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Effects of Mn addition on the microstructure and mechanical properties of Mg–Zn–Sn alloys
Journal Pre-proof Effects of Mn addition on the microstructure and mechanical properties of Mg–Zn–Sn alloys Caihong Hou, Fugang Qi, Zhisong Ye, Nie Zhao, Dingfei Zhang, Xiaoping Ouyang PII:
S0921-5093(20)30024-1
DOI:
https://doi.org/10.1016/j.msea.2020.138933
Reference:
MSA 138933
To appear in:
Materials Science & Engineering A
Received Date: 4 September 2019 Revised Date:
6 January 2020
Accepted Date: 6 January 2020
Please cite this article as: C. Hou, F. Qi, Z. Ye, N. Zhao, D. Zhang, X. Ouyang, Effects of Mn addition on the microstructure and mechanical properties of Mg–Zn–Sn alloys, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2020.138933. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2020 Published by Elsevier B.V.
Effects of Mn addition on the microstructure and mechanical properties of Mg-Zn-Sn alloys﹡ Caihong Houa,b, Fugang Qia,b,*, Zhisong Yea,b, Nie Zhaoa,b,*, Dingfei Zhangc, Xiaoping Ouyanga,b a
School of Materials Science and Engineering, Xiangtan University, Xiangtan, 411105, P.R.China
b
Key Laboratory of Low Dimensional Materials and Application Technology of Ministry of Education, Xiangtan University, Xiangtan, 411105, P.R.China
c
College of Materials Science and Engineering, Chongqing University, Chongqing, 400045, P.R.China * Corresponding Author: E-mail: [email protected], [email protected]
Abstract: The effects of Mn addition on the microstructure and mechanical properties of Mg-Zn-Sn alloys have been investigated by optical microscopy (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM) and uniaxial tensile tests. The results show that the addition of Mn can significantly improve the mechanical properties of the as-extruded and aged Mg-Zn-Sn alloys, which is mainly due to the grain refinement and precipitation strengthening. For the as-extruded alloys, the Mn element mainly exists in the form of
﹡Corresponding
author. Tel.: +86 731 58298119
E-mail address: [email protected] (F.G. Qi), [email protected](N. Zhao) 1
α-Mn particle phase, which is dispersed in the matrix and play the role of fine grain strengthening and dispersion strengthening. For the single-aged and double-aged alloys, the dispersed α-Mn particle phases can also serve as the heterogeneous precipitates during the ageing treatment, which is
nucleation cores of rod-like
beneficial to the nucleation rate of the precipitates. The crystallographic characteristics research shows that the directional relationship between α-Mg, α-Mn is 2110
∕∕ 0001
’
∕∕ 012
, i.e., the
and
phase can form a coherent
interface on the α-Mn phase.
Keywords: Mg-Zn-Sn alloy; Mn; Precipitate; Microstructure; Mechanical properties
1. Introduction Magnesium alloys have been widely used in the aerospace, transportation and automobile industries, due to their excellent properties, such as low density, high specific strength, easy recovery and good casting ability [1, 2]. However, due to their low absolute strength, high production cost and poor corrosion resistance, the practical applications of magnesium alloys are limited [3-5]. Therefore, the development of new low-cost high-performance magnesium alloys has attracted the widespread attention. The Mg-Zn-Sn alloys are a new type of high-strength magnesium alloy system. The content of Sn and Zn elements in such alloys, is relatively high, usually between 3 and 9 wt.% ( hereafter, all compositions are in weight percent unless stated otherwise), so 2
it can be expected that such alloys have obvious response to age hardening [6-8]. The researchers have found that Mg-Zn-Sn alloys have better mechanical properties than Mg-Zn and Mg-Sn alloys, mainly due to the synergistic precipitation strengthening of MgZn2 and Mg2Sn precipitates during the aging treatment [9-11]. Tang et al. [7] investigated the effect of Zn content on the microstructure and mechanical properties of Mg-5Sn-XZn (X = l~4 %) alloys. It was found that the number of fine second phases increased during the extrusion process, and the comprehensive mechanical properties of the alloys were improved with increasing Zn. Cheng et al. [12] studied the strengthening effect of the Mg2Sn phase in the as-extruded Mg-8Sn-2Zn-2Al alloy. It was reported that the as-extruded alloy exhibited higher strength and elongation after homogenization treatment. Sasaki et al. [13] prepared high-strength Mg-Sn-Zn-Al alloy by the extrusion, solid solution and two-stage aging treatment, and its highest mechanical properties were 399 MPa, 370 MPa and 14%, which was attributed to the combined precipitation strengthening of MgZn2 and Mg2Sn precipitates. These studies show that the Mg-Zn-Sn alloys are promising high-strength wrought magnesium alloys, and it is expected to further improve the mechanical properties of Mg-Zn-Sn wrought alloys through alloying and micro-alloying. Mn is an important alloying element in magnesium alloys, and its maximum solid solubility in magnesium is close to 2.2 %. In general, the primary role of Mn is to remove impurity elements when the Mn content in the magnesium alloy is less than 0.3%. However, in recent years, some researchers have found that higher Mn content can refine grains and improve mechanical properties [14-16]. Our research group has 3
previously studied the effects of different Mn contents on the microstructure and properties of Mg-Zn-Mn alloys, and found that the mechanical properties of the alloys were the best when the Mn content was 1%. It is of great interest to explore the possible fine grain strengthening of Mn addition in the high alloying Mg-Zn-Sn alloy, to develop a new high-strength wrought magnesium alloy. Therefore, we have investigated the effects of 1% Mn addition on the microstructure and mechanical properties of Mg-6Zn-4Sn alloy subjected to cast, homogenization, extrusion, solid solution and aging treatment in the present work. 2. Experiment The Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys were prepared with pure Mg (>99.9%, mass fraction), pure Zn (>99.95%), pure Sn (>99.9%) and Mg-4.10% Mn master alloys. All these alloys were melted at 690~750 °C in a ZG-00L vacuum induction melting furnace under an Ar atmosphere. The chemical compositions were analysed with an XRF-1800 CCDE X-ray fluorescence spectrometer, and the results are listed in Table 1. Table 1 The chemical compositions of the studied alloys (wt.%) Actual compositions (wt. %)
No.
Mg
Zn
Sn
Mn
Mg-6Zn-4Sn
Bal.
6.17
4.36
0
Mg-6Zn-4Sn-1Mn
Bal.
5.91
4.46
0.93
The ingots were homogenized at 330 °C for 24 h followed by air cooling. Then, the ingots were extruded on an XJ-500 horizontal extrusion machine, and the extrusion process parameters are shown in Table 2. The as-extruded samples were 4
solution-treated at 440 °C for 2 h, followed by water quenching. After solution treatment, the artificial aging treatments are carried out. In detail, the aging treatments were divided into two types, one was single aging, that is, the samples were aged at 180 °C for 12 h, and the other was double aging, that is, the samples were pre-aged at 90 °C for 24 h, followed by secondary aging at 180 °C for 8 h. Table 2 Extrusion parameters for the studied alloys Test materials
Mg-6Zn-4Sn-xMn(x=0,1)
Billet tempera ture (°C)
Extrusion chamber temperatu re (°C)
Mold hole diameter (mm)
Extrusion speed (m/min)
Extrusi on ratio
360
350
16
2
25
The mechanical properties of the samples were evaluated on an electrical universal material testing machine CMT-5105 (tensile displacement rate of 3 mm/s) at room temperature, and the schematic diagram of tensile testing samples are shown in Fig. 1. Then, the structural constituents of the alloys were analysed with a Rigaku D/MAX-2500 PC X-ray diffractometer using Cu-Kα radiation with a scanning angle from 10° to 90° and a scanning rate of 4°/min. The microstructures of the samples were characterized with a BX53M optical microscope, a TESCAN VEGA II scanning electron microscope equipped with an Oxford INCA Energy 350 energy-dispersive spectrometer and a FEI Tecnai G2 F20 transmission electrical microscope.
