Microstructural evolution and enhanced mechanical properties of Mg-Gd-Y-Zn-Zr alloy via centrifugal casting, ring-rolling and aging

Zhenduo Ma, Guo Li, Qiang Peng, Xiaodong Peng, Daolun Chen,Hanzhu Zhang, Yan Yang, Guobing Wei, Weidong Xie

a International Joint Laboratory for Light Alloys (Ministry of Education), Chongqing University, Chongqing 400044, China

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

c National Engineering Research Center for Magnesium Alloys, Chongqing University, Chongqing 400044, China

d Department of Mechanical and Industrial Engineering, Ryerson University, Toronto, Ontario M5B 2K3, Canada

e Industrial engineering institute, Ningxia Polytechnic, Ningxia 750021, China

Abstract A ring-shaped Mg-8.5Gd-4Y-1Zn-0.4Zr (wt%) alloy was manufactured via centrifugal casting and ring-rolling process.The effects of accumulative ring-rolling reduction amount on the microstructure, texture, and tensile properties of the alloy were investigated.The results indicate that the microstructure of centrifugal cast alloy consists of equiaxed grains and network-like eutectic structure present at grain boundaries.The ring-rolled alloy exhibits a characteristic bimodal microstructure composed of fine dynamic recrystallized (DRXed) grains with weak basal texture and coarse un-DRXed grains with strong basal texture, along with the presence of LPSO phase.With increasing amount of accumulative ring-rolling reduction, the coarse un-DRXed grains are refined via the formation of increasing amount of fine DRXed grains.Meanwhile, the dynamic precipitation of Mg5RE phase occurs, generating a dispersion strengthening effect.A superior combination of strength and ductility is achieved in the ring-rolled alloy after an accumulative rolling reduction of 80%.The tensile strength of this ring-rolled alloy after peak aging is further enhanced, reaching 511MPa, while keeping a reasonable ductility.The salient strengthening mechanisms identified include the grain boundary strengthening of fine DRXed grains, dispersion strengthening of dynamic precipitated Mg5RE phase, short fiber strengthening of LPSO lamellae/rods, and precipitation strengthening of nano-sized prismatic β＇ precipitates and basal γ＇ precipitates.

Keywords: High strength magnesium alloy; Centrifugal casting; Ring rolling; Aging; Mechanical property.

1.Introduction

The development of high-strength magnesium alloys opens the way for their industrial applications [1-5].The study on high-strength magnesium (Mg)-gadolinium (Gd) based alloys with an excellent age hardening response has recently become one of the most promising directions [6,7].Several studies have pointed out that the addition of yttrium (Y) and zinc(Zn) is conducive to further improving the age hardening response of Mg-Gd based alloys [2,3,8].Besides, the addition of Zn leads to the formation of long-period stacking ordered(LPSO) phase, which can effectively hinder the movement of dislocations and improve the strength of magnesium alloys due to its high hardness [9].On the one hand, the blockshaped LPSO phase with a high Young’s modulus promotes the dynamic recrystallization(DRX)through the particle stimulated nucleation (PSN) mechanism during hot rolling [9].On the other hand, the formation of kink bands connected with lamellar LPSO phase can promote uniform deformationand increase the critical resolved shear stress (CRSS) of basal slip and activate the non-basal slip, which improve the strength and ductility simultaneously [10,11].Adding Zr is one of the effective ways to refine the grain size of magnesium alloys, thus Mg-Gd-Y-Zn-Zr alloy is considered to be a high strength magnesium alloy with great potential for the further development.In particular, the composite strengthening of LPSO phase along with the presence of precipitates after aging treatment has been realized to be highly beneficial to achieve the high strength in such Mg-Gd-Y-Zn-Zr alloys[12].Xu et al.[13]reported that the strengthening mechanism of the high-strength Mg-Gd-Y-Zn-Zr wrought alloy is the presence of LPSO phase in the matrix and the precipitation ofβ＇phase at the boundaries of dynamically recrystallized grains.After aging treatment at a certain temperature and time, a large number ofγ＇phase is precipitated in the matrix of deformed alloy, and the strength is significantly improved through the joint strengthening of LPSO+β＇+γ＇.In addition, Yu et al.[14]prepared an ultra-high strength Mg-11Gd-4.5Y-1Nd-1.5Zn-0.5Zr alloy by hot extrusion, cold rolling and aging treatment, and observed that cold rolling can enhance the age hardening effect and promote the generation of high-density dislocations in the matrix, thus stimulating a large number ofβ＇nucleation sites during aging treatment.

