Texture development and tensile properties of Mg-Yb binary alloys during hot extrusion and subsequent annealing

Lu Li, Cuncai Zhang, Hao Lv, Chunrong Liu, Zhuozhang Wen, Jingwen Jiang

School of Materials and Energy, Southwest University, Chongqing, China

Abstract Ytterbium (Yb) containing magnesium alloys have aroused extensive interest due to their excellent mechanical properties after thermomechanical processing and heat treatment.Unfortunately, the sole effect of Yb addition on the microstructure and mechanical properties of pure Mg matrix remains uncertain to date.In this work, the effects of Yb concentration on the texture development and tensile properties of pure Mg matrix during hot extrusion and the subsequent annealing were systematically investigated.The results revealed that the constitutional supercooling induced by Yb addition refined the as-cast microstructure but exerted a negligible effect on the original columnar grain morphology.When extruded at 300 °C, the dynamic recrystallization (DRX) process was considerably retarded.The in-grain misorientation axes (IGMA) analysis combined with TEM observation indicated that non-basal slips operated with increasing Yb concentration.Specifically,the prismatic ＜a＞slip should be robustly activated in Mg-1.0 Yb extrudate, promoting the formation of the texture with {10-10} plane normal to the extrusion direction (ED), while for the Mg-2.0 Yb counterpart, the increased activity of pyramidal ＜c + a＞slip and the relaxation of basal/＜c + a＞dislocations generated an ED-tilted texture component.The preferential grain growth dominated the subsequent annealing texture development at 400 °C when a comparable grain size was achieved.An obvious ED-tilted texture intensity with the peak around ＜-12-13＞was observed in Mg-2.0 Yb alloy, which was primarily caused by grains with the basal orientation vanished and with the non-basal orientations intensified due to a higher concentration of Yb solute.Favored by the grain refinement, the Mg-2.0 Yb extrudate exhibited a high tensile yield strength of 304 ± 3.5 MPa, while the subsequently annealed counterpart presented a favorable elongation to failure of 14.8 ± 1.2%, which mainly due to the homogeneous grain structure, weak ED-tilted texture, and dissolution of coarse phases after high-temperature annealing.

Keywords: Magnesium alloy; Ytterbium; Texture development; Tensile properties.

1.Introduction

The ever-increasing demand for improving fuel economy and decreasing harmful emissions inspires researchers to develop new wrought magnesium (Mg) alloys with enhanced properties [1].Hot extrusion has been widely used to fabricate Mg alloys as dynamic recrystallization is readily initiated in the triaxial compressive stress state, by which the matrix can be significantly refined and thereby,the improved strength can be achieved.However, as the limited number of active deformation systems in hcp magnesium, a strong crystallographic texture with the (0001) plane parallel to the extrusion direction (ED) is prone to be developed upon hot extrusion[2-4].The basal texture with most grains in “hard orientation” is detrimental for the ductility when stretched along the extrusion axis subsequently.As a result,the enhanced strength seems to be a consumption of ductility, and the trade-off between them is still a concern in magnesium extrusion.To this end, it is of significance to tailor the deformation texture of extruded Mg alloys for enhancing strength and ductility synergy.

To date, many efforts have been made to weaken/alter the deformation texture in extruded Mg alloys, where one of the promising attempts is to incorporate a trace amount of rare earth (RE) elements into pure Mg matrix.A large body of studies relating to extruded Mg-RE alloys have shown that a substantial basal texture weakening can be achieved in terms of decreased peak texture intensity and tilted or split basal poles towards the ED [5].For example, Wu et al.[6]reported that more random components of texture were formed in the Mg-Y extrudates, promoting the activation of non-basal slips and therefore an increase in the elongation-to-failure.Kim et al.[7]showed that with increasing Gd addition, the texture peak orientation was changed from the＜2-1-11＞to＜0001＞direction, leading to an extraordinary texture pattern with basal poles parallel to the ED in the Mg-15 Gd alloy and thus a significant improvement in the tensile strength.Wu et al.[8]reported that the addition of Er acted a significant role in texture change and the reduction inc/aratio diversified the deformation mechanisms in the Mg-Er alloys, which facilitated homogeneous deformation during room temperature tension and compression.Moreover, Stanford [9]identified that La, Ce, and Gd were all effective texture modifiers during the extrusion of Mg-based alloys, favored by the formation of“Rare-Earth”(RE)texture at the low alloying levels of 300, 400, and 600 ppm, respectively.Despite these efforts,however, there is still no open reported literature relating to the effects of Yb sole addition on texture development and mechanical properties of pure Mg matrix.

Ytterbium (Yb) is chemically classified by the ionic radii as the group of heavy RE elements.In recent years, Yb has aroused intense interest as a promising alloying additive to modify the microstructure and improve the mechanical performance of traditional magnesium systems.The excellent effects on grain refinement and dense Yb-containing precipitates combine to strengthen the matrix via grain boundary and precipitation hardening.Nevertheless, it was worth noting that the available investigations mainly focused on the Mg-Yb-Zn [10-17]or Mg-RE-Yb-Zn [18-21]multicomponent Mg alloys.As alloying elements such as Zinc and REs (Gd,Sm) also contribute to the improvement of microstructure and mechanical performance of Mg-base alloys, it is critical to clarify the sole effect of Yb addition on the recrystallization behavior, texture development, and mechanical properties of Mg matrix during thermomechanical processing and the subsequent annealing.However, there is no open reported literature involving this topic.In this study, the Mg-Yb binary alloys with different Yb concentrations (0, 0.5, 1.0,and 2.0 wt.%) were prepared.The effects of Yb concentration on the texture development and tensile properties of pure Mg matrix during hot extrusion and the subsequent annealing were systematically investigated.Specifically,the mechanisms underlying the increased activity of non-basal slips during hot extrusion, the preferential grain growth during annealing, and the development of mechanical properties with increasing Yb concentration were discussed in detail.These findings help to develop a thorough understanding of Yb addition alone on the microstructure and mechanical properties of pure Mgand provide valuable insight to design novel Mg-Yb series alloys with favorable microstructure and properties.

