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Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition
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Review Article
Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition Qiyu Liao a , Qichi Le a,∗ , Xingrui Chen a , Xiaoqiang Li a , YanChao Jiang a , Lifeng Ma b , Wenxin Hu c a
Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang 110819, China Taiyuan University Science and Technology, Heavy Machinery Engineering Research Center, Ministry Education, Taiyuan, Shanxi, 030024, China c Baotou Research Institute of Rare Earths, Baotou, 014000, China b
a r t i c l e
i n f o
a b s t r a c t
Article history:
Tensile test at high temperature and electron backscatter diffraction (EBSD) techniques were
Received 10 November 2019
combined to investigate the deformation behavior and microstructure evolution of the as-
Accepted 17 January 2020
extruded AZ110 magnesium alloy with La-rich Mish Metal addition (AZ110LC for short). The
Available online xxx
stretching temperature and strain rate have great effect on the flow stress of the alloy. The flow stress rapidly reaches the peak stress and then decreases with further deformation until
Keywords:
the fracture. The maximum elongation of the alloy is 840% under 573 K and 1.7 × 10−4 s−1 . The
AZ110 magnesium alloy
excellent superplasticity ascribes to the refinement of microstructure and the distribution of
La-rich Mish metal
sub-structures of the as-extruded AZ110LC alloy. Moreover, the strain rate sensitivity coeffi-
Superplastic deformation behavior
cient of the as-extruded AZ110LC alloy is 0.41 and the active energy is 87 kJ/mol under 523 K
Microstructure evolution
and 4.2 × 10−4 s-1 indicating that the deformation mechanism of the alloy is GBS. Finally, the
Cavities
shapes and fraction of cavities closed to tensile fracture have a great influence on the elongation of the tensile samples at different strain rates under 573 K. The formation and growth of the cavities can effectively relax the stress concentration and can effectively coordinate and compensate the GBS mechanism in superplastic deformation of the alloy. © 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
∗
Corresponding author. E-mails: [email protected] (Q. Liao), [email protected] (Q. Le). https://doi.org/10.1016/j.jmrt.2020.01.063 2238-7854/© 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/). Please cite this article in press as: Liao Q, et al. Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.063
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Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. Qiyu Liao received his B.S. and M.S. degree in Material processing engineering from the Northeastern University, Shenyang, China. He is currently with analysis of Plastic deformation of magnesium alloy and Development and application of new magnesium alloys. Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of Melting and casting of magnesium alloys; Plastic deformation of magnesium alloy; Light metal solidification in physical field (electromagnetic field and ultrasonic field); Development and application of new magnesium alloys; Mg alloys corrosion and battery; oxidation and combustion of Mg alloys. Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of application of ultrasonics on preparation of Mg alloys and Mg alloys corrosion and battery.
Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of Mg alloy composite material.
Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of numerical simulation of wrought magnesium alloys.
Institutional: Taiyuan University Science and Technology, Heavy Machinery Engineering Research Center, Ministry Education, Taiyuan, Shanxi, China. He is currently with analysis of optimization design theory and key technology of rolling equipment; theory and key technology of composite technology; nonferrous metal deep processing key technology.
Institutional: Baotou Research Institute of Rare Earths, Baotou, China
1.
Introduction
Wrought magnesium alloys are very attractive in various industrial products because of their low density, excellent mechanical properties, good damping, strong electromagnetic shielding ability and recyclability [1]. In spite of these advantages, the poor formability at room temperature largely limits the application of wrought magnesium alloys due to its hexagonal close packed (HCP) structure [2]. Therefore, utilizing superplastic deformation to enhance their formability is necessary in order to fabricate light structural parts with some complex geometry of thin-walled components. Generally, there are three basic requirements for the occurrence of superplasticity: fine microstructure of equiaxed grains with grain size typically less than 10 m; tensile temperature not exceeding 0.5 TM (TM is the liquidus temperature of metal materials); strain rate higher than 10−6 s-1 but less than 10-2 s-1 [3]. In order to acquire fine grains, several severe plastic deformation (SPD) technologies are employed to fulfill the requirement which are time-consuming and costly [4]. Hence, it is worth to consider alloying with appropriate elements and conventional deformation methods for fine-grain microstructure. Among various magnesium alloys, those alloys based on Mg-Al-Zn system are the most widely used wrought magnesium, due to their excellent mechanical properties and high productivity and good product qualities. However, the content of Al element affects the ductility of the Mg-Al-Zn alloys. Generally, the alloys containing high content of Al element such as AZ80, AZ91 have respectively low ductility at room temperature [5]. Their formability can be further improved by enhancing their ductility through incorporating superplastic behavior. Good formability can be obtained by adding rare earth elements to some Mg-Al-Zn alloys. When alloying with RE elements in Mg-Al-Zn alloys, some RE phases with good thermal stability precipitate in the matrix can effectively refine the initial microstructure. It was reported that the softening behavior of the low-melting-point phase Mg17 Al12 has contribution to the interphase boundary sliding at elevated temperatures for exhibiting superplasticity, but the high volume fraction of the coarse Mg17 Al12 phase leads to low room-temperature ductility [6]. Furthermore, some RE elements can change the texture type of magnesium alloys and enhance its plasticity. Zhong et al. found that alloying Y and Ce elements changes the lattice parameters and make is possible for activation of non-basal slip [7]. On the other hand, the initial microstructure of the
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Fig. 1 – The microstructure characteristic of the as-extruded AZ110LC alloy. (a) IPF (Inverse pole figure) map; (b) SEM micrograph; (c) grain misorientation angle distribution map; (d) average grain size distribution map; (e) pole figures of (0002), (10 1¯ 0) and (1 2¯ 10) planes.
alloy is a significantly decisive factor for the complex flow behavior at different deformation conditions. The microstructure evolution may relate to DRX (dynamic recrystallization) and DRV (dynamic recovery) and the deformation mechanism may relevant to GBS (grain boundary sliding), lattice/grain boundary diffusion, dislocation movement, dislocation creep, dislocation climb, mechanical twinning, and so on [7]. Therefore, systematic investigation on the deformation behavior and microstructure evolution are essential for the application of the Mg-Al-Zn-RE systemic alloy. At present, most Mg-Al-Zn-RE systemic alloys possessing with superplasticity in the recent researches had two characteristics: the weight percentage of Al element is not more than 10% and the content of RE elements is not less than 5%. In this work, the percentage of Al element by weight studied in this work is 11% alloying with 1.8 wt. % La-rich Mish Metal. The purpose of this work is to study the deformation behavior of the as-extruded AZ110LC alloy through high temperature tensile test (423 K ∼ 573 K) and strain rate (1.7 × 10-3 s-1 ∼ 1.7 × 10-4 s-1 ). In order to clarify the deformation mechanism, the flow stress behavior, microstructure and structural characteristics, recrystallization behavior and cavitation theory during different deformation conditions were studied. This work can provide guidance for the wide application of this promising magnesium alloy in industry.
2.
