Science & Technology[J], 2020, 54(0): 160-170 DOI:10.1016/j.jmst.2020.04.031

1. Introduction

With the increasing demand for lightweight structural materials, magnesium (Mg) alloys have received significant interests from aircraft, aerospace, automobile, and electronic industries [[1], [2], [3], [4]]. One of the major issues that currently restrict the larger and wider application of Mg alloys is their inadequate strength [[5], [6], [7], [8]]. As such, developing high-strength Mg alloys has been a research focus over the last two decades.

Mg-gadolinium (Gd) based alloys represent a class of high-strength Mg alloys [9]. Since the maximum solubility of Gd in Mg sharply decreases from 23.49 wt% at eutectic temperature (548 °C) to 3.82 wt% at 200 °C, Mg-Gd alloy system forms an ideal model for precipitation hardening [10]. The dominant strengthening precipitates in peak-aged Mg-Gd alloys are β′ phase, which has a base-centred orthorhombic Bravais lattice: a = 0.650 nm, b = 2.272 nm, c = 0.521 nm and an orientation relationship (OR) with α-Mg matrix: [001]β′//[0001]α and (100)β′// $\left\{ 1\bar{2}10 \right\}$α [11]. Micro-alloying of Ag in Mg-Gd alloys induces formation of precipitate phases on (0002)α planes [12]. The dominant basal precipitates in Mg-Gd-Ag alloys are γ′′ phase, which forms as thin plates with a unit thickness [13]. γ′′ phase has a hexagonal structure: a = 0.548 nm and c = 0.417 nm and an OR: $10\bar{1}0$ γ′′// / $2\bar{1}\bar{1}0$ α and (0001)γ′′//(0001)α [13]. Concurrent precipitation of β′ and γ′′ precipitates on prismatic and basal planes, respectively, constructs isolated domains in α-Mg matrix, which provides the most effective obstacles for dislocation sliding and thus significantly improves strength [14]. As such, a peak-aged Mg-15.6Gd-1.8Ag-0.4 Zr (wt%) casting alloy has reached room temperature tensile yield strength (TYS) of 328 MPa and ultimate tensile strength (UTS) of 423 MPa, the highest values on record in Mg casting alloys thus far [15]. On such a basis, Mg-Gd-Ag-Zr alloys exhibit a great potential for achieving ultra-high strength when suitable thermal-mechanical processing is conducted.

Hot extrusion is a feasible and effective way to improve the mechanical properties of Mg alloys by reducing casting defects, refining grain size, and forming favoured textures [16,17]. In general, billet and mould temperatures are designed to be equal. This extrusion strategy is hereinafter referred to iso-thermal extrusion (ITE). Pre-heating temperature of billet and extrusion die/container of most ITE for Mg-Gd based alloys ranges from 350 °C to 450 °C, while homogenisation (solution treatment) temperature for billet is usually no less than 500 °C [[18], [19], [20]]. Since the solubility of Gd in primary Mg matrix sharply decreases with decreasing temperature, relatively lower pre-heating temperature of billet inevitably leads to considerable pre-precipitation of large second-phase particles before extrusion, particularly in Mg-Gd-Ag alloys in which Ag accelerates precipitation kinetics [21]. As a result, precipitation hardening in Mg-Gd-Ag alloys via post-extrusion ageing treatment is greatly weakened, which is detrimental to hardening. While pre-precipitation of large second-phase particles prior to extrusion could be relieved if the pre-heating temperature is increased to homogenisation temperature or near homogenisation temperature, grain growth is aggressive during extrusion process, which is also detrimental to yield high strength. In addition, hot shortness or transverse cracks are prone to form at high extrusion temperatures, leading to poor extrusion quality and thus deteriorating mechanical properties.

In order to get a balanced combination of precipitation hardening, grain refinement and high extrusion quality, a new extrusion strategy has been applied [22]. Billet is pre-heated at homogenisation temperature to restrict pre-precipitation of large second-phase particles prior to extrusion and extrusion die and container are pre-heated at relatively lower temperature to control actual extrusion temperature and avoid aggressive grain coarsening. Such an extrusion strategy (hereinafter designated as differential-thermal extrusion (DTE)) has achieved a significant increment in strength of Mg-15Gd-1Zn-0.4 Zr (wt%) alloy, an alloy with comparable age-hardening response to Mg-Gd-Ag alloys [23]. As such, the DTE approach shows great potential for obtaining high strength in age-hardenable Mg-Gd-Ag alloys. However, attempts for processing Mg-Gd-Ag alloys via DTE have not been reported thus far.

In this work, Mg-13.1Gd-1.6Ag-0.4 Zr (wt%) alloy was ITE extruded at 350 °C or DTE extruded with pre-heated billet at 500 °C and pre-heated die at 350 °C. In addition, the DTE alloy was further cold rolled with a total thickness reduction of 10 %. Microstructure and mechanical properties of the alloys processed by the above three routes were comparatively studied. The alloy achieves TYS of 421 MPa and UTS of 515 MPa via DTE, cold rolling and ageing. The obtained strength of the alloy is noticeably higher than those of other Mg-Gd(-Y)-Ag extruded alloys with similar compositions and is comparable to those of other high-strength Mg wrought alloys, indicating that the combination of DTE, cold work and ageing is a suitable strategy to pursue high-strength in age-hardenable Mg alloys.