In the aerospace sector, high-strength magnesium alloy ring parts which can meet the demand for weight reduction have significant economic and application values.In detail,this kind of ring parts can be mainly used in satellite cabin,shell and other structural parts.The preparation of traditional magnesium alloy ring parts requires large-sized magnesium alloy ingots as raw materials.However, the solidified microstructure of the large-sized ingots exhibits dendrites and severe composition segregation, which damage the subsequent deformation performance.As a result, it is difficult to use traditional techniques to prepare the seamless magnesium alloy rings.In addition, the traditional fabrication of seamless Mg-Gd-Y-Zn-Zr alloy ring parts needs to undergo large deformation processing such as upsetting, reaming and ring rolling.These complex fabrication processes can easily cause cracking of magnesium alloys and the decrease of yield strength due to excessive strain.Meanwhile, the complex fabrication processes require additional equipment and increase the fabrication costs.Thus, it is urgent to develop a new fabrication process for high-strength Mg-Gd-Y-Zn-Zr alloy ring parts with low fabrication costs.The centrifugal casting is an enabling casting technique used to purify the alloy melt efficiently, transform the developed dendrites into equiaxed grains and homogenize the microstructure [15].Meanwhile,the ring-shaped centrifugal cast ingot is easy to be ring-rolled subsequently, and the large deformation required in the traditional processes can be avoided, thus reducing cracking tendency.

Fig.1.Schematic diagram of the preparation process and sample selection[17].

The aim of this study is to identify the effect of different amounts of accumulative rolling reduction on the microstructure and mechanical properties of a centrifugal cast Mg-Gd-Y-Zn-Zr alloy with homogenization.A 80% ringrolled alloy is aged to further improve the strength and the relevant strengthening mechanisms are discussed.The results can provide the guidelines for the fabrication of high-strength Mg-Gd-Y-Zn-Zr alloy ring parts with low fabrication costs,which is very important for the engineering applications.

2.Experimental procedure

2.1.Alloy preparation

Billet with the composition of Mg-8.5Gd-4Y-1Zn-0.4Zr(wt%) alloy was melted in an electrical furnace at 750°C in Ar atmosphere.After the melt was cooled to 725°C, the melt was cast into a rotating steel mold.According to the previous experience of our research group, the rotational velocity of centrifugal casting mold was determined to be 900rpm [16].Before ring rolling, the inner and outer surfaces of the centrifugal cast ring need to be machined.About 5mm material on both sides of the ring was removed by machining.After machining, the outer diameter and thickness of the centrifugal cast ring are 380mm and 21mm, respectively.

The centrifugal cast rings were homogenized at 510°C for 12h, followed by quenching in water at room temperature.The ring rolling was performed with different amounts of accumulative rolling reduction of 40%, 60% and 80% at 450°C.The purpose for the accumulative rolling reduction variances is to study the microstructure evolution and mechanical properties of the alloy during ring rolling.It should be noted that before the ring rolling, the homogenized raw alloys were preheated at 450°C for 35min because the rolling of magnesium alloys at room temperature is difficult.Then the ring rolling was performed with a reduction of 15% per pass.During each pass, the rolled samples were reheated at 450°C for 15min.After rolling to the preset reduction, three samples with different amounts of accumulative rolling reduction were obtained.Fig.1 shows a schematic diagram illustrating the entire process of centrifugal casting and ring rolling [17].The ring-rolled samples with an accumulative rolling reduction of 80% (RR80 alloys) were subjected to aging treatment at 200°C with an aging time from 0.25h to 128h and the peak-aged sample was obtained at 40h.The purpose for the aging time variances is to determine the peak aging point by aging hardening curve.

