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Effects of alloying composition on the microstructures and mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy
Accepted Manuscript Effects of alloying composition on the microstructures and mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy Fang Wang, Wenlong Xiao, Maowen Liu, Jing Chen, Xiang Li, Jiabing Xi, Chaoli Ma PII:
S0042-207X(18)31645-2
DOI:
https://doi.org/10.1016/j.vacuum.2018.10.072
Reference:
VAC 8345
To appear in:
Vacuum
Received Date: 22 August 2018 Revised Date:
24 October 2018
Accepted Date: 25 October 2018
Please cite this article as: Wang F, Xiao W, Liu M, Chen J, Li X, Xi J, Ma C, Effects of alloying composition on the microstructures and mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy, Vacuum (2018), doi: https://doi.org/10.1016/j.vacuum.2018.10.072. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT 1
Effects of alloying composition on the microstructures and
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mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy
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Fang Wang 1, Wenlong Xiao 1,2, *, Maowen Liu 1, Jing Chen 1, Xiang Li 1, Jiabing Xi 1, Chaoli Ma 1,2
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of Materials Science and Engineering, Beihang University, Beijing 100191, China
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Industry and Information Technology, Beihang University, Beijing, 100191, China
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Abstract: With the purpose of increasing our understanding of the microstructural evolution and
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mechanical
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Key Laboratory of Aerospace Advanced Materials and Performance of Ministry of Education, School
properties
of
Mg-Al-Zn-Ca-RE
system
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Key Laboratory of High-Temperature Structural Materials & Coatings Technology, Ministry of
alloy,
Mg-xAl-(8-x)Zn-4Ca
and
Mg-xAl-(8-x)Zn-3Ca-1RE (wt. %) casting alloys were studied in this paper. The α-Mg dendrite
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was refined, and the intermetallic compounds became finer and less divorced by increasing the Al
12
content. During solidification RE addition consumed Al and Zn to form RE-containing phase in
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the initial solidification stage, and then the residual liquid solidified to form Ca-containing phase
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without RE participation. The majority phases, Al11RE3 and C36, were transformed into Al2REZn2
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and Ca2Mg6Zn3 as the Al content was decreasing. The Ca-containing phases tended to form a
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highly interconnected network which would deteriorate the ductility, while the substitution of RE
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for a part of Ca would make the structure less continuous. A range of eutectic morphologies were
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found in Ca-containing phases. C36 phase could shape a lamellar eutectic structure, and
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Ca2Mg6Zn3 phase was prone to exhibit as a divorced morphology. Besides, Ca2Mg6Zn3 formed at a
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relatively low temperature which will bring about microshrinkage. Mg-6Al-2Zn-3Ca-1RE alloy
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containing refined and more evenly distributed C36 with less continuous morphology showed the
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best mechanical performance.
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Key words: Magnesium alloy; Calcium; Rare earth; Microstructures; Intermetallic compound;
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Mechanical properties.
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*
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School of Materials Science and Engineering, Beihang University, Beijing 100191, China. Tel:
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+86-10-8233 8631; Fax: +86-10-8233 8631
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E-mail: [email protected] (Wenlong Xiao)
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Corresponding author.
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1. Introduction Persistent focus towards lightweight in automotive industry stimulates the
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development of magnesium alloys. Among them, heat resistant Mg alloys are deemed
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as a promising structural material in automobile engine. However, the well-developed
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commercially Mg-Al based alloys, such as AZ91 and AM60, show poor heat
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resistance due to the rapid dynamic precipitation of β-Mg17Al12 phase at high
7
temperatures [1]. There is a general consensus that the most effective way to improve
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heat resistance of Mg-Al based alloy is to acquire thermal stable intermetallic
9
compounds by alloying the alloy with Zn, Ca, Si, Sr and rare earth (RE), etc [1-4]. In
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recent years, Mg-Al-Ca and Mg-Zn-Ca magnesium alloys have been attracted
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numerous studies due to their excellent mechanical properties at both room and
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elevated temperature, low price, texture weakening and good ignition resistance [5-8].
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However, calcium generally brings worse die sticking and hot tearing. It has been
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reported that zinc addition would restore the castability of Mg-Al alloys, and RE
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addition would ameliorate hot tearing susceptibility of Mg-Zn-Al alloys by shortening
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the solidification freeze range [9-11].
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The formation of Ca-containing phases and RE-containing phases can suppress
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the formation of low melting point phases such as Mg-Al, Mg-Zn and Mg-Zn-Al
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phases in Mg-Al-Ca, Mg-Zn-Ca, Mg-Al-RE, Mg-Al-Zn-Ca and Mg-Al-Ca-RE system
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alloys, which are capable to obtain alloys with excellent high temperature properties
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[6, 12-16]. Mg-Al-Zn-Ca-RE alloy containing all the mentioned above system alloys
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is therefore considered as a promising heat resistant magnesium alloy system with
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high performance. In our previous work, the microstructures and mechanical
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properties of Mg-Al-Zn-Ca-La alloys at fixed Al and Zn contents has been
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investigated, Al2LaZn2 and Ca2Mg6Zn3 would form as major intermetallic compounds
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[16]. However, the categories of Ca-containing phase in Mg-Al-Ca alloys are largely
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depend on Al/Ca ratio, and Zn addition would also have a significant effect on the
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phase constitution of Mg-Al-Zn-Ca alloy [5, 14, 17, 18]. Our recently study reveals
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that Al and Zn contents have a great influence on the microstructures and mechanical
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ACCEPTED MANUSCRIPT properties of Mg-Al-Zn-Ca quaternary alloy [14]. Two Laves phases, i.e. C36 and
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C15, a quaternary Q phase and/or Ca2Mg6Zn3 isomorphs phase are observed in the
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casting Mg-xAl-(8-x)Zn-2Ca alloys with different Al and Zn contents, and the alloy
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containing the majority phase of C36 exhibits the best combination of strength and
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ductility. Moreover, the category of second phase transforms from C36 to C15 and
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Ca2Mg6Zn3 phase by decreasing the Al/Zn ratio at a constant Ca content. As Q phase
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is harmful to high temperature properties, we consider that the way of increasing Ca
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content and/or alloying with RE might eliminate such phase [14]. In addition, it has
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been reported that magnesium alloys containing thermally stable phases with proper
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morphology and size as well as distribution are beneficial for the desire of high
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mechanical performance [19, 20]. In that case, the aims of this work are to investigate
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the effects of Al and Zn contents on the microstructure and mechanical properties of
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high Ca-containing Mg-Al-Zn-Ca alloys, and the role of a small amount of RE
14
substitution for Ca in the microstructural evolution and mechanical properties. For
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these purposes, a series of Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys
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are studied, and the relationship between the microstructure and mechanical
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properties are discussed.
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2. Experimental procedure The
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actual
compositions
of
three
Mg-xAl-(8-x)Zn-4Ca
and
three
Mg-xAl-(8-x)Zn-3Ca-1RE alloys (weight percent, wt. %, also for hereafter not
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mentioned) are listed in Table 1, and the alloys were prepared from pure (≥99.9%) Mg,
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Al and Zn elements and Mg-20Ca and Mg-25RE (cerium rich misch metal) master
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alloys by electrical resistance melting furnace. Casting process was performed under
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the protection of an anti-oxidizing flux. The melt was refined by BaCl2 refining agent
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and then holding at 740°C for 20min to purify the melt before pouring into a steel
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mold at room temperature. It is noted that Mg-25RE master alloy was prepared in
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vacuum melting furnace (10-3 Pa) because of the poor ignition resistance. The
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obtained as-cast alloys were designated by AZX(E) followed with three or four
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Table 1. The microstructure was observed by optical microscopy (OM) and scanning
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electron microscopy (SEM) equipped with energy dispersive spectrometry (EDS)
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operated under a vacuum condition of 10-5 Pa. The intermetallic compounds were
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identified by a RIGAKU RINT-2000 X-ray diffractometer (XRD) with Cu Kα
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radiation. Differential scanning calorimetry (DSC) was conducted for analyzing the
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solidification behavior with a heating rate of 10 K/min under Ar gas atmosphere
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protection. The flat dog-bone like tensile test plates with gauge length of 15 mm,
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width of 5.0 mm and thickness of 1.5 mm were machined from the same position of
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the ingots, and tensile tests were performed three times on Instron 8801 universal
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testing machine at room temperature (RT) according to Chinese metallic
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materials-tensile testing standard (GB/T228-2002). The fracture surfaces were
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observed using OM.
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3. Results and Discussion
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3.1 Microstructure
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3.1.1 Phase constitution
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Fig. 1 exhibits the XRD patterns of the studied alloys. It is illustrated in Fig. 1a
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that AZX264 alloy is composed of α-Mg, Ca2Mg6Zn3 and C14-Mg2Ca phases.
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Replacing Ca with RE causes the appearance of new peaks in AZXE2631 alloy (Fig.
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1b). The new peaks belong to Al2REZn2 phase and are identified on the basis of
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Al2CeZn2 phase (PDF#29-0013: I4/mmm space group, a = 0.4244nm and c =
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1.0986nm) [21]. Additionally, unknown peaks are detected in both AZX264 and
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AZXE2631 alloys, they might belong to Ca2Mg6Zn3 in a different crystallographic
25
form [22]. Compared with Fig. 1a, the increase of Al content leads to a decrease in the
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amount of Ca2Mg6Zn3 phase as shown in Fig. 1c (AZX444 alloy), and the peaks of
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C36 phase emerge in this alloy instead of that of C14 phase (Fig. 1c). The peaks
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belong to C36 are consistent with our previous work [14]. Similarity with AZXE2631
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shows no difference with the result in our previous study [16]. When the Al content is
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further increased, though the peaks belonging to C36 phase and Ca2Mg6Zn3 phase are
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still remained in AZX624 alloy (Fig. 1d), the amount of C36 phase is dramatically
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increased. By comparison, Al11RE3 phase is formed in AZXE6231 alloy rather than
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Al2REZn2 in AZXE4431 alloy, which is confirmed in Fig.1e based on Al11Ce3 phase
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(PDF#19-0006: Immm space group, a = 0.4395 nm, b = 1.0092 nm, c = 1.3011 nm).
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3.1.2 Observation of primary α-Mg phase
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Fig. 2 shows the optical micrographs of the as-cast alloys, where the typical
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α-Mg dendrites surrounded by second phases can be found in each alloy. It is evident
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that the secondary dendrite arm spacing (SDAS) is changed by varying the alloying
12
composition. The SDAS of each alloy is present in Fig. 3, which was measured by
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averaging over the distance between the adjacent dendrite arm centers. There were at
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least 60 secondary dendrite arms being counted in different optical micrographs for
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each alloy. AZX264 alloy shows the largest SDAS among the studied Mg-Al-Zn-Ca
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alloys as shown in Fig. 3, and the SDAS decreases obviously and then increases
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slightly with the increase of Al content so that AZX444 alloy has the smallest SDAS.
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When RE is substituted for a part of the Ca content, the dendrite arms become more
19
developed and obtain a tiny refinement as compared with the alloy containing the
20
same Al and Zn contents. As shown in Fig. 2, the majority of second phases are
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formed along the dendritic boundaries and distributed as a highly continuous
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interconnected morphology, especially for the high Zn-containing alloys.
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Grain refinement effect generally involves in the addition of nucleants and/or
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alloying elements into a melt before casting. The additions of Al, Zn, Ca and RE to
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Mg alloys have little effect on nucleation of the primary α-Mg phase since these
26
elements are mostly segregated to form second phases well after the nucleation.
