Microstructures and tensile properties of a high-strength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy

    Microstructures and tensile properties of a Mg–4Al–2RE–2Ca–0.3Mn alloy

high-strength die-cast

Qiang Yang, Kai Guan, Fanqiang Bu, Yaqin Zhang, Xin Qiu, Tian Zheng, Xiaojuan Liu, Jian Meng PII: DOI: Reference:

S1044-5803(16)30025-0 doi: 10.1016/j.matchar.2016.01.024 MTL 8169

To appear in:

Materials Characterization

Received date: Revised date: Accepted date:

18 November 2015 21 January 2016 27 January 2016

Please cite this article as: Yang Qiang, Guan Kai, Bu Fanqiang, Zhang Yaqin, Qiu Xin, Zheng Tian, Liu Xiaojuan, Meng Jian, Microstructures and tensile properties of a highstrength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy, Materials Characterization (2016), doi: 10.1016/j.matchar.2016.01.024

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ACCEPTED MANUSCRIPT Microstructures and tensile properties of a high-strength die-cast Mg4Al2RE2Ca0.3Mn alloy

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Xiaojuan Liua,*, Jian Menga,*

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Qiang Yanga, Kai Guana,d, Fanqiang Bua,d, Yaqin Zhangb, Xin Qiu a,c, Tian Zhenga,

State Key Laboratory of Rare Earth Resource Utilization, Changchun Institute of Applied

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Chemistry, Chinese Academy of Sciences, Changchun 130022, P. R. China Kashui International Holdings Limited, Shenzhen, Guangdong 518111, P. R. China

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Yangzhou hongfu Aluminium Co. Ltd, Yangzhou 100049, P. R. China

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University of Chinese Academy of Sciences, Beijing 100049, P. R. China

Abstract

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Microstructures and tensile properties of a high-pressure die-cast Mg4Al2RE2Ca

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based alloy were thoroughly investigated. The results indicate that the combinative addition of Ca and RE has a good cell refinement effect and can clearly change the intermetallic constituents. In the studied alloy, the intermetallic phases mainly consist of (Mg,Al)2Ca, Al11RE3, Al8RE3, Al2.12RE0.88, and Al10RE2Mn7, wherein the Al11RE3 phase has fully disparate morphologies from the conventional one in the MgAlRE based alloys, the Al8RE3 phase is an unknown phase, and the Al2.12RE0.88 phase contains twins and numerous staking faults. However, the Al2RE phase is absent in the studied alloy. On the other hand, substituting half RE with Ca can also obviously further improve the tensile strength at both room temperature and high temperatures,

* Corresponding author. Tel.: +86-431-85262030; Fax: +86-431-85698041 E-mail address: [email protected] (Xiaojuan Liu); [email protected] (Jian Meng).

ACCEPTED MANUSCRIPT but will also result in low plasticity simultaneously at lower temperatures ( 200 C), to which the key is the highly interconnected and strong intermetallic skeleton.

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Keywords: magnesium alloy; rare earth; calcium; intermetallic phase; tensile property;

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intermetallic skeleton 1. Introduction

Because of its outstanding high temperature properties and creep resistance, the

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high-pressure die-cast (HPDC) AE44 (Mg4Al4RE, wt.%, where RE represents rare

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earth) alloy is growingly attractive for automotive industries for improving fuel efficiency [14]. Nonetheless, its inadequate strength and relatively high cost limit its

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application to power train components. Therefore, many investigations were

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conducted to further improve the strength of AE44 alloy by increasing Al and/or RE,

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or adding some other alloying elements or inoculations [46]. However, these methods increased the cost or the processing (or casting) temperature, or resulted in

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unsatisfactory strength increments. Therefore, to facilitate the application of magnesium alloys in automotive powertrain components, it is significant to further drastically improve the strength of AE44 alloy through some methods with lower cost and processing temperature, such as substituting part of RE by some cheap alloying elements. It is reported that addition of Ca to magnesium can enhance the high temperature strength and creep resistance of MgAl-based alloys [79]. Furthermore, Ca has a low density (1.54 g/cm3 [10]), and low cost, and also is one of the promising elements that can improve ignition resistance of magnesium alloys [8]. Therefore, many efforts