5
Figure 1 The schematic diagram of tensile sample. 3. Results 3.1 As-cast microstructure The X-ray diffraction (XRD) patterns of the as-cast Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are shown in Fig. 2. As can be seen from Fig. 2(a), the Mg-6Zn-4Sn alloy consists of α-Mg, Mg7Zn3 and Mg2Sn phases, after the addition of Mn, the α-Mn phase diffraction peaks are also detected in the Mg-6Zn-4Sn-1Mn alloy, indicating that Mn may mainly exist as a pure α-Mn phase.
6
Figure 2 XRD patterns of the as-cast (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys. The
optical
microstructure
images
of
the
as-cast
Mg-6Zn-4Sn
and
Mg-6Zn-4Sn-1Mn alloys are shown in Fig. 3 (a and b). The two alloys exhibit continuous network dendritic structures with α-Mg matrix and eutectic phases. It is found that the dendritic structure is partially refined with Mn addition. Fig. 3 (c and d) show the backscattered electron (BSE) images of the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. According to the energy-dispersive spectroscopy (EDS) analysis, the secondary phase constitutes of the two alloys are identical. Specifically, the bright contrast is determined to be Mg2Sn phase with a face-centred cubic structure, and the grey contrast is determined to be Mg7Zn3 phase with an orthorhombic structure [17]. Moreover, the line scanning in Fig. 3 (d) shows that Mn 7
element is evenly distributed in the matrix and compounds.
Figure 3 (a and b) Optical and (c and d) BSE-SEM microstructures of the as-cast (a and c) Mg-6Zn-4Sn and (b and d) Mg-6Zn-4Sn-1Mn alloys (inset: line scanning of Mn element in (d)). 3.2 As-homogenized microstructure In order to improve the structural non-uniformity of the ingot and greatly reduce the deformation resistance of the alloys, a homogenization treatment is performed before the extrusion process. Fig. 4 (a and b) show the optical microstructure images of the as-homogenized Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. It is impressive to note that the secondary phases are partially dissolved into the matrix after the homogenization treatment, and 8
the network-like compounds apparently change from continuous to discontinuous. In addition, the compounds are identified by means of SEM and EDS, and it is found that the bulk compounds on the grain boundaries are Mg7Zn3 compounds, which are mostly dissolved during the homogenization treatment. The Mg2Sn compounds remain nearly constant. A comparison of the line scanning microstructures of the as-cast (Fig. 3 (d)) and as-homogenized (Fig. 4 (d)) alloys show that the content of the as-homogenized Mn element is increased in the residual Mg-Zn compounds. This is due to that the Mn-rich phase can be regarded as the heterogeneous nucleated core of the Mg7Zn3 phase in the eutectic reaction, which causes the Mn-rich phase to be wrapped by eutectic compounds. Then, the Mn-rich phase is exposed after homogenization treatment. Additionally, the supersaturated Mn in the matrix will be precipitated during the homogenization treatment.
9
Figure 4 (a and b) Optical and (c and d) BSE-SEM microstructures of the as-homogenized (a and c) Mg-6Zn-4Sn and (b and d) Mg-6Zn-4Sn-1Mn alloys (inset: line scanning of Mn element in (d)). Fig. 5 shows a bright-field TEM (BF-TEM) image of the as-homogenized Mg-6Zn-4Sn-1Mn alloy. The partially supersaturated Mn element is precipitated after homogenization treatment [18]. As seen from the image, there are many fine precipitated particles in the matrix after homogenization treatment. According to the EDS analysis, the blue arrow denotes a kind of Mg-Zn phase, which has a completely consistent directional arrangement and is similar to the rod-like
(MgZn2) phase.
The red circle corresponds to α-Mn particles, which have a nondirectional 10
arrangement and are shaped like rods and regular polygons.
Figure 5 Bright-field TEM image of the as-homogenized Mg-6Zn-4Sn-1Mn alloy. 3.3 As-extruded microstructure Fig. 6 (a and b) show the optical microstructure images of the as-extruded Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. The figures show that the eutectic compounds are broken and distributed as streamlines parallel to the extrusion direction. Both alloys have mixed crystal structures. The reason for this phenomenon is that the Mg-Zn and Mg-Sn compounds can effectively pin dislocations during the deformation process, which inhibits dynamic recrystallization. Therefore, the microstructure will remain unrecrystallized grains. Additionally, the grains of the Mg-6Zn-4Sn-1Mn alloys are finer than those of the Mg-6Zn-4Sn alloys. It is due to that the α-Mn phases play a role in inhibiting dynamic recrystallization. The secondary phases can be observed more clearly in the SEM photographs (Fig. 6 (c and d)). The secondary phases are almost broken during the extrusion process, and the strip-shaped secondary phases are much coarser, as shown in Fig. 6 (d). 11
Fig. 6 (e) shows a BF-TEM image of the as-extruded Mg-6Zn-4Sn-1Mn alloy. The figure indicates that the α-Mn particles have different morphologies. The reasons for these different morphologies can be divided into two aspects. On the one hand, the α-Mn particles are precipitated after homogenization treatment, and these α-Mn particles are further grown during the hot extrusion processing. On the other hand, some new fine α-Mn particles continue to be precipitated during the hot extrusion processing. These α-Mn particles are dispersedly distributed in the matrix, and the particles range in size from 10 to 50 nm, with spherical shape and rod shape.
Figure 6 (a and b) Optical microstructures of the as-extruded (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys (the extruded direction is horizontal). (c and d) BSE-SEM microstructures of the as-extruded (c) Mg-6Zn-4Sn and (d) Mg-6Zn-4Sn-1Mn alloys, (e) bright-field TEM image of the as-extruded Mg-6Zn-4Sn-1Mn alloy. 3.4 Solution-treated and aged microstructures 12
Fig. 7 (a and b) show the optical microstructures of the solution-treated Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. In contrast to the as-extruded alloys, the recrystallized grains grow significantly in the solution-treated alloys. Moreover, the grains size of the Mg-6Zn-4Sn-1Mn alloy is much smaller than that of the Mg-6Zn-4Sn alloy. The reason is that the solid solubility of Zn and Sn in magnesium matrix increase at elevated temperature, so the Mg7Zn3 and Mg2Sn phases are almost dissolved in the matrix. In addition, the α-Mn phases located on the grain boundary can inhibit the grain boundary migration.