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However, the solute elements play a critical role in controlling the growth of the
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nucleated grains and in subsequent nucleation during solidification, which can be
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explained in terms of the growth restriction factor (GRF) [23]. It demonstrates that the
ACCEPTED MANUSCRIPT solute tends to segregate on the ahead of growing dendrite/grain and thus generates
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the constitutional undercooling (△Tc). This constitutional undercooling not only
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suppresses the growth of dendrite/grain, but also promotes nucleation when △Tc
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reaches the undercooling required for nucleation (△Tn). The GRF value, well known
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as Q, is widely used to describe the effect of solutes on dendrite arm spacing/ grain
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size, which can be expressed as
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=∑
,
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−1
,
(1)
where ml,i is the slope of the liquidus line, c0,i is the concentration of the solute, and ki
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is the partition coefficient for the element i in the binary alloy. The relationship
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between grain size and Q for magnesium alloy binary alloyed with Al, Zn and Ca is
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well illustrated in [23, 24], which is plotted in Fig. 4 showing alloying compositions
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calculated by equation (1). It shows a distinct decrease in grain size at a minor
13
addition of either Zn or Ca while only a minimal reduction at a further addition of that
14
exceeding the saturation level. Obviously, the addition level of Al for grain refinement
15
is much higher than that of Zn and Ca. An evident grain/dendrite refinement is
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observed in AZX444 and AZXE4431 alloys compared to AZX264 and AZX2631
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alloys. Since the additions of Zn and Ca in these alloys are larger than the saturation
18
levels and that of Al is less than the level (Fig. 4), the increased Al content should
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account for the refinement effect in the SDAS [23, 24]. When the Al content increases
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from 4 wt% to 6 wt%, it can be inferred from Fig. 4 that the refinement effect should
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still remain. However, only a slightly increase in SDAS can be found in Fig. 3 as
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compared AZX444 alloy to AZX624 alloy and AZX4431 alloy to AZXE6231 alloy. In
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high Al content alloys, the Mg-Al-Ca phase containing a certain amount of Zn would
24
form (section 3.2 in detail), and the contribution of Zn to GRF should be decreased in
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AZX624 and AZX6231 alloys. In that case, the counteracting effect of increasing Al
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solute and decreasing Zn solute makes the SDAS exhibits slightly change when the Al
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content is further increased.
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Xiao, et al deduced that Mg-Ce binary alloy with a low concentration of Ce will
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cause a dramatic effect on grain refinement by using a relative grain size (RGS)
ACCEPTED MANUSCRIPT model proposed by Easton and StJohn [25, 26]. As with similar instance of
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Mg-Al-Zn-RE alloy, the RE elements in Mg-Al-Zn-Ca-RE alloy can easily segregate
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on the solid/liquid interface in the initial stage of solidification, which produces
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barriers between the liquid and solid phases and thus inhibits the diffusion of Al, Zn
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and Ca into the solid phase, leading to the grain/dendrite refinement by raising the
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△Tc. On the other hand, the formed RE-containing phase consumes a large number of
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Al and Zn solutes (section 3.2 in detail), which lowers △Tc. The two factors canceled
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each other out so that there are only a slightly decrease of SDAS in the studied alloys
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with RE addition.
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3.1.3 Identification of intermetallic compounds
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The back-scattered electron images (BEIs) of the studied alloys are shown in Fig.
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5, and there are apparent contrasts among the intermetallic compounds in the studied
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alloys. In general, different phases can be distinguished according to the contrasts in a
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BEI. However, it is hardly probable to differentiate C14-Mg2Ca phase and Ca2Mg6Zn3
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phase by comparing the contrasts in Mg-Zn-Ca alloys [27, 28]. The categories of the
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second phase with different contrasts are further examined by EDS, and the elemental
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compositions of each phase, pointed out by arrows A-J in Fig. 5, are listed in Table 2.
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Fig. 5 shows that Ca-containing phases tend to be a highly interconnected
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microstructure in the studied alloys. The majority phase, Ca2Mg6Zn3 in AZX264 alloy,
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is transformed into C36 phase with a gray contrast by increasing the Al content.
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Whilst in AZX444 and AZX624 alloys, Ca2Mg6Zn3 becomes the minor phase. As
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listed in Table 2, Al and Zn elements are contained in Ca2Mg6Zn3 phase and C36
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phase, respectively. It should be noticed that the composition of C36 phase is different
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between AZX444 and AZX624 alloys, which can be identified as C36-(Mg, Al)2Ca
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and C36-Al2(Mg, Ca), respectively, according to calculating the valence electron
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concentrations [14, 29]. A high content of Al atom is consumed by the formation of
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C36-Al2(Mg, Ca), which can also be found in AZX622 and AZX442 alloys published
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in [14]. But instead of forming Q phase as the low melting phase in AZX622 and
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AZX442 alloys, the Ca2Mg6Zn3 phase is formed in AZX624 and AZX444 alloys
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owing to the increased Ca content. As substituting Ca with RE, RE-containing phase with the brightest contrast is
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visible in the BEIs and locate between the Ca-containing phase, which leads to a less
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continuous morphology of the Ca-containing compounds (Fig. 5b, d and f). There are
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no RE elements detected in Ca-containing phases according to the EDS results (Table
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2), and the RE addition does not change the type of Ca-containing phase except for
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AZXE6231 alloy. In AZXE6231, Al11RE3 phase forms instead of Al2REZn2 phase and
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a few Zn atoms are dissolved in this RE-phase, while Ca2Mg6Zn3 phase which is
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present in AZX624 cannot be detected. As shown in Fig. 5f, a small amount of
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Mg-Al-Zn-Ca quaternary Q phase displaying a medium contrast (arrow I) can be
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identified in this alloy [14]. The amount of Q phase is far less than that in AZX622
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alloy. It seems that adding RE or increasing Ca content can both suppress the
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formation of Q phase.
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3.2 Phase formation character during solidification
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The solidification event, to some extent, determines the size, morphology and
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distribution of second phase and casting defect formation, which in turn to have an
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impact on mechanical properties [30]. In this section, DSC heating curves of the
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representative studied alloys are used to speculate the phase formation during
21
solidification (Fig. 6). As can be seen, each endothermic peak denotes the dissolution
22
of the phase labelled [14, 16]. Table 3 lists the peak temperature of each
23
transformation. For AZXE2631 alloy, the endothermic peaks corresponding to the
24
melting of α-Mg, Al2REZn2, C14 and Ca2Mg6Zn3 are occurred at approximately
25
618°C, 587°C, 489°C and 458°C, respectively. The peaks are shifted when changes
26
are made to the Al and Zn concentrations (AZXE6231 alloy). The melting peak of
27
α-Mg is decreased to 591°C, and the peaks related to C14 and Ca2Mg6Zn3 are
28
disappeared, while a peak that refers to C36 can be detected. It should be mentioned
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in this alloy that the peak of Al11RE3 is hard to distinguish due to the peak is covered
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by α-Mg peak. As for AZX444 alloy, except for the melting peak of α-Mg (593°C),
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two endothermic peaks at approximately 500°C and 445°C are considered to be phase
3
transformation of C36 and Ca2Mg6Zn3, respectively. As mentioned above, it can be known that the formation of RE-containing phase
5
follows on the heels of the nucleation and growth of primary α-Mg in the initial stage
6
of solidification, and which type would be formed is determined by the Al and Zn
7
contents, i.e. Al2REZn2 will transform into Al11RE3 with increasing the Al content.
8
Afterward, the solidification of residual liquid will comply with the solidification
9
event of Mg-Al-Zn-Ca quaternary system. A part of Al and Zn are consumed by the
10
formation of RE-containing phase, and then the Ca/Al and Ca/Zn ratios are changed
11
in the residual liquid, thereby probably changing the type of Ca-containing phase [14,
12
17, 31, 32]. As a summary, Ca2Mg6Zn3 and C14 would transform into C36 with
13
increasing the Al content, which is consistent with our previous work [14]. As listed in
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Table 2, the composition of C36 varies in alloys with different Al contents. It might be
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ascribed to that C36-(Mg, Al)2Ca phase follows the Al2Ca–Mg2Ca pseudo-binary
16
system, while C36-Al2(Mg, Ca) phase complies with the Al2Mg–Al2Ca pseudo-binary
17
system [18, 33].
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As shown in Fig. 5, the morphology of eutectic phases varies in the studied
19
alloys. Owing to the non-equilibrium solidification, both Ca2Mg6Zn3 and Mg-Al-Ca
20
phases form either a typical fully divorced or a partially divorced morphology and
21
exhibits a highly interconnected network. With increasing the Al content and
22
meanwhile decreasing the Zn content for the alloy, an increasing proportion of less
23
divorced morphology and even typical eutectic morphology is present. Fig. 7 shows
24
higher magnification SEM images for a close inspection. As can be seen from Fig.
25
7a-c, a high proportion of granular and fibrous eutectic α-Mg is found, and the
26
Ca2Mg6Zn3 phase becomes less divorced with decreasing the Zn content. This may be
27
attributed to that Zn has a strongly partition effect to eutectic growth [34, 35]. As can
28
be seen in Fig. 7b-d, the C36 phase tends to display as a much less divorced and even
29
lamellar eutectic morphology. It has been reported that Al does not partition as
30
strongly as an equivalent amount of Zn during eutectic phase growth [30]. In addition,
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ACCEPTED MANUSCRIPT primary α-Mg also has a significant impact on eutectic morphology [36]. The
2
relatively low eutectic temperature of Ca2Mg6Zn3 means a wide solidification range
3
as well as a high solid fraction of primary α-Mg during eutectic solidification, which
4
leads to narrow spaces for the eutectic growth and therefore forms more divorced
5
morphology. With increasing the Al content and/or substituting RE for a part amount
6
of Ca, the α-Mg dendrites become much developed (Fig. 2) and the main eutectic
7
reaction occurs at a relatively high temperature. As shown in Fig. 5f, a smaller size,
8
less divorced and more evenly distribution of C36 are thus obtained in AZXE6231
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alloy.
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3.3 The relationship between microstructures and mechanical properties
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3.3.1 Tensile properties
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The typical tensile stress-strain curves of the studied alloys are presented in Fig.
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8, and the average yield strength (YS), ultimate tensile strength (UTS) and elongation
15
(El.) are listed in Table 4. Furthermore, column chart with error bars has been used in
16
Figure 9 to illustrate the effects of alloying compositions. It can be found that the
17
mechanical performance varies with the different alloying compositions. As can be
18
seen in Fig. 8, AZX264 alloy fractured immediately after yielding, showing a limited
19
plastic deformation. Figure 9 shows that the UTS and El. are improved considerably
20
with the increase of the Al content, but the YS has a negligible change when the Al
21
content increases from 2 wt% to 4 wt%. Moreover, the YS is decreased slightly from
22
102±4 MPa of AZX442 alloy to 95±3 MPa of AZX624 alloy. For the RE-containing
23
alloy, the YS firstly decreases from 107±5 MPa to 90±2 MPa and then increases
24
significantly to 146±6 MPa with increasing the Al content. However, the UTS and El.
25
remain nearly constant firstly and, finally, increase substantially. It is noted that the
26
substitution of RE for a part of Ca exhibits different impacts on tensile properties,
27
which depends on Al and Zn concentrations. As compared AZXE2631 alloy with
28
AZX264 alloy, the UTS and El. are remarkably improved, while the YS almost
29
remains unchanged. When the Al content is increased to 4 wt%, the RE addition can
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2
properties are substantially enhanced by replacing a part of Ca with RE when the Al
3
content is further increased to 6 wt%. Among the studied alloys, AZXE6231 alloy has
4
the most optimal strength and ductility, in which the YS, UTS and El. are 146±6 MPa,
5
212±3 MPa and 2.4±0.3%, respectively.
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Fig. 10 displays the optical micrographs of longitudinal section adjacent to the
7
fracture surface. It is clearly seen that the cracks mostly develop along the interface
8
between α-Mg matrix and eutectic compounds where stress concentration is generated,
9
and cracks can readily propagate along the compounds particularly with a divorced
10
morphology, thus deteriorating the ductility. Compared to C36, Ca2Mg6Zn3 in the
11
studied alloys shows a much higher interconnected network and more divorced and
12
coarse morphology, which causes stress concentration and split of α-Mg matrix more
13
readily [37]. As a result, high Zn-containing AZX264 and AZXE2613 alloys with high
14
volume fraction of Ca2Mg6Zn3 exhibit worse ductility. With respect to the less
15
divorced and even lamellar eutectic phase, i.e. C36 phase, the cracks develop not only
16
along the phase/dendrite boundary but also into the eutectic phases (Fig. 10d),
17
resulting in acceptable ductility. Thus, AZX624 and AZXE6231 containing C36 as the
18
majority phase with both refined divorced and lamellar morphologies exhibit
19
relatively high ductility, which is in a good agreement with our previous work [14].
20
Furthermore, RE addition led to the formation of RE-phase in the initial stage of
21
solidification as mentioned above. The RE-phase would occupy between the α-Mg
22
dendrites where the nucleation of Ca-containing phase would then take place. As
23
solidification continued, the growth direction of Ca-containing phase could be
24
restricted by the pre-existed RE-phase. The subsequently formed Ca-containing phase
25
could not link each other as highly as that in the counterpart RE-free alloy. On the
26
other hand, when 1 wt% Ca is replaced by RE, the volume fraction of Ca-containing
27
phase formed is decreased. Consequently, the Ca-containing phases in RE-containing
28
alloy are somehow less continuous (Figure 5), thus the ductility of Mg-Al-Zn-Ca
29
alloy can be slightly improved by RE addition.