Experimental procedures
The original material studied in this work was as-extruded AZ110LC alloy. The chemical composition of the alloy was Mg-11.3Al-0.5Zn-0.15Mn-1.2La-0.6Ce (wt. %). The ingots (Ø60mm × length 180 mm) were melted in a high temperature resistance furnace at 973∼1010 K in the CO2 and SF6 protec-
tive gases atmosphere. And then, the ingots were machined into round bar with a diameter of 70 mm which were further homogenized in a muffle furnace at 673 K for 20 h. These billets were reheated at 573 K for 40minimums and extruded at the same temperature in a reverse extruder with an extrusion ratio of 17:1 and a ram speed of 1.5 mm/s. Tensile specimens (10 mm gauge length and 4 mm gauge diameter) were machined along the extrusion direction. Superplastic tensile tests were conducted at temperature ranging from 423 K to 573 K with an interval of 50 K, together with the strain rates of 1.7 × 10−3 s-1 , 8.4 × 10-4 s-1 , 4.2 × 10−4 s−1 and 1.7 × 10-4 s-1 . Furthermore, the tensile tests interrupted at different engineering strain (0%, 4%, 40% and 90%) were also performed at 523 K and 4.2 × 10-4 s-1 in order to investigate the superplastic deformation mechanism of the alloy. Special to note was that the sample of 0% strain was reheated at 523 K for 30 min in order to investigate the initial microstructure under high temperature tensile deformation. It took 5–10 minimums for the equipped heating furnace to reach the setting temperature and the other 30 minimums was given for the tensile specimens to reach and equilibrate the setting temperature. The microstructure characteristics of the extruded alloys and superplastic deformation samples were observed with optical microscope (OM) and scanning electron microscopy (SEM) equipped with an energy dispersive X-ray spectrometer (EDS). The texture analysis was carried by applying XRD test. The pole figure was measured contained (0002), (10 1¯ 0) and (1 2¯ 10) three crystallographic planes. Electron backscattered diffraction (EBSD) test was implemented on a JEOL 733 electron probe equipped with HKL-EBSD system. Transmission electron microscopy (TEM) was used to image microscopic features. The tensile mechanical property of the extrusion bars was tested at room temperature using a SANS CMT-5105 material tensile prototype.
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Fig. 2 – XRD patterns of the as-extruded AZ110LC alloys.
3.
Results
3.1.
Microstructure of the as-extruded AZ110LC alloy
The initial microstructure features of the as-extruded AZ110LC alloy were characterized by SEM and EBSD, as shown in Fig. 1. The Fig. 1a and d display a fine-grained microstructure with an average grain size of 2.9 m and a typical double-peak structure. Fig. 1b shows the morphology and distribution of
the second phases. Numerous lumpy and punctate precipitates (white contrasts in Fig. 1b) disperse uniformly paralleled to the extrusion direction. To further identify compositions of these precipitates, XRD and EDS analysis were conducted and presented in Figs. 2 and 3. As we can see in Fig. 2 that there are three main precipitates in the as-extruded AZ110LC alloy which are eutectic phase Mg17 Al12 and two kinds of Al-RE phases, Al11 (La/Ce)3 and Al2 (La/Ce). Corresponding elements scanning results (see Fig. 3a), Al, La and Ce elements both exist in the lumpy and punctate precipitates. Furthermore, the EDS results in Fig. 3b are the extensive evidence to identify the rare earth precipitates. The refined equiaxed grains are emerged by DRX undoubtedly. It is obvious that the grainrefining effect attributes to the extrusion deformation and the rare earth phases. In general, with the increasing extrusion temperature, the grain size may increase simultaneously due to the strong temperature sensitivity of magnesium alloy [2]. Hence, the lower extrusion temperature in this work can refrain from grain growth. These grains are almost filled with high dense high angle grain boundaries (HAGB, >15◦ ) which in favor of grain boundary slipping and grain rotation. Fig. 1e shows the (0002), (10 1¯ 0) and (1 2¯ 10) pole figures of the asextruded AZ110LC alloy. From Fig. 1e, the texture intensity distributes along the TD (transverse direction) and implies a typical hot-extrusion fiber texture in the ED-TD plane with relatively randomized basal plane orientations.
3.2.
The tensile properties of the alloy
Fig. 4a and b shows the true stress-true strain curves of the extruded AZ110LC alloy under different deformation tem-
Fig. 3 – (a) Elements scanning figures of the extruded sample of AZ110LC alloy in the ED-TD plane; (b) and (c) TEM micrographs of Al-RE phases of AZ110LC alloy; (d) corresponding EDS results of secondary phases.
Fig. 4 – Effects of (a) test temperature and (b) initial strain rate on the flow behavior of the AZ110LC alloy.
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Fig. 5 – Peak stresses of the AZ110LC alloy under different temperatures and strain rates.
Fig. 6 – Elongation to fracture of the AZ110LC alloy under different tensile temperatures and strain rates.
peratures and strain rates. Obviously, the flow curves are significantly affected by deformation conditions. As shown in Fig. 4, the flow stress decreases with increasing stretching temperatures and decreasing strain rates. The flow stress rapidly reaches peak stress and then decreases with further deformation until the fracture. After peak stress, the test alloy starts to reach the softening stage. The increasing deformation temperature and decreasing strain rate obviously slow the softening rate at the softening stage. When the stretching temperature and strain rate reach 523 K and 8.4 × 10−4 s-1 , the work hardening phenomenon can be observed in the later deformation stage and the reason may be closely related to grain growth. Additionally, when tested at high temperature and slow strain rate, there is a long stage of stable flow stress until fracture. The stable flow stress indicates that the DRX or DRV behaviors and work hardening may occur simultaneously at a balanced condition in the soft stage. Fig. 5 shows the peak stresses of the extruded AZ110LC alloy at different stretching temperatures and strain rates. The peak stress monotonically increases with the decreasing tensile temperature and increasing strain rate. The reason is that the dislocation motion and annihilation during the deformation process are accelerated and the corresponding dislocation multiplication and dislocation density are decreased under the high tensile temperature and high strain rate. Fig. 6 shows elongations of the AZ110LC alloy under different tensile temperatures and strain rates. The alloy exhibits obvious superplasticity within a wide temperature range from 473 K to 573 K. It can be seen that the elongation increases with the rise of tensile temperature except for the condition of exceeding 523 K at the strain rate of 8.4 × 10−4 s-1 . It can be seen in Fig. 6 that the maximum elongation of the alloy is 840% under 573 K and 1.7 × 10−4 s-1 . At the deformation temperature of 523 K, an excellent elongation (640%) also can be obtained at a relatively rapidly strain rate (8.4 × 10−4 s-1 ). Fig. 7 shows all the tensile samples tested at different testing temperatures and strain rates. It can be seen in Fig. 7(c) that the elongation of the alloy under 523 K and 4.2 × 10-4 s-1 is less than that under
523 K and 8.4 × 10-3 s-1 . From Fig. 4(b), the flow stress of the alloy under 523 K and 8.4 × 10-3 s-1 rises when the true strain exceeds 0.9 because of the work hardening behavior. However, the DRV effect is greater than the work hardening effect when deforming under 523 K and 4.2 × 10-4 s-1 . All of those illustrate that the alloy is equipped with good superplastic deformation capability at the given temperatures and strain rates. Base on the similar configuration of all the true stress and true strain curves under different testing temperatures and strain rates, the flow stress curves under 523 K and 4.2 × 10−4 s-1 were shown in Fig. 8. Citing a deformation condition of 523 K and 4.2 × 10−4 s-1 for instance, the flow curve can be divided into four tensile stages, i.e., (I) elastic deformation stage, (II) yielding deformation stage, (III) homogeneous deformation stage, (IV) localized necking stage [8,9]. At the first stage, the flow stress rises linearly and quickly with the increase of strain at the elastic deformation stage. As shown in Fig. 8, the yielding deformation stage is extremely transient and the dislocation accumulates rapidly in this plastic stage which brings about intensely work hardening. When the flow stress reaches the peak stress, it starts to decrease within the further straining. At the homogeneous stage, the accumulating dislocation energy can provide sufficient driving force for flow softening via dynamic softening behaviors such as DRX and DRV behaviors [10]. The steady state of the homogenous deformation stage is a consequence of dynamic equilibrium of work hardening and dynamic softening process [11]. At the last stage, the slope of the flow curve decreases rapidly and localized necking occurs resulting in drastic dropping of flow stress until to fracture. All above, the AZ110LC alloy is equipped with good superplastic deformation capability at the given testing temperatures and strain rates. Moreover, the flow stress curves are different from other magnesium alloys, especially in the homogenous deformation stage. For many other magnesium alloys, the flow stress increases rapidly to the yielding point and reaches to the peak stress slowly. However, the AZ110LC alloy studied in this work has a shorter yielding deformation
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Fig. 7 – Testing samples of different tensile temperatures and strain rates.