2. Experimental procedures

The alloy was prepared using pure Mg, pure Ag, Mg-90Gd (wt%) and Mg-30Zr (wt%) master alloys. Casting process was carried out using an electrical resistant furnace and was under the protection of a mixture atmosphere of 99 vol.% CO2 and 1 vol.% SF6. Firstly, pure Mg was melted at 700 °C. Then, Mg-90Gd mater alloy was added at 750 °C. Thirdly, pure Ag and Mg-30Zr mater alloy were added at 780 °C. Finally, temperature of the melt was decreased to ∼730 °C and the melt was poured in a steel mould pre-heated at 200 °C followed by natural cooling. The actual composition of the alloy was measured using induced coupled plasma atomic emission spectroscopy (ICP) to be Mg-13.1Gd-1.6Ag-0.4 Zr (wt%).

The cast ingots were 65 mm in diameter and 100 mm in length. Firstly, the cast ingots were homogenised at 500 °C for 6 h followed by quenching in hot water of ∼90 °C. After that, the homogenised ingots were machined into cylinder billets of 60 mm diameter and 40 mm length. Hot extrusion was carried out using an extrusion machine with a nominal load force of 315 tone. The extrusion ratio and ram speed were 9 and 2 mm/s, respectively. Two sets of billet and die temperatures were designed. For ITE, the pre-heating temperature of billet, die and container was 350 °C; for DTE, the billet was pre-heated to 500 °C while the die and container were pre-heated to 350 °C. Prior to extrusion, billets were kept at the pre-heating temperature for 30 min in order to reach a uniform temperature distribution. Extruded bars were immediately quenched in cold water when excessing the die exit. For clarity, extruded bars via ITE or DTE are designated as ITE or DTE samples, respectively. In addition, some DTE samples were further cold rolled along the extrusion direction (ED) with a total thickness reduction of 10 % (denoted as DTER samples). Post-extrusion ageing treatment was carried at 200 °C for 32 h using an oil bath. The aged ITE, DTE and DTER samples were named as ITEA, DTEA and DTERA samples, respectively.

Room temperature tensile tests were carried out using a Zwick/Roell Z20 machine. The testing sample has a dog-bone shape with a gauge section of 10 mm length, 3.5 mm width and 2 mm thickness. The testing speed was 1 × 10-3 s-1 and the applied force direction was parallel to ED or the rolling direction (RD). YS was specified as 0.0002 offset strength. Each reported value was the average of three parallel tests.

Microstructure of the samples were examined using JOEL 7800 F field emission scanning electron microscope (SEM) equipped with a back scattered electron (BSE) detector, an Oxford X-ray energy dispersive spectroscopy (EDS) detector and an electron backscatter diffraction (EBSD) detector. EBSD was performed at a step size of 0.5 μm with a scanning area of ∼ 0.3 mm2. Raw EBSD data were collected using Oxford AZtect software and off-line EBSD date processing was carried out using Channel 5 software. Samples for BSE-SEM imaging were prepared by mechanical grinding using SiC sandpapers to 2000 grit surface finish, polishing using 0.2 μm diamond paste and etching using an acetic-picral solution (4.2 g picric acid, 10 ml acetic acid, 80 ml ethyl alcohol and 10 ml water). Samples for EBSD were prepared by firstly mechanical grinding using SiC sandpapers to 2000 grit surface finish and secondly polishing using Struers 0.04 μm colloidal silica suspension (OPS) on a Struers MD-Chem cloth followed by cleaning in pure ethanol for 10 min.

Bright-field scanning transmission electron microscopy (BF-STEM), high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) and selected area electron diffraction (SAED) were performed using a FEI Tecnai F20 S-TWIN transmission electron microscope (TEM) operated at 200 kV. TEM foils were prepared by cutting 0.5 mm thick slices, mechanical grinding to 50 μm thick, punching into 3 mm discs and finally thinned using ion-milling. To evaluate the number densities of precipitate phases, thicknesses of TEM foils were evaluated by matching position averaged convergent beam electron diffraction (PACBED) patterns with simulated patterns [24].