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3.3.2 Strengthening mechanism In general, the yield strength is mainly related to SDAS and intermetallic
3
compound for Mg casting alloys. Based on the well-known Hall-Petch relationship,
4
smaller SDAS has more grain boundaries which can improve the strength. However,
5
AZX444 and AZXE4431 alloys have relatively small SDAS but low yield strength.
6
Additionally, It has been reported previously that the strength would monotonic
7
increase with increasing the volume fraction of intermetallic compounds in Mg-Al
8
and Mg-RE binary alloys [38, 39]. Fig. 11 shows the volume fractions of the
9
intermetallic compounds for the studied alloys. It is measured to be 9.1 vol%, 9.5
10
vol %, 9.6 vol %, 8.0 vol %, 7.2 vol %, and 7.8 vol. % for AZX264，AZX444,
11
AZX624, AZXE2631, AZXE4431 and AZXE6231 alloys, respectively. There are not
12
too much differences between either the three Mg-Al-Zn-Ca alloys or the other three
13
Mg-Al-Zn-Ca-RE alloys, but the yield strength varies. It seems that other factors
14
should be responsible for the strengthening mechanism.
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It is figured out that highly interconnected and strong intermetallic skeleton can
16
form robust encasements for α-Mg cells, thus improving strength [40, 41]. In this
17
study, the high Zn containing alloys have more continuous skeleton, and the yield
18
strength of Mg-Al-Zn-Ca alloy does increase with increasing the Zn content. While
19
for the RE-containing alloys, AZXE6231 alloy shows the highest yield strength
20
although less continuous skeleton is formed. Apart from the above factors, the size
21
and distribution of intermetallic compound also play critical roles in determining the
22
strength, and the well-developed models for interpreting the strengthening mechanism
23
also reveal such importance [38, 42-45]. It can be found in Fig. 2 and Fig. 5 that the
24
size and distribution of the intermetallic compounds in AZXE6231 alloy are
25
remarkably distinguished from other alloys. AZXE6231 alloy has the finest and most
26
evenly distributed intermetallic compounds among the studied alloys. During tensile
27
test, the more evenly distributed C36 and the relative small α-Mg dendrites not only
28
can inhibit the grain boundary slippage and dislocation motion, but can transfer the
29
applied load with the help of good α-Mg reinforcement interfacial integrity [37].
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ACCEPTED MANUSCRIPT Moreover, considering that the bulk modulus of RE phase is much higher than that of
2
Ca phase, the yield strength of RE-containing alloy should be higher as compared
3
with the RE free alloy at the same Al and Zn contents [42]. Nevertheless, AZXE4431
4
alloy shows lower yield strength than AZX444 alloy, and the reason need to be further
5
investigated. In summary, it can be deduced that the highest strength of AZXE6231
6
alloy is mainly attributed to the dispersion strengthening as a result of the refined and
7
more evenly distributed intermetallic compounds caused by the refined dendrites and
8
the less divorced eutectic morphology [36, 43].
Fig. 12 depicts the work-hardening rate (Θ) and true strain curves of AZXE2631,
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AZX444 and AZXE6231 alloys, and Θ equals to
, where σ is true stress, ε is
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true strain and
12
beginning of plastic deformation and then decreased gradually until fracture for all the
13
alloys. AZX6231 alloy with the majority of C36 phase displays the highest work
14
hardening rate, while AZX2631 alloy with the majority of Ca2Mg6Zn3 phase has the
15
lowest one, and AZX444 alloy with the both phases owns the moderate one, which is
16
in a good agreement with our previous work [14]. Combined with the obtained better
17
ductility resulting from the less continuous structure, AZXE6231 alloy therefore
18
shows the highest UTS among the studied alloys. However, the failure still occurs in
19
the stage of uniform deformation during the tensile tests.
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is strain rate. The work hardening rates rapidly decreased at the
The last but not the least, casting defects of these alloys are considered as well.
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Fig. 13 displays optical macrographs of the studied alloys. It is obvious in Fig. 13a
22
and b that numerous micro-shrinkages are present in AZX264 alloy, therefore, it has
23
the worst mechanical properties among the studied alloys, and the failure strength of
24
the samples are even occurred at around the 0.2% yield strength. Either increasing Al
25
content or RE addition can suppress the micro-shrinkage as can be detected in Fig.
26
13b and c. The formation of micro-shrinkage is involved with the final stage of the
27
solidification process for eutectic growth. The formation of divorced morphology
28
which need independent nucleation and growth will provide a much greater resistance
29
to feeding [30, 46]. Additionally, a low eutectic point can extend the solidification
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2
containing divorced Ca2Mg6Zn3 as the majority phase has the worst feedability.
3
Increasing Al content and/or RE addition can reduce and even suppress the formation
4
of Ca2Mg6Zn3 as well as modify the eutectic morphology. In that case, it is expected
5
that AZXE6231 has the best castability, so that micro-shrinkage can hardly be
6
observed (Fig. 13d).
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4. Conclusions
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Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys were prepared by
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gravity casting to examine the microstructure and the resultant mechanical
11
performance, which could give a guidance for alloy design of this five-element
12
system heat resistance magnesium alloy. The main conclusions could be drawn
13
hereafter.
14
(1) The α-Mg dendrites became more developed and the second dendrite arm spacing
15
decreased evidently with increasing the Al content in Mg-xAl-(8-x)Zn-4Ca alloy,
16
and substituting 1wt% RE for the equivalent amount of Ca enabled the dendrites
17
slightly refined. The intermetallic compounds were located at the dendrite
18
boundaries, which became much finer and more evenly distributed as the
19
dendrites were refined.
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(2) RE addition consumed the Al and Zn contents to form Al11RE3 and Al2REZn2 in
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the initial stage of solidification at high Al alloy and high Zn alloy, respectively,
22 23 24 25
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and it did not take part in the subsequently formation of Ca-containing phases which would comply with the solidification event of Mg-Al-Zn-Ca quaternary alloy system. In the quaternary alloy, the majority phase transformed from Ca2Mg6Zn3 into C36 with the increase of the Al content.
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(3) Ca-containing phases tended to form a highly interconnected continuous network
27
which became less continuous because of RE addition. Among the Ca-containing
28
phases, Ca2Mg6Zn3 phase had a more divorced morphology and solidified at a
29
relatively low temperature which would cause microshrinkage, while C36 phase
ACCEPTED MANUSCRIPT 1
formed at a higher temperature owned a much less divorced and even lamellar
2
eutectic morphology and a better castability. (4) The size, distribution and category of intermetallic compound played important
4
roles in determining the mechanical properties. Mg-6Al-2Zn-3Ca-1RE alloy
5
showed the best mechanical performance among the studied alloys due to the less
6
continuous, much refined and more evenly distributed C36 phase as well as the
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obtained soundness casting.
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Acknowledgements
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The authors are grateful to the financial support by National Key Research and
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Development Program of China (No. 2016YFB0301103) and National Natural Science
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Foundation of China (NSFC, No. 51401010, 51671007 and 51671012). I (Fang Wang)
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want to thank the corresponding author and prof. Ma in particular for their patient
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guidance and help.
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References
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[1] A.A. Luo, Recent magnesium alloy development for elevated temperature applications, Int. Mater. Rev. 49(1) (2004) 13-30. 197-217.
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[2] M. Pekguleryuz, M. Celikin, Creep resistance in magnesium alloys, Int. Mater. Rev. 55(4) (2010) [3] S. Zhu, M.A. Easton, T.B. Abbott, J.-F. Nie, M.S. Dargusch, N. Hort, M.A. Gibson, Evaluation of
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Magnesium Die-Casting Alloys for Elevated Temperature Applications: Microstructure, Tensile Properties, and Creep Resistance, Metallurgical and Materials Transactions A 46(8) (2015) 3543-3554. [4] W. Xiao, M.A. Easton, M.S. Dargusch, S. Zhu, M.A. Gibson, The influence of Zn additions on the microstructure and creep resistance of high pressure die cast magnesium alloy AE44, Materials Science and Engineering: A 539 (2012) 177-184. [5] R. Ninomiya, T. Ojiro, K. Kubota, Improved heat resistance of Mg-Al alloys by the Ca addition, Acta Metallurgica et Materialia 43(2) (1995) 669-674. [6] A.A. Luo, B.R. Powell, M.P. Balogh, Creep and microstructure of magnesium-aluminum-calcium based alloys, Metallurgical and Materials Transactions A 33(3) (2002) 567-574. [7] Y. Du, M. Zheng, B. Jiang, Comparison of microstructure and mechanical properties of Mg-Zn microalloyed with Ca or Ce, Vacuum 151 (2018) 221-225. [8] M. Sakamoto, S. Akiyama, K. Ogi, Suppression of ignition and burning of molten Mg alloys by Ca bearing stable oxide film, J. Mater. Sci. Lett. 16(12) (1997) 1048-1050.
ACCEPTED MANUSCRIPT [9] I.A. Anyanwu, Y. Gokan, S. Nozawa, A. Suzuki, S. Kamado, Y. Kojima, S. Takeda, T. Ishida, Development of New Die-castable Mg-Zn-Al-Ca-RE Alloys for High Temperature Applications, MATERIALS TRANSACTIONS 44(4) (2003) 562-570. [10] W. Xiao, S. Zhu, M.A. Easton, M.S. Dargusch, M.A. Gibson, J.-f. Nie, Microstructural characterization of high pressure die cast Mg-Zn-Al-RE alloys, Mater. Charact. 65 (2012) 28-36. [11] J. Song, F. Pan, B. Jiang, A. Atrens, M.-X. Zhang, Y. Lu, A review on hot tearing of magnesium alloys, Journal of Magnesium and Alloys 4(3) (2016) 151-172.
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[12] A. Suzuki, N.D. Saddock, J.R. TerBush, B.R. Powell, J.W. Jones, T.M. Pollock, Precipitation Strengthening of a Mg-Al-Ca–Based AXJ530 Die-cast Alloy, Metallurgical and Materials Transactions A 39(3) (2008) 696-702.
[13] Z. Zhang, R. Tremblay, D. Dubé, Microstructure and mechanical properties of ZA104 (0.3–0.6Ca) die-casting magnesium alloys, Materials Science and Engineering: A 385(1–2) (2004) 286-291.
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[14] F. Wang, T. Hu, Y. Zhang, W. Xiao, C. Ma, Effects of Al and Zn contents on the microstructure and mechanical properties of Mg-Al-Zn-Ca magnesium alloys, Materials Science and Engineering: A 704 (2017) 57-65.
[15] K.N. Braszczyńska-Malik, A. Grzybowska, Microstructure of Mg-5Al-0.4Mn-xRE (x = 3 and
M AN U
5 wt.%) alloys in as-cast conditions and after annealing, J. Alloys Compd. 663 (2016) 172-179. [16] W. Zhang, W. Xiao, F. Wang, C. Ma, Development of heat resistant Mg-Zn-Al-based magnesium alloys by addition of La and Ca: Microstructure and tensile properties, J. Alloys Compd. 684 (2016) 8-14.
[17] L. Zhang, K.-k. Deng, K.-b. Nie, F.-j. Xu, K. Su, W. Liang, Microstructures and mechanical properties of Mg–Al–Ca alloys affected by Ca/Al ratio, Materials Science and Engineering: A 636 (2015) 279-288.
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[18] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Solidification paths and eutectic intermetallic phases in Mg–Al–Ca ternary alloys, Acta Mater. 53(9) (2005) 2823-2834. [19] X. Lu, G. Zhao, J. Zhou, C. Zhang, L. Sun, Effect of extrusion speeds on the microstructure, texture and mechanical properties of high-speed extrudable MgZnSnMnCa alloy, Vacuum 157 (2018) 180-191.
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[20] H. Pan, G. Qin, M. Xu, H. Fu, Y. Ren, F. Pan, Z. Gao, C. Zhao, Q. Yang, J. She, B. Song, Enhancing mechanical properties of Mg–Sn alloys by combining addition of Ca and Zn, Materials & Design 83 (2015) 736-744.