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ACCEPTED MANUSCRIPT were devoted to investigating the microstructures and the properties of the MgAlCa based alloys [817], or using Ca to further improve the high temperature properties of

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the classical magnesium alloy systems such as AZ91 (Mg9Al1Zn, wt.%) [1820],

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ZA62 (Mg6Zn2Al, wt.%) [21], and AZ31 (Mg3Al1Zn, wt.%) [22]. Recently, some investigators have also paid attention to the modification of Ca on the properties of the MgAlRE based alloys. For example, Bakke et al. [4] reported that Ca

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addition noticeably enhanced the yield strength and the hardness of the conventional

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HPDC AE44 alloy, although drastically reduced its ultimate strength, elongation, and impact energy absorption simultaneously. Subsequently, Zhang et al. [23] reported

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that Ca addition clearly improved the creep resistance of the metal mold-casted AE41

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(Mg4Al1RE, wt.%) at 150 C and 100 MPa. However, the intermetallic phases in

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the corresponding Ca-modified MgAlRE based alloys were not investigated in detail. Additionally, few investigations dealing with further improving the strength of

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the HPDC AE44 alloy by substituting part of rare earth by cheap alloying element(s) were conducted. Thus, it is very necessary that thoroughly investigating the microstructures and the tensile properties of the HPDC AE44 based alloy with part RE substituted by Ca, to provide new insight into alloy design principles for further development of the HPDC MgAlRE based alloys. In this work, the HPDC Mg4Al2RE2Ca0.3Mn (AEX422) alloy was prepared, where RE represents misch metals of La and Ce. Then, its microstructures and tensile properties at various temperatures (20250 C) were thoroughly investigated. Finally, the underlying mechanisms were analyzed and discussed.

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ACCEPTED MANUSCRIPT 2. Experimental material and procedures The AEX422 alloy was prepared from high-purity magnesium and aluminum, and

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Mg20 wt.% LaCe, Mg28 wt.% Ca, and Mg2 wt.% Mn master alloys. All the

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metals preheated to approximately 300 C were melted in an electric resistance furnace and protected by CO2 + 1 vol% SF6 mixed gas during the whole melting process. After being heated to approximately 745 C, the melt was fully stirred for 8

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min, followed by keeping static for 40 minutes and cooling down to 700 5 C.

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Finally, tensile test bars with size of 60 mm in length and 6 mm in diameter were casted using a 280 ton clamping force cold chamber die-cast machine and a die

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equipped with an oil heating/cooling system. The temperature of the oil heater was set

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to 280 ºC previously. The chemical compositions of the obtained alloy were

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determined by inductivity coupled plasma atomic emission spectroscopy (ICP-AES) apparatus, and were Mg4.17Al1.79RE1.98Ca0.24Mn in wt.%, where La and Ce

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are approximately 0.81 wt.% and 0.98 wt.%, respectively. The microstructures of the studied alloy were characterized using optical microscopy (OM), X-ray diffraction (XRD) with Cu Kradiation (= 1.5418 Å), backscatter scanning electron microscopy (SEM), and transmission electron microscopy (TEM) equipped with an energy dispersive spectrometer (EDS). Since the microstructures near the core region are significantly distinct from those close to the skin [24], all micrographs presented in this work respond to ones of the middle portions on the cross-sections of tensile bars. The samples for OM and SEM observations were grounded using different grades of polishing papers and polished

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ACCEPTED MANUSCRIPT with Al2O3 paste. The TEM foils in the form of 3 mm diameter discs were mechanically grounded to approximately 30 m and then ion-milled by a precision

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ion polishing system (PIPS Gatan) equipped with cooling system by liquid nitrogen.

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Some tensile samples were heat treated at 400 C for 16 hours. The tensile tests were performed at room temperature (20 ºC), 75 ºC, 150 ºC, 200 ºC, and 250 ºC, with a strain rate of 1.0 × 10-3 s-1. The values presented in this work correspond to the

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average of at least three effective measurements for every condition to confirm

before each high temperature test.