Figure 7 (a and b) Optical microstructures of the solution treated (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys. Fig. 8 shows the BF-TEM images and corresponding fast Fourier transform (FFT) patterns of the double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys taken from the 1120
zone axis. Based on the EDS analysis and previous studies, three different
precipitates are observed in the Mg-6Zn-4Sn alloy, as shown in Fig. 8 (a). One is a rod-like
phase parallel to 0001 , the other is a disc-like
phase lying on the
0001 basal plane and the last is a Mg2Sn phase without particular morphology or direction [19-21]. After the addition of Mn element, the precipitates of the 13
Mg-6Zn-4Sn-1Mn alloy mainly consists of
,
, Mg2Sn and α-Mn phases, as
shown in Fig. 8 (b). In addition, it is observed that the number density of the precipitates increase after the addition of Mn. At the same time, it is also found that the rod-like
phases are significantly shorter and coarser in Fig. 8 (b). The reason is
possible that the Mn element exists in the form of the α-Mn phases in the matrix, which promote the precipitation of the
phases. Thereby, the
phases are
formed much earlier, and grow up more fully.
Figure 8 Bright-field TEM images of the double aged (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys (inset: 1120
diffraction pattern).
Mn element is mainly dissolved in the matrix during the melting process and then continuously precipitates and grows during the homogenization, extrusion and solution treatment. Fig.9 are the typical α-Mn particles in the investigated alloys. Most of the supersaturated Mn elements have precipitated after the hot extrusion processing. Therefore, during the subsequent solution treatment, most of the α-Mn phases coarsen and grow to 100 nm (Fig. 9 (b and c)). During the subsequent aging 14
treatment, the size of the α-Mn phases doesn’t change due to the lower temperature (Fig. 9 (a)). At the same time, it can be seen that the α-Mn phases have many different morphologies from the image (Fig. 9), mainly including spherical shapes, rod shapes and regular polygons. Moreover, a part of the α-Mn phases can be used as the nucleation core of rod-like β1' phases during the aging treatment, as shown in Fig. 9 (d and e).
Figure 9 (a, b and c) Morphologies of the α-Mn precipitates for the (b and c) solution treated and (a) pre-aged Mg-6Zn-4Sn-1Mn alloys. (d and e) Morphologies and function of the α-Mn precipitates for the aged Mg-6Zn-4Sn-1Mn alloys. (d) 440 °C/2 h+180 °C/2 h, (e) 440 °C/2 h+180 °C/12 h. 3.5 Mechanical properties The room-temperature mechanical properties of the as-extruded, single-aged and 15
double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are summarized in Fig. 10. For the investigated alloys, the aged samples exhibit significantly higher strengths than the as-extruded samples, and the double-aged alloys exhibit higher strengths than the single-aged alloys. During the solution treatment, the Mg-Zn and Mg-Sn compounds are dissolved in the matrix, and then those compounds will precipitate at the subsequent aging treatment. Precipitates can provide a precipitation enhancing effect, so the properties of the as-extruded alloys are greatly improved after the aging treatments. For the double-stage aging treatment, the Guinier Preston (G.P.) zone is generated during the pre-aging treatment, which can provide the nucleation cores of
phases during the second high-temperature aging treatment [22]. Therefore, the
double-aged alloys have better mechanical properties than the single-aged alloys. In addition, Mn addition can improve the room-temperature tensile properties of the Mg-6Zn-4Sn alloy. As seen from Fig. 10, regardless of the extruded state or aged state, the mechanical properties of the Mg-6Zn-4Sn-1Mn alloy are better than those of the Mg-6Zn-4Sn alloy. In the as-extruded state, the Mn is mainly distributed around the α-Mg matrix in the form of dispersed particles, which can refine the grains and enhance dispersibility. Therefore, the mechanical properties of the as-extruded Mg-6Zn-4Sn-1Mn alloys are greater than those of the as-extruded Mg-6Zn-4Sn alloys, and the ultimate tensile strength (UTS) and yield strength (YS) are increased by 41MPa and 59MPa, respectively. In the aged state, the α-Mn precipitated phases can inhibit the grain growth during the solid solution process and play a role in the grain refinement strengthening. Furthermore, the dispersed distribution of α-Mn 16
precipitated phases provide precipitation strengthening. Therefore, the mechanical properties of the aged Mg-6Zn-4Sn-1Mn alloys are greater than those of the aged Mg-6Zn-4Sn alloys. It is impressive to note that the UTS and YS of the single-aged alloys improve 65MPa and 95MPa, respectively, and the UTS and YS of the double-aged alloys improve 68MPa and 80MPa, respectively.
Figure 10 Mechanical properties of (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys at different states. 4. Discussions The evolution of grain size and secondary phases will affect the mechanical properties of magnesium alloys after the addition of 1% Mn. Based on the current research, the corresponding mechanisms will be discussed as follows: alloy melt purification, grain refinement and heterogeneous nucleated core. 17
4.1 Alloy melt purification When the magnesium alloy is melted, impurity elements such as Fe and Si are interfused by the iron mould and raw materials, and these impurity elements will reduce the corrosion resistance of the alloys. According to the literature [23], a small amount of Mn can improve the quality of the alloy by decreasing the corrosion rate. Fig. 11 (a) shows that a small amount of Mn-rich phases have been found in the extruded structures. The line scans show that these Mn-rich phases contain Mn, Fe, Al and other elements, which may form during the casting process. Moreover, Mn-rich phases containing Al element are also observed in the solution-treated Mg-6Zn-4Sn-1Mn alloy (Fig. 11 (b)). These indicate that Mn can be combined with impurity elements and then remove the impurities. It should be noted that the Mn element can form compounds with impurity elements such as Fe and Si during the melting process, and then play a role in removing impurity elements. However, it is possible that a few compounds have the same density as the matrix and remain suspended in the melt without separation and precipitation. Therefore, small amounts of Mn-rich compounds are occasionally observed in experimental alloy structures.
18
Figure 11 (a) SE-SEM of the as-extruded Mg-6Zn-4Sn-1Mn alloy. (b) SE-SEM of the solution treated Mg-6Zn-4Sn-1Mn alloy. 4.2 Grain refinement Based on current microstructural analysis, it is known that the addition of Mn can refine grains. For the as-cast structure, most of the Mn is dissolved in the matrix. However, some α-Mn phases accumulate in the front of the α-Mg phases, which results in compositional undercooling that inhibits grain growth. Therefore, the dendritic structures are refined, which has an effect on grain refinement strengthening. For the as-extruded and heat-treated structures, the α-Mn phases precipitate and grow further during the homogenization, thermal deformation and high-temperature solution treatment. There are many fine α-Mn phases diffusely distributed in the matrix, which hinders dynamic recrystallization and grain boundary migration. In this process, Mn plays the role of grain refinement strengthening and precipitation strengthening. Therefore, the room-temperature mechanical properties of the alloys can be improved by Mn addition, which can be confirmed from Fig. 10. 19
According to the research [24], the total yield strength can be estimated as where
=
+
+
+
is the intrinsic (Peierls) stress required to move a dislocation,
grain refinement strengthening,
is the solid solution strengthening, and
(1) is the is
the precipitation strengthening. SIMAR et al. [25] proposed that the effect of grain size on alloy strength meets the Hall–Petch relationship. !