[21] G. Cordier, E. Czech, H. Schäfer, P. Woll, Structural characterization of new ternary compounds of
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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44
uranium and cerium, Journal of the Less Common Metals 110(1) (1985) 327-330. [22] P.M. Jardim, G. Solórzano, J.B.V. Sande, Precipitate Crystal Structure Determination in Melt Spun Mg-1.5wt%Ca-6wt%Zn Alloy, Microsc. Microanal. 8(6) (2002) 487-496. [23] D.H. StJohn, M. Qian, M.A. Easton, P. Cao, Z. Hildebrand, Grain refinement of magnesium alloys, Metallurgical and Materials Transactions A 36(7) (2005) 1669-1679. [24] Y.C. Lee, A.K. Dahle, D.H. StJohn, The role of solute in grain refinement of magnesium, Metallurgical and Materials Transactions A 31(11) (2000) 2895-2906. [25] W. Xiao, Y. Shen, L. Wang, Y. Wu, Z. Cao, S. Jia, L. Wang, The influences of rare earth content on the microstructure and mechanical properties of Mg–7Zn–5Al alloy, Materials & Design 31(7) (2010) 3542-3549. [26] M.A. Easton, D.H. StJohn, A model of grain refinement incorporating alloy constitution and potency of heterogeneous nucleant particles, Acta Mater. 49(10) (2001) 1867-1878.
ACCEPTED MANUSCRIPT 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17
[27] G. Levi, S. Avraham, A. Zilberov, M. Bamberger, Solidification, solution treatment and age
18
[34] M. Nave, A. Dahle, D. StJohn, The Role of Zinc in the Eutectic Solidification of Magnesium‐
19
Aluminium‐Zinc Alloys, Magnesium Technology 2000
20
[35] M. Nave, A. Dahle, D. StJohn, Eutectic Growth Morphologies in Magnesium‐Aluminium Alloys,
21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41
Magnesium technology 2000
hardening of a Mg–1.6 wt.% Ca–3.2 wt.% Zn alloy, Acta Mater. 54(2) (2006) 523-530. [28] P. Yin, N.F. Li, T. Lei, L. Liu, C. Ouyang, Effects of Ca on microstructure, mechanical and corrosion properties and biocompatibility of Mg–Zn–Ca alloys, J. Mater. Sci. Mater. Med. 24(6) (2013) 1365-1373. [29] T. Rzychoń, Characterization of Mg-rich clusters in the C36 phase of the Mg–5Al–3Ca–0.7Sr– 0.2Mn alloy, J. Alloys Compd. 598 (2014) 95-105.
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[30] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D.H. StJohn, Development of the as-cast microstructure in magnesium–aluminium alloys, Journal of Light Metals 1(1) (2001) 61-72.
[31] Z.T. Li, X.D. Zhang, M.Y. Zheng, X.G. Qiao, K. Wu, C. Xu, S. Kamado, Effect of Ca/Al ratio on microstructure and mechanical properties of Mg-Al-Ca-Mn alloys, Materials Science and Engineering: A 682 (2017) 423-432.
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[32] S. Wasiur-Rahman, M. Medraj, Critical assessment and thermodynamic modeling of the binary Mg–Zn, Ca–Zn and ternary Mg–Ca–Zn systems, Intermetallics 17(10) (2009) 847-864. [33] Y. Zhong, J. Liu, R.A. Witt, Y.-h. Sohn, Z.-K. Liu, Al2(Mg,Ca) phases in Mg–Al–Ca ternary 573-576.
M AN U
system: First-principles prediction and experimental identification, Scripta Mater. 55(6) (2006)
(2000) 243-250.
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(2000) 233-242.
[36] M.S. Dargusch, M. Nave, S.D. McDonald, D.H. StJohn, The effect of aluminium content on the eutectic morphology of high pressure die cast magnesium–aluminium alloys, J. Alloys Compd. 492(1) (2010) L64-L68.
[37] M. Wang, D.H. Xiao, W.S. Liu, Effect of Si addition on microstructure and properties of
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magnesium alloys with high Al and Zn contents, Vacuum 141 (2017) 144-151. [38] C.H. Cáceres, C.J. Davidson, J.R. Griffiths, C.L. Newton, Effects of solidification rate and ageing on the microstructure and mechanical properties of AZ91 alloy, Materials Science and Engineering: A
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325(1–2) (2002) 344-355.
[39] T.L. Chia, M.A. Easton, S.M. Zhu, M.A. Gibson, N. Birbilis, J.F. Nie, The effect of alloy composition on the microstructure and tensile properties of binary Mg-rare earth alloys, Intermetallics 17(7) (2009) 481-490.
[40] D. Amberger, P. Eisenlohr, M. Göken, On the importance of a connected hard-phase skeleton for the creep resistance of Mg alloys, Acta Mater. 60(5) (2012) 2277-2289. [41] Q. Yang, K. Guan, F. Bu, Y. Zhang, X. Qiu, T. Zheng, X. Liu, J. Meng, Microstructures and tensile properties of a high-strength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy, Mater. Charact. [42] Q. Yang, K. Guan, F. Bu, Y. Zhang, X. Qiu, T. Zheng, X. Liu, J. Meng, Microstructures and tensile properties of a high-strength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy, Mater. Charact. 113 (2016) 180-188. [43] C.D. Lee, K.S. Shin, Effects of precipitate and dendrite arm spacing on tensile properties and fracture behavior of As-Cast magnesium-aluminum alloys, Metals and Materials International 9(1)
ACCEPTED MANUSCRIPT 1 2 3 4 5 6 7
(2003) 21-27. [44] M. Mabuchi, K. Higashi, Strengthening mechanisms of Mg-Si alloys, Acta Mater. 44(11) (1996) 4611-4618. [45] J.W. Luster, M. Thumann, R. Baumann, Mechanical properties of aluminium alloy 6061–Al2O3 composites, Mater. Sci. Technol. 9(10) (1993) 853-862. [46] D.M. Stefanescu, Science and engineering of casting solidification, third edition ed., Springer2015.
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ACCEPTED MANUSCRIPT Table captions: Table 1. Actual compositions of the studied alloys. Table 2. EDS results of intermetallic compounds in each alloy. Table 3. Endothermic peak temperatures for the DSC heating curves in Fig. 6.
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Table 4. Tensile properties of the studied alloys.
ACCEPTED MANUSCRIPT Table 1. Actual compositions of the studied alloys. Actual composition (wt. %) Alloy Zn
Ca
Ce
La
Pr
Nd
Si
Fe
Mn
Mg
AZX264
2.279
6.373
4.048
/
/
/
/
0.008
0.002
0.003
Bal.
AZX444
4.250
4.119
3.726
/
/
/
/
0.007
0.002
0.002
Bal.
AZX624
6.342
2.117
4.066
/
/
/
/
0.009
0.003
0.004
Bal.
AZXE2631
2.280
6.205
3.186
0.537
0.255
0.027
0.139
0.015
0.004
0.005
Bal.
AZXE4431
4.232
3.958
3.098
0.539
0.234
0.025
0.128
0.013
0.003
0.006
Bal.
AZXE6231
6.442
2.046
3.039
0.506
0.226
0.022
0.119
0.018
0.004
0.007
Bal.
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Al
ACCEPTED MANUSCRIPT Table 2. EDS results of intermetallic compounds in each alloy. Elemental composition (at. %) Arrow
Al
Zn
Ca
RE
Mg
Ca2Mg6Zn3
8.70
12.62
14.63
/
64.05
B
Al2REZn2
22.95
15.67
1.22
8.19
51.97
C
Ca2Mg6Zn3
3.86
19.05
11.85
/
65.24
D
C36-(Mg, Al)2Ca
13.74
1.78
E
C36-Al2(Mg, Ca)
25.60
0.66
F
Ca2Mg6Zn3
7.61
9.28
G
C36-Al2(Mg, Ca)
28.57
0.92
H
Al11RE3
36.15
I
Q
23.97
AZX444
AZX624
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A AZXE2631
10.76
/
73.72
8.34
/
65.40
5.72
/
77.39
8.51
/
62.00
2.24
1.05
8.80
51.76
9.18
3.17
/
63.68
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AZXE6231
Intermetallic compound
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Alloy
ACCEPTED MANUSCRIPT Table 3. Endothermic peak temperatures for the DSC heating curves in Fig. 6. Temperature of endothermic peak/ Alloy
Primary phase
Intermetallic compound Al11RE3/Al2REZn2
C36
C14
Ca2Mg6Zn3
AZXE2631
618
587
\
489
458
AZX444
593
\
500
\
445
AZXE6231
591
submerged
517
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α-Mg
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\
\
ACCEPTED MANUSCRIPT Table 4. Tensile properties of the studied alloys. YS/MPa
UTS/MPa
El./%
AZX264
103±6
103±6
0.2±0.0
AZX444
102±4
155±5
1.4±0.2
AZX624
95±3
172±2
2.1±0.2
AZXE2631
107±5
138±2
1.0±0.2
AZXE4431
90±2
139±4
1.1±0.1
AZXE6231
146±6
212±3
2.4±0.3
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Alloy
ACCEPTED MANUSCRIPT Figure captions: Fig. 1 XRD patterns of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZX624 and (e) AZXE6231 alloys. Fig. 2 Optical micrographs of (a) AZX264, (b) AZXE2613, (c) AZX444, (d) AZXE4431, (e)
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AZX624 and (f) AZXE6231 alloys. Fig. 3 Second dendrite arm spacing of the studied alloys.
Fig. 4 Grain size of the binary magnesium alloys with a range of Al, Zn and Ca alloying contents plotted against the growth restriction factor Q. The data is taken from [23, 24].
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Fig. 5 BEIs of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 6 DSC heating curves of (a) AZXE2631, (b) AZX444 and (c) AZXE6231 alloys. Fig. 7 Eutectic morphologies of (a) Ca2Mg6Zn3 and (b) C36 in AZX444 alloy, (c) Ca2Mg6Zn3 and C36 in AZX624 alloy and (d) C36 in AZXE6231 alloy.
Fig. 8 Tensile stress-strain curves of AZX264, AZXE2631, AZX444, AZXE4431, AZX624 and AZXE6231 alloys.
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Fig. 9 Column chart of the mechanical properties of the as-cast alloys.
Fig. 10 Fracture lateral surfaces of (a) AZX264, (b) AZEX2613, (c) AZX444 and (d) AZEX6213 alloys.
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Fig. 11 Volume fractions of the intermetallic compounds in each alloy. Fig. 12 Work-hardening rate-true strain curves of AZXE2631, AZX444 and AZXE6231 alloys.
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Fig. 13 Optical macrographs of (a) AZX264, (b) AZXE2631, (c) AZX444 and (d) AZEX6231 alloys.
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Fig. 1 XRD patterns of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZX624 and (e) AZXE6231
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alloys.
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Fig. 2 Optical micrographs of (a) AZX264, (b) AZXE2613, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 3 Second dendrite arm spacing of the studied alloys.
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Fig. 4 Grain size of the binary magnesium alloys with a range of Al, Zn and Ca alloying contents
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plotted against the growth restriction factor Q. The data is taken from [23, 24].
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Fig. 5 BEIs of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 6 DSC heating curves of (a) AZXE2631, (b) AZX444 and (c) AZXE6231 alloys.
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Fig. 7 Eutectic morphologies of (a) Ca2Mg6Zn3 and (b) C36 in AZX444 alloy, (c) Ca2Mg6Zn3 and
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C36 in AZX624 alloy and (d) C36 in AZXE6231 alloy.
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Fig. 8 Tensile stress-strain curves of AZX264, AZXE2631, AZX444, AZXE4431, AZX624 and
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Fig. 9 Column chart of the mechanical properties of the as-cast alloys.
b
c
d
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Fig. 10 Fracture lateral surfaces of (a) AZX264, (b) AZEX2613, (c) AZX444 and (d) AZEX6213
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Fig. 11 Volume fractions of the intermetallic compounds in each alloy.
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Fig. 12 Work-hardening rate-true strain curves of AZXE2631, AZX444 and AZXE6231 alloys.
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c
d
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Fig. 13 Optical macrographs of (a) AZX264, (b) AZXE2631, (c) AZX444 and (d) AZEX6231
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ACCEPTED MANUSCRIPT Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys were prepared by gravity casting, and the microstructures and mechanical properties were studied. The category, size, morphology and distribution of the intermetallic compounds were significantly influenced by alloying composition.
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C36 formed as a majority compound instead of Ca2Mg6Zn3 with the increase of Al content, which was much finer and more evenly distributed in high Al content alloy.
Replacing a part of Ca by RE could further refine the microstructure, thus
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Mg-6Al-2Zn-3Ca-1RE obtained the best mechanical property.