3.3. The microstructure evolution during superplastic tensile deformation
Fig. 8 – Four stages of the deformation process of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 . (the samples after different deformation stages marked S1, S2, S3 and S4.).
stage and has a large strain at the homogenous stage. It is obvious that the mechanical properties are largely related to the microstructure evolution and deformation behavior, which will be taken into consideration in the later chapters.
Fig. 9 depicts the IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages calculated from the EBSD tests. Fig. 9a shows the IPF map of the initial as-extruded AZ110LC alloy undergone 30 minimum heating. Compared with the initial as-extruded AZ110LC alloy, the fraction of sub-structures increases from 0.07 to 0.15 which results from the thermal effect. There is no doubt that the microstructure evolution is closely related to deformation degree and thermal effect during high temperature tensile tests. Fig. 10 shows grain boundaries distribution maps of the AZ110LC alloy under 523 K and 4.2 × 10-4s-1 and Fig. 11 displays variation in average grain size and average aspect ratio of the AZ110LC alloy at different deformation stages. At the elastic deformation stage (Fig. 2b), the average grain size decreases slightly, but the fraction of coarse grains (d>3 m) and fine grains increase. Besides, the fraction of sub-structures has a slight decrease. Therefore, due to deformation and thermal effect, some sub-structures evolve into HAGBs shaped with fine grains. With the further straining, the sub-structures in the AZ110LC alloy are well developed. The average grain shape aspect ratio decreases greatly and the average grain size decreases slightly. In this stage, the accumulated dislocations in the yielding deformation stage gradually release. This
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Fig. 9 – IPF maps showing grain morphology distribution of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 10 – Grain boundaries distribution maps of the AZ110LC alloy under 523 K and 4.2 × 10-4 s−1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
brings about that the distribution of grain size are much more homogenous and grain shape are more equiaxed, which may ascribe to the migration and rotation of grains and develop excellent ductility finally. However, at the end of the homogenous deformation stage, the average grain size increases slowly and the average grain shape aspect ratio increases greatly. This is a result of dynamic softening behavior due to DRV and DRX. Therefore, the microstructure characteristic in
the gauge section of the AZ110LC alloy is intimately related to the fluctuation of flow stress during tensile test at high temperature. The increasing strain can provide DRX and DRV, which could lead to the occurrence of softening behavior. Fig. 12 shows IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. It can be seen that the maximum texture intensity gradually decreases from 4.28 to 2.12 after different deformation strains. From Fig. 12a,
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Fig. 11 – Variation in average grain size and average aspect ratio of the AZ110LC alloy at different deformation stages.
the texture peaks are contracted around region A (arrows in Fig. 12a). This texture type was reported by Hu et al., which was a kind of rare earth texture [12]. At the first deformation stage, the texture peaks transform to revolve around (01 1¯ 0) (region B in Fig. 12b). With the further tensile deformation, the texture peaks are contracted around (01 1¯ 0) and (11 2¯ 0) (region C in Fig. 12c), which results in the weakening and randomized basal texture. At the end of the third stage, the texture intensity decreases to a minimum value and the texture peaks change to around region D (arrow in Fig. 12d), which depicts texture weakening and grain rotation during the homogeneous deformation stage. The evolution of texture during the tensile deformation is intimately related to grain rotation and GBS which can weaken the texture intensity. The texture analysis can contribute to investigating the deformation mechanism and it can be inferred that the GBS is the major deformation mechanism and grain rotation acts as an accommodation mechanism to GBS during superplastic deformation [13].
4.
Discussion
4.1.
Superplastic deformation behavior
Generally, the power-law relationship is used to describe the flow process of metals, and the strain rate sensitivity coefficient (m) reflects the capability of plastic deformation. The
Fig. 13 – The strain rate sensitivity coefficient m decimated by the curve of ln versus ln ε˙ .
value of m can be evaluated as the slope of ln and ln ε˙ , which is defined as [7,14]: m=
∂ ln |ε, T ∂ ln ε˙
(1)
where is the flow stress(MPa), ε˙ is the strain rate (s−1 ), T is the deformation temperature (K). As shown in Eq. 1, the value of m reflects the variation of flow stress and strain rate under different tensile temperatures and strains. The strain rate sensitivity embodies the capability of the transfer and diffusion of necking which resulted in stability of the homogenous stage [15]. Corresponding the value of peak stresses of different deformation conditions, the value of m can be calculated by the slope of ln and ln ε˙ , seen in Fig. 13. By good linear fitting, the value of m is approximately equal to 0.41. Furthermore, the elongation to fracture is 450% which indicates an excellent superplasticity under the deformation condition of 523 K and 4.2 × 10-4 s−1 . Meanwhile, the slight work hardening in the end of the homogenous stage can slow down the development in necking [16]. In order to investigate the superplastic deformation behaviors dominated by the thermal activity, the stress-strain behavior is described with the constitute equation:
Q
ε˙ = A n exp −
RT
(2)
where Q is the activation energy of flow process (J mol−1 ), A is material parameter, n is the stress index (1/m), T is the abso-
Fig. 12 – IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4. Please cite this article in press as: Liao Q, et al. Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.063
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fine grains with HAGBs and also promote the sub-structures transform to LAGBs grains [20]. There are two kinds of nucleation patterns of DRX, CDRX (continue dynamic recrystallization) and DDRX (discontinue dynamic recrystallization) respectively [21]. Generally, sub-structures transform to LAGBs by merging and reorganizing lattice dislocations and further change into HAGBs which are conducive to GBS [21]. Moreover, the recrystallization behavior is closely related to the dislocation merging and interlacing. Therefore, the investigation of dislocation density evolution is important to depict the recrystallization behavior during different superplastic deformation stages. To quantitatively study the dislocation density evolution, the KAM (kernel average misorientation) method was adopted to determine the local misorientation using EBSD data. The local misorientation at a point of 90 nm × 90 nm was calculated by its surrounding 24 points [22]: Fig. 14 – The activation energy Q decimated by the curve of ln versus 104 /T.