3. Results

3.1. As-extruded microstructure

After homogenisation treatment at 500 °C for 6 h, grains in the alloy are approximately equiaxed in shape and no apparent eutectic compounds are observed along grain boundaries (Fig. 1(a)). Careful examining the BSE-SEM image of as-homogenised alloy finds a minor fraction of small particles showing brighter contrast disperse in the microstructure, as indicated by the orange arrows in Fig. 1(a). Since the intensity in BSE-SEM image is proportional to the atomic number of element, these particles are rich in solute element(s). The inset in Fig. 1(a) indicates that the bright particles are cuboidal. EDS-SEM measurements of α-Mg matrix (Fig. 1(c)) and the cuboidal particle (Fig. 1(d)) indicate that the composition of α-Mg matrix is close to the overall composition of the alloy while the cuboidal particle contains more than 80 at.% of Gd. Referring to the EDS spectrum of α-Mg matrix, it is clear that Ag is barely detected in the cuboidal particle (Fig. 1(c) and (d)). Based on the composition and unique shape of the cuboidal particles, they are identified as so-called cuboid phase, which is frequently found in Mg-rare earth (RE) alloys after homogenisation or solution treatments [[25], [26], [27]]. The cuboid phase is generally recognised as REH2 compound [28] or RE solid solution [29]. Given the small size and quite low volume fraction, the cuboid phase is likely to have negligible effects on the microstructure and mechanical properties of the alloy and thus is not considered in the following. As shown by the orientation map in Fig. 1(b), grains in the as-homogenised alloy are randomly oriented and have an averaged grain size of 44 ± 28 μm.

Fig. 1.

Fig. 1.   (a) BSE-SEM image and (b) orientation map showing the grain shape, size and orientation of the alloy after homogenisation treatment at 500 °C for 6 h. Representative EDS spectra collected from α-Mg matrix (c) and cuboid phase (d), respectively. The inset in (a) shows an enlarge image of a cuboid phase. The orientation map in (b) is coloured using an inverse pole figure (IPF) colouring scheme.


In contrast to the homogenised microstructure, a large number of second-phase particles disperse in the ITE sample after extrusion (Fig. 2(a)). These second-phase particles tend to form strings that are roughly parallel to ED and occupy a volume fraction of ∼8%. The enlarged image of second-phase particles indicate that they are generally micro-sized and predominantly distribute along grain boundaries (Fig. 2(b)). Compared with the ITE sample, the volume fraction of second-phase particles in the DTE sample is noticeably lower, accounting for a volume fraction of ∼2% (Fig. 2(c)). Since second-phase particles mainly distribute along grain boundaries (Fig. 2(d)), they display an inhomogeneous distribution in the DTE sample due to the existence of coarse elongated grains (Fig. 2(c)). BF-STEM image of second-phase particles at triple conjunctions of grain boundaries is shown in Fig. 2(e). The SAED pattern obtained from the particle indicated by the blue arrow in Fig. 2(e) can be indexed according to a face centred cubic (fcc) structure with the lattice parameter a = 2.23 nm in the $1\bar{1}\bar{2}$ zone axis. A representative EDS-SEM spectrum collected from a second-phase particle is provided in Fig. 2(f) and the averaged composition of second-phase particles is Mg-(13.7 ± 0.5)Gd-(2.4 ± 0.2)Ag (at.%), showing a stoichiometry approximating Mg5(Gd, Ag). According to Mg-Gd binary phase diagram [10], the equilibrium phase at the Mg-rich side is Mg5Gd compound having an fcc structure with the lattice parameter a = 2.23 nm. Thus, based on the diffraction and EDS results of the second-phase particles, they are likely to be Mg5Gd type compound, in which a certain amount of Gd is substituted by Ag (hereafter denoted as Mg5(Gd, Ag) compound). Since BSE-SEM images cannot clearly reveal the shape and orientation of grains, microstructure of the ITE and DTE samples are further examined using EBSD.

Fig. 2.

Fig. 2.   (a-d) BSE-SEM images showing microstructure of the alloy after (a, b) ITE and (c, d) DTE; (e) BF-STEM image showing second-phase particles located at triple conjunctions of grain boundaries; (f) a representative EDS spectrum collected from a second-phase particle. (b) and (d) are enlarged images of particle-rich regions in the ITE and DTE samples, respectively, and the orange arrows indicate that second-phase particles mainly distribute along grain boundaries. The inset in (e) is the corresponding SAED pattern of the particle indicate by the blue arrow. The extrusion direction in (a-d) is horizontal.


The orientation map and inverse pole figures (IPFs) of the ITE sample are shown in Fig. 3. The majority of grains in the ITE sample are approximately equiaxed in shape and their sizes are obviously smaller than those of grains in the as-homogenised sample, indicating that the fine equiaxed grains in the ITE sample are formed via dynamic recrystallization (DRX) during extrusion. Furthermore, the equiaxed grains with relatively smaller grain sizes appear like to distribute in the form of strings roughly parallel to ED, which is similar to the distribution of Mg5(Gd, Ag) compounds mentioned above. The averaged grain size of DRXed grains is 2.8 ± 1.6 μm. The elongated unDRXed grains only occupy a volume fraction of ∼1.0 % (Fig. S1(a)), so that the ITE sample is near fully recrystallized. As shown in Fig. 3(b), the ITE sample overall has a weak [0001]//ED texture of 2.7 mud (multiple of uniform distribution). This [0001]//ED texture is intimately related to the DRXed grains, which also have a preferred orientation of [0001]//ED (Fig. 3(c)). In contrast to equiaxed DRXed grains, the elongated unDRXed grains are strongly textured with a preferred orientation of $01\bar{1}0$ // ED (Fig. 3(d)).

Fig. 3.