2.Materials and experimental method

The alloy ingot with a nominal composition(wt.%)of MgxYb (wt.%,x= 0.5, 1.0, and 2.0) was melted in an electrical resistance furnace using pure Mg and Mg-15% Yb, under fluxing protection.The chemical compositions of as-cast billets were analyzed using the inductively coupled plasma optical emission spectroscopy (ICP-OES) method, as shown in Table 1.The as-cast ingot was soluted at 400 °C for 24 h and then hot extruded at 300 °C at an extrusion ratio of 15:1, followed by a cold-water quench.Afterward, the samples of hot extruded pure Mg, Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb alloys were annealed at 400 °C for 15, 30, 60, and 100 min respectively to obtain a similar average grain size.

Table 1Chemical composition of Mg-Yb alloys in wt.% measured by ICP-OES method.

The samples were mechanically polished and then etched with an etchant of 5 g picric acid, 6 ml acetic acid, 80 ml alcohol, and 10 ml water for optical microscope observation(OM, Zeiss MAT200).The microstructures were further characterized via scanning electron microscope(SEM,JSM-6610),transmission electron microscopy (TEM, FEI Talos F200X operated at 200 kV), and electron backscatter diffraction(EBSD, Oxford Instruments AZtecHKL).For the EBSD observation, the samples were sectioned from the cross-section of the extrudates and then electropolished in an aqueous solution of 5.2 g sodium thiocyanate, 9.4 g citric acid, 1.2 g 8-hydroxyquinoline, 100 ml ethanol, 10 ml isopropanol, 1.8 ml perchloric acid and 2.3 ml distilled water.The grain orientation spread (GOS) value smaller than 2° to isolate DRX grains and calculate the DRX fraction.The average grain size estimates were achieved by the linear intercept procedure.Tensile samples were cut into the dog-bone shape with 5 mm in gage diameter and 25 mm in gage length along the ED.Tensile tests were conducted at an initial strain rate of 10?3s?1at ambient temperature by using a CMT5504 tensile test machine and each record was the average of 3 individual measurements.

3.Results and discussion

3.1.As-cast microstructure

Fig.1 shows the optical micrographs of as-cast pure Mg and Mg-Yb binary alloys with different Yb concentrations.It can be observed that all alloys show typical columnargrain morphology and the columnar width is decreased from 384 μm in the case of pure Mg sample to 296 μm, 245 μm,and 209 μm in the cases of Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb alloys, respectively.Besides, Fig.1b manifests that the microstructure of the as-cast Mg-0.5 Yb sample is free of large compounds either in grains or at grain boundaries,indicating that Yb is uniformly dissolved in the Mg matrix.When a higher concentration of Yb is alloyed, as shown in Fig.1c and d, a large number of second phases discontinuously distribute in the matrix and their density and dimension gradually increase with increasing Yb concentration.According to X-ray diffraction data and the Mg-Yb binary phase diagram, these compounds can be identified as the Mg2Yb phase.

In this investigation, it was found that the microstructure was refined with increasing Yb concentration.This should correlate with the change of solidification behavior induced by Yb addition.According to the phases present in binary Mg-Yb alloy, the equilibrium partition coefficient (k) of Yb in Mg is less than 1 (k≈0.17).Therefore, Yb atoms tend to squeeze into the solid-liquid interface during solidification.As the solubility and diffusion rate of Yb in the Mg matrix are fairly low, a large number of Yb atoms will segregate to the solid-liquid interface, suppressing the grain growth during solidification.Meanwhile, the solute segregation during solidification resulted in the occurrence of constitutional supercooling.According to the literature [22], the parameter which quantifies the capacity of constitutional supercooling on nucleation for each alloy is the growth restriction factor,Q, defined as wheremis the liquidus slope,kis the equilibrium partition coefficient, andc0is the alloy composition.Thus, when 2.0 wt.% Yb is alloyed in pure Mg, theQvalue of about 5.1 is much lower than that of ZK60-2.0 Yb alloy (about 56.1),indicating a remarkably weak capacity of constitutional supercooling for the Mg-2.0 Yb alloy possesses.

It is generally accepted that the driving force for nucleation in the melt of alloy is generated by the formation of constitutional supercooling, and once it exceeds the undercooling required for nucleation on a potent particle, then nucleation of the equiaxed grain initiates [23].Therefore, the sufficient constitutional supercooling combined with the potent nucleation agent (Zr) rendered the as-cast ZK60-2.0 Yb alloy an equiaxed grain structure [10,14], whereas the Mg matrix with a high concentration of 2.0 wt.% Yb incorporation still retained the columnar-type grain morphology.In other words,the constitutional supercooling effect induced by 2.0 wt.%Yb alloying alone can only play a role in hindering the growth of original columnar grains rather than promoting the formation of a nucleus.Accordingly, it is suggested to incorporate other powerful nucleating agents such as Zr or Al to improve the modification effect of Yb in magnesium alloys.

3.2.As-extruded microstructure

3.2.1.Texture characteristics

The representative EBSD inverse pole figure (IPF) maps,(0001) pole figures (PFs), and the inverse pole figures for ED (ED IPFs) of as-extruded pure Mg, Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb extrudates in cross-section are shown in Fig.2a, b, c, and d, respectively.It is found that the initial columnar type microstructure has transformed into equiaxed grains despite some coarse deformed grains remained,indicating that incomplete dynamic recrystallization occurs during hot extrusion.With increasing Yb concentration, the average grain size and the DRX fraction (fDRX) decrease from ?18.1 μm and 98.9% in the case of pure Mg alloy to ?10.0 μm and 97.3%, ?3.4 μm and 52.1%, ?2.2 μm and 46.7% in the cases of 0.5, 1.0, and 2.0 wt.% Yb alloyed samples, respectively.This heterogeneous microstructure with gradually refined recrystallized grains suggested that the addition of Yb retarded the DRX progress of the Mg matrix,which was in good agreement with the observation in the hot extrusion of Mg-Zn-Yb-Zr alloys [15].According to our previous investigations [24,25], the variation in DRX kinetics of Yb-containing Mg alloys should be primarily associated with the solute drag effect induced by Yb alloying.Moreover, it is obvious that with increasing Yb concentration, the texture randomization considerably increased accompanied by a notable decrease in the maximum texture intensity from 11.9 m.u.d.(the overall multiples of the uniform distribution,m.u.d.) in the case of pure Mg to 8.3, 7.8 and 5.2 m.u.d.in the cases of Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb samples, respectively.