Fig.2.OM micrographs of experimental alloys with different position: (a) Inside; (b) Middle; (c) Outside.

2.2.Microstructural characterization

The microstructures were observed using an optical microscope (OM), a VegaIILMU scanning electron microscope equipped with an energy-dispersive X-ray spectrometer(EDS)and a FEI Tecnai G2 F20 transmission electron microscope(TEM).The electron backscatter diffraction (EBSD) was conducted on a FEI NOVA 400 FG SEM using the step size of 1.5μm.The samples with different centrifugal radius were selected as shown in Fig.1.Meanwhile, the chemical composition of the three samples was analyzed by inductively coupled plasma emission spectrometer (ICP-OES).The OM,SEM and EBSD observations of the ring-rolled samples were measured on the RD-ND plane, as specified by a black rectangle in Fig.1.The phases in the experimental alloys were identified via a Rigaku D/ Max2500PC XRD with a diffraction angle range from 20° to 90° at a scanning rate of 5°/min.In order to determine the phase transformation temperature of the second phase in centrifugal casting alloy,the samples were analyzed by Netzsch STA44PF3 differential thermal analysis(DTA) tester.The sample is a wafer with a diameter of 3mm and a thickness of 1mm.The temperature range of the test is 25 ～600 °C and the heating rate is 10K/min.

2.3.Mechanical testing

Tensile tests were performed on a WDW-100D machine at a strain rate of 1 mm·min?1.The tensile samples with a gage size of 10mm×3mm×2mm were cut along a direction parallel to the RD direction.The location of the tensile sample selection is specified by a black rectangle in Fig.1.Tensile tests were conducted on alloys with different accumulative rolling reductions and peak aged RR80 alloy.Hardness tests were conducted by HV-30 Vickers at a load of 100g and a dwell time of 10s.Five indentations were made for each material condition to obtain an average hardness value.

3.Results and discussion

3.1.Microstructure

3.1.1.Microstructure before hot rolling

Fig.2 shows the OM images of the sample with different centrifugal radius.It can be seen that the microstructure characteristic of the samples with different centrifugal radius is almost the same.Meanwhile, the actual chemical composi-tion is also very close (as shown in Table 1).The reasons for this phenomenon are as follows.The outer and inner sides of the alloy ring are in contact with the mold and air, respectively.Thus, the solidification rate of melt in both sides of the ring is close.Besides, the wall thickness of centrifugal ring is relatively thin, so there is almost no gradient distribution of composition in this centrifugal ring.Therefore, this paper does not analyze the influence of gradient distribution on the subsequent rolling.

Table 1Actual chemical composition of experimental alloys with different centrifugal radius.

Table 2EDS analysis at typical positions in the alloy in Fig.6(b) (at%).

The centrifugal cast alloy was homogenized before hot rolling to dissolve the reticulated eutectic structure located at the grain boundaries and avoid the occurrence of cracking during hot rolling [18].Based on the DTA curve of the experimental alloy, shown in Fig.3(a), the eutectic temperature is about 520°C.In general, the homogenization temperature is 5-10°C lower than the eutectic temperature [19], thus the homogenization temperature in the present study is selected to be 510°C.XRD patterns of the centrifugal cast and homogenized alloys are shown in Fig.3(b).It can be seen that the centrifugal cast alloy consists ofα-Mg matrix,Mg12REZn and Mg3RE phases.The Mg12REZn phase belongs to LPSO phase while Mg3RE phase is a network eutectic phase.After the homogenization treatment,the diffraction peaks of Mg3RE phase disappear, and a diffraction peak of Mg12REZn phase intensifies.As observed in the OM images (shown in Fig.3c and d), the centrifugal cast alloy includes equiaxed grains and network eutectic structure distributed at the grain boundaries.However, the homogenization treatment gives rise to the dissolution of the network eutectic phases into theα-Mg matrix and the formation of block-shaped LPSO phases at the grain boundaries.In addition, the average grain size of the experimental alloy increases from ～53μm to ～96μm after homogenization.