= where
"
(2)
is the locking parameter and d is the average grain size. Due to the low
solubility of Mn in magnesium for the investigated alloys, we assume that the effect of Mn on
is negligible [26]. Combined with the references [26, 27], we use
= 0.64 MPa m" . The grain sizes (d) for the double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are 108.2 µm and 37.1µm, respectively. Calculated with Eq. (2),
values of Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are 61.5 MPa and105.1
MPa, respectively. The precipitation strengthening stems from the interaction of dislocation, which can be expressed as follows: =
$%& '
= ()*
(3)
where M is an orientation factor related to the basal texture in the material, )* is the critical resolved shear stress (CRSS) for the operative slip system, + is the shear modulus, b is the magnitude of the Burgers vector, and L is the spacing between obstacles. Combined with the previous study [24, 28-30], the value of M is 4.5 for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. Meanwhile, the value of )* is 45.6 MPa 20
and 56.2 MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys, respectively. Therefore, the calculated
is 205.2 MPa and 252.9 MPa for the Mg-6Zn-4Sn and
Mg-6Zn-4Sn-1Mn alloys, respectively. This indicates that the precipitation strengthening is the largest contribution to strengthening. We neglect the contribution of solid solution strengthening ( small effect present in the investigated alloy [31, 32]. We use on the references [27, 33]. Summing all contributions,
) due to the very =10.7 MPa based = 277.4 MPa and
=368.7 MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys, are approximately consistent with the measured yield strength of 298 MPa and 378MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. 4.3 Heterogeneous nucleated core A peculiar phenomenon is observed in the microstructural analysis of Mg-6Zn-4Sn-1Mn alloy, in which the growth of the rod-like
phases is dependent
on the α-Mn phases during the aging treatments. This finding indicates that the α-Mn phases can be regarded as the heterogeneous nucleated cores of the rod-like phases, which can be confirmed from Fig. 9. To clearly study the orientation relationship between α-Mn and
, a pre-aged (440
°C/2 h+180 °C/2 h) structure is selected for observation. Fig. 12 (a) shows the high-resolution TEM (HRTEM) image (taken from 2110 , zone axis) of the rod-like
phases in the pre-aged Mg-6Zn-4Sn-1Mn alloy. Moreover, the diffraction
spots obtained by FFT transformation from the HRTEM image are shown in Fig. 12 (b). A set of spots can be separated and calibrated after finding the matrix and the 21
rod-like
phase, which indicates that the orientation relationship between rod-like
phase and matrix is
0001
’
∕∕ 2110
, 1010
’
phases and rod-like
∕∕ 100
, 2110
’
∕∕ 0002
,
, and the
phases and α-Mn phases is 0001
orientation relationship between rod-like 012
,
’
∕∕
. Therefore, it is concluded that the interface of α-Mn
phases are fully coherent boundaries, and the rod-like
phases can be attached to α-Mn phase nucleation. According to the different precipitate arrangements simulated with corresponding energy density [34], the simulations show preferential localization of energies at corners and interfaces. Nucleation and growth of the new phase would be favored at such location, due to the such event is associated with an overall reduction in strain energy [35]. And the total change in free energy required for nucleation of a spherical solid precipitate from the matrix is: /
∆. = 0 12 0 Δ.4
!
+ 412
/
+ 0 12 0 6
(4)
The first two terms include the free energy change per unit volume (∆.7 ), and the energy change needed to create the unit area of the interface (
). The third term
takes into account the strain energy per unit volume (6), the energy required to permit a precipitate to fit into the surrounding matrix during the nucleation and growth of the precipitate, introduced when the precipitate forms in a solid, rigid matrix. The precipitate does not occupy the same volume that is displaced, so additional energy is required to accommodate the precipitate in the matrix [36]. This explains the phenomenon that the rod-like
phases are attached to the α-Mn phases
nucleationfrom the perspective of strain drive, as shown in Fig.12 (a). 22
The edge-to-edge matching model proposed by Zhang and Kelly [37, 38] requires that the interatomic spacing misfit along matching directions is less than 10%, and the d-value mismatch between matching planes is normally less than 6%. Moreover, the corresponding formulas are as follows: 89 =
|;< !;= | ;=
, 8> =
|
(5)
=
In specific, the interatomic spacing of Mn in the direction of [012] is 0.89 nm and that of
in the direction of [0001] is 0.86 nm. The interplanar spacing of Mn in the in the direction of 1010 is 0.90 nm.
direction of (100) is 0.89 nm and that of
Substituting the value into the formula to obtain the r-value misfit of the direction pair 0001 1010
’
’
∕∕ 012
is 3.37%, and the d-value mismatch of the plane pair
∕∕ 100
is 1.12%, which further shows that the α-Mn precipitated
phases can be regarded as the heterogeneous core of the rod-like A part of the the Fig.12 (a), the found that the
phase.
phases are nucleated and grown on the α-Mn phases. As seen in phases are rod-shaped. Combined with the research [39], it is phases attached to the nucleation of α-Mn phases are all rod-shaped,
and this phenomenon can also be observed in Fig. 9 (d) and (e). When the
phases
are attached to the α-Mn phases and nucleate, the microstructure has changed, and the lattice constants have changed. In this case, in order to maintain the coherent or semi-coherent state between the α-Mn phase and
phase, the α-Mn crystal lattice
will be distorted, and then the elastic strain energy and interface energy anisotropy will be generated. Due to the interaction between the interface energy anisotropy and the elastic strain energy of α-Mn phases, the shape and volume of the 23
phases are
restricted, and the morphology of the morphology of the
phases are determined. The specific
phases may be calculated by the phase field method [40, 41].
The phase field method is a microstructure simulation method developed from the Ginzburg-Landau theory. Combining the effect of ordering potential and thermodynamic driving force, it can predict the morphology of particles through field variables.
Figure 12 (a) HRTEM image of the rod-like β'1 phase nucleated in α-Mn particle phase for the underaged (440°C/2h+180°C/2h) Mg-6Zn-4Sn-1Mn alloy, taken from 2110
zone axis. (b) FFT pattern from (a).
5. Conclusion In summary, the current work systematically investigates the microstructure and mechanical properties of the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys via OM, SEM, TEM and tensile tests. The results are as follows: (1) Mn mainly exists in the form of α-Mn particle phases, which have no effect on the original phase composition of the Mg-Zn-Sn ternary alloy. The morphologies of 24
the α-Mn phases are mainly spherical shape, rod shape and regular polygon. (2) Mn has a significant effect on the microstructure of the Mg-Zn-Sn alloy. The α-Mn phases can inhibit dynamic recrystallization and hinder grain boundary migration. (3) The α-Mn precipitate phases can be regarded as the heterogeneous nucleation cores of the rod-like
phases during the aging treatment process, which
promotes the precipitation of the rod-like between α-Mg,
and α-Mn is 2110
phases. The directional relationship ∕∕ 0001
’
∕∕ 012
.
(4) The strengths of the Mg-6Zn-4Sn-1Mn alloy are much better than that of the Mg-6Zn-4Sn alloy regardless of the as-extruded or aged state. The room-temperature mechanical properties are improved by grain refinement strengthening and precipitation strengthening after the addition of Mn. Acknowledgements This work was supported by National Natural Science Foundation of China (51701172), China Postdoctoral Science Foundation (2018M632977), Natural Science Foundation of Hunan Province (2018JJ3504), Educational Commission of Hunan Province of China (16C1527), Foundation of Xiangtan University (KZ08034), and Natural Science Foundation of Xiangtan University (KZ03014). The authors acknowledge Dr. Xiaoming Zhang from ZKKF (Beijing) Science & Technology Co., Ltd. for TEM observations.
References 25
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28
CRediT Author Statement Caihong Hou: Formal analysis, Data Curation, Writing-Original Draft. Fugang Qi: Conceptualization, Writing-Review & Editing, Project administration. Zhisong Ye: Investigation. Nie Zhao: Methodology. Dingfei Zhang: Resources. Xiaoping Ouyang: Supervision, Funding acquisition.