1 sur j − i n n
i =
j=1
lute temperature (K), R is the universal gas constant (8.314 J mol−1 K−1 ). The stress can be regarded as peak stress. In order to calculate the value of Q, the Eq. (2) can be transferred as: ln =
ln˙ε − lnA 1 Q + · n T R, n
(3) GND =
Thus, Eq. 3 can be given to the following form: Q =R·n·
∂ln ε˙ = const ∂1/T
(4)
The value of Q is measured from the slope in a plot of ln as a function of 1/T, as shown in Eq. 4 and Fig. 14. The active energy can be calculated to be ∼87 kJ/mol. This value is close to the expected energy for grain boundary diffusion in magnesium alloys (∼92 kJ/mol) [17]. Usually, the active energy can determine the deformation mechanism of conventional metals [18]. Therefore, in the present study, according to the above analysis, it can be concluded that GBS is the primary deformation mechanism of the AZ110LC alloy. In general, GBS occurs easily in magnesium alloys, but becomes blocked or inhibited from further sliding at regions such as triple points or grain boundary ledges. The occurrence of the RE phases could block the GBS at the triple junctions and boundaries due to the dislocation sliding is hard to get across the phase boundaries. Because of the impending effect of RE phases, the strength of the alloy could be enhanced during the hot deformation process [19].
4.2.
where i is the value of the local misorientation at point ‘i’ and jsur represents the misorientation angle at its neighboring point ‘j’. Strain gradient theory method is used to deduce the GND (geometrically necessary dislocations) density information [23].
Recrystallization mechanism
The microstructure evolution during the superplastic deformation process is closely related to DRX and DRV behaviors which could refine the grains and increase the deformation patency. Fig. 15 shows the distribution maps of three types of grains and structures, DRX grains (blue color regions), substructures (yellow color regions) and deformed grains (red color regions). DRX behavior can generate a large number of
2i = Ai ub
where GND is the mean density of GND and i can be regarded as the mean local misorientation angle (mean ); u is the unit length of the point which is 90 nm; b is the Burgers 2 vector. A= ub is a constant for magnesium alloy. The local misorientation angle distribution maps of the AZ110LC alloy at different deformation stages were shown in Figs. 16 and 17, and the mean local misorientation angles (mean ) are also calculated in Fig. 17. As shown in Figs. 16 and 17, the GND density decreases in stage I and increases rapidly in the further stages which are consistent with the evolution of sub-structures. Instable sub-structures absorb in-grain slipping and climbing dislocation during superplastic deformation and finally transform to HAGBs [24]. Plastic deformation is carried out by the grain slipping and climbing dislocation, and the migration, sliding and rotation of sub-grain boundaries play an important role in coordinating deformation [25]. The deformation behavior in the homogenous stage mainly contains DRV and DRX softening behaviors and work hardening behavior. The homogenous stage is a long process of dynamic equilibrium between strain hardening and strain softening. The rising of GND density mains increasing dislocation accumulation and accelerate the effect of work hardening. According to the Ball-Hutchison theory, dislocations accumulate on the adjacent grain boundaries by grain interior, resulting in stress concentration and when the stress reaches a certain value, the front dislocation of plugging will disappear along the grain boundary, and they will occur again [26]. That mains, the superplastic deformation of the AZ110LC alloy is strongly relevant to the density of dislocation evolution, specific performance: the rising of the density
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Fig. 15 – The DRX distribution maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 16 – Local misorientation maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
of dislocation in the homogenous stage is instrumental in enhancing the ductility in the superplastic deformation.
4.3.
Cavitation
The development of cavities is a significant characteristic of the failure of the superplastic alloys. The cavitation development reflects the failure behavior of tensile tests [27,28].
Fig. 18 shows the distribution of cavities in the tensile samples closed to tensile fracture at different strain rates under 473 K. As shown in Fig. 18, the cavities distribute dispersively in the gauge length and present two kinds of shapes: punctaform and strip-form (red arrows in Fig. 18). The shapes and the fraction of cavities have a great influence on the elongation of the tensile samples at different strain rates under a certain temperature. From Figs. 6 and 18, cavities with high round-
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Fig. 17 – Local misorientation angle distribution maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 18 – Cavities in AZ110LC alloy deformed at 523 K and constant strain rate of (a) 1.7 × 10−3 s-1 ; (b) 8.4 × 10-4 s-1 ; (c) 4.2 × 10-4 s-1 and (d) 1.7 × 10-4 s-1 .
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ness coefficients lead to relatively lower elongation. Moreover, the lower roundness coefficients tend to be recorded at a lower strain rate [29]. This also demonstrates that, when the strain rate is large, the diffusion of vacancies near the grain boundaries cannot perform fully, resulting in lower elongation. Nevertheless, the vacancies near the grain boundaries can diffuse fully under lower strain rate and the vacancies have enough time to form the cavities, which leads to large elongation percentage of the alloy. The formation and growth of cavities can effectively relax the stress concentration in triple junctions and boundaries between matrix and particles [30,31]. The stress concentration is easily developed near the grain boundary and particles which impedes the grain boundary slipping. As shown in Fig. 18, the cavities with high roundness coefficients form and grow near the rare earth particles. As the active energy of grain boundary slip is slightly higher than that of grain boundary diffusion and lattice diffusion during superplastic deformation of magnesium alloys, the grain boundary slip always produces stress concentration at the triple junctions or at the phase boundary between material reinforcement phase and matrix, which results in the in-grain slipping being hindered, resulting in vacancies and overlap [32]. Therefore, there are mechanisms for coordinating and compensating grain boundary sliding in superplastic deformation of magnesium alloys, such as the diffusion of atoms and cavities and the movement of dislocations.
Conflict of interest We the undersigned declare that this manuscript entitled “Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish Metal addition” is original, has not been published before and is not currently being considered for publication elsewhere. We confirm that the manuscript has been read and approved by all named authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us. We understand that the corresponding author is the sole contact for the Editorial process. He is responsible for communicating with the other authors about progress, submissions of revisions and final approval of proofs.
Acknowledgments This research was financially supported by the National Key Research and Development Program of China (Grant No. 2016YFB0301104) and National Natural Science Foundation of China grant number 51771043 and The authors are also grateful for supported by State Key Laboratory of Baiyunobo Rare Earth Resource Researches and Comprehensive Utilization.
references
5.
Conclusions
In summary, the superplastic behavior and microstructure evolution of the as-extruded AZ110LC alloy during different tensile temperatures and strain rates were investigated in this article. The main outcomes could be concluded as follow: (1) The AZ110LC alloy reaches the optimal elongation of 840% under 573 K and 1.7 × 10−4 s-1 , and exhibits excellent superplastic forming capability even at lower stretching temperatures and higher strain rates. (2) The flow stress decreases with increasing testing temperatures and decreasing strain rates. The peak stress is strongly dependent on the deformation temperatures and strain rates. The flow stress rapidly reaches peak stress and then decreases with further deformation until the fracture. (3) The excellent superplasticity of the as-extruded AZ110LC alloy ascribes to the refinement of microstructure and the distribution of sub-structures. (4) The strain rate sensitivity coefficient of the as-extruded AZ110LC alloy is 0.41 and the active energy is 87 kJ/mol under 523 K and 4.2 × 10−4 s-1 indicating that the deformation mechanism of the alloy is GBS accommodated by grain rotation and dislocation diffusion. (5) The shapes and fraction of cavities have a great influence on the elongation of the AZ110LC alloy. The formation and growth of cavities can effectively relax the stress concentration and the cavities could effectively coordinate and compensate the GBS in superplastic deformation of the AZ110LC alloy.