Fig. 3.   (a) Orientation map of the ITE sample; (b-d) IPFs of the whole data set, the DRXed grain subset and the unDRXed grain subset, respectively. The schematic diagrams at the bottom right corner of (a) show the observation plane, the sample coordination system and the legend for (a). The orientation map (a) and IPFs (b-d) are constructed referring to ED and (a) is coloured based on an IPF colouring scheme.


Fig. 4 shows the orientation map and IPFs of the DTE sample. Compared with the ITE sample, a much larger proportion of unDRXed grains elongating roughly along ED are observed, indicating that the DTE sample is partially recrystallized (Fig. 4(a)). The volume fraction of elongated unDRXed grains is ∼27 % (Fig. S1(b)). Different from the ITE sample, the DTE sample overall exhibits a $01\bar{1}0$//ED texture of 4.9 mud (Fig. 4(b)). As seen from Fig. 4(c), the DRXed grains is near randomly oriented, having a texture maximum of only 1.3 mud. In contrast, the unDRXed grains are strongly textured, showing a $01\bar{1}0$//ED texture of 13.6 mud (Fig. 4(d)). Thus, the overall $01\bar{1}0$//ED texture in the DTE sample is attributed to the increased volume fraction of unDRXed grains. The averaged grain size of DRXed grains in the DTE sample is 2.6 ± 1.5 μm, which is slightly smaller than that of DRXed grains in the ITE sample.

Fig. 4.

Fig. 4.   (a) Orientation map of the DTE sample; (b-d) IPFs of the whole data set, the DRXed grain subset and the unDRXed grain subset, respectively. The schematic diagrams at the bottom right corner of (a) illustrate the observation plane, the sample coordination system and the legend for (a). The orientation map (a) and IPFs (b-d) are constructed referring to ED and (a) is coloured based on an IPF colouring scheme.


Since the orientation maps and IPFs in Fig. 3, Fig. 4 only reveal the preferred orientations of grains referring to ED, pole figures (PFs) of the ITE and DTE samples with the projection plane perpendicular to ED are constructed to further analyse texture features in the ITE and DTE samples (Fig. 5). Two texture components are noticed in the ITE sample: [0001]//ED and $01\bar{1}0$ //ED (Fig. 5(a-d)). The former is relatively stronger and is associated with DRXed grains based on the IPF shown in Fig. 3(c), indicating that the [0001] axis of DRXed grains tend to be parallel to ED meanwhile the $01\bar{1}0$ axis rotate about ED (Fig. 5(a) and (b)). The latter is relatively weaker and arises from unDRXed grains as shown in Fig. 5(c) and (d), indicating that the $01\bar{1}0$ axis of unDRXed grains is aligned with ED while the [0001] axis rotate about ED. Different from that of the ITE sample, the DTE sample exhibits a stronger $01\bar{1}0$ //ED texture (Fig. 5(e-h)). As shown in Fig. 5(g) and (h), the $01\bar{1}0$ //ED texture in the DTE sample is caused by unDRXed grains. It is worth noticing that the texture of unDRXed grains in the DTE sample and that of unDRXed grains in the ITE sample are similar (Fig. 5(c, d, g, h)), but their contributions to the overall textures of ITE and DTE samples are different.

Fig. 5.

Fig. 5.   $\left\{ 0001 \right\}$ and $01\bar{1}0$ PFs of the ITE (a-d) and DTE (e-h) samples. The number at the bottom of each PF indicates the corresponding maximum intensity. (c) and (d) are the unDRXed grain subset extracted from (a) and (b), respectively and (g) and (h) are the unDRXed grain subset extracted from (e) and (f), respectively.


3.2. As-cold-rolled microstructure

Microstructure and texture of the DTER sample appear like those of the DTE sample when the orientation map and IPFs are constructed refereeing to RD (//ED) (Fig. S2). To distinguish the effect of cold rolling on microstructure of the DTER sample, the orientation map and IPFs of the DTER sample are also constructed referring to ND and shown in Fig. 6. Compared with the ITE and DTE samples, the DTER sample has a higher proportion of dark (unindexed) regions, due to a higher number density of crystal defects (e.g. dislocations) introduced by cold rolling. Most of grains, including DRXed and unDRXed grains, display in red or near red colour in the orientation map, indicating that grains in the DTER sample have a relatively narrower distribution of orientations (Fig. 6(a)). The DTER sample overall has a [0001]//ND texture of 2.7 mud (Fig. 6(b)). This [0001]//ND texture with a slightly weaker intensity is also found in the IPF of DRXed grains (Fig. 6(c)). The unDRXed grains occupy a volume fraction of ∼ 20 % (Fig. S1(c)) and exhibit a texture maximum of 3.6 mud slightly deviated from the [0001] direction (Fig. 6(d)). The averaged grain size of DRXed grains in the DTER sample is 2.5 ± 1.6 μm, close to that of DRXed grains in the DTE sample.

Fig. 6.

Fig. 6.   (a) Orientation map of the DTER sample; (b-d) IPFs of the whole data set, the DRXed grain subset and the unDRXed grain subset, respectively. The schematic diagrams at the bottom right corner of (a) illustrate the observation plane, the sample coordination system and the legend for (a). The orientation map (a) and IPFs (b-d) are constructed referring to ND and (a) is coloured based on an IPF colouring scheme.