3.1.2.Microstructure after hot rolling

Figs.4 and 5 show the OM images and inverse pole figures(IPF) maps of ring-rolled alloys with different amounts of accumulative rolling reduction.As shown in Figs.4(a) and 5(a),the un-DRXed grains occupy the majority of microstructure in the ring-rolled 40% alloy, and only a few DRXed grains appear at the boundaries of un-DRXed grains.As seen from Fig.5(a), there are obvious color contrast changes inside the un-DRXed grains because of the existence of subgrains and a large number of dislocations, which indicate the changes of grain orientation in these regions.This EBSD observation corresponds to the change of the orientations of lamellar LPSO phase in the OM observation (shown in Fig.4(a)).The accumulated strain energy stored in the un-DRXed grains plus grain boundary energy provides energy for the occurrence of DRX.Moreover, the lamellar and blocky LPSO phase is located in the interior and at the boundary of un-DRXed grains, respectively.A large number of lamellar LPSO phase in the un-DRXed grains can inhibit the occurrence of DRX and the growth of DRXed grains.On the contrary, the blocky LPSO phase distributed at un-DRXed grain boundaries can promote the occurrence of DRX via PSN mechanism.Thus, the DRX occurs preferentially at un-DRXed grain boundaries.As shown in Fig.4(b) and (c),the degree of kinked lamellar LPSO phase and the volume fraction of “necklace” microstructure increase with increasing accumulative rolling reduction from 60% to 80%.The“necklace” microstructure is presumed to be DRXed grains.With the kinking of LPSO phase,the surroundingα-Mg phase generates a kink-like plastic deformation by the slip of dislocations [20], to ensure the continuity of deformation.According to the Friedel-Escaig mechanism of dislocation cross-slip[21,22], a transition from a spiral dislocation to a blade dislocation occurs in the kink bands.Blade dislocations gradually form low-angle grain boundaries through climbing [22-24].Finally, the sub-grain boundaries are gradually transformed into high-angle grain boundaries by continuously absorbing dislocations in the small-angle grains.As a result, the kinked grain boundaries with a large difference in orientations can be formed [25].The kink band boundary can become favorable nucleation sites for continuous dynamic recrystallization(CDRX).The “necklace” microstructure forms at the kinked grain boundaries as indicated by the white circle in Fig.4(b)and (c).As the rolling reduction amount and rolling stress increase, the kink deformation of lamellar LPSO phase and the accumulated strain energy increase,thereby promoting the nucleation of “necklace” microstructure at the kink band boundary.As shown in Fig.5(b) and (c), the (0001) plane of coarse un-DRXed grains is almost parallel to the rolling surface,while the fine DRXed grains show a more random orientation.The grain boundaries in the DRX region are clearly observed, so that the “necklace” microstructure is confirmed to be fine DRXed grains.In a word, with increasing amount of accumulative rolling reduction, the coarse un-DRXed grains are fragmented and refined.The volume fraction of the fine DRXed grains around them increases significantly.

Fig.3.(a) DTA curve, (b) XRD patterns, and OM micrographs of (c) centrifugal cast and (d) homogenized Mg-Gd-Y-Zn-Zr alloy.

Fig.4.Optical microstructure of the ring-rolled alloy with different amounts of accumulative rolling reduction at (a) 40%, (b) 60% and (c) 80%, respectively.

Fig.5.IPF maps of the ring-rolled alloy with different amounts of accumulative rolling reductions at (a) 40%, (b) 60% and (c) 80%, respectively.

Fig.6.Texture evolution of the ring-rolled alloys with different amounts of accumulative rolling reduction.

Fig.6 shows the texture evolution of the ring-rolled samples with different amounts of accumulative rolling reduction.The obvious rolling basal texture with the (0001) basal plane of magnesium grains perpendicular to ND is formed in all the ring-rolled samples.With increasing rolling reduction amount,the basal pole spreads from ND toward RD to some extent.Meanwhile, the distribution of basal pole inclines from ND to TD, and the degree of expansion to TD and RD increases gradually.In addition, the decrease of the basal texture intensity with increasing amount of accumulative rolling reduction is directly related to the occurrence of DRX to a greater extent with increasing volume fraction of DRXed grains (Fig.5).