S0921-5093(20)30024-1
DOI:
https://doi.org/10.1016/j.msea.2020.138933
Reference:
MSA 138933
To appear in:
Materials Science & Engineering A
Received Date: 4 September 2019 Revised Date:
6 January 2020
Accepted Date: 6 January 2020
Please cite this article as: C. Hou, F. Qi, Z. Ye, N. Zhao, D. Zhang, X. Ouyang, Effects of Mn addition on the microstructure and mechanical properties of Mg–Zn–Sn alloys, Materials Science & Engineering A (2020), doi: https://doi.org/10.1016/j.msea.2020.138933. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. 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. © 2020 Published by Elsevier B.V.
Effects of Mn addition on the microstructure and mechanical properties of Mg-Zn-Sn alloys﹡ Caihong Houa,b, Fugang Qia,b,*, Zhisong Yea,b, Nie Zhaoa,b,*, Dingfei Zhangc, Xiaoping Ouyanga,b a
School of Materials Science and Engineering, Xiangtan University, Xiangtan, 411105, P.R.China
b
Key Laboratory of Low Dimensional Materials and Application Technology of Ministry of Education, Xiangtan University, Xiangtan, 411105, P.R.China
c
College of Materials Science and Engineering, Chongqing University, Chongqing, 400045, P.R.China * Corresponding Author: E-mail: [email protected], [email protected]
Abstract: The effects of Mn addition on the microstructure and mechanical properties of Mg-Zn-Sn alloys have been investigated by optical microscopy (OM), X-ray diffraction (XRD), scanning electron microscopy (SEM), transmission electron microscopy (TEM) and uniaxial tensile tests. The results show that the addition of Mn can significantly improve the mechanical properties of the as-extruded and aged Mg-Zn-Sn alloys, which is mainly due to the grain refinement and precipitation strengthening. For the as-extruded alloys, the Mn element mainly exists in the form of
﹡Corresponding
author. Tel.: +86 731 58298119
E-mail address: [email protected] (F.G. Qi), [email protected](N. Zhao) 1
α-Mn particle phase, which is dispersed in the matrix and play the role of fine grain strengthening and dispersion strengthening. For the single-aged and double-aged alloys, the dispersed α-Mn particle phases can also serve as the heterogeneous precipitates during the ageing treatment, which is
nucleation cores of rod-like
beneficial to the nucleation rate of the precipitates. The crystallographic characteristics research shows that the directional relationship between α-Mg, α-Mn is 2110
∕∕ 0001
’
∕∕ 012
, i.e., the
and
phase can form a coherent
interface on the α-Mn phase.
Keywords: Mg-Zn-Sn alloy; Mn; Precipitate; Microstructure; Mechanical properties
1. Introduction Magnesium alloys have been widely used in the aerospace, transportation and automobile industries, due to their excellent properties, such as low density, high specific strength, easy recovery and good casting ability [1, 2]. However, due to their low absolute strength, high production cost and poor corrosion resistance, the practical applications of magnesium alloys are limited [3-5]. Therefore, the development of new low-cost high-performance magnesium alloys has attracted the widespread attention. The Mg-Zn-Sn alloys are a new type of high-strength magnesium alloy system. The content of Sn and Zn elements in such alloys, is relatively high, usually between 3 and 9 wt.% ( hereafter, all compositions are in weight percent unless stated otherwise), so 2
it can be expected that such alloys have obvious response to age hardening [6-8]. The researchers have found that Mg-Zn-Sn alloys have better mechanical properties than Mg-Zn and Mg-Sn alloys, mainly due to the synergistic precipitation strengthening of MgZn2 and Mg2Sn precipitates during the aging treatment [9-11]. Tang et al. [7] investigated the effect of Zn content on the microstructure and mechanical properties of Mg-5Sn-XZn (X = l~4 %) alloys. It was found that the number of fine second phases increased during the extrusion process, and the comprehensive mechanical properties of the alloys were improved with increasing Zn. Cheng et al. [12] studied the strengthening effect of the Mg2Sn phase in the as-extruded Mg-8Sn-2Zn-2Al alloy. It was reported that the as-extruded alloy exhibited higher strength and elongation after homogenization treatment. Sasaki et al. [13] prepared high-strength Mg-Sn-Zn-Al alloy by the extrusion, solid solution and two-stage aging treatment, and its highest mechanical properties were 399 MPa, 370 MPa and 14%, which was attributed to the combined precipitation strengthening of MgZn2 and Mg2Sn precipitates. These studies show that the Mg-Zn-Sn alloys are promising high-strength wrought magnesium alloys, and it is expected to further improve the mechanical properties of Mg-Zn-Sn wrought alloys through alloying and micro-alloying. Mn is an important alloying element in magnesium alloys, and its maximum solid solubility in magnesium is close to 2.2 %. In general, the primary role of Mn is to remove impurity elements when the Mn content in the magnesium alloy is less than 0.3%. However, in recent years, some researchers have found that higher Mn content can refine grains and improve mechanical properties [14-16]. Our research group has 3
previously studied the effects of different Mn contents on the microstructure and properties of Mg-Zn-Mn alloys, and found that the mechanical properties of the alloys were the best when the Mn content was 1%. It is of great interest to explore the possible fine grain strengthening of Mn addition in the high alloying Mg-Zn-Sn alloy, to develop a new high-strength wrought magnesium alloy. Therefore, we have investigated the effects of 1% Mn addition on the microstructure and mechanical properties of Mg-6Zn-4Sn alloy subjected to cast, homogenization, extrusion, solid solution and aging treatment in the present work. 2. Experiment The Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys were prepared with pure Mg (>99.9%, mass fraction), pure Zn (>99.95%), pure Sn (>99.9%) and Mg-4.10% Mn master alloys. All these alloys were melted at 690~750 °C in a ZG-00L vacuum induction melting furnace under an Ar atmosphere. The chemical compositions were analysed with an XRF-1800 CCDE X-ray fluorescence spectrometer, and the results are listed in Table 1. Table 1 The chemical compositions of the studied alloys (wt.%) Actual compositions (wt. %)
No.
Mg
Zn
Sn
Mn
Mg-6Zn-4Sn
Bal.
6.17
4.36
0
Mg-6Zn-4Sn-1Mn
Bal.
5.91
4.46
0.93
The ingots were homogenized at 330 °C for 24 h followed by air cooling. Then, the ingots were extruded on an XJ-500 horizontal extrusion machine, and the extrusion process parameters are shown in Table 2. The as-extruded samples were 4
solution-treated at 440 °C for 2 h, followed by water quenching. After solution treatment, the artificial aging treatments are carried out. In detail, the aging treatments were divided into two types, one was single aging, that is, the samples were aged at 180 °C for 12 h, and the other was double aging, that is, the samples were pre-aged at 90 °C for 24 h, followed by secondary aging at 180 °C for 8 h. Table 2 Extrusion parameters for the studied alloys Test materials
Mg-6Zn-4Sn-xMn(x=0,1)
Billet tempera ture (°C)
Extrusion chamber temperatu re (°C)
Mold hole diameter (mm)
Extrusion speed (m/min)
Extrusi on ratio
360
350
16
2
25
The mechanical properties of the samples were evaluated on an electrical universal material testing machine CMT-5105 (tensile displacement rate of 3 mm/s) at room temperature, and the schematic diagram of tensile testing samples are shown in Fig. 1. Then, the structural constituents of the alloys were analysed with a Rigaku D/MAX-2500 PC X-ray diffractometer using Cu-Kα radiation with a scanning angle from 10° to 90° and a scanning rate of 4°/min. The microstructures of the samples were characterized with a BX53M optical microscope, a TESCAN VEGA II scanning electron microscope equipped with an Oxford INCA Energy 350 energy-dispersive spectrometer and a FEI Tecnai G2 F20 transmission electrical microscope.