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Review Article
Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition Qiyu Liao a , Qichi Le a,∗ , Xingrui Chen a , Xiaoqiang Li a , YanChao Jiang a , Lifeng Ma b , Wenxin Hu c a
Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang 110819, China Taiyuan University Science and Technology, Heavy Machinery Engineering Research Center, Ministry Education, Taiyuan, Shanxi, 030024, China c Baotou Research Institute of Rare Earths, Baotou, 014000, China b
a r t i c l e
i n f o
a b s t r a c t
Article history:
Tensile test at high temperature and electron backscatter diffraction (EBSD) techniques were
Received 10 November 2019
combined to investigate the deformation behavior and microstructure evolution of the as-
Accepted 17 January 2020
extruded AZ110 magnesium alloy with La-rich Mish Metal addition (AZ110LC for short). The
Available online xxx
stretching temperature and strain rate have great effect on the flow stress of the alloy. The flow stress rapidly reaches the peak stress and then decreases with further deformation until
Keywords:
the fracture. The maximum elongation of the alloy is 840% under 573 K and 1.7 × 10−4 s−1 . The
AZ110 magnesium alloy
excellent superplasticity ascribes to the refinement of microstructure and the distribution of
La-rich Mish metal
sub-structures of the as-extruded AZ110LC alloy. Moreover, the strain rate sensitivity coeffi-
Superplastic deformation behavior
cient of the as-extruded AZ110LC alloy is 0.41 and the active energy is 87 kJ/mol under 523 K
Microstructure evolution
and 4.2 × 10−4 s-1 indicating that the deformation mechanism of the alloy is GBS. Finally, the
Cavities
shapes and fraction of cavities closed to tensile fracture have a great influence on the elongation of the tensile samples at different strain rates under 573 K. The formation and growth of the cavities can effectively relax the stress concentration and can effectively coordinate and compensate the GBS mechanism in superplastic deformation of the alloy. © 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).
∗
Corresponding author. E-mails: [email protected] (Q. Liao), [email protected] (Q. Le). https://doi.org/10.1016/j.jmrt.2020.01.063 2238-7854/© 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/). Please cite this article in press as: Liao Q, et al. Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.063
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Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. Qiyu Liao received his B.S. and M.S. degree in Material processing engineering from the Northeastern University, Shenyang, China. He is currently with analysis of Plastic deformation of magnesium alloy and Development and application of new magnesium alloys. Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of Melting and casting of magnesium alloys; Plastic deformation of magnesium alloy; Light metal solidification in physical field (electromagnetic field and ultrasonic field); Development and application of new magnesium alloys; Mg alloys corrosion and battery; oxidation and combustion of Mg alloys. Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of application of ultrasonics on preparation of Mg alloys and Mg alloys corrosion and battery.
Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of Mg alloy composite material.
Institutional: Key Lab of Electromagnetic Processing of Materials, Ministry of Education, Northeastern Univeristy, Shenyang, China. He is currently with analysis of numerical simulation of wrought magnesium alloys.
Institutional: Taiyuan University Science and Technology, Heavy Machinery Engineering Research Center, Ministry Education, Taiyuan, Shanxi, China. He is currently with analysis of optimization design theory and key technology of rolling equipment; theory and key technology of composite technology; nonferrous metal deep processing key technology.
Institutional: Baotou Research Institute of Rare Earths, Baotou, China
1.
Introduction
Wrought magnesium alloys are very attractive in various industrial products because of their low density, excellent mechanical properties, good damping, strong electromagnetic shielding ability and recyclability [1]. In spite of these advantages, the poor formability at room temperature largely limits the application of wrought magnesium alloys due to its hexagonal close packed (HCP) structure [2]. Therefore, utilizing superplastic deformation to enhance their formability is necessary in order to fabricate light structural parts with some complex geometry of thin-walled components. Generally, there are three basic requirements for the occurrence of superplasticity: fine microstructure of equiaxed grains with grain size typically less than 10 m; tensile temperature not exceeding 0.5 TM (TM is the liquidus temperature of metal materials); strain rate higher than 10−6 s-1 but less than 10-2 s-1 [3]. In order to acquire fine grains, several severe plastic deformation (SPD) technologies are employed to fulfill the requirement which are time-consuming and costly [4]. Hence, it is worth to consider alloying with appropriate elements and conventional deformation methods for fine-grain microstructure. Among various magnesium alloys, those alloys based on Mg-Al-Zn system are the most widely used wrought magnesium, due to their excellent mechanical properties and high productivity and good product qualities. However, the content of Al element affects the ductility of the Mg-Al-Zn alloys. Generally, the alloys containing high content of Al element such as AZ80, AZ91 have respectively low ductility at room temperature [5]. Their formability can be further improved by enhancing their ductility through incorporating superplastic behavior. Good formability can be obtained by adding rare earth elements to some Mg-Al-Zn alloys. When alloying with RE elements in Mg-Al-Zn alloys, some RE phases with good thermal stability precipitate in the matrix can effectively refine the initial microstructure. It was reported that the softening behavior of the low-melting-point phase Mg17 Al12 has contribution to the interphase boundary sliding at elevated temperatures for exhibiting superplasticity, but the high volume fraction of the coarse Mg17 Al12 phase leads to low room-temperature ductility [6]. Furthermore, some RE elements can change the texture type of magnesium alloys and enhance its plasticity. Zhong et al. found that alloying Y and Ce elements changes the lattice parameters and make is possible for activation of non-basal slip [7]. On the other hand, the initial microstructure of the
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Fig. 1 – The microstructure characteristic of the as-extruded AZ110LC alloy. (a) IPF (Inverse pole figure) map; (b) SEM micrograph; (c) grain misorientation angle distribution map; (d) average grain size distribution map; (e) pole figures of (0002), (10 1¯ 0) and (1 2¯ 10) planes.
alloy is a significantly decisive factor for the complex flow behavior at different deformation conditions. The microstructure evolution may relate to DRX (dynamic recrystallization) and DRV (dynamic recovery) and the deformation mechanism may relevant to GBS (grain boundary sliding), lattice/grain boundary diffusion, dislocation movement, dislocation creep, dislocation climb, mechanical twinning, and so on [7]. Therefore, systematic investigation on the deformation behavior and microstructure evolution are essential for the application of the Mg-Al-Zn-RE systemic alloy. At present, most Mg-Al-Zn-RE systemic alloys possessing with superplasticity in the recent researches had two characteristics: the weight percentage of Al element is not more than 10% and the content of RE elements is not less than 5%. In this work, the percentage of Al element by weight studied in this work is 11% alloying with 1.8 wt. % La-rich Mish Metal. The purpose of this work is to study the deformation behavior of the as-extruded AZ110LC alloy through high temperature tensile test (423 K ∼ 573 K) and strain rate (1.7 × 10-3 s-1 ∼ 1.7 × 10-4 s-1 ). In order to clarify the deformation mechanism, the flow stress behavior, microstructure and structural characteristics, recrystallization behavior and cavitation theory during different deformation conditions were studied. This work can provide guidance for the wide application of this promising magnesium alloy in industry.
2.