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3.1 Microstructures

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3. Results and discussion

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reproducibility, and a 30 min holding was applied to balance the testing temperature

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Fig. 1a shows the OM of the HPDC AEX422 alloy, wherein the alloy consists of continuous or semi-continuous reticular cell boundaries and α-Mg cells. The average

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cell size measured using linear intercept method is approximately 7 m, which is much smaller than those of the conventional HPDC AE44 alloy (approximately 14 m) [25, 26] and AX53 alloy (approximately 10 m) [27]. Therefore, the combinative addition of RE and Ca has a better cell refinement effect on microstructures of the HPDC Mg4Al-based alloys than the respective RE or Ca addition. Fig. 1b represents the backscatter SEM image of the studied alloy. It can be seen that most intermetallic particles locate on cell boundaries while only a few of small blocky particles in α-Mg cells. In addition, the intermetallic phases can be simply classified into two groups: the dark phase and the bright phase. The dark phase mainly presents a lamellar

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ACCEPTED MANUSCRIPT structure consisting of single layer between two cells or alternative layers of magnesium and an intermetallic phase, which is very similar to that in the HPDC

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MgAlCa based alloys [27]. The bright phase has no coincident morphology and can

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be classified into five types: the lath-like phase, the rod-like phase, the small blocky phase, the lumpy phase with large volume, and the irregular-shaped phase, which are significantly different from the AlRE intermetallic phases in the conventional HPDC

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MgAlRE based alloys [2830], and were investigated in detail by TEM and EDS in

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this work.

Fig. 2 shows the XRD pattern of the studied alloy. The result indicates that there are

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mainly three intermetallic phases, namely Al11RE3, (Mg,Al)2Ca, and Al2.12RE0.88. In

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addition, there are also some peaks labeled as “unknown phase” since they do not

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match any known related phases in the MgAlRE/Ca based alloys to our best knowledge. Fig. 3ac show the bright-field TEM images of the dark phase with

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alternating layers, single layer, and triple junction, respectively, in the AEX422 alloy. Based on the corresponding selected area electron diffraction (SAED) patterns (Figs. 3d-e) and the EDS analysis results, the dark phase was confirmed as (Mg,Al)2Ca with a dihexagonal C36 crystal structure [11]. The experimental lattice parameters of a and c are approximately 0.594 nm and 1. 971 nm, respectively, which are well consistent with the reported lattice parameters of a = 0.596 nm, and c = 1.979 nm for the C36 phase in the as-cast MgAlCa ternary system [12, 16]. Furthermore, the EDS analysis results also illustrate that there is little RE and no manganese in the C36 phase. Luo et al. [27] reported that the C36 phase and the eutectic magnesium have a

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ACCEPTED MANUSCRIPT certain crystallographic orientation relationship as: (0001)Mg//(0001)C36 and [ 1 2 1 0]Mg//[11 2 0]C36. However, there was no regularity within the various regions

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examined in this work. In addition, since the unit cell of the C36 phase is twice that of

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the Mg2Ca (C14) phase along the c-axis [12], it is possible that the C14 phase could coexist with the C36 phase. Fig. 4 shows the magnified TEM micrograph of the C36 phase with B = [2 1 1 0], with no discernable phase boundary inside the compound but

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uniform lattice fringes perpendicular to [0001]. Thus, the C14 phase is absent in the

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AEX422 alloy.

Fig. 5a shows the bright-field TEM micrographs of the dominant bright phase, the

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lath-like phase, in the studied alloy. Based on the corresponding SAED patterns (Fig.

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5b and c) and EDS analysis results, the lath-like phase was identified as Al11RE3

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(body-centered orthorhombic structure, a = 0.443 nm, b = 1.314 nm, and c = 1.013 nm) [29, 31], in which there are approximately 2 at.% Ca segregated. Fig. 5d shows

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the bright-field TEM image of the irregular-shaped phase. The corresponding SAED patterns (Figs. 5e and f) demonstrate a same structure to the lath-like phase, also being Al11RE3, although the Ca concentration (approximately 0.6 at.%) in this phase is much lower. In addition, the corresponding EDS analysis results also illustrate that some manganese segregated in the A part while no manganese in B part. Therefore, it is deduced that the morphology differences of the Al11RE3 phase may be caused by the differences of Mn/Ca segregations during formation. Fig. 6ac represent the bright-field TEM image, the corresponding SAED pattern, and the representative EDS pattern, respectively, of the rod-like phase. Although having a similar morphology to

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ACCEPTED MANUSCRIPT the Al11RE3 phase in the conventional AE44 alloy, the rod-like phase is not Al11RE3. In addition, since the SAED patterns cannot be reasonably indexed by any known