5
Figure 1 The schematic diagram of tensile sample. 3. Results 3.1 As-cast microstructure The X-ray diffraction (XRD) patterns of the as-cast Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are shown in Fig. 2. As can be seen from Fig. 2(a), the Mg-6Zn-4Sn alloy consists of α-Mg, Mg7Zn3 and Mg2Sn phases, after the addition of Mn, the α-Mn phase diffraction peaks are also detected in the Mg-6Zn-4Sn-1Mn alloy, indicating that Mn may mainly exist as a pure α-Mn phase.
6
Figure 2 XRD patterns of the as-cast (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys. The
optical
microstructure
images
of
the
as-cast
Mg-6Zn-4Sn
and
Mg-6Zn-4Sn-1Mn alloys are shown in Fig. 3 (a and b). The two alloys exhibit continuous network dendritic structures with α-Mg matrix and eutectic phases. It is found that the dendritic structure is partially refined with Mn addition. Fig. 3 (c and d) show the backscattered electron (BSE) images of the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. According to the energy-dispersive spectroscopy (EDS) analysis, the secondary phase constitutes of the two alloys are identical. Specifically, the bright contrast is determined to be Mg2Sn phase with a face-centred cubic structure, and the grey contrast is determined to be Mg7Zn3 phase with an orthorhombic structure [17]. Moreover, the line scanning in Fig. 3 (d) shows that Mn 7
element is evenly distributed in the matrix and compounds.
Figure 3 (a and b) Optical and (c and d) BSE-SEM microstructures of the as-cast (a and c) Mg-6Zn-4Sn and (b and d) Mg-6Zn-4Sn-1Mn alloys (inset: line scanning of Mn element in (d)). 3.2 As-homogenized microstructure In order to improve the structural non-uniformity of the ingot and greatly reduce the deformation resistance of the alloys, a homogenization treatment is performed before the extrusion process. Fig. 4 (a and b) show the optical microstructure images of the as-homogenized Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. It is impressive to note that the secondary phases are partially dissolved into the matrix after the homogenization treatment, and 8
the network-like compounds apparently change from continuous to discontinuous. In addition, the compounds are identified by means of SEM and EDS, and it is found that the bulk compounds on the grain boundaries are Mg7Zn3 compounds, which are mostly dissolved during the homogenization treatment. The Mg2Sn compounds remain nearly constant. A comparison of the line scanning microstructures of the as-cast (Fig. 3 (d)) and as-homogenized (Fig. 4 (d)) alloys show that the content of the as-homogenized Mn element is increased in the residual Mg-Zn compounds. This is due to that the Mn-rich phase can be regarded as the heterogeneous nucleated core of the Mg7Zn3 phase in the eutectic reaction, which causes the Mn-rich phase to be wrapped by eutectic compounds. Then, the Mn-rich phase is exposed after homogenization treatment. Additionally, the supersaturated Mn in the matrix will be precipitated during the homogenization treatment.
9
Figure 4 (a and b) Optical and (c and d) BSE-SEM microstructures of the as-homogenized (a and c) Mg-6Zn-4Sn and (b and d) Mg-6Zn-4Sn-1Mn alloys (inset: line scanning of Mn element in (d)). Fig. 5 shows a bright-field TEM (BF-TEM) image of the as-homogenized Mg-6Zn-4Sn-1Mn alloy. The partially supersaturated Mn element is precipitated after homogenization treatment [18]. As seen from the image, there are many fine precipitated particles in the matrix after homogenization treatment. According to the EDS analysis, the blue arrow denotes a kind of Mg-Zn phase, which has a completely consistent directional arrangement and is similar to the rod-like
(MgZn2) phase.
The red circle corresponds to α-Mn particles, which have a nondirectional 10
arrangement and are shaped like rods and regular polygons.
Figure 5 Bright-field TEM image of the as-homogenized Mg-6Zn-4Sn-1Mn alloy. 3.3 As-extruded microstructure Fig. 6 (a and b) show the optical microstructure images of the as-extruded Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. The figures show that the eutectic compounds are broken and distributed as streamlines parallel to the extrusion direction. Both alloys have mixed crystal structures. The reason for this phenomenon is that the Mg-Zn and Mg-Sn compounds can effectively pin dislocations during the deformation process, which inhibits dynamic recrystallization. Therefore, the microstructure will remain unrecrystallized grains. Additionally, the grains of the Mg-6Zn-4Sn-1Mn alloys are finer than those of the Mg-6Zn-4Sn alloys. It is due to that the α-Mn phases play a role in inhibiting dynamic recrystallization. The secondary phases can be observed more clearly in the SEM photographs (Fig. 6 (c and d)). The secondary phases are almost broken during the extrusion process, and the strip-shaped secondary phases are much coarser, as shown in Fig. 6 (d). 11
Fig. 6 (e) shows a BF-TEM image of the as-extruded Mg-6Zn-4Sn-1Mn alloy. The figure indicates that the α-Mn particles have different morphologies. The reasons for these different morphologies can be divided into two aspects. On the one hand, the α-Mn particles are precipitated after homogenization treatment, and these α-Mn particles are further grown during the hot extrusion processing. On the other hand, some new fine α-Mn particles continue to be precipitated during the hot extrusion processing. These α-Mn particles are dispersedly distributed in the matrix, and the particles range in size from 10 to 50 nm, with spherical shape and rod shape.
Figure 6 (a and b) Optical microstructures of the as-extruded (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys (the extruded direction is horizontal). (c and d) BSE-SEM microstructures of the as-extruded (c) Mg-6Zn-4Sn and (d) Mg-6Zn-4Sn-1Mn alloys, (e) bright-field TEM image of the as-extruded Mg-6Zn-4Sn-1Mn alloy. 3.4 Solution-treated and aged microstructures 12
Fig. 7 (a and b) show the optical microstructures of the solution-treated Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. In contrast to the as-extruded alloys, the recrystallized grains grow significantly in the solution-treated alloys. Moreover, the grains size of the Mg-6Zn-4Sn-1Mn alloy is much smaller than that of the Mg-6Zn-4Sn alloy. The reason is that the solid solubility of Zn and Sn in magnesium matrix increase at elevated temperature, so the Mg7Zn3 and Mg2Sn phases are almost dissolved in the matrix. In addition, the α-Mn phases located on the grain boundary can inhibit the grain boundary migration.
Figure 7 (a and b) Optical microstructures of the solution treated (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys. Fig. 8 shows the BF-TEM images and corresponding fast Fourier transform (FFT) patterns of the double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys taken from the 1120
zone axis. Based on the EDS analysis and previous studies, three different
precipitates are observed in the Mg-6Zn-4Sn alloy, as shown in Fig. 8 (a). One is a rod-like
phase parallel to 0001 , the other is a disc-like
phase lying on the
0001 basal plane and the last is a Mg2Sn phase without particular morphology or direction [19-21]. After the addition of Mn element, the precipitates of the 13
Mg-6Zn-4Sn-1Mn alloy mainly consists of
,
, Mg2Sn and α-Mn phases, as
shown in Fig. 8 (b). In addition, it is observed that the number density of the precipitates increase after the addition of Mn. At the same time, it is also found that the rod-like
phases are significantly shorter and coarser in Fig. 8 (b). The reason is
possible that the Mn element exists in the form of the α-Mn phases in the matrix, which promote the precipitation of the
phases. Thereby, the
phases are
formed much earlier, and grow up more fully.