Typical SEM images of the experimental alloy with different amounts of accumulative ring rolling reduction are shown in Fig.7.As the accumulative rolling reduction increases,the volume fraction and size of the lamellar LPSO phase decrease.The decomposition of the LPSO phase is an endothermic phase transition process, and the stress along the phase interface provides a driving force for the decomposition of the LPSO phase.With increasing amount of rolling reduction, the accumulated strain near the interface also increases,which facilitates the decomposition and breaking of the LPSO phase.Along with the dissolution of the LPSO phase in the un-DRXed region, a large number of RE atoms are dissolved as well, which provides a condition for the dynamic precipitates in the DRXed region [26].

The EDS results in Fig.7(b) are shown in Table 2.The composition of a LPSO lamella at point A in the un-DRXed region is about 95.5Mg-1.7Gd-1.3Y-1.5Zn (at%).As shown by the red arrows in Fig.7(b) and (c), the rod-shaped phase distributed along the RD can be observed at the grain boundaries in the DRXed region.The EDS result of point B is 95.8Mg1.6Gd-1.2Y-1.4Zn (at%), indicating that the rod-shaped phase is also LPSO phase.Besides, since the basal plane of rod-shaped phase was observed to be parallel to the rolling surface [27], the rod-shaped LPSO phase should be left after the breaking of some LPSO lamellae in the un-DRXed region.Meanwhile, many granular precipitates can also be observed, as shown by the yellow arrows in Fig.7(b).According to the EDS result of point C, the composition of the granular precipitated phase is 83.7Mg-11.4Gd-4.9Y-0.1Zn (at%), suggesting that the granular precipitated phase is Mg5RE phase.This is a result of dynamic precipitation of supersaturated RE atoms in the DRXed region.The number of rod-shaped LPSO phase and granular Mg5RE phase in the DRXed region increases significantly with increasing amount of accumulative rolling reduction.

Fig.7.SEM micrographs of the ring-rolled alloy with different amounts of cumulative rolling reduction at (a) 40%, (b) 60% and (c) 80%, respectively.

Fig.8.TEM images of the RR80 alloy: (a) bright-field image of LPSO phase, (b) bright-field image of Mg5RE phase, (c) SAED pattern of region A and (d) SAED pattern of region B.

Bright-field TEM images of the RR80 alloy are shown in Fig.8.As shown in Fig.8(a), a large number of LPSO lamellae are observed along the [110]Mgdirection.According to the selected area electron diffraction (SAED) pattern of region A (Fig.8(c)), these lamellar precipitates should belong to 14H-LPSO phase.After homogenization treatment, a large number of RE and Zn atoms are dissolved inα-Mg matrix,which provides the chemical order condition for the precipitation of LPSO phase during hot rolling.In the subsequent hot rolling process, a large number of dislocations would be produced in the grains.Some dislocations would decompose to generate stacking faults; meanwhile, the RE and Zn solute atoms in the matrix would rearrange stacking, so that the chemical conditions and structural conditions for the formation of LPSO phase could be satisfied simultaneously, thus promoting the precipitation of LPSO phase [28].

In addition, it can be seen that several granular precipitates with a size of 100-200nm are distributed in grains and at grain boundaries.According to the SAED pattern of region B (Fig.8(d)), these granular precipitates are confirmed to be Mg5RE phase with a fcc structure.The precipitation of Mg5RE phase in the process of hot rolling belongs to the heterogeneous nucleation solid-state phase transformation.Since crystal defects such as grain boundaries and dislocations have a higher energy, it is easier for the second phase to nucleate at these sites.That is, the higher stored energy at these defects can reduce the energy required for nucleation of the second phases, thus these positions will become the preferred positions for the precipitation [29].When the rolling reduction reaches 80%, sufficient dislocations are stored in the grains, leading to the entanglement of dislocations and the formation of high stress/strain regions, which provide the nucleation sites for the dynamic precipitation.The rare earth atoms spread rapidly along with the movement of dislocations.When the composition condition is satisfied, the Mg5RE phase is formed.At the same time, with increasing number and entanglement of dislocations, new sub-grains are formed with small angle grain boundaries.Rare earth atoms on dislocations would be concentrated at these grain boundaries.With continuous accumulation of rare earth atoms, the nucleation and growth of the Mg5RE phase begin, and a large number of precipitated particles are finally formed on the small angle grain boundaries.These Mg5RE precipitates can effectively pin the grain boundary, hinder the movement of the grain boundary and dislocation slip, thus improving the strength [30].