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related phases and its atomic ratio of Al:RE is approximately 2.7, the rod-like phase

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was thus marked by Al8RE3 with unknown crystal structure and lattice parameters in this work. Furthermore, the corresponding EDS results also represent that there are some calcium and manganese segregated in this phase. Fig. 7a shows the bright-field

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TEM micrograph of the small blocky phase. From the corresponding SAED patterns

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(Fig. 7b), it was confirmed as Al2.12RE0.88 (hexagonal crystal structure, a = 0.4478 nm, and c = 0.4347 nm [3]), and some twins existed in this phase. In addition, numerous

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stacking faults can be found in this phase based on the high resolution TEM (HRTEM)

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image (Fig. 7c). Moreover, the EDS analysis result indicates that there is some

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calcium segregated in this phase, but no manganese. Fig. 7de show the bright-field TEM image, the corresponding SAED pattern, and the representative EDS pattern,

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respectively, of the lumpy phase with large volume. Since the ratio of Al:RE:Mn is approximately equal to 10:2:7 and the corresponding SAED pattern matches well with the Al10RE2Mn7 phase (hexagonal crystal structure, a = 0.90 nm, c = 1.31 nm [29, 32]), the lumpy phase was thus identified as Al10RE2Mn7. The EDS analysis results demonstrate no calcium in this phase. Fig. 8a shows the bright-field TEM image of a large plate phase consisted of two parts marked by A and B. It can be seen that part A and part B have fully different SAED patterns (Fig. 8b and c) and chemical components (Fig. 8d and e), being confirmed as Al2.12RE0.88 and Al11RE3, respectively. Additionally, the two phases have

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ACCEPTED MANUSCRIPT the following crystallographic orientation relationship as: (10 1 0)A 2.12RE0.88//(002) Al11RE3, and [ 1 2 1 0]Al2.12RE0.88//[100] Al11RE3.

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It is reported that the Al2.12RE0.88 phase was unstable and undergone a transition into

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the Al2RE phase after long-term annealing at 175 C [3]. Nonetheless, the coexistence of Al2.12RE0.88 and Al2RE was not observed in this work. The EDS analysis results indicate that both Al2.12RE0.88 and Al11RE3 phases contain some Ca and Mn, being

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approximately 2.1 at.% and 1.1 at.% in Al11RE3, respectively, and approximately 0.4

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at.% and 0.3 at.%, respectively, in Al2.12RE0.88. It should be noted that these results are distinctly deviated from the above results for the same phases with different

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morphologies, which further confirmed that the phase morphologies for a certain

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phase are closely related to particularly the segregated alloying elements. Furthermore,

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the results also suggest that the phase transformation way or production of the Al2.12RE0.88 phase could be conditional, but this need more elaborate investigation.

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3.2 Tensile properties

Fig. 9 shows the representative tensile curves of the AEX422 alloy at 20 ºC, 75 ºC, 150 ºC, 200 ºC, and 250 ºC. It is obvious that this alloy exhibits very high yield strength and obvious strain hardening behavior at room temperature, although both of them decrease as testing temperature rising. The values of the tensile properties for this alloy, including the ultimate tensile strength (UTS), the tensile yield strength (TYS), and the elongation to failure (ε), were summarized in Table 1 and compared with the similar conventional heat-resistant HPDC alloys, such as AE44, AX52 (Mg5Al2Ca, wt.%), AXJ531 (Mg5Al3Ca1Sr, wt.%). It can be seen that both

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ACCEPTED MANUSCRIPT UTS and TYS of the studied alloy decrease with the increase of testing temperature while ε rises up through the testing temperatures in this work. In addition, the TYS of

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the AEX422 alloy is much higher than those of the conventional heat-resistant HPDC

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alloys at room temperature. For instance, the TYS for the studied alloy is 204 MPa, which is greater by 38%, 27%, and 7% than those of the conventional AE44 alloy (148 MPa) [25, 28], AX52 alloy (161 MPa) [32, 33], and AXJ531 alloy (190 MPa)

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[32, 33], respectively, although with a slightly lower or comparative UTS and a

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distinctly smaller ε. Furthermore, the studied alloy also exhibits very outstanding high temperature properties. The UTS, TYS and ε of the AEX422 alloy are 164 MPa, 126

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MPa, and 17%, respectively, at 200 C while those of the conventional AE44 alloy are

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120 MPa, 100 MPa, and 24% [25, 28], respectively.

