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Effect of semi-solid isothermal heat treatment on the microstructure of Mg–6A1–1Zn–0.7Si alloy
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 6 ( 2 0 0 8 ) 374–381
journal homepage: www.elsevier.com/locate/jmatprotec
Effect of semi-solid isothermal heat treatment on the microstructure of Mg–6A1–1Zn–0.7Si alloy Yang Mingbo a,b,∗ , Pan Fusheng b , Cheng Renju b , Bai Liang b a b
Materials Science & Engineering College, Chongqing Institute of Technology, Chongqing 400050, PR China Materials Science & Engineering College, Chongqing University, Chongqing 400030, PR China
a r t i c l e
i n f o
a b s t r a c t
Article history:
In this paper, the effect of semi-solid isothermal heat treatment on the microstructure
Received 2 August 2007
of Mg–6A1–1Zn–0.7Si alloy, especially on the semi-solid microstructural evolution and the
Received in revised form
modification of Chinese script shaped Mg2 Si phase, are first investigated. The research
29 September 2007
results indicate that it is possible to produce the Mg–6A1–1Zn–0.7Si alloy with non-dendritic
Accepted 14 December 2007
microstructure by the semi-solid isothermal heat treatment. After treated at 575–585 ◦ C for 120 min, the experimental alloy can obtain a non-dendritic microstructure with a 12–21% liquid content and an average size range of 48–67 m of the unmelted primary solid parti-
Keywords:
cles. In addition, the semi-solid isothermal heat treatment can modify the Chinese script
Mg–A1–Si-based alloys
shaped Mg2 Si phases in experimental alloy. After treated at 580 ◦ C or 585 ◦ C for 120 min,
Semi-solid
the morphology of Mg2 Si phases in the experimental alloy changes from the initial Chinese
Semi-solid isothermal heat
script shape to granule and/or polygon shapes.
treatment
© 2007 Elsevier B.V. All rights reserved.
Mg2 Si phase
1.
Introduction
Magnesium alloys are the lightest structural alloys commercially available and have great potential for applications in automotive, aerospace industries and others. However, in recent years, improving the elevated temperature properties has become a critical issue for the further application of magnesium alloys. Previous investigations (Luo and Pekguleryuz, 1994; Dargusch et al., 2004) showed that the Mg–Al–Si-based alloys are a potential elevated temperature magnesium alloys, because the Mg2 Si phase in the Mg–Al–Si-based alloys has high melting point, high hardness, low density, high elastic modulus and low thermal expansion coefficient, and the Mg2 Si phase is very stable and can impede grain boundary sliding at elevated temperature (Lu et al., 2001). However, the challenge of Mg–Al–Si-based alloys is that the Mg2 Si phase is prone
to forming undesirable, coarse Chinese script shape under low solidification rate, which would damage the mechanical properties of the magnesium alloys. Due to the above-mentioned reason, the research about modification of Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys by microalloying, has received much attention all over the world, and consequently many researches have been carried out. It had been reported that the Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys could be modified by Sb (Yuan et al., 2002; Srinivasan et al., 2005; Nam et al., 2006), Ca and P additions (Quimby et al., 2006; Kim et al., 1999; Yoo et al., 2004). However, some investigations also indicated that Sb addition was not an effective modifier of Mg2 Si phase (Quimby et al., 2006), Ca addition could result in cast defects such as a hot-crack (Tang et al., 2005), P addition could produce ignition and the amount of P addition was also diffi-
∗ Corresponding author at: Materials Science & Engineering College, Chongqing Institute of Technology, Chongqing 400050, PR China. Tel.: +86 23 68667455; fax: +86 23 68667714. E-mail address: [email protected] (Y. Mingbo). 0924-0136/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.12.041
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 6 ( 2 0 0 8 ) 374–381
cult to control (Yuan et al., 2002). Therefore, further research about the modification of Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys needs to be considered. It is well known that the semi-solid metal (SSM) processing which has low processing temperature that can effectively resolve the problem of oxidation and burning of magnesium alloys, now is becoming an important forming process for the fabrication of magnesium alloys parts. In general, the semi-solid metal (SSM) processing is mainly composed of semi-solid material production, partial remelting and thixoforming, thereinto the semi-solid materials production is the most important one. At present, the methods of obtaining semi-solid materials with non-dendritic microstructure include mechanical stirring (MS), electromagnetic stirring (ES), strain-induced melt activation (SIMA), spray deposition (SD), liquidus cast and semi-solid isothermal heat treatment (Chen et al., 2001; Kim et al., 2002). Among these methods, the semi-solid isothermal heat treatment is a new way being found in the 1990s, which omits the special procedure to fabricate the semi-solid materials but fulfils the semisolid non-dendritic microstructure during heating prior to thixoforming. In addition, recent results indicated that the dendtrtic Mg2 Si phases in Al–Si alloys could be modified to globular and/or elliptic shapes by semi-solid isothermal heat treatment (Qin et al., 2005). Therefore, the semi-solid isothermal heat treatment is thought as a potential method for the modification of Chinese script shaped Mg2 Si phase in Mg–Al–Si-based magnesium alloys. But up to now, the investigation about the effect of semi-solid isothermal treatment on the microstructure of Mg–Al–Si-based alloys has not been carried out. In order to provide a theoretical guide for the properties improving of Mg–Al–Si-based magnesium alloys by SSM processing, the present work investigates the effect of semi-solid isothermal heat treatment on the microstructure of Mg–6Al–1Zn–0.7Si alloy, and particular attention is paid to the semi-solid microstructural evolution and the modification of Chinese script shaped Mg2 Si phase.