Fig.9.Bright-field TEM images and the corresponding SAED patterns of the peak-aged RR80 alloy: (a, c) β＇ precipitates +LPSO phase; (b, d) β＇precipitates+γ＇ precipitates.

3.1.3.Microstructure of peak-aged RR80 alloy

Fig.9 shows the bright-field TEM images and the corresponding SAED patterns of the peak-aged RR80 alloy.As shown in Fig.9(a) taken from [110]Mgdirection, numerous fine lenticular precipitates and lamellae coexist in the peakaged alloy.As seen from Fig.9(c), there are two sets of diffraction spots in region A, which are confirmed the coexistence ofβ＇phase and 14H-type LPSO phase.The fine lenticular precipitates (marked by yellow arrows) are identified to beβ＇precipitates according to the additional diffraction spots observed at 1/4(010)α?Mg, 2/4(010)α?Mgand 3/4(010)α?Mg[31].The lamellar phases (marked by red arrows) are confirmed to be 14H-type LPSO lamellae.

As shown in Fig.9(b), the peak aging indeed results in the presence of bothβ＇phase precipitated on the prismatic plane andγ＇phase precipitated on the basal plane.According to the SAED pattern of region B (Fig.9(d)), the coexistence of prismaticβ＇precipitates (marked by yellow arrows)and basalγ＇precipitates (marked by white arrows) is confirmed.When Zn is added to Mg-RE alloy with a high rare earth content, not only LPSO phase but also solute segregation faults will be formed on the basal plane ofα-Mg matrix.The so-called solute segregation faults are actually theγ＇precipitates on the basal plane [32].As reported in the study of Mg-6.5Gd-2.5Dy-1.8Zn alloy[33],theγ＇precipitates have a better strengthening effect than the LPSO phase.Theβ＇precipitates andγ＇precipitates thus become the main strengthening phases after peak aging.As a result, the synergistic strengthening effect of dense nano-sized prismaticβ＇precipitates and basalγ＇precipitates is expected to contribute to the enhanced age hardening response and strength significantly.

Fig.10.Age-hardening curve of the RR80 alloy at 200°C.

3.2.Mechanical properties and strengthening mechanisms

The RR80 alloy is selected for the subsequent aging because of its better comprehensive performance.Fig.10 shows the age hardening curve of the ringrolled 80% alloy at 200°C.The change trend of the hardness corresponds to the precipitation sequence of precipitates during aging.As reported by Wang et al.[34], the precipitation sequence of Mg-Gd-Y-Zn alloys is:α-Mg (SSSS)→β＇＇(DO19)→β＇(bco)→β1(fcc)→β(fcc) and SSSS→I2stacking fault→γ＇(hcp)→LPSO (ordered hcp).At the initial stage of aging, the hardness increases slowly.This is attributed to the formation ofβ＇＇precipitates and I2stacking faults, which have a coherent relationship with the matrix.While they can impede the movement of dislocations with a certain strengthening effect, their ability to impede dislocations is relatively weak and the strengthening effect is limited because of their small size [35].After aging for 8h,the hardness of the alloy increases rapidly, exhibiting obvious age hardening characteristics.Due to their poor thermal stability, theβ＇＇precipitates and I2stacking faults are subsequently transformed intoβ＇precipitates andγ＇precipitates,respectively.The volume fraction ofβ＇precipitates andγ＇precipitates increases significantly.After aging for 40h, the number density ofβ＇precipitates andγ＇precipitates reaches the maximum value, leading to a peak hardness of 124 HV for the alloy.The hardness decreases gradually during subsequent aging from 40h to 128h, which corresponds to the over-aging stage.Theβ＇precipitates andγ＇precipitates becomeβphase and LPSO phase, respectively, along with their coarsening in the later stage.