Fig.2.EBSD IPF maps and their corresponding (0001) PFs and ED IPFs obtained from cross-sections of as-extruded (a) pure Mg, (b) Mg-0.5 Yb, (c)Mg-1.0 Yb, and (d) Mg-2.0 Yb alloys. Aver shows the average grain size and ?DRX indicates the area fraction of recrystallized grains.

As shown in Fig.2b, the ED IPF of the Mg-0.5 Yb sample exhibits a typical＜01-10＞-＜?12-10＞// ED fiber texture, which is analogous to the pure Mg extrudate(Fig.2a), indicating that a trace amount of Yb addition exerts limited influence on the texture pattern, despite a pronounced decrease in intensity.In the case of Mg-1.0 Yb sample(Fig.2c), a weaker＜01-10＞// ED texture is generated suggesting that the＜01-10＞direction of most grains is nearly parallel to the ED.Subsequently increasing Yb concentration to 2.0 wt.% (Fig.2d), some new characteristics appear in the PF including more secondary poles and the increased tilt angle from TD to ED of the poles (from ?5.7° to ?31.5°).The corresponding ED IPF of the Mg-2.0 Yb sample exhibits a dominant orientation around＜04-43＞, suggesting that thec-axes of most grains incline to the ED, which is different from the pure Mg and lower Yb-containing counterparts mainly with theirc-axes perpendicular to the ED after extrusion.The mechanisms underlying the aforementioned texture development will be discussed in the next section.

3.2.2.Dominant deformation mode

Available research indicates that the addition of RE elements into Mg alloys can change the critical resolved shear stress (CRSS) and flow stresses of various deformation modes, which will certainly change the activities of various slip modes during plastic deformation and, as a result,affect the final deformation texture [26-28].Therefore, it is of importance to understand the effect of Yb addition on the dominant deformation mode operated in the extrusion of Mg-Yb binary alloys.To clarify this, the in-grain misorientation axes (IGMA) analysis and TEM analysis were performed to evaluate the active slips during straining.

(1) IGMA analysis

The IGMA analysis is obtained from the EBSD measurements, based on the primary assumption that bending of crystal lattice under the action of slip takes place about a certain Taylor axis.Therefore, the dominant slip mode in a deformed grain can be determined by simply matching the Taylor axis for a given deformation mechanism (Table 2) to the experimentally measured IGMA [29].On this basis, it can be concluded that the domination of prismatic＜a＞slip triggers a lattice rotation about＜0001＞axis and coactivation of slip variants belonging to basal＜a＞slip or the second-order pyramidal＜c+a＞slips, both of which have＜10-10＞as their Taylor axes, can develop the＜uvt0＞-type IGMA distribution.In this work, the misorientation axes of material-pointpairs with angles between 1.2 and 2.0° are plotted since the available study reveals that increasing the upper cutoff misorientation angle above 2.0° exerted a negligible effect on the qualitative distribution features of IGMA [30].

Table 2Deformation Modes available in Magnesium and Corresponding Taylor Axes[29].

The representative IGMA distributions developed in the extruded Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb alloys are shown with their kernel average misorientation (KAM) maps in Fig.3.As the KAM map is indicative of the local misorientation level,an increased density of dislocations is observed in samples with increasing Yb concentration.For the 0.5 wt.%Yb alloyed specimen, the IGMA distributions of the selected deformed grains with higher dislocation densities (marked by capital letters A-F in Fig.3a) are depicted in Fig.3d.In grains A-E, weak in-grain orientation spreads (IGOS) with the maximum intensity (MI) of IGMA lower than 2.0 m.u.d.are observed,indicating a small extent of lattice rotation takes place, and thus they are considered to have a uniform IGMA distribution [30].However, it is interesting to find that the distribution of IGMA in grain F is concentrated around the＜0001＞axis with the MI higher than 2.0 m.u.d.,which is regarded to develop preferential IGMA and the prismatic＜a＞slip is dominantly operative.Notably, the number of grains with this IGMA distribution is very limited, thereby, exerting a negligible effect on the overall texture development.That is why the fiber texture with an＜01-10＞-＜?12-10＞arc in the case of Mg-0.5 Yb extrudate is analogous to that of Yb-free counterpart.

In the case of Mg-1.0 Yb counterpart (Fig.3b and e), a substantially different characteristic of IGMA distribution is found: strong intensities with the MI higher than 2.0 m.u.d.around the＜0001＞axis are observed in all selected grains.This＜0001＞-type IGMA distribution combined with the development of the＜01-10＞// ED texture indicates that the prismatic＜a＞slip should be robustly activated in the pathway of Mg-1.0 Yb extrusion.Moreover, when 2.0 wt.% Yb is alloyed (Fig.3c and f), besides the pronounced spread of prismatic poles observed in grains B and D, a different distribution pattern with diffused intensity (MI＞2.0 m.u.d.) of IGMA around the＜0001＞axis is found in grains A, C, E,and F.This phenomenon may be caused by the minor coactivation of other slip modes including pyramidal＜a＞slip orpyramidal＜c+a＞slip, which in turn results in slight rotations of the lattice about different Taylor axes [30].To clarify this, TEM observation was performed to determine the dominant slip mode based on theg?b= 0 invisible criterion.