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It is reported that the mechanical properties of alloys are mainly related to cell size, intermetallic particles, and solid solution [3438]. Consequently, the underlying

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reasons for why the AEX422 alloy performs considerably higher TYS but distinctly lower ε than the conventional AE44, AX52, and AXJ531 alloys were revealed based on these factors. Firstly, the AEX422 alloy has much finer cells than the conventional AE44 alloy. Smaller cells demonstrate more cell boundaries to volume, thus more obstacles for dislocation motion, and improved strength and hardness [39, 40]. Based on HallPetch relationship, the yield strength (y) can be calculated as follow [41]:

y = 0 + Kd1/2, where 0 is a constant, K a parameter mirroring the influence of cell boundary on yield strength, and d the average cell diameter. Since magnesium alloys ordinarily

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ACCEPTED MANUSCRIPT have a great value of K due to the limited glide systems, smaller cells can thus obviously improve the yield strength. Consequently, the AEX422 alloy exhibits higher

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yield strength than the conventional AE44 alloy. However, since the AX52 and the

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AXJ531 alloys have a comparative cell size (510 m [9, 32]) to the studied alloy, the contribution of cell refinement to strength increments was thus not considered. Secondly, each of AEX422, AE44, AX52, and AXJ531 alloys contains numerous

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intermetallic particles located on cell boundaries. However, their morphologies, types

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and distribution are of a great difference. Amberger et al. [42] reported that the highly interconnected and strong intermetallic skeleton can form robust encasements for

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-Mg cells, thus improved strength. With respect to the MgAlRE based and the

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MgAlCa based alloys, the mainly intermetallic compounds are Al11RE3 and C14,

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C15, or C36, respectively. Based on bulk modulus (B), it can be deduced that the Al11RE3 phase (B = 75 GPa [43]) is more robust than the C14, C15 and C36 phases (B

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< 58 GPa [44]). Therefore, the Al11RE3 skeleton is stronger than the C14, C15, and C36 skeletons. However, the adjacent Al11RE3 particles are ordinarily discrete to each other while the C14, C15, and C36 skeletons are integrated based on the microstructures of the conventional AE44 [9], AX52 [28], and AXJ531 alloys [32]. Thus, the C14, C15, and C36 skeletons are more highly interconnected than the Al11RE3 skeleton. Nonetheless, the intermetallic skeleton in the AX52 alloy is disconnected due to the low concentration of the alloying elements. Taken the above factors together, it can be deduced that the strengthening effect of the intermetallic skeletons in these alloys follows: AEX422 > AXJ531 > AE44 > AX52. Nevertheless,

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ACCEPTED MANUSCRIPT the highly interconnected and strong intermetallic skeleton can also result in low plasticity. Fig. 10 shows the representative tensile curves along with the backscatter

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SEM micrographs of the as-cast and the heat-treated AEX422 samples. It is obvious

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that the as-cast AEX422 sample has much higher yield strength but significantly lower plasticity than the heat treated one. From SEM images, it can be seen that the highly interconnected intermetallic skeleton changed isolated or disconnected after

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heat treatment, although the cell size was not obviously changed. Therefore, the

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highly interconnected and strong intermetallic skeleton is the key to the very high yield strength but distinctly low plasticity.

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Finally, calcium has relatively higher solid solubility (approximately 16 wt.% at

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517 C) in magnesium [45] than lanthanum and cerium. Due to the great degree of

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supercooling for HPDC method, there are possibly some supersaturated solute calcium atoms in Mg matrix, which will modify the bonding force to interact with the

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moving dislocations when deformation [6]. Therefore, the AEX422 alloy has higher tensile yield strength than the conventional AE44 alloy. 4. Conclusion

The microstructures and tensile properties of the HPDC AEX422 alloy have been thoroughly investigated, and the following conclusions can be drawn: 1. The combinative addition of Ca and RE has better cell refinement effect and can clearly change the intermetallic constituents. The intermetallic phases in the HPDC AEX422 alloy can be simply grouped into the dark phase and the bright phase based on backscatter SEM image. The dark phase is the C36 phase and the