Fig.11 shows the tensile stress-strain curves of the alloys with different amounts of accumulative ring rolling reduction and peak-aged RR80 alloy with ring rolling 80% tested at room temperature.The corresponding tensile properties of the studied alloys are summarized in Table 3.The results reveal that the ultimate tensile strength (UTS) and yield strength(YS) are remarkably enhanced with increasing amount of rolling reduction.A UTS of 390MPa, a YS of 330MPa and a fracture elongation (EL) of 12.2% were achieved after ring rolling with an accumulative reduction of 80%.Furthermore,the peak-aged RR80 alloy exhibits a value of UTS and YS of 511MPa and 435MPa, respectively, along with an EL of 5.3%.

Table 3Tensile properties of the alloys with different amounts of accumulative rolling reduction and peak-aged RR80 alloy tested at room temperature.

Fig.11.Tensile stress-strain curves of the alloys with different amounts of accumulative rolling reduction and peak-aged RR80 alloy tested at room temperature.

The improvement of tensile properties with increasing amount of accumulative rolling reduction up to 80% is due to the breaking of LPSO phase, grain refinement and dynamically precipitated Mg5RE phase.The deformation kinking of LPSO phase can coordinate the deformation of matrix grains,which is helpful for the improvement of plasticity.Due to the presence of deformation kinks, the stress concentration is easy to occur at the interface ofα-Mg/LPSO phase, which promotes the refinement ofα-Mg grains near LPSO phase via DRX[36].In addition,LPSO phase lamellae are arranged parallel to the rolling direction after hot rolling,producing a short fiber strengthening effect which is similar to that in composite materials [20].The lamellar/rod-shaped LPSO phase can effectively impede the slip of dislocations, the growth of twins and the growth of microcracks [37].The LPSO phase lamellae in the 40% ring-rolled alloy are relatively coarse, which are insufficient to exert a strong strengthening and toughening effect.With increasing amount of rolling reduction, both the extent of broken LPSO lamellae and the number of rodshaped LPSO phase increase, which is equivalent to refine the LPSO phase.Moreover, the kink deformation of LPSO phase intensifies, increasing the volume fraction of DRXed grains.Therefore, the strength and plasticity of the alloy simultaneously increase as the amount of ring rolling reduction increases.

The volume fraction of fine recrystallized grains increases significantly with increasing amount of accumulative rolling reduction, leading to a decrease in the average grain size(as if the coarse un-DRXed grains are refined).According to the Hall-Petch relationship, the strengthening stemming from grain refinement gives rise to the increase of yield strength of alloy [38].Besides, the weakened texture of DRXed grains is also beneficial for facilitating the basal slip, releasing the stress concentration transferred from the un-DRXed grains and promoting the homogeneous deformation [13].Furthermore, the refinement of the DRXed grains can significantly shorten the dislocation slip distance and relieve the stress concentration caused by the accumulation of too many dislocations, thus improving the plasticity.As a result, the increasing volume fraction of fine DRXed grains with increasing amount of accumulative rolling reduction improves the strength and ductility simultaneously (Table 3).

The accumulated strain near the interface increases with increasing amount of rolling reduction, which enhances the precipitation driving force and promotes the dynamic precipitation of Mg5RE phase.These dynamically precipitated Mg5RE particles can effectively pin the grain boundaries and impede the slip of dislocations, thus playing a role of dispersion strengthening.In turn, the dynamically precipitated Mg5RE particles can promote the DRX via PSN mechanism,enhancing the grain boundary strengthening effect.