S0042-207X(18)31645-2
DOI:
https://doi.org/10.1016/j.vacuum.2018.10.072
Reference:
VAC 8345
To appear in:
Vacuum
Received Date: 22 August 2018 Revised Date:
24 October 2018
Accepted Date: 25 October 2018
Please cite this article as: Wang F, Xiao W, Liu M, Chen J, Li X, Xi J, Ma C, Effects of alloying composition on the microstructures and mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy, Vacuum (2018), doi: https://doi.org/10.1016/j.vacuum.2018.10.072. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT 1
Effects of alloying composition on the microstructures and
2
mechanical properties of Mg-Al-Zn-Ca-RE magnesium alloy
3
Fang Wang 1, Wenlong Xiao 1,2, *, Maowen Liu 1, Jing Chen 1, Xiang Li 1, Jiabing Xi 1, Chaoli Ma 1,2
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1
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of Materials Science and Engineering, Beihang University, Beijing 100191, China
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2
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Industry and Information Technology, Beihang University, Beijing, 100191, China
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Abstract: With the purpose of increasing our understanding of the microstructural evolution and
9
mechanical
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Key Laboratory of Aerospace Advanced Materials and Performance of Ministry of Education, School
properties
of
Mg-Al-Zn-Ca-RE
system
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Key Laboratory of High-Temperature Structural Materials & Coatings Technology, Ministry of
alloy,
Mg-xAl-(8-x)Zn-4Ca
and
Mg-xAl-(8-x)Zn-3Ca-1RE (wt. %) casting alloys were studied in this paper. The α-Mg dendrite
11
was refined, and the intermetallic compounds became finer and less divorced by increasing the Al
12
content. During solidification RE addition consumed Al and Zn to form RE-containing phase in
13
the initial solidification stage, and then the residual liquid solidified to form Ca-containing phase
14
without RE participation. The majority phases, Al11RE3 and C36, were transformed into Al2REZn2
15
and Ca2Mg6Zn3 as the Al content was decreasing. The Ca-containing phases tended to form a
16
highly interconnected network which would deteriorate the ductility, while the substitution of RE
17
for a part of Ca would make the structure less continuous. A range of eutectic morphologies were
18
found in Ca-containing phases. C36 phase could shape a lamellar eutectic structure, and
19
Ca2Mg6Zn3 phase was prone to exhibit as a divorced morphology. Besides, Ca2Mg6Zn3 formed at a
20
relatively low temperature which will bring about microshrinkage. Mg-6Al-2Zn-3Ca-1RE alloy
21
containing refined and more evenly distributed C36 with less continuous morphology showed the
22
best mechanical performance.
23
Key words: Magnesium alloy; Calcium; Rare earth; Microstructures; Intermetallic compound;
24
Mechanical properties.
25
*
26
School of Materials Science and Engineering, Beihang University, Beijing 100191, China. Tel:
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+86-10-8233 8631; Fax: +86-10-8233 8631
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E-mail: [email protected] (Wenlong Xiao)
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Corresponding author.
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1. Introduction Persistent focus towards lightweight in automotive industry stimulates the
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development of magnesium alloys. Among them, heat resistant Mg alloys are deemed
4
as a promising structural material in automobile engine. However, the well-developed
5
commercially Mg-Al based alloys, such as AZ91 and AM60, show poor heat
6
resistance due to the rapid dynamic precipitation of β-Mg17Al12 phase at high
7
temperatures [1]. There is a general consensus that the most effective way to improve
8
heat resistance of Mg-Al based alloy is to acquire thermal stable intermetallic
9
compounds by alloying the alloy with Zn, Ca, Si, Sr and rare earth (RE), etc [1-4]. In
10
recent years, Mg-Al-Ca and Mg-Zn-Ca magnesium alloys have been attracted
11
numerous studies due to their excellent mechanical properties at both room and
12
elevated temperature, low price, texture weakening and good ignition resistance [5-8].
13
However, calcium generally brings worse die sticking and hot tearing. It has been
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reported that zinc addition would restore the castability of Mg-Al alloys, and RE
15
addition would ameliorate hot tearing susceptibility of Mg-Zn-Al alloys by shortening
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the solidification freeze range [9-11].
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The formation of Ca-containing phases and RE-containing phases can suppress
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the formation of low melting point phases such as Mg-Al, Mg-Zn and Mg-Zn-Al
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phases in Mg-Al-Ca, Mg-Zn-Ca, Mg-Al-RE, Mg-Al-Zn-Ca and Mg-Al-Ca-RE system
20
alloys, which are capable to obtain alloys with excellent high temperature properties
21
[6, 12-16]. Mg-Al-Zn-Ca-RE alloy containing all the mentioned above system alloys
22
is therefore considered as a promising heat resistant magnesium alloy system with
23
high performance. In our previous work, the microstructures and mechanical
24
properties of Mg-Al-Zn-Ca-La alloys at fixed Al and Zn contents has been
25
investigated, Al2LaZn2 and Ca2Mg6Zn3 would form as major intermetallic compounds
26
[16]. However, the categories of Ca-containing phase in Mg-Al-Ca alloys are largely
27
depend on Al/Ca ratio, and Zn addition would also have a significant effect on the
28
phase constitution of Mg-Al-Zn-Ca alloy [5, 14, 17, 18]. Our recently study reveals
29
that Al and Zn contents have a great influence on the microstructures and mechanical
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ACCEPTED MANUSCRIPT properties of Mg-Al-Zn-Ca quaternary alloy [14]. Two Laves phases, i.e. C36 and
2
C15, a quaternary Q phase and/or Ca2Mg6Zn3 isomorphs phase are observed in the
3
casting Mg-xAl-(8-x)Zn-2Ca alloys with different Al and Zn contents, and the alloy
4
containing the majority phase of C36 exhibits the best combination of strength and
5
ductility. Moreover, the category of second phase transforms from C36 to C15 and
6
Ca2Mg6Zn3 phase by decreasing the Al/Zn ratio at a constant Ca content. As Q phase
7
is harmful to high temperature properties, we consider that the way of increasing Ca
8
content and/or alloying with RE might eliminate such phase [14]. In addition, it has
9
been reported that magnesium alloys containing thermally stable phases with proper
10
morphology and size as well as distribution are beneficial for the desire of high
11
mechanical performance [19, 20]. In that case, the aims of this work are to investigate
12
the effects of Al and Zn contents on the microstructure and mechanical properties of
13
high Ca-containing Mg-Al-Zn-Ca alloys, and the role of a small amount of RE
14
substitution for Ca in the microstructural evolution and mechanical properties. For
15
these purposes, a series of Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys
16
are studied, and the relationship between the microstructure and mechanical
17
properties are discussed.
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2. Experimental procedure The
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actual
compositions
of
three
Mg-xAl-(8-x)Zn-4Ca
and
three
Mg-xAl-(8-x)Zn-3Ca-1RE alloys (weight percent, wt. %, also for hereafter not
22
mentioned) are listed in Table 1, and the alloys were prepared from pure (≥99.9%) Mg,
23
Al and Zn elements and Mg-20Ca and Mg-25RE (cerium rich misch metal) master
24
alloys by electrical resistance melting furnace. Casting process was performed under
25
the protection of an anti-oxidizing flux. The melt was refined by BaCl2 refining agent
26
and then holding at 740°C for 20min to purify the melt before pouring into a steel
27
mold at room temperature. It is noted that Mg-25RE master alloy was prepared in
28
vacuum melting furnace (10-3 Pa) because of the poor ignition resistance. The
29
obtained as-cast alloys were designated by AZX(E) followed with three or four
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Table 1. The microstructure was observed by optical microscopy (OM) and scanning
3
electron microscopy (SEM) equipped with energy dispersive spectrometry (EDS)
4
operated under a vacuum condition of 10-5 Pa. The intermetallic compounds were
5
identified by a RIGAKU RINT-2000 X-ray diffractometer (XRD) with Cu Kα
6
radiation. Differential scanning calorimetry (DSC) was conducted for analyzing the
7
solidification behavior with a heating rate of 10 K/min under Ar gas atmosphere
8
protection. The flat dog-bone like tensile test plates with gauge length of 15 mm,
9
width of 5.0 mm and thickness of 1.5 mm were machined from the same position of
10
the ingots, and tensile tests were performed three times on Instron 8801 universal
11
testing machine at room temperature (RT) according to Chinese metallic
12
materials-tensile testing standard (GB/T228-2002). The fracture surfaces were
13
observed using OM.
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3. Results and Discussion
16
3.1 Microstructure
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3.1.1 Phase constitution
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Fig. 1 exhibits the XRD patterns of the studied alloys. It is illustrated in Fig. 1a
19
that AZX264 alloy is composed of α-Mg, Ca2Mg6Zn3 and C14-Mg2Ca phases.
20
Replacing Ca with RE causes the appearance of new peaks in AZXE2631 alloy (Fig.
21
1b). The new peaks belong to Al2REZn2 phase and are identified on the basis of
22
Al2CeZn2 phase (PDF#29-0013: I4/mmm space group, a = 0.4244nm and c =
23
1.0986nm) [21]. Additionally, unknown peaks are detected in both AZX264 and
24
AZXE2631 alloys, they might belong to Ca2Mg6Zn3 in a different crystallographic
25
form [22]. Compared with Fig. 1a, the increase of Al content leads to a decrease in the
26
amount of Ca2Mg6Zn3 phase as shown in Fig. 1c (AZX444 alloy), and the peaks of
27
C36 phase emerge in this alloy instead of that of C14 phase (Fig. 1c). The peaks
28
belong to C36 are consistent with our previous work [14]. Similarity with AZXE2631
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2
shows no difference with the result in our previous study [16]. When the Al content is
3
further increased, though the peaks belonging to C36 phase and Ca2Mg6Zn3 phase are
4
still remained in AZX624 alloy (Fig. 1d), the amount of C36 phase is dramatically
5
increased. By comparison, Al11RE3 phase is formed in AZXE6231 alloy rather than
6
Al2REZn2 in AZXE4431 alloy, which is confirmed in Fig.1e based on Al11Ce3 phase
7
(PDF#19-0006: Immm space group, a = 0.4395 nm, b = 1.0092 nm, c = 1.3011 nm).
8
3.1.2 Observation of primary α-Mg phase
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Fig. 2 shows the optical micrographs of the as-cast alloys, where the typical
10
α-Mg dendrites surrounded by second phases can be found in each alloy. It is evident
11
that the secondary dendrite arm spacing (SDAS) is changed by varying the alloying
12
composition. The SDAS of each alloy is present in Fig. 3, which was measured by
13
averaging over the distance between the adjacent dendrite arm centers. There were at
14
least 60 secondary dendrite arms being counted in different optical micrographs for
15
each alloy. AZX264 alloy shows the largest SDAS among the studied Mg-Al-Zn-Ca
16
alloys as shown in Fig. 3, and the SDAS decreases obviously and then increases
17
slightly with the increase of Al content so that AZX444 alloy has the smallest SDAS.
18
When RE is substituted for a part of the Ca content, the dendrite arms become more
19
developed and obtain a tiny refinement as compared with the alloy containing the
20
same Al and Zn contents. As shown in Fig. 2, the majority of second phases are
21
formed along the dendritic boundaries and distributed as a highly continuous
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interconnected morphology, especially for the high Zn-containing alloys.
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Grain refinement effect generally involves in the addition of nucleants and/or
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alloying elements into a melt before casting. The additions of Al, Zn, Ca and RE to
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Mg alloys have little effect on nucleation of the primary α-Mg phase since these
26
elements are mostly segregated to form second phases well after the nucleation.
27
However, the solute elements play a critical role in controlling the growth of the
28
nucleated grains and in subsequent nucleation during solidification, which can be
29
explained in terms of the growth restriction factor (GRF) [23]. It demonstrates that the
ACCEPTED MANUSCRIPT solute tends to segregate on the ahead of growing dendrite/grain and thus generates
2
the constitutional undercooling (△Tc). This constitutional undercooling not only
3
suppresses the growth of dendrite/grain, but also promotes nucleation when △Tc
4
reaches the undercooling required for nucleation (△Tn). The GRF value, well known
5
as Q, is widely used to describe the effect of solutes on dendrite arm spacing/ grain
6
size, which can be expressed as
7
=∑
,
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−1
,
(1)
where ml,i is the slope of the liquidus line, c0,i is the concentration of the solute, and ki
9
is the partition coefficient for the element i in the binary alloy. The relationship
10
between grain size and Q for magnesium alloy binary alloyed with Al, Zn and Ca is
11
well illustrated in [23, 24], which is plotted in Fig. 4 showing alloying compositions
12
calculated by equation (1). It shows a distinct decrease in grain size at a minor
13
addition of either Zn or Ca while only a minimal reduction at a further addition of that
14
exceeding the saturation level. Obviously, the addition level of Al for grain refinement
15
is much higher than that of Zn and Ca. An evident grain/dendrite refinement is
16
observed in AZX444 and AZXE4431 alloys compared to AZX264 and AZX2631
17
alloys. Since the additions of Zn and Ca in these alloys are larger than the saturation
18
levels and that of Al is less than the level (Fig. 4), the increased Al content should
19
account for the refinement effect in the SDAS [23, 24]. When the Al content increases
20
from 4 wt% to 6 wt%, it can be inferred from Fig. 4 that the refinement effect should
21
still remain. However, only a slightly increase in SDAS can be found in Fig. 3 as
22
compared AZX444 alloy to AZX624 alloy and AZX4431 alloy to AZXE6231 alloy. In
23
high Al content alloys, the Mg-Al-Ca phase containing a certain amount of Zn would
24
form (section 3.2 in detail), and the contribution of Zn to GRF should be decreased in
25
AZX624 and AZX6231 alloys. In that case, the counteracting effect of increasing Al
26
solute and decreasing Zn solute makes the SDAS exhibits slightly change when the Al
27
content is further increased.