Figure 8 Bright-field TEM images of the double aged (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys (inset: 1120
diffraction pattern).
Mn element is mainly dissolved in the matrix during the melting process and then continuously precipitates and grows during the homogenization, extrusion and solution treatment. Fig.9 are the typical α-Mn particles in the investigated alloys. Most of the supersaturated Mn elements have precipitated after the hot extrusion processing. Therefore, during the subsequent solution treatment, most of the α-Mn phases coarsen and grow to 100 nm (Fig. 9 (b and c)). During the subsequent aging 14
treatment, the size of the α-Mn phases doesn’t change due to the lower temperature (Fig. 9 (a)). At the same time, it can be seen that the α-Mn phases have many different morphologies from the image (Fig. 9), mainly including spherical shapes, rod shapes and regular polygons. Moreover, a part of the α-Mn phases can be used as the nucleation core of rod-like β1' phases during the aging treatment, as shown in Fig. 9 (d and e).
Figure 9 (a, b and c) Morphologies of the α-Mn precipitates for the (b and c) solution treated and (a) pre-aged Mg-6Zn-4Sn-1Mn alloys. (d and e) Morphologies and function of the α-Mn precipitates for the aged Mg-6Zn-4Sn-1Mn alloys. (d) 440 °C/2 h+180 °C/2 h, (e) 440 °C/2 h+180 °C/12 h. 3.5 Mechanical properties The room-temperature mechanical properties of the as-extruded, single-aged and 15
double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are summarized in Fig. 10. For the investigated alloys, the aged samples exhibit significantly higher strengths than the as-extruded samples, and the double-aged alloys exhibit higher strengths than the single-aged alloys. During the solution treatment, the Mg-Zn and Mg-Sn compounds are dissolved in the matrix, and then those compounds will precipitate at the subsequent aging treatment. Precipitates can provide a precipitation enhancing effect, so the properties of the as-extruded alloys are greatly improved after the aging treatments. For the double-stage aging treatment, the Guinier Preston (G.P.) zone is generated during the pre-aging treatment, which can provide the nucleation cores of
phases during the second high-temperature aging treatment [22]. Therefore, the
double-aged alloys have better mechanical properties than the single-aged alloys. In addition, Mn addition can improve the room-temperature tensile properties of the Mg-6Zn-4Sn alloy. As seen from Fig. 10, regardless of the extruded state or aged state, the mechanical properties of the Mg-6Zn-4Sn-1Mn alloy are better than those of the Mg-6Zn-4Sn alloy. In the as-extruded state, the Mn is mainly distributed around the α-Mg matrix in the form of dispersed particles, which can refine the grains and enhance dispersibility. Therefore, the mechanical properties of the as-extruded Mg-6Zn-4Sn-1Mn alloys are greater than those of the as-extruded Mg-6Zn-4Sn alloys, and the ultimate tensile strength (UTS) and yield strength (YS) are increased by 41MPa and 59MPa, respectively. In the aged state, the α-Mn precipitated phases can inhibit the grain growth during the solid solution process and play a role in the grain refinement strengthening. Furthermore, the dispersed distribution of α-Mn 16
precipitated phases provide precipitation strengthening. Therefore, the mechanical properties of the aged Mg-6Zn-4Sn-1Mn alloys are greater than those of the aged Mg-6Zn-4Sn alloys. It is impressive to note that the UTS and YS of the single-aged alloys improve 65MPa and 95MPa, respectively, and the UTS and YS of the double-aged alloys improve 68MPa and 80MPa, respectively.
Figure 10 Mechanical properties of (a) Mg-6Zn-4Sn and (b) Mg-6Zn-4Sn-1Mn alloys at different states. 4. Discussions The evolution of grain size and secondary phases will affect the mechanical properties of magnesium alloys after the addition of 1% Mn. Based on the current research, the corresponding mechanisms will be discussed as follows: alloy melt purification, grain refinement and heterogeneous nucleated core. 17
4.1 Alloy melt purification When the magnesium alloy is melted, impurity elements such as Fe and Si are interfused by the iron mould and raw materials, and these impurity elements will reduce the corrosion resistance of the alloys. According to the literature [23], a small amount of Mn can improve the quality of the alloy by decreasing the corrosion rate. Fig. 11 (a) shows that a small amount of Mn-rich phases have been found in the extruded structures. The line scans show that these Mn-rich phases contain Mn, Fe, Al and other elements, which may form during the casting process. Moreover, Mn-rich phases containing Al element are also observed in the solution-treated Mg-6Zn-4Sn-1Mn alloy (Fig. 11 (b)). These indicate that Mn can be combined with impurity elements and then remove the impurities. It should be noted that the Mn element can form compounds with impurity elements such as Fe and Si during the melting process, and then play a role in removing impurity elements. However, it is possible that a few compounds have the same density as the matrix and remain suspended in the melt without separation and precipitation. Therefore, small amounts of Mn-rich compounds are occasionally observed in experimental alloy structures.
18
Figure 11 (a) SE-SEM of the as-extruded Mg-6Zn-4Sn-1Mn alloy. (b) SE-SEM of the solution treated Mg-6Zn-4Sn-1Mn alloy. 4.2 Grain refinement Based on current microstructural analysis, it is known that the addition of Mn can refine grains. For the as-cast structure, most of the Mn is dissolved in the matrix. However, some α-Mn phases accumulate in the front of the α-Mg phases, which results in compositional undercooling that inhibits grain growth. Therefore, the dendritic structures are refined, which has an effect on grain refinement strengthening. For the as-extruded and heat-treated structures, the α-Mn phases precipitate and grow further during the homogenization, thermal deformation and high-temperature solution treatment. There are many fine α-Mn phases diffusely distributed in the matrix, which hinders dynamic recrystallization and grain boundary migration. In this process, Mn plays the role of grain refinement strengthening and precipitation strengthening. Therefore, the room-temperature mechanical properties of the alloys can be improved by Mn addition, which can be confirmed from Fig. 10. 19
According to the research [24], the total yield strength can be estimated as where
=
+
+
+
is the intrinsic (Peierls) stress required to move a dislocation,
grain refinement strengthening,
is the solid solution strengthening, and
(1) is the is
the precipitation strengthening. SIMAR et al. [25] proposed that the effect of grain size on alloy strength meets the Hall–Petch relationship. !