(2) TEM identification

Fig.4 shows the TEM images of as-extruded Mg-Yb binary alloys.In the case of Mg-0.5 Yb sample (Fig.4a), a few Yb-containing globular particles (cf.Fig.4b inset, the EDX elemental mapping of Yb) with 50-80 nm in diameter sparsely distributed at grain boundaries.The high magnification images in the conditions ofg= 0001 (Fig.4b) andg= 10-10 (Fig.4c) demonstrate that the existence of residual＜a＞dislocations (as indicated by blue arrows in Fig.4c)in as-deformed grains but no obvious＜c+a＞dislocations can be detected.It was in good agreement with the findings in Fig.3a and d, presenting a low dislocation density with negligible intensities of IGMA distribution.However,compared with the counterpart with a low Yb concentration,Mg-2.0 Yb sample exhibited a different precipitation pattern with massive activated dislocations, i.e., blocky particles with 500-600 nm in length (as indicated by the dashed circle in Fig.4d) mainly precipitated at grain boundaries in companion with fine dot-like particles profusely distributed in grain interiors (cf.Fig.4e, as yellow arrows indicated).It should be related to a higher concentration of Yb incorporated in the Mg matrix and the limited solid solubility of Yb under the deformation temperature.Given the pinning effect of precipitates and higher stored energy, as a result, a large population of dislocations appeared in as-deformed regions,including an increased density of＜c+a＞dislocations (as indicated by red arrows in Fig.4e) and residual＜a＞dislocations (as blue arrows indicated in Fig.4f).As previous investigations indicated[31-33], the elevated deformation temperature combined with a high concentration of RE solute decreased the critical resolved shear stresses (CRSS) of pyramidal＜c+a＞slip,bringing the CRSS values of basal and non-basal slips closer together, which was mainly responsible for the increased activity of basal and＜c+a＞slip observed in the Mg-2.0 Yb extrudate.Furthermore, it is found that the＜c+a＞dislocations mainly lay on the basal plane and these immobile dislocations may lose the ability to accommodate c-axis deformation, resulting in a low ductility during the subsequent stretching [34].

3.2.3.Mechanisms of texture development

From the above analysis, it could be concluded that a continuous and pronounced basal texture weakening was in companion with the change in the dominant deformation mode with increasing Yb addition.Nevertheless, the question was still open as to why the single＜01-10＞// ED texture and the subsequent＜04?43＞// ED component were developed with the increase of Yb concentration.

For the texture transformation from an＜01?10＞-＜?12-10＞fiber in the case of Mg-0.5 Yb extrudate to a weak＜01?10＞// ED texture in the 1.0 wt.% Yb alloyed counterpart, an analogous phenomenon was reported by Honniball et al.[35]who suggested that the stability of＜01-10＞grains towards prismatic slip locked them in place and thus resulted in the development of a single＜01-10＞// ED texture.Besides,Mayama[36]demonstrated that a single slip of the prismatic slip led to the lattice rotation around the [0001]axis,which aligned the＜-12-10＞direction with the compressivedirection, i.e.during extrusion, the {01-10} plane tended to orient in the extrusion direction.This transformation can be schematically depicted as shown in Fig.5.

Fig.3.KAM and IGMA distributions of as-extruded (a, d) Mg-0.5 Yb, (b, e) Mg-1.0 Yb, and (c, f) Mg-2.0 Yb alloys.The IGMA distribution for the deformed grains marked as A through F in (a) Mg-0.5 Yb, (b) Mg-1.0 Yb, and (c) Mg-2.0 Yb alloys.The maximum/minimum intensities of each IGMA distribution are also given at the bottom of each IGMA distribution.

On the other hand, for the ED-titled texture observed in Mg-2.0 Yb sample, it may be attributed to the recovery of basal/＜c+a＞dislocations.According to Imandoust et al.[37], relaxation of basal/＜c+a＞dislocations in REcontaining Mg alloys initiates a “backward rotation” toward the [0001]pole, which in traditional Mg alloys rarely moves the grains more than 15° toward the [0001]pole.It was suggested that the effects induced by RE addition including reducing the grain boundary energy, decreasing the CRSS for＜c+a＞slip, solute-drag effect, and delayed DRX process were in combination to prolong the nucleation stage and that higher dislocation densities, particularly of＜c+a＞dislocations, were present to drive it.These dense dislocations repel one another which induces the backward crystal rotation as they move apart during recovery.

In this work, it was interesting to find that Yb addition exerted similar effects on pure Mg during hot extrusion.The high-angle annular dark-field (HAADF) images and the corresponding EDX elemental mappings demonstrate that the Yb element is potent to segregate to grain boundaries existing in the form of precipitates, and/or solute atoms (cf.Fig.6).The segregation of Yb at grain boundaries may result in reduced grain boundary energy and exert a strong drag effect dur-ing boundary deformation and/or migration [28].Meanwhile,the addition of Yb reduced the CRSS values of non-basal slips, leading to a considerable increase in dislocation density in Mg-2.0 Yb extrudate.As shown in Fig.7a and b,the densities of residual low angle grain boundaries (LAGBs)in as-extruded Mg-Yb binary alloys increase with increasing Yb concentration.TEM micrographs (Fig.7c and d) further demonstrate that a large number of sub-grains and LAGBs(indicated by cyan arrows in Fig.7d) appear in the Mg-2.0 Yb extrudate.Therefore, based on the aforementioned observations, it could be concluded that the continuous DRX mechanism may dominate the Mg-2.0 Yb extrusion and its DRX progress should be considerably retarded with a prolonged nucleation stage as well as a higher dislocation density,which was also in good agreement with the results observed in Fig.2.

Fig.4.Low- and high-magnification TEM images of as-extruded (a-c) Mg-0.5 Yb and (d-f) Mg-2.0 Yb alloys in two-beam conditions of (b, e) g = [0001]and (c, f) g = [10-10], respectively.

Fig.5.Schematic explanation of the stability of the grain orientation with {01-10} plane normal to the extrusion direction induced by prismatic slip.