The remarkable improvement of tensile strength after peakaging treatment (Fig.11 and Table 3) is due to following factors.The lenticularβ＇precipitates distribute in the matrix on the cylinder-like surface in a form of triangles.That is,the denseβ＇precipitates formed on {110} prismatic planes are perpendicular to the basal plane, which impede the basal dislocation slip and improve the strength but reduce the plasticity of the alloy [2].However, the presence of LPSO phase can coordinate the deformation behavior.The kink deformation of LPSO phase can effectively adapt to the deformation process of the alloy.Thus, LPSO phase improves the strength and plasticity of the alloy concurrently, rather than act as a potential initiation point for cracking [39].The numerous LPSO lamellae with a narrow interlamellar spacing can restrain the coarsening of prismaticβ＇precipitates, which are thus beneficial to retain fineβ＇precipitates and increase strength [40].The combined effect of LPSO phase lamellae with fineβ＇precipitates restricts the basal and prismatic slip and hinders the growth of microcracks via the formation of an approximately closed micro-space [36].The strengthening mechanisms of peak-aged RR80 alloy involve directly those of a normal ring-rolled alloy and the strong precipitation strengthening mechanisms.The fiber strengthening of LPSO phase lamellae, the dispersion strengthening of the dynamically precipitated Mg5RE phase, the grain boundary strengthening of the fine DRXed grains and especially the precipitation strengthening of the nano-sizedβ＇andγ＇are operatedall together to make the peak-aged alloy with a high tensile strength.It should be emphasized that the basalγ＇precipitates divide the Mg matrix grain into some triangular networks,while the prismaticβ＇precipitates can be evenly distributed on the three cylindrical surfaces of the networks, thus forming numerous small near-continuous blocks.The twinning and the generation of dislocations would be limited in the blocks.Even though the twins and dislocations are generated under the higher applied stresses in the later stage of tensile deformation, the twin expansion and dislocation slippage would become effectively impeded by the blocks, since they form an ideal co-precipitates strengthening mode [41].Therefore, the comprehensive mechanical properties of the peak-aged alloy are greatly improved, as shown in Fig.11 and Table 3.

4.Conclusions

The microstructures and tensile properties of a centrifugal cast Mg-8.5Gd-4Y-1Zn-0.4Zr (wt.%) alloy with different amounts of accumulative ring rolling reduction were investigated, along with the effect of aging on the microstructures and tensile properties in a ring-rolled part with a rolling reduction of 80%.The following conclusions could be drawn:

1) The microstructure of the alloy fabricated directly by centrifugal casting consists mainly of equiaxed grains and network-like eutectic structure present at grain boundaries.After homogenization treatment at 510°C for 12h, the network eutectic structure are resolved into theα-Mg matrix, and the block-shaped LPSO phase is formed at grain boundaries.Then the average grain size increases from ～53μm to ～96μm after homogenization.

2) The ring-rolled alloy exhibits a characteristic bimodal microstructure composed of fine dynamic recrystallized(DRXed) grains with weak basal texture and coarse un-DRXed grains with strong basal texture.With increasing amount of accumulative rolling reduction, the volume fraction of fine DRXed grains and the breaking degree of LPSO lamellae increase significantly, leading to the simultaneous increase in the strength and plasticity.Meanwhile, the dynamic precipitation of Mg5RE phase occurs, exhibiting dispersion strengthening effect.

3) A UTS of 390MPa, a YS of 330MPa and an EL of 12.2% are achieved in the ring-rolled alloy when the accumulative rolling reduction increases to 80%.The strengthening mechanisms include 1) the grain boundary strengthening of fine DRXed grains, 2) the dispersion strengthening of dynamic precipitated Mg5RE phase, and 3) the short fiber strengthening of LPSO lamellae/rods.

4) The peak-aging of the alloy after the accumulative rolling reduction of 80% is observed to be performed at 200°C for 40h.The peak-aged alloy exhibits significantly high values of UTS and YS of 511MPa and 434MPa, respectively, with a decent EL of 5.1%.In the peak-aging state, the significant increase in the strength is attributed to the additional strong precipitation strengthening of nano-sizedβ＇precipitates present on the prismatic plane andγ＇precipitates on the basal plane, in addition to the above three strengthening mechanisms.

Declaration of Competing Interest

None.

Acknowledgments

The authors would like to acknowledge financial support by Fundamental Research Funds for the National Key Research and Development Program of China (Project No.2016YFB0700403), the Venture & Innovation Support Program for Chongqing Overseas Returnees (Project No.cx2018057), the Chongqing Research Program of Basic Research and Frontier Technology(Project Nos.cstc2019jcyjmsxm0548 and cstc2019jcyj-msxmX0306), the Fundamental Research Funds for the Central Universities (Project No.2021CDJJMRH-001).