= where
"
(2)
is the locking parameter and d is the average grain size. Due to the low
solubility of Mn in magnesium for the investigated alloys, we assume that the effect of Mn on
is negligible [26]. Combined with the references [26, 27], we use
= 0.64 MPa m" . The grain sizes (d) for the double-aged Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are 108.2 µm and 37.1µm, respectively. Calculated with Eq. (2),
values of Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys are 61.5 MPa and105.1
MPa, respectively. The precipitation strengthening stems from the interaction of dislocation, which can be expressed as follows: =
$%& '
= ()*
(3)
where M is an orientation factor related to the basal texture in the material, )* is the critical resolved shear stress (CRSS) for the operative slip system, + is the shear modulus, b is the magnitude of the Burgers vector, and L is the spacing between obstacles. Combined with the previous study [24, 28-30], the value of M is 4.5 for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. Meanwhile, the value of )* is 45.6 MPa 20
and 56.2 MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys, respectively. Therefore, the calculated
is 205.2 MPa and 252.9 MPa for the Mg-6Zn-4Sn and
Mg-6Zn-4Sn-1Mn alloys, respectively. This indicates that the precipitation strengthening is the largest contribution to strengthening. We neglect the contribution of solid solution strengthening ( small effect present in the investigated alloy [31, 32]. We use on the references [27, 33]. Summing all contributions,
) due to the very =10.7 MPa based = 277.4 MPa and
=368.7 MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys, are approximately consistent with the measured yield strength of 298 MPa and 378MPa for the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys. 4.3 Heterogeneous nucleated core A peculiar phenomenon is observed in the microstructural analysis of Mg-6Zn-4Sn-1Mn alloy, in which the growth of the rod-like
phases is dependent
on the α-Mn phases during the aging treatments. This finding indicates that the α-Mn phases can be regarded as the heterogeneous nucleated cores of the rod-like phases, which can be confirmed from Fig. 9. To clearly study the orientation relationship between α-Mn and
, a pre-aged (440
°C/2 h+180 °C/2 h) structure is selected for observation. Fig. 12 (a) shows the high-resolution TEM (HRTEM) image (taken from 2110 , zone axis) of the rod-like
phases in the pre-aged Mg-6Zn-4Sn-1Mn alloy. Moreover, the diffraction
spots obtained by FFT transformation from the HRTEM image are shown in Fig. 12 (b). A set of spots can be separated and calibrated after finding the matrix and the 21
rod-like
phase, which indicates that the orientation relationship between rod-like
phase and matrix is
0001
’
∕∕ 2110
, 1010
’
phases and rod-like
∕∕ 100
, 2110
’
∕∕ 0002
,
, and the
phases and α-Mn phases is 0001
orientation relationship between rod-like 012
,
’
∕∕
. Therefore, it is concluded that the interface of α-Mn
phases are fully coherent boundaries, and the rod-like
phases can be attached to α-Mn phase nucleation. According to the different precipitate arrangements simulated with corresponding energy density [34], the simulations show preferential localization of energies at corners and interfaces. Nucleation and growth of the new phase would be favored at such location, due to the such event is associated with an overall reduction in strain energy [35]. And the total change in free energy required for nucleation of a spherical solid precipitate from the matrix is: /
∆. = 0 12 0 Δ.4
!
+ 412
/
+ 0 12 0 6
(4)
The first two terms include the free energy change per unit volume (∆.7 ), and the energy change needed to create the unit area of the interface (
). The third term
takes into account the strain energy per unit volume (6), the energy required to permit a precipitate to fit into the surrounding matrix during the nucleation and growth of the precipitate, introduced when the precipitate forms in a solid, rigid matrix. The precipitate does not occupy the same volume that is displaced, so additional energy is required to accommodate the precipitate in the matrix [36]. This explains the phenomenon that the rod-like
phases are attached to the α-Mn phases
nucleationfrom the perspective of strain drive, as shown in Fig.12 (a). 22
The edge-to-edge matching model proposed by Zhang and Kelly [37, 38] requires that the interatomic spacing misfit along matching directions is less than 10%, and the d-value mismatch between matching planes is normally less than 6%. Moreover, the corresponding formulas are as follows: 89 =
|;< !;= | ;=
, 8> =
|
(5)
=
In specific, the interatomic spacing of Mn in the direction of [012] is 0.89 nm and that of
in the direction of [0001] is 0.86 nm. The interplanar spacing of Mn in the in the direction of 1010 is 0.90 nm.
direction of (100) is 0.89 nm and that of
Substituting the value into the formula to obtain the r-value misfit of the direction pair 0001 1010
’
’
∕∕ 012
is 3.37%, and the d-value mismatch of the plane pair
∕∕ 100
is 1.12%, which further shows that the α-Mn precipitated
phases can be regarded as the heterogeneous core of the rod-like A part of the the Fig.12 (a), the found that the
phase.
phases are nucleated and grown on the α-Mn phases. As seen in phases are rod-shaped. Combined with the research [39], it is phases attached to the nucleation of α-Mn phases are all rod-shaped,
and this phenomenon can also be observed in Fig. 9 (d) and (e). When the
phases
are attached to the α-Mn phases and nucleate, the microstructure has changed, and the lattice constants have changed. In this case, in order to maintain the coherent or semi-coherent state between the α-Mn phase and
phase, the α-Mn crystal lattice
will be distorted, and then the elastic strain energy and interface energy anisotropy will be generated. Due to the interaction between the interface energy anisotropy and the elastic strain energy of α-Mn phases, the shape and volume of the 23
phases are
restricted, and the morphology of the morphology of the
phases are determined. The specific
phases may be calculated by the phase field method [40, 41].
The phase field method is a microstructure simulation method developed from the Ginzburg-Landau theory. Combining the effect of ordering potential and thermodynamic driving force, it can predict the morphology of particles through field variables.
Figure 12 (a) HRTEM image of the rod-like β'1 phase nucleated in α-Mn particle phase for the underaged (440°C/2h+180°C/2h) Mg-6Zn-4Sn-1Mn alloy, taken from 2110
zone axis. (b) FFT pattern from (a).
5. Conclusion In summary, the current work systematically investigates the microstructure and mechanical properties of the Mg-6Zn-4Sn and Mg-6Zn-4Sn-1Mn alloys via OM, SEM, TEM and tensile tests. The results are as follows: (1) Mn mainly exists in the form of α-Mn particle phases, which have no effect on the original phase composition of the Mg-Zn-Sn ternary alloy. The morphologies of 24
the α-Mn phases are mainly spherical shape, rod shape and regular polygon. (2) Mn has a significant effect on the microstructure of the Mg-Zn-Sn alloy. The α-Mn phases can inhibit dynamic recrystallization and hinder grain boundary migration. (3) The α-Mn precipitate phases can be regarded as the heterogeneous nucleation cores of the rod-like
phases during the aging treatment process, which
promotes the precipitation of the rod-like between α-Mg,
and α-Mn is 2110
phases. The directional relationship ∕∕ 0001
’
∕∕ 012
.
(4) The strengths of the Mg-6Zn-4Sn-1Mn alloy are much better than that of the Mg-6Zn-4Sn alloy regardless of the as-extruded or aged state. The room-temperature mechanical properties are improved by grain refinement strengthening and precipitation strengthening after the addition of Mn. Acknowledgements This work was supported by National Natural Science Foundation of China (51701172), China Postdoctoral Science Foundation (2018M632977), Natural Science Foundation of Hunan Province (2018JJ3504), Educational Commission of Hunan Province of China (16C1527), Foundation of Xiangtan University (KZ08034), and Natural Science Foundation of Xiangtan University (KZ03014). The authors acknowledge Dr. Xiaoming Zhang from ZKKF (Beijing) Science & Technology Co., Ltd. for TEM observations.
References 25
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CRediT Author Statement Caihong Hou: Formal analysis, Data Curation, Writing-Original Draft. Fugang Qi: Conceptualization, Writing-Review & Editing, Project administration. Zhisong Ye: Investigation. Nie Zhao: Methodology. Dingfei Zhang: Resources. Xiaoping Ouyang: Supervision, Funding acquisition.