It is worth illustrating that sub-grains formation generally minimizes strain energy and is a potential mechanism by which a crystal rotation can occur without twinning or phase transformation [38].However, the angle of typical sub-grain boundaries is often small and it is hard to account for the larger misorientation observed.Fig.8 shows the subsets ofKAM maps with dynamic recrystallized (DRXed) grains differentiated from deformed grains, and the corresponding ED IPFs for Mg-1.0 Yb and Mg-2.0 Yb extrudates.It is interesting to find that the prismatic＜a＞slip dominated Mg-1.0 Yb sample develops the primary＜01-10＞// ED texture both in the subsets of DRXed and deformed grains,while the pyramidal＜c+a＞slip operated Mg-2.0 Yb counterpart generates two different texture components, i.e.,deformed grains exhibiting a strong orientation around＜04-43＞and the DRXed grains presenting a dominant orientation distribution along＜03?34＞-＜?12-12＞arc.It is clear that thec-axes of the DRXed and deformed grains incline to the ED to some extent.Given that an increased density of basal＜a＞and pyramidal＜c+a＞dislocations are appeared in Mg-2.0 Yb extrudate(cf.Figs.4e and 7b), the development of ED-tilted texture component may be attributed to the backward crystal rotation toward the [0001]pole during relaxation of basal/＜c+a＞dislocations.Moreover, a larger rotation angle exhibited by DRX grains may be due to the segregation of Yb atoms to grain boundaries that led to the oriented growth of recrystallized grains.

Fig.6.HAADF-STEM images and EDXS maps showing grain boundaries, dislocations, precipitates and the distribution of Mg and Yb elements in Mg-2.0 Yb extrudate.

Based on the above-mentioned discussion, it could be concluded that the basal＜a＞slip should be operated in the extruded Mg-0.5 Yb alloy, generating a typical＜01-10＞-＜?12-10＞fiber texture;the prismatic＜a＞slip was robustly activated in as-extruded Mg-1.0 wt.% Yb alloy, which resulted in the lattice rotation around the [0001]axis and stabilized the＜-12-10＞direction aligned to the compressive di-rection, promoting the formation of a single＜10-10＞// ED texture; when a high concentration of 2.0 wt.% Yb alloyed,the increased activity of pyramidal＜c+a＞slip and the relaxation of basal/＜c+a＞dislocations resulted in the lattice rotation withc-axes of most grains inclined to the ED, developing the subsequent＜04-43＞// ED texture component.

Fig.7.EBSD results showing HAGBs and LAGBs in (a) Mg-0.5 Yb and (b) Mg-2.0 Yb extrudates, and TEM bright-field images of Mg-2.0 Yb alloy showing (c) sub-grains and (d) LAGBs.

3.3.As-annealed microstructure

3.3.1.Texture characteristics

According to preliminary investigations,Yb-containing Mg alloys exhibited good thermal stability, especially for the alloys with high Yb concentrations.A long time of heating with a relatively low annealing temperature could result in the coarsening of precipitates, which was detrimental to the ductility of the as-annealed samples.Therefore, a relatively high temperature with short durations was adopted in subsequent annealing treatments.Meanwhile,a comparable average grain size was achieved in the Yb-containing samples after annealing to decrease the influence of grain size and highlight the effect of annealing texture on the resulting tensile properties.

Fig.9 illustrates the representative EBSD IPF maps of the pure Mg, Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb extrudates annealed at 400 °C for different durations of 15, 30,60, and 100 min, respectively.The average grain size of each sample was labeled in figures.In the Mg-Yb binary alloys,the homogeneity of grain size decreased with increasing Ybconcentration during annealing despite a similar average grain size of 13.0-14.0 μm was achieved, indicating that the preferential grain growth was enhanced with the increase of Yb content.Another microstructure characteristic worthy of note was that after prolonging the incubation time to 100 min, almost 3 times more than that of 0.5 Yb alloyed counterpart,Mg-2.0 Yb alloy could just obtain a comparable average grain size, further verifying that a high concentration of Yb alloying resulted in the low mobility of grain boundaries during the grain growth process.

Fig.8.Subsets of KAM maps with recrystallized grains differentiated from deformed grains, and the corresponding ED IPFs for (a, b, and c) Mg-1.0Yb and(d, e, and f) Mg-2.0 Yb extrudates.

Fig.9.EBSD IPF maps and their corresponding (0001) PFs and ED IPFs obtained from cross-sections of as-annealed (a) pure Mg, (b) Mg-0.5 Yb, (c)Mg-1.0 Yb, and (d) Mg-2.0 Yb alloys. Aver shows the average grain size.

The (0001) PFs and the corresponding ED IPFs in Fig.9ad exhibit the texture development of pure Mg, Mg-0.5 Yb,Mg-1.0 Yb, and Mg-2.0 Yb samples during annealing.Basal texture weakening including the dispersed texture and the decreased maximum intensity was observed in the samples with increasing Yb concentration, i.e., more secondary poles with an increased tilt angle from TD to ED of the poles (from ?7.6° to ?52.2°) were evident and the maximum pole intensity of (0001) PF reduced from 14.1 m.u.d.in the case of pure Mg to 9.2, 4.8 and 3.8 m.u.d.in the cases of Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb samples, respectively.Compared with the as-extruded counterparts, the maximum pole intensities were weakened both in as-annealed Mg-1.0 Yb (7.8 to 4.8 m.u.d.) and Mg-2.0 Yb (5.2 to 3.8 m.u.d.)samples except for pure Mg and Mg-0.5 Yb alloys, contrarily, exhibiting an enhanced intensity from 11.9 to 14.1 and from 8.3 to 9.2 m.u.d.respectively, which should be related to the preferential growth in the direction of＜-12-10＞during annealing, as the peak shown in ED IPF in Fig.9a and b.When further increased the Yb concentration, there was a tendency of shifting from the＜-12-10＞peak in the case of pure Mg and Mg-0.5 Yb samples to the spreading of the＜-24-23＞and＜-12-13＞peaks in the cases of 1.0 and 2.0 wt.% Yb alloyed counterparts, respectively.The movement of texture peak to the [0001]axis along the [0001]-＜?12-10＞side of the stereographic triangle suggested that thec-axes of most grains inclined to the ED and, interestingly, a higher Yb concentration corresponded to a larger inclination angle.The development of ED-tilted texture in as-annealed Mg-Yb binary alloys will discuss in the next section.