2.
Experimental procedure
The Mg–6Al–1Zn–0.7Si alloy was prepared by adding the following materials: commercial AM60 alloy, pure Al, Mg and Zn (>99.9 wt%), and Al–30 wt%Si master alloy. The experimental alloy was melted in a crucible resistance furnace and protected by a flux addition. After treated by C2 Cl6 , the melt was held at 740 ◦ C for 20 min and then poured into a permanent mould whose size of mould cavity is ø8 mm × 120 mm. The actual composition of experimental alloy was analyzed and given as following (wt%): 92.37 Mg, 5.92 Al, 0.79 Zn, 0.24 Mn and 0.68 Si. The semi-solid isothermal treatment experiment was carried out in a box-like resistance furnace. Firstly, the differential scanning calorimetry (DSC) was carried out by using a NETZSCH STA 449C system in order to obtain the temperature of semi-solid isothermal heat treatment. Sample weighted around 30 mg was heated in a flowing argon atmosphere from 30 ◦ C to 700 ◦ C for 5 min before being cooled down to 100 ◦ C. The heating curve was recorded at a controlling speed of 15 ◦ C/min. Fig. 1 shows the DSC heating curve of the experimental alloy. As shown in Fig. 1, there are two peaks in
375
Fig. 1 – DSC heating curve of the as-cast experimental alloy.
heating curve, respectively corresponding to the melting of ␣-Mg matrix and second phase transformation, and the onset and peak temperatures of the matrix melting are 518.8 ◦ C and 612.4 ◦ C, respectively. Therefore, in order to obtain relatively high liquid fraction and reduce the oxidation and burning, the temperatures of semi-solid isothermal heat treatment, 575 ◦ C, 580 ◦ C and 585 ◦ C which approximately correspond to a range of 35–43% liquid content (the detailed calculation is given in the following section), were selected for the experimental alloy. Samples with the dimensions of 15 mm long and 8 mm in diameter were cut from the above-mentioned alloy and put into a box-like chamber of resistance furnace. The specimen temperature was measured by means of a thermocouple which was fixed on the specimen surface. The heating rate and temperature fluctuating range of the resistance furnace were 1 ◦ C/min and ±1 ◦ C, respectively. After samples reached the experimental temperature, they were respectively held for 30 min, 60 min, 90 min and 120 min, and then taken out for water quenching quickly. Then the specimens for the following semi-solid microstructural analysis were obtained. Fig. 2 shows the schematic drawing of the sample location for the semi-solid microstructural analysis. After the as-cast and/or semi-solid samples were etched in a 8% nitric acid solution in distilled water, the centeral microstructures of these samples were examined by Olympus Optical microscope and JOEL JSM-6460LV type scanning
Fig. 2 – Schematic drawing of the sample location for semi-solid microstructural analysis.
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Table 1 – EDS results of the as-cast experimental alloy (at.%) Positions
Fig.4b-A Fig.4b-B Fig.4b-C Fig.4b-D
Elements
Total (%)
Mg
Al
Zn
Si
92.26 67.29 64.16 77.00
4.31 2.02 32.50 21.69
3.43 – 3.34 1.31
0.11 30.69 – –
100 100 100 100
3.