PFs of the DTER sample, with the projection plane perpendicular to RD or ND, respectively, are shown in Fig. 7. Examining from the ND-TD projection, a $\left\langle 01\bar{1}0 \right\rangle $//RD texture is observed in the DTER sample (Fig. 7(a) and (b)). This texture component in the DTER is geometrically equivalent to the $\left\langle 01\bar{1}0 \right\rangle $//ED texture component in the DTE sample, because ED is parallel to RD. Compared with (0001) poles in the DTE sample which rotate about ED, (0001) poles in the DTER sample concentrate around ND and somewhat spread towards TD (Fig. 7(c) and (d)). The combined observations based on IPFs and PFs indicate that the $\left\langle 01\bar{1}0 \right\rangle $//ED texture component in the DTE sample is preserved and grains in the DTE sample rotate toward ND during cold rolling process (Fig. 6, Fig. 7). As a result, the DTER sample has a [0001]//ND and $\left\langle 01\bar{1}0 \right\rangle $//RD (//ED) texture.

Fig. 7.

Fig. 7.   $~\left\{ 0001 \right\}$ and $\left\{ 01\bar{1}0 \right\}$ P PFs of the ITE and DTE samples. The number under each PF indicates the maximum intensity. (a and b) are constructed with the projection plane perpendicular to RD and (c and d) are constructed with the projection plane perpendicular to ND.


3.3. As-aged microstructure

Orientation maps and IPFs of the ITEA, DTEA and DTERA samples are shown in Fig. 8. The ITEA sample is near fully recrystallized, with a volume fraction of unDRXed grains under 2.0 % (Figs. 8(a) and S1(d)). Besides, the averaged grain size of DRXed grains in the ITEA sample is 2.7 ± 1.5 μm. Therefore, both the volume fraction of unDRXed grains and the averaged grain size of DRXed grains in the ITEA sample are close to those of the ITE counterpart. The DTEA and DTERA samples preserve the bimodal-grained structure of the DTE and DTER counterparts, both containing ∼20 % volume fraction of coarse elongated unDRXed grains (Fig. S1(e) and (f)). In addition, the averaged grain sizes of equiaxed DRXed grains in the DTEA and DTERA samples are 2.5 ± 1.5 μm and 2.3 ± 1.6 μm, respectively, both are near equal to those of DRXed grains in the DTE and DTER counterparts. It is also noted that textures in the ITEA, DTEA and DTERA samples maintain like those in the ITE, DTE and DTER counterparts, respectively. In a word, ageing treatment at 200 °C has negligible effects on the grain shape, size and texture of the samples.

Fig. 8.

Fig. 8.   Orientation maps showing the grain shape and orientation in the (a) ITEA, (b) DTEA and (c) DTERA samples. The legends for the orientation maps and the corresponding IPFs are shown at the top-right and bottom-left corners of (a-c), respectively. (a) and (b) are constructed referring to ED while (c) is constructed referring to ND. The observation plane and sample coordination system of the ITEA, DTEA and DTERA samples are the same to their ITE, DTE and DTER counterparts shown in Figs. 3, 4 and 6.


Precipitates in the ITEA, DTEA and DTERA samples after ageing treatment are shown in Fig. 9. In $\left\langle \bar{2}110 \right\rangle $ α zone axis images (Fig. 9(a, c, e)), plate-like precipitates lying on (0002)α planes are observed. These basal precipitates are γ′′ phase as reported in our previous study [13,30]. To evaluate the size and number density of precipitates, quantitative measurements of precipitates are summarised in Fig. 10. The averaged diameters of γ′′ precipitates in the ITEA and DTEA samples are 16.7 ± 6.2 nm and 15.7 ± 5.2 nm, respectively, which is slightly larger than that of γ′′ precipitates (13.2 ± 5.6 nm) in the DTERA sample (Fig. 10(a)). The number density of γ′′ increases from 7.4 ± 0.5 × 1022/m3 in the ITEA sample to 10.3 ± 1.2 × 1022/m3 in the DTEA sample and finally achieves 11.7 ± 0.5 × 1022/m3 in the DTERA sample (Fig. 10(d)).

Fig. 9.

Fig. 9.   HAADF-STEM images showing morphology and distribution of precipitates in the (a, b) ITEA, (c, d) DTEA and (e, f) DTERA samples. The electron beam is parallel to 2¯110 α in (a, c, e) and [0001]α in (b, d, f), respectively.


Fig. 10.

Fig. 10.   Quantitative measurements of precipitates in the ITEA, DTEA and DTERA samples; (a) the averaged diameter of γ′′ precipitates; (b) the averaged length and width of β′ precipitates; (c) the averaged length of chain-like structures; (d) the number densities of γ′′, β′ and chain-like structures. The diameter of γ′′ precipitates is measured along $\left\langle 10\bar{1}0 \right\rangle $ α in $\left\langle \bar{2}110 \right\rangle $ α HAADF-STEM images; the length and width of β′ precipitates are measured along $\left\langle 10\bar{1}0 \right\rangle $ α and $\left\langle \bar{2}110 \right\rangle $ α, respectively, in [0001]α HAADF-STEM images; and the length of chain-like structures is measured along $\left\langle 10\bar{1}0 \right\rangle $α in [0001]α HAADF-STEM images.