3.3.2.Preferential grain growth

It was interesting to find that with increasing Yb concentration, the texture development of Mg-Yb binary alloys after annealing exhibited an opposite shift direction on the stereographic triangle to the [0001]axis with respect to that evolution during the extrusion.The above-mentioned texture development in companion with a pronounced basal texture weakening during high-temperature annealing should be intimately correlated with the preferential growth of grains with non-basal or near-random orientations.Therefore, preferential grain growth was adopted to elaborate on the mechanisms underlying texture development after high-temperature annealing.To clarify this, the as-annealed microstructure of Mg-Yb binary alloys was separated into two categories based on the grain size: smaller and larger than 1.5 times the average grain size (?20 μm), and they were analyzed by EBSD,respectively.

Fig.10 depicts the subsets of IPF maps and the corresponding (0001) PFs and ED IPFs for small (≤20 μm) and large (＞20 μm) grains of Mg-0.5 Yb, Mg-1.0 Yb, and Mg-2.0 Yb alloys from Fig.9b, c, and d, respectively.Compared with the large grains, the small ones exhibited a more dispersed distribution.As illustrated in Fig.10a, c, and e, the(0001) PFs of the small grains show relatively weak basal poles with much broader orientation distributions, especially for the Mg-2.0 Yb sample.On the contrary,the(0001)PFs of the large grains (Fig.10b, d, and f) exhibit a more intensive distribution of basal poles with a larger maximum pole intensity compared with that of the small grains, which demonstrates the occurrence of preferential grain growth during annealing.The ED IPFs may provide a direct suggestion for the specific growth direction in Mg alloys with different Yb concentrations.As shown in Fig.10a and b, when 0.5 wt.%Yb is alloyed, a broad orientation distribution around＜-12-10＞pole displayed by small grains gradually evolves to an intensive one during grain growth, which is analogous to the preferential grain growth in pure Mg.In the case of Mg-1.0 Yb sample(Fig.10c and d),the grain growth changes the texture pattern from a broad orientation distribution along＜02-21＞-＜?12-11＞arc to a strong＜＜-24-23＞peak, while the addition of 2.0 wt.% Yb leads to a transition in the texture pattern from a dispersed＜02-21＞-＜?24-23＞arc to an enhanced＜?12-13＞peak (Fig.10e and f).

According to the recent study [39], the preferential growth direction ofα-Mg corresponds to the＜2-1-10＞and＜2-1-1x＞directions, where＜2-1-10＞is the preferential growth direction of the basal plane, and＜2-1-1x＞is for the nonbasal planes.In the current investigation, Mg-0.5 Yb alloy exhibited a typical oriented grain growth towards the＜2-1-10＞direction, which was analogous to the pure magnesium (cf.Figs.2a and 9a), suggesting that a trace amount of Yb addition exerted limited influence on the preferential growth of matrix.However, when higher concentrations of Yb were alloyed, the specific growth orientations changed from＜2-1-10＞to＜4-2-23＞and＜2-1-13＞in the cases of Mg-1.0 Yb and Mg-2.0 Yb samples, respectively.This preferred growth direction of non-basal planes should be related to the change in surface energy and the crystallographic anisotropy triggered by sufficient Yb addition.Notably, on the other hand, the more the Yb alloyed, the closer the preferred growth direction to the [0001]on the stereographic triangle, which indicated that the preferential growth direction of non-basal planes was highly dependent on the solution concentration.A similar phenomenon was also observed in Mg-Al and Mg-Zn alloys.Du et al.[40]reported that with increasing Al concentration, the preferential growth direction of the non-basal plane changed from＜11-23＞to＜22-45＞, and for Mg-Zn alloys, this direction changes from＜11-23＞to＜22-45＞or＜11-22＞as the Zn concentration increased.

Therefore, It can be concluded that the texture weakening/altering in Mg-Yb binary alloys during high-temperature annealing was caused by grains with the basal orientation vanished and with the non-basal orientations intensified due to an increase in Yb addition.Importantly, the root of this tendency, as Bohlenet al.indicated [41], may be linked with the boundary pinning caused by RE segregation and/or REcontaining particles.

Fig.10.EBSD IPF maps, (0001) PFs, and the corresponding ED IPFs of the grains (a, c, and e) smaller or (b, d, and f) larger than 1.5 times the average grain size (?20 μm) of as-annealed (a, b) Mg-0.5 Yb, (c, d) Mg-1.0 Yb, and (e, f) Mg-2.0 Yb alloys.

3.4.Mechanical properties and fracture characteristics

3.4.1.Mechanical properties of as-extruded alloys

The true stress-strain curves of as-extruded pure Mg and Mg-Yb binary alloys are depicted in Fig.11a.The tensile yield strength (TYS), ultimate tensile strength (UTS), and elongation to fracture (EL) of samples along the extrusion direction are listed in Table 3.It is clear that with increasing Yb concentration, the TYS and UTS significantly improved in companion with a decrease in EL except for the pure Mg extrudate.Specifically, the TYS increased from ?78 MPa in the case of pure Mg sample to ?153, ?244, and ?304 MPa in the cases of 0.5, 1.0, and 2.0 wt.% Yb alloyed counterparts,while the EL firstly increased from ?5.6% in the case of pure Mg to ?9.8% when 0.5 wt.% Yb alloyed and then decreased to ?5.7 and ?4.8% in the cases of Mg-1.0 Yb and Mg-2.0 Yb counterparts.

Table 3Mechanical properties of as-extruded and the subsequently annealed Mg-Yb binary alloys when stretched along the extrusion direction.