Results and discussion
fore, under the permanent moulding casting of this paper, the Mg2 Si phases in the experimental alloy exhibit typical Chinese script shaped morphology. It is well known that the presence of fine and uniform phases distributed along the grain boundaries in microstructure is easier to act as an effective straddle to the dislocation motion thus improving the properties of engineering alloys (Balasubramani et al., in press). Apparently, the coarse Chinese script shaped Mg2 Si phases in the as-cast experimental alloy would give a detrimental effect on the mechanical properties of the alloy since the long cracks can easily nucleate along the interface between Chinese script shaped Mg2 Si particles and ␣-Mg matrix (Kim et al., 1999).
3.1.
As-cast microstructure
3.2.
Fig. 3 – XRD result of the as-cast experimental alloy.
electron microscope (SEM) equipped with Oxford energy dispersive X-ray spectrometer (EDS). The grain size was analyzed by the standard linear intercept method using a Olympus stereomicroscope. The phases in the experimental alloy were also analyzed by D/Max-1200X type analyzer operated at 40 kV and 30 mA.
Fig. 3 shows the XRD result of the as-cast experimental alloy. It is found from Fig. 3 that, similar to the common Mg–Al–Sibased magnesium alloys (Dargusch et al., 2004; Srinivasan et al., 2005; Nam et al., 2006; Quimby et al., 2006; Kim et al., 1999), the main phases in as-cast experimental alloy are ␣Mg, Mg17 Al12 and Mg2 Si. Fig. 4 shows the optical and SEM images of the as-cast experimental alloy. As shown in Fig. 4, the experimental alloy exhibits typical dendritic microstructure (Fig. 4a), and in Fig. 4b the coarse Chinese script shaped phases and white second phases are also observed. According to the EDS result of the as-cast experimental alloy (Table 1), the Chinese script shaped phases and white second phases in Fig. 4b are Mg2 Si and Mg17 Al12 , respectively. In general, the Mg2 Si phases in Mg–Al–Si-based alloys are prone to forming coarse Chinese script shape under low solidification rate (Luo and Pekguleryuz, 1994; Dargusch et al., 2004). There-
Semi-solid microstructural evolution
Fig. 5 shows the semi-solid microstructural evolution of the experimental alloy treated at 580 ◦ C for different holding times from 30 min to 120 min. It is found from Fig. 5 that during the initial period of holding (30 min), the liquid is dispersed discontinuously, and the “entrapped liquid” pool is also present inside the primary solids. With the holding time increasing from 60 to 120 min, the amount of the liquid phase both inside the grains and between the grains increases, and the liquid is dispersed continuously along grain boundaries. The mechanism of partial melting in the initial stage of the semi-solid isothermal holding, could be related to the dissolutions of the last solidified phases with low melting temperatures (Wang et al., 1997; Zoqui and Robert, 2001). Fig. 6 shows the semi-solid microstructures of the experimental alloy treated at different temperatures for 120 min. Comparing Figs. 4–6, it is found that the semi-solid microstruc-
Fig. 4 – As-cast microstructures of the experimental alloy for (a) optical image; (b) SEM image.
j o u r n a l o f m a t e r i a l s p r o c e s s i n g t e c h n o l o g y 2 0 6 ( 2 0 0 8 ) 374–381
377
Fig. 5 – Semi-solid microstructures of the experimental alloy treated at 580 ◦ C for different times for: (a) 30 min; (b) 60 min; (c) 90 min; (d) 120 min.
ture, which comprises of unmelted primary solid particles surrounded by a transformation product of the remelted liquid, is essentially different with the as-cast microstructure. As shown in Figs. 5 and 6, the morphology of the unmelted primary solid particles is approximately globular shape, and with the increasing of the holding time and semi-solid isothermal temperature, the size and globular trendance of unmelted primary solid particles decrease and become more obvious, respectively. For example, after treated at 575 ◦ C, 580 ◦ C and 585 ◦ C for 120 min, the measured average size of unmelted primary solid particles in semi-solid experimental alloy is 67 m, 53 m and 48 m, respectively. In addition, it is surprising that the coarsening of unmelted primary solid particles by the mechanisms of coalescence and Ostwald ripening (Czerwinski, 2005; Wang et al., in press; Chen et al., 2003), is not observed under the experimental condition of this paper. At the same time, another surprising result is found from Figs. 5 and 6, namely the volume fraction of liquid phase in the semi-solid microstructures of experimental alloy is relatively low. According to the Scheil equation (Eq. (1)) (Zhang et al., 2007), the volume fraction of liquid fL in the semi-solid experimental alloy should be a constant value at a given temperature of semi-solid isothermal heat treatment, with the assumption that the homogenization of the liquid is complete and no diffusion is in the solid.