In addition to basal γ′′ precipitates, a large number of prismatic precipitates are also found in [0001]α zone axis images (Fig. 9(b, d, f)). The main prismatic precipitates are β′ phase, showing bright lattice fringes parallel to $\left\{ 10\bar{1}0 \right\}$ α. The size of β′ precipitates is largest in the ITEA sample and smallest in the DTERA sample (Fig. 10(b)). On the contrary, β′ precipitates have the lowest number density (7.8 ± 0.6 × 1022/m3) in the ITEA sample but the highest number density (15.8 ± 0.8 × 1022/m3) in the DTERA sample (Fig. 10(d)). Besides β′ phase, many protrusions showing a chain-like shape are noticed around β′ precipitates (Fig. 9(b, d, f)). These chain-like structures are made up of hexagon rings and/or single-layered zig-zag plates and roughly extend along $\left\langle 01\bar{1}0 \right\rangle $ α directions. While hexagon rings and single-layered zig-zig plates are recognised as solute clusters or Guinier-Preston (GP) zones [11,31], the crystal structures of chain-like structures have not been well-defined in the literature. The lengths of chain-like structures in all samples are close (Fig. 10(c)). While the number densities of chain-like structures in the ITEA and DTEA samples are near equal (i.e. 5.0 ± 0.2 × 1022/m3 and 5.3 ± 0.5 × 1022/m3, respectively), that of chain-like structures in the DTERA sample is near doubled (9.0 ± 0.5 × 1022/m3). The above results indicate that precipitation in the DTEA sample is stronger than that in the ITEA sample and cold rolling prior to ageing further enhances precipitation in the DTERA sample.

3.4. Mechanical properties

Mechanical properties of the alloy in different processing conditions are summarised in Fig. 11. The ITE sample has tensile yield strength (TYS) of 236 MPa, ultimate tensile strength (UTS) of 323 MPa and elongation of 17.6 %. Compared with the ITE sample, the DTE sample achieves higher TYS (282 MPa) and UTS (333 MPa) but lower elongation (9.2 %). After cold rolling, TYS and UTS of the DTER sample are further increased to 313 MPa and 384 MPa, respectively, accompanying with decreasing of elongation to 4.6 %.

Fig. 11.

Fig. 11.   Room temperature tensile properties of the ITE, DTE, DTER, ITEA, DTEA and DTERA samples.


The strength of the alloy is significantly improved after ageing treatment. Specifically, TYS and UTS of the ITEA sample become 318 MPa and 404 MPa, respectively, which are 82 MPa and 81 MPa, respectively, higher than those of the ITE counterpart. The DTEA sample has TYS of 386 MPa and UTS of 450 MPa, showing an increment of 104 MPa and 117 MPa, respectively, but elongation decreases to 3%. The DTERA sample has the highest strength, of which TYS and UTS are 421 MPa and 515 MPa, respectively. The achieved strength of the DTERA sample is much higher than those of other Mg-Gd(-Y)-Ag extruded alloys with similar compositions developed previously [32,33] and is comparable to other ultra-high strength Mg wrought alloys reported in the literature [[34], [35], [36], [37]].

4. Discussion

4.1. Texture evolution

Higher extrusion temperature generally accelerates recrystallization kinetics and thus leads to a higher fraction of recrystallized grains. Based on such assumption, it is expected that the DTE sample with pre-heated billet at 500 °C should have a higher fraction of recrystallized grains than the ITE sample with pre-heated billet at 350 °C. On the contrary, it is found that the ITE sample is near fully recrystallized with the volume fraction of unrecrystallized grains under 2%, while the DTE sample contains unrecrystallized grains with the volume fraction of over 20 % (Fig. 3, Fig. 4). These results indicate that recrystallization kinetics in the ITE samples should be accelerated by some other factors. It is noticed that the ITE sample has near 4-fold volume fraction of Mg5(Gd, Ag) second-phase particles than the DTE sample (Fig. 2). These Mg5(Gd, Ag) second-phase particles are generally micro-sized and mainly form via dynamic precipitation during extrusion. Furthermore, the recrystallized grains with relatively smaller grain sizes appear like to distribute in the form of strings parallel to ED, which is similar to the distribution of Gd5(Mg, Ag) particles, indicating that nucleation of recrystallized grains is intimately related to Gd5(Gd, Ag) particles. In addition, it has been found that Mg5Gd particles stimulated nucleation of recrystallized grains in many Mg-Gd based alloys [38,39]. Given the relatively large volume fractions of Mg5(Gd, Ag) particles, particle stimulation nucleation (PSN) is likely to be a major mechanism for dynamic recrystallization in the ITE and DTE samples. Increasing the fraction of Mg5(Gd, Ag) particles accelerates recrystallization kinetics. As a result, the ITE sample with a larger volume fraction of Mg5(Gd, Ag) particles is near fully recrystallized while the DTE sample containing a less volume fraction of Mg5(Gd, Ag) particles develops the bimodal-grained structure, in which unrecrystallized grains account for a volume fraction of over 20 %.