Experimental procedures
The original material studied in this work was as-extruded AZ110LC alloy. The chemical composition of the alloy was Mg-11.3Al-0.5Zn-0.15Mn-1.2La-0.6Ce (wt. %). The ingots (Ø60mm × length 180 mm) were melted in a high temperature resistance furnace at 973∼1010 K in the CO2 and SF6 protec-
tive gases atmosphere. And then, the ingots were machined into round bar with a diameter of 70 mm which were further homogenized in a muffle furnace at 673 K for 20 h. These billets were reheated at 573 K for 40minimums and extruded at the same temperature in a reverse extruder with an extrusion ratio of 17:1 and a ram speed of 1.5 mm/s. Tensile specimens (10 mm gauge length and 4 mm gauge diameter) were machined along the extrusion direction. Superplastic tensile tests were conducted at temperature ranging from 423 K to 573 K with an interval of 50 K, together with the strain rates of 1.7 × 10−3 s-1 , 8.4 × 10-4 s-1 , 4.2 × 10−4 s−1 and 1.7 × 10-4 s-1 . Furthermore, the tensile tests interrupted at different engineering strain (0%, 4%, 40% and 90%) were also performed at 523 K and 4.2 × 10-4 s-1 in order to investigate the superplastic deformation mechanism of the alloy. Special to note was that the sample of 0% strain was reheated at 523 K for 30 min in order to investigate the initial microstructure under high temperature tensile deformation. It took 5–10 minimums for the equipped heating furnace to reach the setting temperature and the other 30 minimums was given for the tensile specimens to reach and equilibrate the setting temperature. The microstructure characteristics of the extruded alloys and superplastic deformation samples were observed with optical microscope (OM) and scanning electron microscopy (SEM) equipped with an energy dispersive X-ray spectrometer (EDS). The texture analysis was carried by applying XRD test. The pole figure was measured contained (0002), (10 1¯ 0) and (1 2¯ 10) three crystallographic planes. Electron backscattered diffraction (EBSD) test was implemented on a JEOL 733 electron probe equipped with HKL-EBSD system. Transmission electron microscopy (TEM) was used to image microscopic features. The tensile mechanical property of the extrusion bars was tested at room temperature using a SANS CMT-5105 material tensile prototype.
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Fig. 2 – XRD patterns of the as-extruded AZ110LC alloys.
3.
Results
3.1.
Microstructure of the as-extruded AZ110LC alloy
The initial microstructure features of the as-extruded AZ110LC alloy were characterized by SEM and EBSD, as shown in Fig. 1. The Fig. 1a and d display a fine-grained microstructure with an average grain size of 2.9 m and a typical double-peak structure. Fig. 1b shows the morphology and distribution of
the second phases. Numerous lumpy and punctate precipitates (white contrasts in Fig. 1b) disperse uniformly paralleled to the extrusion direction. To further identify compositions of these precipitates, XRD and EDS analysis were conducted and presented in Figs. 2 and 3. As we can see in Fig. 2 that there are three main precipitates in the as-extruded AZ110LC alloy which are eutectic phase Mg17 Al12 and two kinds of Al-RE phases, Al11 (La/Ce)3 and Al2 (La/Ce). Corresponding elements scanning results (see Fig. 3a), Al, La and Ce elements both exist in the lumpy and punctate precipitates. Furthermore, the EDS results in Fig. 3b are the extensive evidence to identify the rare earth precipitates. The refined equiaxed grains are emerged by DRX undoubtedly. It is obvious that the grainrefining effect attributes to the extrusion deformation and the rare earth phases. In general, with the increasing extrusion temperature, the grain size may increase simultaneously due to the strong temperature sensitivity of magnesium alloy [2]. Hence, the lower extrusion temperature in this work can refrain from grain growth. These grains are almost filled with high dense high angle grain boundaries (HAGB, >15◦ ) which in favor of grain boundary slipping and grain rotation. Fig. 1e shows the (0002), (10 1¯ 0) and (1 2¯ 10) pole figures of the asextruded AZ110LC alloy. From Fig. 1e, the texture intensity distributes along the TD (transverse direction) and implies a typical hot-extrusion fiber texture in the ED-TD plane with relatively randomized basal plane orientations.
3.2.
The tensile properties of the alloy
Fig. 4a and b shows the true stress-true strain curves of the extruded AZ110LC alloy under different deformation tem-
Fig. 3 – (a) Elements scanning figures of the extruded sample of AZ110LC alloy in the ED-TD plane; (b) and (c) TEM micrographs of Al-RE phases of AZ110LC alloy; (d) corresponding EDS results of secondary phases.
Fig. 4 – Effects of (a) test temperature and (b) initial strain rate on the flow behavior of the AZ110LC alloy.
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Fig. 5 – Peak stresses of the AZ110LC alloy under different temperatures and strain rates.
Fig. 6 – Elongation to fracture of the AZ110LC alloy under different tensile temperatures and strain rates.
peratures and strain rates. Obviously, the flow curves are significantly affected by deformation conditions. As shown in Fig. 4, the flow stress decreases with increasing stretching temperatures and decreasing strain rates. The flow stress rapidly reaches peak stress and then decreases with further deformation until the fracture. After peak stress, the test alloy starts to reach the softening stage. The increasing deformation temperature and decreasing strain rate obviously slow the softening rate at the softening stage. When the stretching temperature and strain rate reach 523 K and 8.4 × 10−4 s-1 , the work hardening phenomenon can be observed in the later deformation stage and the reason may be closely related to grain growth. Additionally, when tested at high temperature and slow strain rate, there is a long stage of stable flow stress until fracture. The stable flow stress indicates that the DRX or DRV behaviors and work hardening may occur simultaneously at a balanced condition in the soft stage. Fig. 5 shows the peak stresses of the extruded AZ110LC alloy at different stretching temperatures and strain rates. The peak stress monotonically increases with the decreasing tensile temperature and increasing strain rate. The reason is that the dislocation motion and annihilation during the deformation process are accelerated and the corresponding dislocation multiplication and dislocation density are decreased under the high tensile temperature and high strain rate. Fig. 6 shows elongations of the AZ110LC alloy under different tensile temperatures and strain rates. The alloy exhibits obvious superplasticity within a wide temperature range from 473 K to 573 K. It can be seen that the elongation increases with the rise of tensile temperature except for the condition of exceeding 523 K at the strain rate of 8.4 × 10−4 s-1 . It can be seen in Fig. 6 that the maximum elongation of the alloy is 840% under 573 K and 1.7 × 10−4 s-1 . At the deformation temperature of 523 K, an excellent elongation (640%) also can be obtained at a relatively rapidly strain rate (8.4 × 10−4 s-1 ). Fig. 7 shows all the tensile samples tested at different testing temperatures and strain rates. It can be seen in Fig. 7(c) that the elongation of the alloy under 523 K and 4.2 × 10-4 s-1 is less than that under
523 K and 8.4 × 10-3 s-1 . From Fig. 4(b), the flow stress of the alloy under 523 K and 8.4 × 10-3 s-1 rises when the true strain exceeds 0.9 because of the work hardening behavior. However, the DRV effect is greater than the work hardening effect when deforming under 523 K and 4.2 × 10-4 s-1 . All of those illustrate that the alloy is equipped with good superplastic deformation capability at the given temperatures and strain rates. Base on the similar configuration of all the true stress and true strain curves under different testing temperatures and strain rates, the flow stress curves under 523 K and 4.2 × 10−4 s-1 were shown in Fig. 8. Citing a deformation condition of 523 K and 4.2 × 10−4 s-1 for instance, the flow curve can be divided into four tensile stages, i.e., (I) elastic deformation stage, (II) yielding deformation stage, (III) homogeneous deformation stage, (IV) localized necking stage [8,9]. At the first stage, the flow stress rises linearly and quickly with the increase of strain at the elastic deformation stage. As shown in Fig. 8, the yielding deformation stage is extremely transient and the dislocation accumulates rapidly in this plastic stage which brings about intensely work hardening. When the flow stress reaches the peak stress, it starts to decrease within the further straining. At the homogeneous stage, the accumulating dislocation energy can provide sufficient driving force for flow softening via dynamic softening behaviors such as DRX and DRV behaviors [10]. The steady state of the homogenous deformation stage is a consequence of dynamic equilibrium of work hardening and dynamic softening process [11]. At the last stage, the slope of the flow curve decreases rapidly and localized necking occurs resulting in drastic dropping of flow stress until to fracture. All above, the AZ110LC alloy is equipped with good superplastic deformation capability at the given testing temperatures and strain rates. Moreover, the flow stress curves are different from other magnesium alloys, especially in the homogenous deformation stage. For many other magnesium alloys, the flow stress increases rapidly to the yielding point and reaches to the peak stress slowly. However, the AZ110LC alloy studied in this work has a shorter yielding deformation
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Fig. 7 – Testing samples of different tensile temperatures and strain rates.