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ACCEPTED MANUSCRIPT bright one is mainly the Al-RE based phases, including Al11RE3, Al8RE3, Al2.12RE0.88, and Al10RE2Mn7, wherein the Al11RE3 phase in the studied alloy

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presents disparate morphologies from that in the conventional AE44 alloy, being

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lath-like, irregular-shaped, and plate-like, the rod-like Al8RE3 phase is an unknown phase, and the Al2.12RE0.88 phase contains twins and stacking faults. 2. The EDS analysis results illustrate that little calcium and manganese segregate in

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the C36 phase in the studied alloy and the morphology for a certain Al-RE phase

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are closely related to particularly the segregated alloying elements. 3. The AEX422 alloy has clearly higher tensile strength than the similar

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conventional heat-resistant magnesium alloys (AE44, AX52, and AXJ531), which

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is mainly attributed to the highly interconnected and strong intermetallic skeleton

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that can form robust encasements for -Mg cells. However, it is also the key to the low plasticity at lower temperatures for the studied alloy.

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Acknowledgements

This work is supported by the National Key Technologies R&D Program 20130305007GX, 20130305011GX, 2014-GX-216A, 2015DFH50210, the National Natural Science Foundation of China under grants no. 20921002, and by the Joint research program of the Chinese Academy of Science and the Japan Society for the Promotion of Science (JSPS) GJHZ1413. References [1] Pekguleryuz M, Celikin M. Creep resistance in magnesium alloys. Int Mater Rev 2010; 55: 197217.

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The values of UTS (MPa), TYS (MPa), and  (%) of the HPDC AEX422 alloy and the similar

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conventional heat-resistant HPDC magnesium alloys (AE44, AX52, and AXJ531) at room temperature (20 C), 75 C, 150 C, 200 C, and 250 C.

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Fig. 1. OM image and backscatter SEM image of the HPDC AEX422 alloy.

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Fig. 2. XRD pattern of the HPDC AEX422 alloy.

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Fig. 3. Bright-field TEM images and the corresponding SAED patterns of the dark phase with (a)

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and (b) alternating layers, (c) and (d) single layer, and (e) and (f) triple junction, respectively, in

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the HPDC AEX422 alloy with B = [2 1 1 0].

Fig. 4. Bright-field TEM images of the C36 phase (B = [2 1 1 0]).

Fig. 5. (a) and (d) Bright-field TEM images and (b), (c), (e), and (f) the corresponding SAED patterns the lath-like phase (ac) and the irregular-shaped phase (df).

Fig. 6. (a) Bright-field TEM image, (b) the corresponding SAED pattern, and (c) the representative EDS pattern for the rod-like phase.

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yellow arrows, and SF represents stacking fault in figure (c).

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lumpy phase. In figure (b), the diffraction contrast spots that come from twins were marked by

Fig. 8. (a) Bright field TEM image of the large plate phase with A and B parts, (b) and (c) the

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corresponding SAED patterns for part A and part B, respectively, and (d) and (e) the representative

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EDS patterns for part A and part B, respectively.

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250 C in this work.

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Fig. 9. Representative tensile curves of the HPDC AEX422 alloy at 20 C, 75 C, 150 C, 200 C,

Fig. 10. Representative tensile curves at 20 C along with the respective backscatter SEM images

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ACCEPTED MANUSCRIPT Table 1 20 C Alloys

75/150 C

200/250 C

UTS

TYS



UTS

TYS



UTS

234 11 245

204 4 148

4 0.6 11

2209/

1823/

72/

1997

1483

93

-/157

-/123

-/23

AX52c,d

228

161

13

-/-

-/-

-/-

AXJ531c,d

238

190

8

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Ref. [25],b Ref. [28], c Ref. [32], d Ref. [33]

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a

1262/

173/

1234

1122

345

120/105

100/89

24/16

-/-

-/-

-/-

-/-

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1646/

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a,b

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AEX422

TYS

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ACCEPTED MANUSCRIPT Highlights 1. Combinative addition of Ca and RE has a very good cell refinement effect.

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2. The Al-RE phases are disparate from those in the conventional AE44 alloy.

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3. The studied alloy has much higher yield strength than the similar magnesium alloys.

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4. The intermetallic skeleton plays a key role on the tensile properties.

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