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Xiao, et al deduced that Mg-Ce binary alloy with a low concentration of Ce will
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cause a dramatic effect on grain refinement by using a relative grain size (RGS)
ACCEPTED MANUSCRIPT model proposed by Easton and StJohn [25, 26]. As with similar instance of
2
Mg-Al-Zn-RE alloy, the RE elements in Mg-Al-Zn-Ca-RE alloy can easily segregate
3
on the solid/liquid interface in the initial stage of solidification, which produces
4
barriers between the liquid and solid phases and thus inhibits the diffusion of Al, Zn
5
and Ca into the solid phase, leading to the grain/dendrite refinement by raising the
6
△Tc. On the other hand, the formed RE-containing phase consumes a large number of
7
Al and Zn solutes (section 3.2 in detail), which lowers △Tc. The two factors canceled
8
each other out so that there are only a slightly decrease of SDAS in the studied alloys
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with RE addition.
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3.1.3 Identification of intermetallic compounds
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The back-scattered electron images (BEIs) of the studied alloys are shown in Fig.
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5, and there are apparent contrasts among the intermetallic compounds in the studied
14
alloys. In general, different phases can be distinguished according to the contrasts in a
15
BEI. However, it is hardly probable to differentiate C14-Mg2Ca phase and Ca2Mg6Zn3
16
phase by comparing the contrasts in Mg-Zn-Ca alloys [27, 28]. The categories of the
17
second phase with different contrasts are further examined by EDS, and the elemental
18
compositions of each phase, pointed out by arrows A-J in Fig. 5, are listed in Table 2.
19
Fig. 5 shows that Ca-containing phases tend to be a highly interconnected
20
microstructure in the studied alloys. The majority phase, Ca2Mg6Zn3 in AZX264 alloy,
21
is transformed into C36 phase with a gray contrast by increasing the Al content.
22
Whilst in AZX444 and AZX624 alloys, Ca2Mg6Zn3 becomes the minor phase. As
23
listed in Table 2, Al and Zn elements are contained in Ca2Mg6Zn3 phase and C36
24
phase, respectively. It should be noticed that the composition of C36 phase is different
25
between AZX444 and AZX624 alloys, which can be identified as C36-(Mg, Al)2Ca
26
and C36-Al2(Mg, Ca), respectively, according to calculating the valence electron
27
concentrations [14, 29]. A high content of Al atom is consumed by the formation of
28
C36-Al2(Mg, Ca), which can also be found in AZX622 and AZX442 alloys published
29
in [14]. But instead of forming Q phase as the low melting phase in AZX622 and
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AZX442 alloys, the Ca2Mg6Zn3 phase is formed in AZX624 and AZX444 alloys
2
owing to the increased Ca content. As substituting Ca with RE, RE-containing phase with the brightest contrast is
4
visible in the BEIs and locate between the Ca-containing phase, which leads to a less
5
continuous morphology of the Ca-containing compounds (Fig. 5b, d and f). There are
6
no RE elements detected in Ca-containing phases according to the EDS results (Table
7
2), and the RE addition does not change the type of Ca-containing phase except for
8
AZXE6231 alloy. In AZXE6231, Al11RE3 phase forms instead of Al2REZn2 phase and
9
a few Zn atoms are dissolved in this RE-phase, while Ca2Mg6Zn3 phase which is
10
present in AZX624 cannot be detected. As shown in Fig. 5f, a small amount of
11
Mg-Al-Zn-Ca quaternary Q phase displaying a medium contrast (arrow I) can be
12
identified in this alloy [14]. The amount of Q phase is far less than that in AZX622
13
alloy. It seems that adding RE or increasing Ca content can both suppress the
14
formation of Q phase.
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3.2 Phase formation character during solidification
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The solidification event, to some extent, determines the size, morphology and
18
distribution of second phase and casting defect formation, which in turn to have an
19
impact on mechanical properties [30]. In this section, DSC heating curves of the
20
representative studied alloys are used to speculate the phase formation during
21
solidification (Fig. 6). As can be seen, each endothermic peak denotes the dissolution
22
of the phase labelled [14, 16]. Table 3 lists the peak temperature of each
23
transformation. For AZXE2631 alloy, the endothermic peaks corresponding to the
24
melting of α-Mg, Al2REZn2, C14 and Ca2Mg6Zn3 are occurred at approximately
25
618°C, 587°C, 489°C and 458°C, respectively. The peaks are shifted when changes
26
are made to the Al and Zn concentrations (AZXE6231 alloy). The melting peak of
27
α-Mg is decreased to 591°C, and the peaks related to C14 and Ca2Mg6Zn3 are
28
disappeared, while a peak that refers to C36 can be detected. It should be mentioned
29
in this alloy that the peak of Al11RE3 is hard to distinguish due to the peak is covered
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by α-Mg peak. As for AZX444 alloy, except for the melting peak of α-Mg (593°C),
2
two endothermic peaks at approximately 500°C and 445°C are considered to be phase
3
transformation of C36 and Ca2Mg6Zn3, respectively. As mentioned above, it can be known that the formation of RE-containing phase
5
follows on the heels of the nucleation and growth of primary α-Mg in the initial stage
6
of solidification, and which type would be formed is determined by the Al and Zn
7
contents, i.e. Al2REZn2 will transform into Al11RE3 with increasing the Al content.
8
Afterward, the solidification of residual liquid will comply with the solidification
9
event of Mg-Al-Zn-Ca quaternary system. A part of Al and Zn are consumed by the
10
formation of RE-containing phase, and then the Ca/Al and Ca/Zn ratios are changed
11
in the residual liquid, thereby probably changing the type of Ca-containing phase [14,
12
17, 31, 32]. As a summary, Ca2Mg6Zn3 and C14 would transform into C36 with
13
increasing the Al content, which is consistent with our previous work [14]. As listed in
14
Table 2, the composition of C36 varies in alloys with different Al contents. It might be
15
ascribed to that C36-(Mg, Al)2Ca phase follows the Al2Ca–Mg2Ca pseudo-binary
16
system, while C36-Al2(Mg, Ca) phase complies with the Al2Mg–Al2Ca pseudo-binary
17
system [18, 33].
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As shown in Fig. 5, the morphology of eutectic phases varies in the studied
19
alloys. Owing to the non-equilibrium solidification, both Ca2Mg6Zn3 and Mg-Al-Ca
20
phases form either a typical fully divorced or a partially divorced morphology and
21
exhibits a highly interconnected network. With increasing the Al content and
22
meanwhile decreasing the Zn content for the alloy, an increasing proportion of less
23
divorced morphology and even typical eutectic morphology is present. Fig. 7 shows
24
higher magnification SEM images for a close inspection. As can be seen from Fig.
25
7a-c, a high proportion of granular and fibrous eutectic α-Mg is found, and the
26
Ca2Mg6Zn3 phase becomes less divorced with decreasing the Zn content. This may be
27
attributed to that Zn has a strongly partition effect to eutectic growth [34, 35]. As can
28
be seen in Fig. 7b-d, the C36 phase tends to display as a much less divorced and even
29
lamellar eutectic morphology. It has been reported that Al does not partition as
30
strongly as an equivalent amount of Zn during eutectic phase growth [30]. In addition,
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ACCEPTED MANUSCRIPT primary α-Mg also has a significant impact on eutectic morphology [36]. The
2
relatively low eutectic temperature of Ca2Mg6Zn3 means a wide solidification range
3
as well as a high solid fraction of primary α-Mg during eutectic solidification, which
4
leads to narrow spaces for the eutectic growth and therefore forms more divorced
5
morphology. With increasing the Al content and/or substituting RE for a part amount
6
of Ca, the α-Mg dendrites become much developed (Fig. 2) and the main eutectic
7
reaction occurs at a relatively high temperature. As shown in Fig. 5f, a smaller size,
8
less divorced and more evenly distribution of C36 are thus obtained in AZXE6231
9
alloy.
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3.3 The relationship between microstructures and mechanical properties
12
3.3.1 Tensile properties
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The typical tensile stress-strain curves of the studied alloys are presented in Fig.
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8, and the average yield strength (YS), ultimate tensile strength (UTS) and elongation
15
(El.) are listed in Table 4. Furthermore, column chart with error bars has been used in
16
Figure 9 to illustrate the effects of alloying compositions. It can be found that the
17
mechanical performance varies with the different alloying compositions. As can be
18
seen in Fig. 8, AZX264 alloy fractured immediately after yielding, showing a limited
19
plastic deformation. Figure 9 shows that the UTS and El. are improved considerably
20
with the increase of the Al content, but the YS has a negligible change when the Al
21
content increases from 2 wt% to 4 wt%. Moreover, the YS is decreased slightly from
22
102±4 MPa of AZX442 alloy to 95±3 MPa of AZX624 alloy. For the RE-containing
23
alloy, the YS firstly decreases from 107±5 MPa to 90±2 MPa and then increases
24
significantly to 146±6 MPa with increasing the Al content. However, the UTS and El.
25
remain nearly constant firstly and, finally, increase substantially. It is noted that the
26
substitution of RE for a part of Ca exhibits different impacts on tensile properties,
27
which depends on Al and Zn concentrations. As compared AZXE2631 alloy with
28
AZX264 alloy, the UTS and El. are remarkably improved, while the YS almost
29
remains unchanged. When the Al content is increased to 4 wt%, the RE addition can
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ACCEPTED MANUSCRIPT result in the decrease of both strength and ductility. However, the mechanical
2
properties are substantially enhanced by replacing a part of Ca with RE when the Al
3
content is further increased to 6 wt%. Among the studied alloys, AZXE6231 alloy has
4
the most optimal strength and ductility, in which the YS, UTS and El. are 146±6 MPa,
5
212±3 MPa and 2.4±0.3%, respectively.
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Fig. 10 displays the optical micrographs of longitudinal section adjacent to the
7
fracture surface. It is clearly seen that the cracks mostly develop along the interface
8
between α-Mg matrix and eutectic compounds where stress concentration is generated,
9
and cracks can readily propagate along the compounds particularly with a divorced
10
morphology, thus deteriorating the ductility. Compared to C36, Ca2Mg6Zn3 in the
11
studied alloys shows a much higher interconnected network and more divorced and
12
coarse morphology, which causes stress concentration and split of α-Mg matrix more
13
readily [37]. As a result, high Zn-containing AZX264 and AZXE2613 alloys with high
14
volume fraction of Ca2Mg6Zn3 exhibit worse ductility. With respect to the less
15
divorced and even lamellar eutectic phase, i.e. C36 phase, the cracks develop not only
16
along the phase/dendrite boundary but also into the eutectic phases (Fig. 10d),
17
resulting in acceptable ductility. Thus, AZX624 and AZXE6231 containing C36 as the
18
majority phase with both refined divorced and lamellar morphologies exhibit
19
relatively high ductility, which is in a good agreement with our previous work [14].
20
Furthermore, RE addition led to the formation of RE-phase in the initial stage of
21
solidification as mentioned above. The RE-phase would occupy between the α-Mg
22
dendrites where the nucleation of Ca-containing phase would then take place. As
23
solidification continued, the growth direction of Ca-containing phase could be
24
restricted by the pre-existed RE-phase. The subsequently formed Ca-containing phase
25
could not link each other as highly as that in the counterpart RE-free alloy. On the
26
other hand, when 1 wt% Ca is replaced by RE, the volume fraction of Ca-containing
27
phase formed is decreased. Consequently, the Ca-containing phases in RE-containing
28
alloy are somehow less continuous (Figure 5), thus the ductility of Mg-Al-Zn-Ca
29
alloy can be slightly improved by RE addition.