3.3. The microstructure evolution during superplastic tensile deformation
Fig. 8 – Four stages of the deformation process of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 . (the samples after different deformation stages marked S1, S2, S3 and S4.).
stage and has a large strain at the homogenous stage. It is obvious that the mechanical properties are largely related to the microstructure evolution and deformation behavior, which will be taken into consideration in the later chapters.
Fig. 9 depicts the IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages calculated from the EBSD tests. Fig. 9a shows the IPF map of the initial as-extruded AZ110LC alloy undergone 30 minimum heating. Compared with the initial as-extruded AZ110LC alloy, the fraction of sub-structures increases from 0.07 to 0.15 which results from the thermal effect. There is no doubt that the microstructure evolution is closely related to deformation degree and thermal effect during high temperature tensile tests. Fig. 10 shows grain boundaries distribution maps of the AZ110LC alloy under 523 K and 4.2 × 10-4s-1 and Fig. 11 displays variation in average grain size and average aspect ratio of the AZ110LC alloy at different deformation stages. At the elastic deformation stage (Fig. 2b), the average grain size decreases slightly, but the fraction of coarse grains (d>3 m) and fine grains increase. Besides, the fraction of sub-structures has a slight decrease. Therefore, due to deformation and thermal effect, some sub-structures evolve into HAGBs shaped with fine grains. With the further straining, the sub-structures in the AZ110LC alloy are well developed. The average grain shape aspect ratio decreases greatly and the average grain size decreases slightly. In this stage, the accumulated dislocations in the yielding deformation stage gradually release. This
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Fig. 9 – IPF maps showing grain morphology distribution of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 10 – Grain boundaries distribution maps of the AZ110LC alloy under 523 K and 4.2 × 10-4 s−1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
brings about that the distribution of grain size are much more homogenous and grain shape are more equiaxed, which may ascribe to the migration and rotation of grains and develop excellent ductility finally. However, at the end of the homogenous deformation stage, the average grain size increases slowly and the average grain shape aspect ratio increases greatly. This is a result of dynamic softening behavior due to DRV and DRX. Therefore, the microstructure characteristic in
the gauge section of the AZ110LC alloy is intimately related to the fluctuation of flow stress during tensile test at high temperature. The increasing strain can provide DRX and DRV, which could lead to the occurrence of softening behavior. Fig. 12 shows IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. It can be seen that the maximum texture intensity gradually decreases from 4.28 to 2.12 after different deformation strains. From Fig. 12a,
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Fig. 11 – Variation in average grain size and average aspect ratio of the AZ110LC alloy at different deformation stages.
the texture peaks are contracted around region A (arrows in Fig. 12a). This texture type was reported by Hu et al., which was a kind of rare earth texture [12]. At the first deformation stage, the texture peaks transform to revolve around (01 1¯ 0) (region B in Fig. 12b). With the further tensile deformation, the texture peaks are contracted around (01 1¯ 0) and (11 2¯ 0) (region C in Fig. 12c), which results in the weakening and randomized basal texture. At the end of the third stage, the texture intensity decreases to a minimum value and the texture peaks change to around region D (arrow in Fig. 12d), which depicts texture weakening and grain rotation during the homogeneous deformation stage. The evolution of texture during the tensile deformation is intimately related to grain rotation and GBS which can weaken the texture intensity. The texture analysis can contribute to investigating the deformation mechanism and it can be inferred that the GBS is the major deformation mechanism and grain rotation acts as an accommodation mechanism to GBS during superplastic deformation [13].
4.
Discussion
4.1.
Superplastic deformation behavior
Generally, the power-law relationship is used to describe the flow process of metals, and the strain rate sensitivity coefficient (m) reflects the capability of plastic deformation. The
Fig. 13 – The strain rate sensitivity coefficient m decimated by the curve of ln versus ln ε˙ .
value of m can be evaluated as the slope of ln and ln ε˙ , which is defined as [7,14]: m=
∂ ln |ε, T ∂ ln ε˙
(1)
where is the flow stress(MPa), ε˙ is the strain rate (s−1 ), T is the deformation temperature (K). As shown in Eq. 1, the value of m reflects the variation of flow stress and strain rate under different tensile temperatures and strains. The strain rate sensitivity embodies the capability of the transfer and diffusion of necking which resulted in stability of the homogenous stage [15]. Corresponding the value of peak stresses of different deformation conditions, the value of m can be calculated by the slope of ln and ln ε˙ , seen in Fig. 13. By good linear fitting, the value of m is approximately equal to 0.41. Furthermore, the elongation to fracture is 450% which indicates an excellent superplasticity under the deformation condition of 523 K and 4.2 × 10-4 s−1 . Meanwhile, the slight work hardening in the end of the homogenous stage can slow down the development in necking [16]. In order to investigate the superplastic deformation behaviors dominated by the thermal activity, the stress-strain behavior is described with the constitute equation:
Q
ε˙ = A n exp −
RT
(2)
where Q is the activation energy of flow process (J mol−1 ), A is material parameter, n is the stress index (1/m), T is the abso-
Fig. 12 – IPF maps of the AZ110LC alloy under 523 K and 4.2 × 10−4 s-1 at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4. Please cite this article in press as: Liao Q, et al. Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish metal addition. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.063
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fine grains with HAGBs and also promote the sub-structures transform to LAGBs grains [20]. There are two kinds of nucleation patterns of DRX, CDRX (continue dynamic recrystallization) and DDRX (discontinue dynamic recrystallization) respectively [21]. Generally, sub-structures transform to LAGBs by merging and reorganizing lattice dislocations and further change into HAGBs which are conducive to GBS [21]. Moreover, the recrystallization behavior is closely related to the dislocation merging and interlacing. Therefore, the investigation of dislocation density evolution is important to depict the recrystallization behavior during different superplastic deformation stages. To quantitatively study the dislocation density evolution, the KAM (kernel average misorientation) method was adopted to determine the local misorientation using EBSD data. The local misorientation at a point of 90 nm × 90 nm was calculated by its surrounding 24 points [22]: Fig. 14 – The activation energy Q decimated by the curve of ln versus 104 /T.