The results revealed that Yb addition posed a significant effect on the mechanical properties of as-extruded Mg alloys.In the cases of Mg-Yb binary alloys, although the texture weakening with increasing Yb concentration was observed,the grain refinement still dominated the strength improvement,and the local stress concentration induced by the increased density of coarse precipitates and undissolved phases (as indicated by black arrows in Fig.12) was primarily responsible for the ductility deterioration.However, it was interesting to find that a distinct improvement in both strength and ductility was observed in the Mg-0.5 Yb extrudate compared with that of pure Mg.The enhancement in strength should be attributed to the pronounced grain refinement and solution strengthening induced by a trace amount of Yb addition, and meanwhile,texture weakening may account for the ductility improvement.

Fig.11.(a) Stress-strain curves and (b) work-hardening rate curves of as-extruded pure Mg and Mg-Yb alloys.

Fig.11b shows the correspondingΘ-(σ-σ0.2) curves,whereΘ= dσ/ dεrepresents the strain-hardening rate.It is observed that the initial hardening rate increases with increasing Yb concentration and the work hardening behaviors exhibited two distinct characteristics.In the cases of Mg-0.5 Yb and Mg-1.0 Yb samples, the strain-hardening curve consisted of an extended elastic-plastic transition (EPT) stage and a steadily decreased stage III.The absence of stage II in these curves suggested that prismatic＜a＞slips may be dominantly operative when stretched at ambient temperature.However, in the case of Mg-2.0 Yb sample, a short stage with a relatively constant work hardening rate after the EPT,corresponding to stage II, was observed, followed by a steep decrease in the strain-hardening rate.The presence of stage II in the 2.0 wt.% Yb alloyed sample indicated that the basal＜a＞slip was activated, but the suppression in this stage to some extent also suggested a rapid transition of the dominant deformation mode.

It was worth mentioning that the aforementioned features in strain-hardening curves were well coincident with the texture characteristics discussed in Section 3.2.1.The basal textured pure Mg, Mg-0.5 Yb, and Mg-1.0 Yb alloys with fairly low values of the Schmid factor for basal slip, as shown in Fig.13a, b, and c, are unfavorable for the activity of basal＜a＞slip, may promote the activation of prismatic＜a＞slip to accommodate strains and lead to a direct transition from EPT to stage III in the strain-hardening curve.On the contrary,the 2.0 wt.%Yb alloyed counterpart withc-axes of most of the grains tilted towards the ED exhibited a relatively high Schmid factor for the basal slip of ?0.3 (Fig.13d), which was favorable for the activation of basal＜a＞slip at the initial straining and thus a flat stage II appeared.

3.4.2.Mechanical properties of as-annealed Mg-Yb alloys

As the grain size of pure Mg after extrusion (?18.1 μm)was much larger than that of Yb-containing Mg extrudates after annealing (?13.0 μm), therefore in this part of the investigation, more attention was paid to the as-annealed Mg-Yb binary alloys with a similar average grain size to highlight the effect of annealing texture on the resulting tensile properties.Fig.14a and b depict the true stress-strain curves and the correspondingΘ-(σ-σ0.2)curves of as-annealed Mg-Yb binary alloys with different Yb concentrations.It is interesting to find that an opposite trend in mechanical properties is observed in these curves compared with those of as-extruded Mg-Yb binary samples, i.e.with increasing Yb concentration, the strength is considerably decreased accompanied by an increase in elongation to failure (cf.Table 3).Considering the fact that all the samples achieved a comparable grain size after annealing, the duration of heating and the resultant variation in grain growth, precipitates, and undissolved phases may account for this reversed effect.Moreover, it was also noticed that the strength in the as-annealed condition was much lower than that of as-extruded counterparts (except for the 0.5 wt.% Yb alloyed sample), which should be attributed to apparent grain growth during high-temperature annealing and thus weakening the effect of grain refinement strengthening dominated in the extruded samples.

After annealing, the Mg-0.5 Yb alloy displayed a limited grain growth with an increase in the average grain size from ?10.0 to ?13.5 μm, and an intensified texture intensity with the maximum texture intensity increasing from 8.3 to 9.2 m.u.d., which was due to a transition from a typical＜01-10＞-＜-12-10＞texture to a concentrated＜-12-10＞fiber.The basal texture intensification combined with the solid solute strengthening due to phase dissolution during annealing (Fig.12a and d) resulted in an increase in strength and deterioration in elongation.The strain-hardening curve(black curve in Fig.14b) rationalized it with a direct transition from EPT to stage III.The initial “hard” orientation unfavorable for the basal＜a＞slip (a low Schmid factor for basal slip,cf.Fig.13f) promoted＜a＞screw dislocations to enter the prismatic plane through cross-slip [42], which resulted in a high dynamic recovery rate, corresponding to a low slope of stage III (dΘ/dσ,cf.Fig.14b).

On the other hand, the ductility improvement and the strength deterioration observed in as-annealed Mg-1.0 Yb and Mg-2.0 Yb samples were attributed to the homogeneous grain structure, weak ED-tilted texture, and dissolution of coarse phases.As shown in Figs.2 and 9, the average grain sizes in extruded and annealed conditions were considerably increased from ?3.4 to 13.3 μm in the case of Mg-1.0 Yb sample, and from ?2.2 to ?13.0 μm in the case of Mg-2.0 Yb alloy, respectively.An increase in the recrystallization extent combined with the preferential grain growth resulted in the development of ED-tilted texture with weaker intensities(MI from 7.8 to 4.8 m.u.d., and 5.2 to 3.8 m.u.d., respectively), leading to an increase in the Schmid factor for basalslip and elongations.Moreover, the annealing treatment at a high temperature of 400 °C for a long duration promoted the dissolution of coarse precipitates and undissolved phases, especially for the Mg-2.0 Yb alloy (cf.Fig.12c and f), which effectively reduced the tendency of stress concentration at interfaces.Although the solid solute strengthening induced by phase dissolution may appear, this effect should be counterweighted by a loss in grain refinement due to the pronounced grain growth during high-temperature annealing.