The fractions of recrystallized or unrecrystallized grains further affect the texture of the ITE and DTE samples. Both unrecrystallized grains in the ITE and DTE samples are strongly textured with the $\left\langle 01\bar{1}0 \right\rangle $ axis parallel to ED (Fig. 5), which is the most commonly observed texture in extruded rods of hexagonal materials [[40], [41], [42], [43]]. Since unrecrystallized grains occupy a volume fraction of over 20 % and recrystallized grains show near random texture (Fig. 4), the DTE sample overall exhibits a $\left\langle 01\bar{1}0 \right\rangle $ //ED fibre texture, which is largely determined by the preferred $\left\langle 01\bar{1}0 \right\rangle $ // ED orientation of unrecrystallized grains. In contrast, the ITE sample overall shows a weak [0001]//ED texture. As illustrated in Fig. 3(c), such unusual [0001]//ED texture arises from recrystallized grains. [0001]//ED texture is generally found in Mg-RE extruded alloys, particularly in those heavily alloyed with Gd [[44], [45], [46], [47]]. Very recently, it is found that the intensity of [0001]//ED texture component in a Mg-14Gd-2Ag-0.5 Zr (wt%) extruded alloy increased during annealing at 400 °C with the preferential growth of grains having [0001]//ED orientation [46]. As mentioned above, the ITE sample is near fully recrystallized. Since grains nucleated via PSN are near randomly oriented [48], a small change in the portion of the grains having a certain preferred orientation will change the overall texture. Considering grain nucleation and growth undergo simultaneously, it is reasonable to believe a certain portion of recrystallized grains have finished nucleation and have underwent grain growth. Thus, the [0001]//ED texture in the ITE sample is presumably associated with the preferred growth of the recrystallized grains with the [0001] axis parallel to ED.

Different from the fact that (0001) poles of unrecrystallized grains rotate about ED, (0001) poles of both recrystallized and unrecrystallized grains concentrate near ND in the DTER sample. As shown in Fig. 6(a), only limited twins are observed in the DTER sample, indicating that dislocation slip is the dominant deformation mode during cold rolling process. Theoretically, the normal of slip plane tend to rotate towards to the axis of compression stress during cold deformation [49]. Since basal slip is the main deformation mode for Mg alloys at room temperature, basal planes of most grains tend to be perpendicular to ND after cold-rolling process.

In general, grain shape and grain size of the ITEA, DTEA and DTERA samples appear like those of their ITE, DTE and DTER counterparts, indicating that ageing treatment at 200 °C dose not significantly arouse grain growth and static recrystallization. As such, the textures in the ITEA, DTEA and DTERA samples maintain similar to those in the ITE, DTE and DTER samples.

4.2. Strengthening mechanism

The DTERA sample obtains remarkably high strength after differential-thermal extrusion, cold rolling and ageing. TYS, UTS and elongation of the DTERA sample are 421 MPa, 515 MPa and 3.1 %, respectively (Fig. 11). Compared with those of other Mg-Gd(-Y)-Ag extruded alloys with similar compositions reported previously [32,33], the DTERA sample has significantly higher strength. In addition, strength of the DTERA sample is comparable with those of other high-strength Mg wrought alloys reported thus far [[34], [35], [36], [37]]. As Mg-13.1Gd-1.6Ag-0.4 Zr alloy is heavily alloyed (the total alloying concentration being over 15 wt%) and the DTERA sample is processed via extrusion, cold rolling and ageing, high strength of the DTERA sample comes from the composite strengthening from solid solution strengthening, grain boundary strengthening, work hardening, texture strengthening and precipitation hardening. Base on the detailed microstructural characterisation in this study, two factors are most significant in controlling strength of the DTERA sample, i.e. bimodal-grained structure and enhanced precipitation.

Bimodal-grained structure is a kind of heterogeneous structure [50]. The key element for heterogenous materials to achieve a good combination of strength and plasticity is that they contain domains with sharp strength differences [51]. In our case, the DTERA sample exhibits a bimodal-grained structure composed of fine equiaxed DRXed grains and coarse elongated unDRXed grains (Fig. 8(c)). On one hand, the unDRXed grains are strongly textured with the basal planes parallel to ED/RD. Such preferred orientation of the unDRXed grains is unfavourable for activating basal slip or tensile twinning during tension along ED/RD. On the other hand, the DRXed grains show a randomised orientation and have an averaged grain size of 2.3 ± 1.6 μm. Thus, the fined and randomised DRXed grains not only improve strength of the alloy by grain boundary strengthening but also effectively accommodate plastic deformation between grains. As a result, the bimodal grained structure provides a strong and ductile framework for Mg-13.1Gd-1.6Ag-0.4 Zr alloy.