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3.3.2 Strengthening mechanism In general, the yield strength is mainly related to SDAS and intermetallic
3
compound for Mg casting alloys. Based on the well-known Hall-Petch relationship,
4
smaller SDAS has more grain boundaries which can improve the strength. However,
5
AZX444 and AZXE4431 alloys have relatively small SDAS but low yield strength.
6
Additionally, It has been reported previously that the strength would monotonic
7
increase with increasing the volume fraction of intermetallic compounds in Mg-Al
8
and Mg-RE binary alloys [38, 39]. Fig. 11 shows the volume fractions of the
9
intermetallic compounds for the studied alloys. It is measured to be 9.1 vol%, 9.5
10
vol %, 9.6 vol %, 8.0 vol %, 7.2 vol %, and 7.8 vol. % for AZX264，AZX444,
11
AZX624, AZXE2631, AZXE4431 and AZXE6231 alloys, respectively. There are not
12
too much differences between either the three Mg-Al-Zn-Ca alloys or the other three
13
Mg-Al-Zn-Ca-RE alloys, but the yield strength varies. It seems that other factors
14
should be responsible for the strengthening mechanism.
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It is figured out that highly interconnected and strong intermetallic skeleton can
16
form robust encasements for α-Mg cells, thus improving strength [40, 41]. In this
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study, the high Zn containing alloys have more continuous skeleton, and the yield
18
strength of Mg-Al-Zn-Ca alloy does increase with increasing the Zn content. While
19
for the RE-containing alloys, AZXE6231 alloy shows the highest yield strength
20
although less continuous skeleton is formed. Apart from the above factors, the size
21
and distribution of intermetallic compound also play critical roles in determining the
22
strength, and the well-developed models for interpreting the strengthening mechanism
23
also reveal such importance [38, 42-45]. It can be found in Fig. 2 and Fig. 5 that the
24
size and distribution of the intermetallic compounds in AZXE6231 alloy are
25
remarkably distinguished from other alloys. AZXE6231 alloy has the finest and most
26
evenly distributed intermetallic compounds among the studied alloys. During tensile
27
test, the more evenly distributed C36 and the relative small α-Mg dendrites not only
28
can inhibit the grain boundary slippage and dislocation motion, but can transfer the
29
applied load with the help of good α-Mg reinforcement interfacial integrity [37].
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ACCEPTED MANUSCRIPT Moreover, considering that the bulk modulus of RE phase is much higher than that of
2
Ca phase, the yield strength of RE-containing alloy should be higher as compared
3
with the RE free alloy at the same Al and Zn contents [42]. Nevertheless, AZXE4431
4
alloy shows lower yield strength than AZX444 alloy, and the reason need to be further
5
investigated. In summary, it can be deduced that the highest strength of AZXE6231
6
alloy is mainly attributed to the dispersion strengthening as a result of the refined and
7
more evenly distributed intermetallic compounds caused by the refined dendrites and
8
the less divorced eutectic morphology [36, 43].
Fig. 12 depicts the work-hardening rate (Θ) and true strain curves of AZXE2631,
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AZX444 and AZXE6231 alloys, and Θ equals to
, where σ is true stress, ε is
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true strain and
12
beginning of plastic deformation and then decreased gradually until fracture for all the
13
alloys. AZX6231 alloy with the majority of C36 phase displays the highest work
14
hardening rate, while AZX2631 alloy with the majority of Ca2Mg6Zn3 phase has the
15
lowest one, and AZX444 alloy with the both phases owns the moderate one, which is
16
in a good agreement with our previous work [14]. Combined with the obtained better
17
ductility resulting from the less continuous structure, AZXE6231 alloy therefore
18
shows the highest UTS among the studied alloys. However, the failure still occurs in
19
the stage of uniform deformation during the tensile tests.
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is strain rate. The work hardening rates rapidly decreased at the
The last but not the least, casting defects of these alloys are considered as well.
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Fig. 13 displays optical macrographs of the studied alloys. It is obvious in Fig. 13a
22
and b that numerous micro-shrinkages are present in AZX264 alloy, therefore, it has
23
the worst mechanical properties among the studied alloys, and the failure strength of
24
the samples are even occurred at around the 0.2% yield strength. Either increasing Al
25
content or RE addition can suppress the micro-shrinkage as can be detected in Fig.
26
13b and c. The formation of micro-shrinkage is involved with the final stage of the
27
solidification process for eutectic growth. The formation of divorced morphology
28
which need independent nucleation and growth will provide a much greater resistance
29
to feeding [30, 46]. Additionally, a low eutectic point can extend the solidification
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2
containing divorced Ca2Mg6Zn3 as the majority phase has the worst feedability.
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Increasing Al content and/or RE addition can reduce and even suppress the formation
4
of Ca2Mg6Zn3 as well as modify the eutectic morphology. In that case, it is expected
5
that AZXE6231 has the best castability, so that micro-shrinkage can hardly be
6
observed (Fig. 13d).
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4. Conclusions
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Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys were prepared by
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gravity casting to examine the microstructure and the resultant mechanical
11
performance, which could give a guidance for alloy design of this five-element
12
system heat resistance magnesium alloy. The main conclusions could be drawn
13
hereafter.
14
(1) The α-Mg dendrites became more developed and the second dendrite arm spacing
15
decreased evidently with increasing the Al content in Mg-xAl-(8-x)Zn-4Ca alloy,
16
and substituting 1wt% RE for the equivalent amount of Ca enabled the dendrites
17
slightly refined. The intermetallic compounds were located at the dendrite
18
boundaries, which became much finer and more evenly distributed as the
19
dendrites were refined.
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(2) RE addition consumed the Al and Zn contents to form Al11RE3 and Al2REZn2 in
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the initial stage of solidification at high Al alloy and high Zn alloy, respectively,
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and it did not take part in the subsequently formation of Ca-containing phases which would comply with the solidification event of Mg-Al-Zn-Ca quaternary alloy system. In the quaternary alloy, the majority phase transformed from Ca2Mg6Zn3 into C36 with the increase of the Al content.
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(3) Ca-containing phases tended to form a highly interconnected continuous network
27
which became less continuous because of RE addition. Among the Ca-containing
28
phases, Ca2Mg6Zn3 phase had a more divorced morphology and solidified at a
29
relatively low temperature which would cause microshrinkage, while C36 phase
ACCEPTED MANUSCRIPT 1
formed at a higher temperature owned a much less divorced and even lamellar
2
eutectic morphology and a better castability. (4) The size, distribution and category of intermetallic compound played important
4
roles in determining the mechanical properties. Mg-6Al-2Zn-3Ca-1RE alloy
5
showed the best mechanical performance among the studied alloys due to the less
6
continuous, much refined and more evenly distributed C36 phase as well as the
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obtained soundness casting.
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Acknowledgements
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The authors are grateful to the financial support by National Key Research and
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Development Program of China (No. 2016YFB0301103) and National Natural Science
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Foundation of China (NSFC, No. 51401010, 51671007 and 51671012). I (Fang Wang)
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want to thank the corresponding author and prof. Ma in particular for their patient
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guidance and help.
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References
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[1] A.A. Luo, Recent magnesium alloy development for elevated temperature applications, Int. Mater. Rev. 49(1) (2004) 13-30. 197-217.
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[2] M. Pekguleryuz, M. Celikin, Creep resistance in magnesium alloys, Int. Mater. Rev. 55(4) (2010) [3] S. Zhu, M.A. Easton, T.B. Abbott, J.-F. Nie, M.S. Dargusch, N. Hort, M.A. Gibson, Evaluation of
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Magnesium Die-Casting Alloys for Elevated Temperature Applications: Microstructure, Tensile Properties, and Creep Resistance, Metallurgical and Materials Transactions A 46(8) (2015) 3543-3554. [4] W. Xiao, M.A. Easton, M.S. Dargusch, S. Zhu, M.A. Gibson, The influence of Zn additions on the microstructure and creep resistance of high pressure die cast magnesium alloy AE44, Materials Science and Engineering: A 539 (2012) 177-184. [5] R. Ninomiya, T. Ojiro, K. Kubota, Improved heat resistance of Mg-Al alloys by the Ca addition, Acta Metallurgica et Materialia 43(2) (1995) 669-674. [6] A.A. Luo, B.R. Powell, M.P. Balogh, Creep and microstructure of magnesium-aluminum-calcium based alloys, Metallurgical and Materials Transactions A 33(3) (2002) 567-574. [7] Y. Du, M. Zheng, B. Jiang, Comparison of microstructure and mechanical properties of Mg-Zn microalloyed with Ca or Ce, Vacuum 151 (2018) 221-225. [8] M. Sakamoto, S. Akiyama, K. Ogi, Suppression of ignition and burning of molten Mg alloys by Ca bearing stable oxide film, J. Mater. Sci. Lett. 16(12) (1997) 1048-1050.
ACCEPTED MANUSCRIPT [9] I.A. Anyanwu, Y. Gokan, S. Nozawa, A. Suzuki, S. Kamado, Y. Kojima, S. Takeda, T. Ishida, Development of New Die-castable Mg-Zn-Al-Ca-RE Alloys for High Temperature Applications, MATERIALS TRANSACTIONS 44(4) (2003) 562-570. [10] W. Xiao, S. Zhu, M.A. Easton, M.S. Dargusch, M.A. Gibson, J.-f. Nie, Microstructural characterization of high pressure die cast Mg-Zn-Al-RE alloys, Mater. Charact. 65 (2012) 28-36. [11] J. Song, F. Pan, B. Jiang, A. Atrens, M.-X. Zhang, Y. Lu, A review on hot tearing of magnesium alloys, Journal of Magnesium and Alloys 4(3) (2016) 151-172.
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[12] A. Suzuki, N.D. Saddock, J.R. TerBush, B.R. Powell, J.W. Jones, T.M. Pollock, Precipitation Strengthening of a Mg-Al-Ca–Based AXJ530 Die-cast Alloy, Metallurgical and Materials Transactions A 39(3) (2008) 696-702.
[13] Z. Zhang, R. Tremblay, D. Dubé, Microstructure and mechanical properties of ZA104 (0.3–0.6Ca) die-casting magnesium alloys, Materials Science and Engineering: A 385(1–2) (2004) 286-291.
SC
[14] F. Wang, T. Hu, Y. Zhang, W. Xiao, C. Ma, Effects of Al and Zn contents on the microstructure and mechanical properties of Mg-Al-Zn-Ca magnesium alloys, Materials Science and Engineering: A 704 (2017) 57-65.
[15] K.N. Braszczyńska-Malik, A. Grzybowska, Microstructure of Mg-5Al-0.4Mn-xRE (x = 3 and
M AN U
5 wt.%) alloys in as-cast conditions and after annealing, J. Alloys Compd. 663 (2016) 172-179. [16] W. Zhang, W. Xiao, F. Wang, C. Ma, Development of heat resistant Mg-Zn-Al-based magnesium alloys by addition of La and Ca: Microstructure and tensile properties, J. Alloys Compd. 684 (2016) 8-14.
[17] L. Zhang, K.-k. Deng, K.-b. Nie, F.-j. Xu, K. Su, W. Liang, Microstructures and mechanical properties of Mg–Al–Ca alloys affected by Ca/Al ratio, Materials Science and Engineering: A 636 (2015) 279-288.
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[18] A. Suzuki, N.D. Saddock, J.W. Jones, T.M. Pollock, Solidification paths and eutectic intermetallic phases in Mg–Al–Ca ternary alloys, Acta Mater. 53(9) (2005) 2823-2834. [19] X. Lu, G. Zhao, J. Zhou, C. Zhang, L. Sun, Effect of extrusion speeds on the microstructure, texture and mechanical properties of high-speed extrudable MgZnSnMnCa alloy, Vacuum 157 (2018) 180-191.
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[20] H. Pan, G. Qin, M. Xu, H. Fu, Y. Ren, F. Pan, Z. Gao, C. Zhao, Q. Yang, J. She, B. Song, Enhancing mechanical properties of Mg–Sn alloys by combining addition of Ca and Zn, Materials & Design 83 (2015) 736-744.