Furthermore,it was worth noting that Mg-1.0 Yb and Mg-2.0 Yb annealed samples exhibited a similar characteristic to the 0.5 wt.% Yb alloyed counterpart in the strain-hardening behavior and higher slopes of stage III were achieved (cf.Fig.14b).The suppression of stage II in the samples with high Schmid factors for basal slip may be due to the reduced discrepancy in the critical resolved shear stress (CRSS) of basal and non-basal slip systems induced by Yb solute, which promoted the activity of multiple slip systems and thus reduced the contribution from the basal slip.Previous studies [33,43]reported that the CRSS values of non-basal slips considerably decreased with RE solute, which contributed to a decrease in plastic anisotropy.

3.4.3.Fracture characteristics

Fig.15 depicts the fracture morphologies of as-extruded and as-annealed Mg-Yb binary alloys after tensile tests.An opposite trend in fracture characteristics was also observed in the samples with different conditions.In the cases of as-extruded samples, the fracture mechanism transited from the mixed brittle and ductile fracture to the brittle fracture mode with increasing Yb concentration.It should be related to the coarsening of precipitates and an increased density of undissolved phases with increasing Yb addition (cf.Fig.12a,b, and c) since Yb has a low solubility and diffusion rate in Mg [44].The blocky phases may hinder dislocation slips and thus strengthen the matrix to some extent, but cracks are more readily initiated at the interfaces between coarse phases and the Mg matrix [45].As Fig.16 demonstrated,a large number of fractured particles were observed in the as-extruded Mg-2.0 Yb fracture surface, indicating the stress concentration was evident.What’s more, the typical of twin-induced crack planes (large cleavage planes) are observed in the fracture surfaces of Mg-1.0 Yb and Mg-2.0 Yb extrudates (Fig.15b and c), which indicates that {10-11}contraction twins and {10-11}-{10-12} double twins may be formed during tension.As shown in Fig.2, the area fraction of unDRXed grains increases because of the suppression of DRX induced by increasing Yb concentration.Twins were easily formed in these coarse unDRXed grains during the subsequent tension along the ED and then acting as crack initiation sites during fracture.It was rationalized the lowest elongation to failure of ?4.8% of the as-extruded Mg-2.0 Yb sample obtained.On the other hand, for the annealed sample,the fracture mechanism transited from the ductile-brittle fracture mode with dimples and cleavage planes to the ductile fracture mode with increasing Yb concentration.This should be attributed to the development of ED-titled texture (cf.Fig.9c and d) and apparent dissolution of coarse precipitates and undissolved phases after high-temperature annealing (cf.Fig.12e and f), which effectively reduced the tendency of stress concentration at interfaces.The fracture characteristics were well consistent with the results of mechanical properties.

From the above discussion, it could be concluded that with increasing Yb concentration, the CRSS of non-basal slips decreased, which increased the activities of non-basal slips and randomized/weakened the basal texture.It improved the ductility of Mg-Yb binary alloys when stretched along the ED, however, the increased density of coarse precipitates and undissolved phases due to a high concentration of Yb addition resulted in the occurrence of localized deformation and premature failure.Therefore, it was suggested that other approaches such as alloying, additional homogenization, or high-temperature deformation could be incorporated in the Mg-Yb base alloys to refine precipitates and promote the dis-solution of second phases, and by these means, the improved strength may be achieved without the loss of plasticity.The above-mentioned attempts will be reported in our forthcoming paper.

Fig.15.SEM images of the tensile fracture surfaces of (a-c) as-extruded and (d-f) as-annealed (a, d) Mg-0.5 Yb, (b, e)Mg-1.0 Yb, and (c, f) Mg-2.0 Yb alloys.

Fig.16.(a) SEM image of the tensile fracture surface of as-extruded Mg-2.0 Yb alloy and (b) its corresponding EDS Mapping.

4.Conclusion

In this study, the effects of Yb concentration on the texture development and mechanical properties of Mg-Yb binary alloys during hot extrusion and the subsequent annealing were systematically investigated, and key conclusions were drawn:

(1) Yb addition effectively refines the as-cast grains but exerts a negligible effect on the transition from a typical columnar to the equiaxed grain structure.The constitutional supercooling induced by a high concentration of Yb addition alone can only play a role in hindering the growth of original columnar grains rather than promoting the formation of a nucleus.

(2) The IGMA analysis combined with TEM observation indicate that non-basal slips are operated with increasing Yb concentration.For the Mg-1.0 Yb alloy, the prismatic＜a＞slip is robustly activated, promoting the formation of the＜10-10＞// ED texture; when 2.0 wt.% Yb alloyed, the increased activity of pyramidal＜c+a＞slip and the relaxation of basal/＜c+a＞dislocations generate the ED-tilted texture.

(3) The annealed texture intensities shift from the＜-12-10＞peak to the spreading of the＜-24-23＞and＜-12-13＞peaks with increasing Yb concentration.The texture weakening/altering in Mg-Yb binary alloys during high-temperature annealing is caused by grains with thebasal orientation vanished and with the non-basal orientations intensified due to the increased concentration of Yb solute.

(4) An opposite trend in mechanical properties is observed in the samples in extruded and annealed conditions.The Mg-2.0 Yb extrudate exhibits a high TYS of 304 ± 3.5 MPa while the subsequently annealed counterpart presents a favorable EL of 14.8 ± 1.2% when stretched along the ED.Yb-induced grain refinement improves the strength after extrusion, and the combination of homogeneous grain structure, weak ED-tilted texture, and dissolution of coarse phases after hightemperature annealing accounts for the elongation enhancement.

Declaration of Competing Interest

The authors declare that there is no conflict of interest.

Acknowledgments

This work was financially supported by the National Natural Science Foundation of China (Grant Nos.51975484 and 51605392), the Natural Science Foundation Project of CQ CSTC (Grant No.cstc2020jcyj-msxmX0170), and the Fundamental Research Funds for the Central Universities (Grant No.XDJK2020B001).