1 sur j − i n n
i =
j=1
lute temperature (K), R is the universal gas constant (8.314 J mol−1 K−1 ). The stress can be regarded as peak stress. In order to calculate the value of Q, the Eq. (2) can be transferred as: ln =
ln˙ε − lnA 1 Q + · n T R, n
(3) GND =
Thus, Eq. 3 can be given to the following form: Q =R·n·
∂ln ε˙ = const ∂1/T
(4)
The value of Q is measured from the slope in a plot of ln as a function of 1/T, as shown in Eq. 4 and Fig. 14. The active energy can be calculated to be ∼87 kJ/mol. This value is close to the expected energy for grain boundary diffusion in magnesium alloys (∼92 kJ/mol) [17]. Usually, the active energy can determine the deformation mechanism of conventional metals [18]. Therefore, in the present study, according to the above analysis, it can be concluded that GBS is the primary deformation mechanism of the AZ110LC alloy. In general, GBS occurs easily in magnesium alloys, but becomes blocked or inhibited from further sliding at regions such as triple points or grain boundary ledges. The occurrence of the RE phases could block the GBS at the triple junctions and boundaries due to the dislocation sliding is hard to get across the phase boundaries. Because of the impending effect of RE phases, the strength of the alloy could be enhanced during the hot deformation process [19].
4.2.
where i is the value of the local misorientation at point ‘i’ and jsur represents the misorientation angle at its neighboring point ‘j’. Strain gradient theory method is used to deduce the GND (geometrically necessary dislocations) density information [23].
Recrystallization mechanism
The microstructure evolution during the superplastic deformation process is closely related to DRX and DRV behaviors which could refine the grains and increase the deformation patency. Fig. 15 shows the distribution maps of three types of grains and structures, DRX grains (blue color regions), substructures (yellow color regions) and deformed grains (red color regions). DRX behavior can generate a large number of
2i = Ai ub
where GND is the mean density of GND and i can be regarded as the mean local misorientation angle (mean ); u is the unit length of the point which is 90 nm; b is the Burgers 2 vector. A= ub is a constant for magnesium alloy. The local misorientation angle distribution maps of the AZ110LC alloy at different deformation stages were shown in Figs. 16 and 17, and the mean local misorientation angles (mean ) are also calculated in Fig. 17. As shown in Figs. 16 and 17, the GND density decreases in stage I and increases rapidly in the further stages which are consistent with the evolution of sub-structures. Instable sub-structures absorb in-grain slipping and climbing dislocation during superplastic deformation and finally transform to HAGBs [24]. Plastic deformation is carried out by the grain slipping and climbing dislocation, and the migration, sliding and rotation of sub-grain boundaries play an important role in coordinating deformation [25]. The deformation behavior in the homogenous stage mainly contains DRV and DRX softening behaviors and work hardening behavior. The homogenous stage is a long process of dynamic equilibrium between strain hardening and strain softening. The rising of GND density mains increasing dislocation accumulation and accelerate the effect of work hardening. According to the Ball-Hutchison theory, dislocations accumulate on the adjacent grain boundaries by grain interior, resulting in stress concentration and when the stress reaches a certain value, the front dislocation of plugging will disappear along the grain boundary, and they will occur again [26]. That mains, the superplastic deformation of the AZ110LC alloy is strongly relevant to the density of dislocation evolution, specific performance: the rising of the density
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Fig. 15 – The DRX distribution maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 16 – Local misorientation maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
of dislocation in the homogenous stage is instrumental in enhancing the ductility in the superplastic deformation.
4.3.
Cavitation
The development of cavities is a significant characteristic of the failure of the superplastic alloys. The cavitation development reflects the failure behavior of tensile tests [27,28].
Fig. 18 shows the distribution of cavities in the tensile samples closed to tensile fracture at different strain rates under 473 K. As shown in Fig. 18, the cavities distribute dispersively in the gauge length and present two kinds of shapes: punctaform and strip-form (red arrows in Fig. 18). The shapes and the fraction of cavities have a great influence on the elongation of the tensile samples at different strain rates under a certain temperature. From Figs. 6 and 18, cavities with high round-
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Fig. 17 – Local misorientation angle distribution maps of the AZ110LC alloy at different deformation stages. (a), (b), (c) and (d) respectively for the gauge of S1, S2, S3 and S4.
Fig. 18 – Cavities in AZ110LC alloy deformed at 523 K and constant strain rate of (a) 1.7 × 10−3 s-1 ; (b) 8.4 × 10-4 s-1 ; (c) 4.2 × 10-4 s-1 and (d) 1.7 × 10-4 s-1 .
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ness coefficients lead to relatively lower elongation. Moreover, the lower roundness coefficients tend to be recorded at a lower strain rate [29]. This also demonstrates that, when the strain rate is large, the diffusion of vacancies near the grain boundaries cannot perform fully, resulting in lower elongation. Nevertheless, the vacancies near the grain boundaries can diffuse fully under lower strain rate and the vacancies have enough time to form the cavities, which leads to large elongation percentage of the alloy. The formation and growth of cavities can effectively relax the stress concentration in triple junctions and boundaries between matrix and particles [30,31]. The stress concentration is easily developed near the grain boundary and particles which impedes the grain boundary slipping. As shown in Fig. 18, the cavities with high roundness coefficients form and grow near the rare earth particles. As the active energy of grain boundary slip is slightly higher than that of grain boundary diffusion and lattice diffusion during superplastic deformation of magnesium alloys, the grain boundary slip always produces stress concentration at the triple junctions or at the phase boundary between material reinforcement phase and matrix, which results in the in-grain slipping being hindered, resulting in vacancies and overlap [32]. Therefore, there are mechanisms for coordinating and compensating grain boundary sliding in superplastic deformation of magnesium alloys, such as the diffusion of atoms and cavities and the movement of dislocations.
Conflict of interest We the undersigned declare that this manuscript entitled “Superplastic deformation behavior of the as-extruded AZ110 magnesium alloy with La-rich Mish Metal addition” is original, has not been published before and is not currently being considered for publication elsewhere. We confirm that the manuscript has been read and approved by all named authors and that there are no other persons who satisfied the criteria for authorship but are not listed. We further confirm that the order of authors listed in the manuscript has been approved by all of us. We understand that the corresponding author is the sole contact for the Editorial process. He is responsible for communicating with the other authors about progress, submissions of revisions and final approval of proofs.
Acknowledgments This research was financially supported by the National Key Research and Development Program of China (Grant No. 2016YFB0301104) and National Natural Science Foundation of China grant number 51771043 and The authors are also grateful for supported by State Key Laboratory of Baiyunobo Rare Earth Resource Researches and Comprehensive Utilization.
references
5.
Conclusions
In summary, the superplastic behavior and microstructure evolution of the as-extruded AZ110LC alloy during different tensile temperatures and strain rates were investigated in this article. The main outcomes could be concluded as follow: (1) The AZ110LC alloy reaches the optimal elongation of 840% under 573 K and 1.7 × 10−4 s-1 , and exhibits excellent superplastic forming capability even at lower stretching temperatures and higher strain rates. (2) The flow stress decreases with increasing testing temperatures and decreasing strain rates. The peak stress is strongly dependent on the deformation temperatures and strain rates. The flow stress rapidly reaches peak stress and then decreases with further deformation until the fracture. (3) The excellent superplasticity of the as-extruded AZ110LC alloy ascribes to the refinement of microstructure and the distribution of sub-structures. (4) The strain rate sensitivity coefficient of the as-extruded AZ110LC alloy is 0.41 and the active energy is 87 kJ/mol under 523 K and 4.2 × 10−4 s-1 indicating that the deformation mechanism of the alloy is GBS accommodated by grain rotation and dislocation diffusion. (5) The shapes and fraction of cavities have a great influence on the elongation of the AZ110LC alloy. The formation and growth of cavities can effectively relax the stress concentration and the cavities could effectively coordinate and compensate the GBS in superplastic deformation of the AZ110LC alloy.
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