Enhanced precipitation in the DTERA sample relates to two key factors. Firstly, the DTE sample has an increased saturation of solute atoms in α-Mg matrix prior to ageing, compared with the ITE sample. This is because the billet for DTE was pre-heated at homogenisation temperature (500 °C), at which precipitation of Mg5(Gd, Ag) particles during pre-heating process of extrusion is maximumly relieved. Therefore, the DTE sample has a less volume fraction of Mg5(Gd, Ag) second-phase particles compared with the ITE sample. Since nucleation rate of precipitates is proportional to saturation of solutes, the number densities of β′, γ′′ and chain-like structures increase in the DTEA sample. As a result, the DTEA sample has a stronger age-hardening response, compared with the ITE sample. Secondly, precipitation in the DTERA sample is further enhanced by cold rolling. Compared with the DTEA sample, the number densities of precipitates further increase while their sizes slightly decrease. Particularly, the number density of chain-like structures is near doubled after cold rolling (Fig. 10(d)). It is believed that crystal defects (e.g. dislocations) introduced by cold rolling may provide more heterogenous nucleation sites and thus enhances nucleation rates of precipitates. The concurrent precipitation of a large number density of both basal and prismatic precipitates forms near-closed blocks in α-Mg matrix, which provides effective obstacles for dislocation sliding and twining. Thus, the combination of DTE and cold rolling make the most out of precipitation hardening, which is another key factor for achieving high strength in Mg-13.1Gd-1.6Ag-0.4 Zr alloy.

While having achieved considerably high strength, elongation of Mg-13.1Gd-1.6Ag-0.4 Zr alloy drops below 5%, which may somewhat limit its applications. It is believed that the limited plasticity of Mg-13.1Gd-1.6Ag-0.4 Zr alloy is largely attributed to Mg5(Gd, Ag) particles that form during extrusion process, which decrease cohesion of grain boundaries and cause local stress concentration. Reducing the volume fraction and refining the size of Mg5(Gd, Ag) particles via microalloying (e.g. slowing down precipitation kinetics) or optimizing processing parameters (e.g. extrusion at higher temperatures but at lower speeds) may contribute to improve the plasticity of Mg-13.1Gd-1.6Ag-0.4 Zr alloy, deserving further attempts.

5. Conclusions

In this work, Mg-13.1Gd-1.6Ag-0.4 Zr alloy was extruded by iso-thermal extrusion or differential-thermal extrusion. Microstructural evolution and mechanical properties of the alloy subjected to different processing conditions, i.e. as-extruded, as-cold rolled or peak-aged, were comparatively studied. The main findings and conclusions are summarised in the following:

(1) Mg-13.1Gd-1.6Ag-0.4 Zr alloy exhibits a nearly fully recrystallized structure after iso-thermal extrusion at 350 °C and has an unusual [0001]//ED texture. Such a texture is related to the growth advantage of recrystallized grains with the [0001] axis aligned with ED.

(2) Mg-13.1Gd-1.6Ag-0.4 Zr alloy shows a bimodal-grained structure that is composed of fine equiaxed recrystallized grains and coarse elongated unrecrystallized grains after differential-thermal extrusion with billet pre-heated to 500 °C and die pre-heated to 350 °C. In contrast to the [0001]//ED texture forms after iso-thermal extrusion, Mg-13.1Gd-1.6Ag-0.4 Zr alloy develops a typical basal texture of $\left\langle 01\bar{1}0 \right\rangle $//ED after differential-thermal extrusion. Such texture change is due to the increased volume fraction of unrecrystallized grains which are strongly textured with the $\left\langle 10\bar{1}0 \right\rangle $ axis parallel to ED.

(3) Differential-thermal extrusion leads to a less volume fraction of Mg5(Gd, Ag) second-phase particles after extrusion, which improves the supersaturation of solutes in α-Mg matrix. In addition, cold rolling prior to ageing causes refinement of precipitates and increment of the number density. As a result, the combination of differential-thermal extrusion and cold work greatly enhances the age-hardening response of Mg-13.1Gd-1.6Ag-0.4 Zr alloy during post-extrusion ageing.

(4) Mg-13.1Gd-1.6Ag-0.4 Zr alloy fabricated via differential-thermal extrusion, cold rolling and ageing displays excellent mechanical properties, reaching YS of 421 MPa and UTS of 515 MPa with acceptable elongation of 3.1 %. The achieved strength of Mg-13.1Gd-1.6Ag-0.4 Zr alloy is much higher than those of other Mg-Gd(-Y)-Ag extruded alloys with similar compositions reported previously. The findings suggest that differential-thermal extrusion plus cold work prior to ageing is a suitable processing route for achieving ultra-high strength in age-hardenable Mg alloys.

Data availability

The authors declare that the main data supporting the findings and conclusions of this study are available from the corresponding author upon reasonable request.

eclaration of Competing Interest

The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

Acknowledgements

This work was supported financially by the National Natural Science Foundation of China (Nos. 51901027, 51971130, 51771113 and 51671128), the China Postdoctoral Science Foundation (No. 2018M643408), and the Natural Science Foundation of Chongqing, China (No. XmT2019012).

Appendix A. Supplementary data

Supplementary material related to this article can be found, inthe online version, at doi:https://doi.org/10.1016/j.jmst.2020.04.031.

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