[21] G. Cordier, E. Czech, H. Schäfer, P. Woll, Structural characterization of new ternary compounds of
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1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44
uranium and cerium, Journal of the Less Common Metals 110(1) (1985) 327-330. [22] P.M. Jardim, G. Solórzano, J.B.V. Sande, Precipitate Crystal Structure Determination in Melt Spun Mg-1.5wt%Ca-6wt%Zn Alloy, Microsc. Microanal. 8(6) (2002) 487-496. [23] D.H. StJohn, M. Qian, M.A. Easton, P. Cao, Z. Hildebrand, Grain refinement of magnesium alloys, Metallurgical and Materials Transactions A 36(7) (2005) 1669-1679. [24] Y.C. Lee, A.K. Dahle, D.H. StJohn, The role of solute in grain refinement of magnesium, Metallurgical and Materials Transactions A 31(11) (2000) 2895-2906. [25] W. Xiao, Y. Shen, L. Wang, Y. Wu, Z. Cao, S. Jia, L. Wang, The influences of rare earth content on the microstructure and mechanical properties of Mg–7Zn–5Al alloy, Materials & Design 31(7) (2010) 3542-3549. [26] M.A. Easton, D.H. StJohn, A model of grain refinement incorporating alloy constitution and potency of heterogeneous nucleant particles, Acta Mater. 49(10) (2001) 1867-1878.
ACCEPTED MANUSCRIPT 1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17
[27] G. Levi, S. Avraham, A. Zilberov, M. Bamberger, Solidification, solution treatment and age
18
[34] M. Nave, A. Dahle, D. StJohn, The Role of Zinc in the Eutectic Solidification of Magnesium‐
19
Aluminium‐Zinc Alloys, Magnesium Technology 2000
20
[35] M. Nave, A. Dahle, D. StJohn, Eutectic Growth Morphologies in Magnesium‐Aluminium Alloys,
21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41
Magnesium technology 2000
hardening of a Mg–1.6 wt.% Ca–3.2 wt.% Zn alloy, Acta Mater. 54(2) (2006) 523-530. [28] P. Yin, N.F. Li, T. Lei, L. Liu, C. Ouyang, Effects of Ca on microstructure, mechanical and corrosion properties and biocompatibility of Mg–Zn–Ca alloys, J. Mater. Sci. Mater. Med. 24(6) (2013) 1365-1373. [29] T. Rzychoń, Characterization of Mg-rich clusters in the C36 phase of the Mg–5Al–3Ca–0.7Sr– 0.2Mn alloy, J. Alloys Compd. 598 (2014) 95-105.
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[30] A.K. Dahle, Y.C. Lee, M.D. Nave, P.L. Schaffer, D.H. StJohn, Development of the as-cast microstructure in magnesium–aluminium alloys, Journal of Light Metals 1(1) (2001) 61-72.
[31] Z.T. Li, X.D. Zhang, M.Y. Zheng, X.G. Qiao, K. Wu, C. Xu, S. Kamado, Effect of Ca/Al ratio on microstructure and mechanical properties of Mg-Al-Ca-Mn alloys, Materials Science and Engineering: A 682 (2017) 423-432.
SC
[32] S. Wasiur-Rahman, M. Medraj, Critical assessment and thermodynamic modeling of the binary Mg–Zn, Ca–Zn and ternary Mg–Ca–Zn systems, Intermetallics 17(10) (2009) 847-864. [33] Y. Zhong, J. Liu, R.A. Witt, Y.-h. Sohn, Z.-K. Liu, Al2(Mg,Ca) phases in Mg–Al–Ca ternary 573-576.
M AN U
system: First-principles prediction and experimental identification, Scripta Mater. 55(6) (2006)
(2000) 243-250.
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(2000) 233-242.
[36] M.S. Dargusch, M. Nave, S.D. McDonald, D.H. StJohn, The effect of aluminium content on the eutectic morphology of high pressure die cast magnesium–aluminium alloys, J. Alloys Compd. 492(1) (2010) L64-L68.
[37] M. Wang, D.H. Xiao, W.S. Liu, Effect of Si addition on microstructure and properties of
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magnesium alloys with high Al and Zn contents, Vacuum 141 (2017) 144-151. [38] C.H. Cáceres, C.J. Davidson, J.R. Griffiths, C.L. Newton, Effects of solidification rate and ageing on the microstructure and mechanical properties of AZ91 alloy, Materials Science and Engineering: A
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325(1–2) (2002) 344-355.
[39] T.L. Chia, M.A. Easton, S.M. Zhu, M.A. Gibson, N. Birbilis, J.F. Nie, The effect of alloy composition on the microstructure and tensile properties of binary Mg-rare earth alloys, Intermetallics 17(7) (2009) 481-490.
[40] D. Amberger, P. Eisenlohr, M. Göken, On the importance of a connected hard-phase skeleton for the creep resistance of Mg alloys, Acta Mater. 60(5) (2012) 2277-2289. [41] Q. Yang, K. Guan, F. Bu, Y. Zhang, X. Qiu, T. Zheng, X. Liu, J. Meng, Microstructures and tensile properties of a high-strength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy, Mater. Charact. [42] Q. Yang, K. Guan, F. Bu, Y. Zhang, X. Qiu, T. Zheng, X. Liu, J. Meng, Microstructures and tensile properties of a high-strength die-cast Mg–4Al–2RE–2Ca–0.3Mn alloy, Mater. Charact. 113 (2016) 180-188. [43] C.D. Lee, K.S. Shin, Effects of precipitate and dendrite arm spacing on tensile properties and fracture behavior of As-Cast magnesium-aluminum alloys, Metals and Materials International 9(1)
ACCEPTED MANUSCRIPT 1 2 3 4 5 6 7
(2003) 21-27. [44] M. Mabuchi, K. Higashi, Strengthening mechanisms of Mg-Si alloys, Acta Mater. 44(11) (1996) 4611-4618. [45] J.W. Luster, M. Thumann, R. Baumann, Mechanical properties of aluminium alloy 6061–Al2O3 composites, Mater. Sci. Technol. 9(10) (1993) 853-862. [46] D.M. Stefanescu, Science and engineering of casting solidification, third edition ed., Springer2015.
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ACCEPTED MANUSCRIPT Table captions: Table 1. Actual compositions of the studied alloys. Table 2. EDS results of intermetallic compounds in each alloy. Table 3. Endothermic peak temperatures for the DSC heating curves in Fig. 6.
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Table 4. Tensile properties of the studied alloys.
ACCEPTED MANUSCRIPT Table 1. Actual compositions of the studied alloys. Actual composition (wt. %) Alloy Zn
Ca
Ce
La
Pr
Nd
Si
Fe
Mn
Mg
AZX264
2.279
6.373
4.048
/
/
/
/
0.008
0.002
0.003
Bal.
AZX444
4.250
4.119
3.726
/
/
/
/
0.007
0.002
0.002
Bal.
AZX624
6.342
2.117
4.066
/
/
/
/
0.009
0.003
0.004
Bal.
AZXE2631
2.280
6.205
3.186
0.537
0.255
0.027
0.139
0.015
0.004
0.005
Bal.
AZXE4431
4.232
3.958
3.098
0.539
0.234
0.025
0.128
0.013
0.003
0.006
Bal.
AZXE6231
6.442
2.046
3.039
0.506
0.226
0.022
0.119
0.018
0.004
0.007
Bal.
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Al
ACCEPTED MANUSCRIPT Table 2. EDS results of intermetallic compounds in each alloy. Elemental composition (at. %) Arrow
Al
Zn
Ca
RE
Mg
Ca2Mg6Zn3
8.70
12.62
14.63
/
64.05
B
Al2REZn2
22.95
15.67
1.22
8.19
51.97
C
Ca2Mg6Zn3
3.86
19.05
11.85
/
65.24
D
C36-(Mg, Al)2Ca
13.74
1.78
E
C36-Al2(Mg, Ca)
25.60
0.66
F
Ca2Mg6Zn3
7.61
9.28
G
C36-Al2(Mg, Ca)
28.57
0.92
H
Al11RE3
36.15
I
Q
23.97
AZX444
AZX624
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A AZXE2631
10.76
/
73.72
8.34
/
65.40
5.72
/
77.39
8.51
/
62.00
2.24
1.05
8.80
51.76
9.18
3.17
/
63.68
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AZXE6231
Intermetallic compound
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Alloy
ACCEPTED MANUSCRIPT Table 3. Endothermic peak temperatures for the DSC heating curves in Fig. 6. Temperature of endothermic peak/ Alloy
Primary phase
Intermetallic compound Al11RE3/Al2REZn2
C36
C14
Ca2Mg6Zn3
AZXE2631
618
587
\
489
458
AZX444
593
\
500
\
445
AZXE6231
591
submerged
517
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α-Mg
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\
\
ACCEPTED MANUSCRIPT Table 4. Tensile properties of the studied alloys. YS/MPa
UTS/MPa
El./%
AZX264
103±6
103±6
0.2±0.0
AZX444
102±4
155±5
1.4±0.2
AZX624
95±3
172±2
2.1±0.2
AZXE2631
107±5
138±2
1.0±0.2
AZXE4431
90±2
139±4
1.1±0.1
AZXE6231
146±6
212±3
2.4±0.3
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Alloy
ACCEPTED MANUSCRIPT Figure captions: Fig. 1 XRD patterns of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZX624 and (e) AZXE6231 alloys. Fig. 2 Optical micrographs of (a) AZX264, (b) AZXE2613, (c) AZX444, (d) AZXE4431, (e)
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AZX624 and (f) AZXE6231 alloys. Fig. 3 Second dendrite arm spacing of the studied alloys.
Fig. 4 Grain size of the binary magnesium alloys with a range of Al, Zn and Ca alloying contents plotted against the growth restriction factor Q. The data is taken from [23, 24].
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Fig. 5 BEIs of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 6 DSC heating curves of (a) AZXE2631, (b) AZX444 and (c) AZXE6231 alloys. Fig. 7 Eutectic morphologies of (a) Ca2Mg6Zn3 and (b) C36 in AZX444 alloy, (c) Ca2Mg6Zn3 and C36 in AZX624 alloy and (d) C36 in AZXE6231 alloy.
Fig. 8 Tensile stress-strain curves of AZX264, AZXE2631, AZX444, AZXE4431, AZX624 and AZXE6231 alloys.
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Fig. 9 Column chart of the mechanical properties of the as-cast alloys.
Fig. 10 Fracture lateral surfaces of (a) AZX264, (b) AZEX2613, (c) AZX444 and (d) AZEX6213 alloys.
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Fig. 11 Volume fractions of the intermetallic compounds in each alloy. Fig. 12 Work-hardening rate-true strain curves of AZXE2631, AZX444 and AZXE6231 alloys.
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Fig. 13 Optical macrographs of (a) AZX264, (b) AZXE2631, (c) AZX444 and (d) AZEX6231 alloys.
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Fig. 1 XRD patterns of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZX624 and (e) AZXE6231
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Fig. 2 Optical micrographs of (a) AZX264, (b) AZXE2613, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 3 Second dendrite arm spacing of the studied alloys.
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Fig. 4 Grain size of the binary magnesium alloys with a range of Al, Zn and Ca alloying contents
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plotted against the growth restriction factor Q. The data is taken from [23, 24].
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Fig. 5 BEIs of (a) AZX264, (b) AZXE2631, (c) AZX444, (d) AZXE4431, (e) AZX624 and (f) AZXE6231 alloys.
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Fig. 6 DSC heating curves of (a) AZXE2631, (b) AZX444 and (c) AZXE6231 alloys.
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Fig. 7 Eutectic morphologies of (a) Ca2Mg6Zn3 and (b) C36 in AZX444 alloy, (c) Ca2Mg6Zn3 and
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Fig. 8 Tensile stress-strain curves of AZX264, AZXE2631, AZX444, AZXE4431, AZX624 and
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Fig. 9 Column chart of the mechanical properties of the as-cast alloys.
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c
d
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Fig. 10 Fracture lateral surfaces of (a) AZX264, (b) AZEX2613, (c) AZX444 and (d) AZEX6213
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Fig. 11 Volume fractions of the intermetallic compounds in each alloy.
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Fig. 12 Work-hardening rate-true strain curves of AZXE2631, AZX444 and AZXE6231 alloys.
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Fig. 13 Optical macrographs of (a) AZX264, (b) AZXE2631, (c) AZX444 and (d) AZEX6231
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alloys.
ACCEPTED MANUSCRIPT Mg-xAl-(8-x)Zn-4Ca and Mg-xAl-(8-x)Zn-3Ca-1RE alloys were prepared by gravity casting, and the microstructures and mechanical properties were studied. The category, size, morphology and distribution of the intermetallic compounds were significantly influenced by alloying composition.
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C36 formed as a majority compound instead of Ca2Mg6Zn3 with the increase of Al content, which was much finer and more evenly distributed in high Al content alloy.
Replacing a part of Ca by RE could further refine the microstructure, thus
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Mg-6Al-2Zn-3Ca-1RE obtained the best mechanical property.