fL =
T − T (1/1−k) M L TM − T
(1)
where TM is the melting point of pure metal, TL the liquidus temperature of the alloy, T the temperature of semi-solid isothermal heat treatment, k is the equilibrium distribution coefficient. Under the assumption that the interactions among various alloying elements are neglected, the k of the multicomponent is treated using the equivalent pseudo-binary method (Li et al., 1998), which is expressed as
k=
mck ii i mi ci
(2)
where ci is the initial composition of the i alloying element, mi and ki are the liquidus slope and partition coefficient of i element in the Mg-i binary alloys, respectively. Under the experimental condition of this paper, i is Zn and Al elements, respectively. Based on the DSC result (Fig. 1), the value of TL is 612.4 ◦ C. In the present paper, the values of TM is 650 ◦ C (Wang et al., in press; Zhang et al., 2007). According to the reference (Zeng et al., 2006), mAl = −6.87; mZn = −6.04; kAl = 0.37; kZn = 0.12. cAl and cZn are 5.92 and 0.79 wt%, respectively. So according to this Eqs. (1) and (2), the liquid fractions of semi-solid experimental alloy at 575 ◦ C, 580 ◦ C and 585 ◦ C are 0.35, 0.39 and 0.43, respectively. While in the present experiment, the practical values of fL measured by the auto image analysis system respectively are 0.12, 0.16 and 0.21 for the samples treated at 575 ◦ C, 580 ◦ C and 585 ◦ C for 120 min. The possible reason is that the liquid–solid diffusion equilibrium is not achieved due to the short isothermal holding time in the present experiment. However, the exact reason is not completely clear. It is
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Fig. 6 – Semi-solid microstructures of the experimental alloy treated at different temperatures for 120 min for: (a) 575 ◦ C; (b) 580 ◦ C; (c) 585 ◦ C.
that little Si dissolves in the matrix after semi-solid isothermal heat treatment, and Si is still in the form of Mg2 Si in the alloy. However, it is surprisingly observed from the Fig. 7 that the Mg2 Si phases in the semi-solid experimental alloy exhibit granule and/or polygon shapes, indicating that the semi-solid isothermal heat treatment could modify the Chinese script shaped Mg2 Si phases in the as-cast alloy.
the subject for further study in our group. In addition, considering that a 40–60% liquid content is generally defined as a working range of SSM in order to obtain adequate flow ability for a semi-solid slurry (Czerwinski, 2005), then in further study a longer isothermal holding time than 120 min and more appropriate temperature range of semi-solid isothermal treatment should be considered for the experimental alloy. Fig. 7 shows the SEM images of the experimental alloy treated at 580 ◦ C and 585 ◦ C for 120 min. In Fig. 7, the morphology of precipitates whose EDS results are listed in Table 2, is clearly exhibited, and in Fig. 7c and d the fine secondary ␣Mg phases surrounded by eutectic Mg17 Al12 phases, are also observed. At the same time, the eutectic Mg17 Al12 phases both inside the grains and between the grains are found to respectively exhibit granule and sawtooth shapes, which is consistent with the results of other Mg–Al-based alloys (Czerwinski and Zielinska-Lipiec, 2005; Chen et al., 2007; Zhang et al., 2008). In addition, it is proved from Tables 1 and 2
3.3.
Modification mechanism of Mg2 Si phase
Since the melting temperature of Mg2 Si phase is 1085 ◦ C (Czerwinski, 2005), the modification of Chinese script shaped Mg2 Si phase during the semi-solid isothermal heat treatment could not be achieved by melted mode. It is well known that the temperature and solute concentration exist fluctuation during the solidification process of as-cast experimental alloy. Therefore, the surface of Mg2 Si phase should exist curvature fluctuation. Based on the above analysis, one pos-
Table 2 – EDS results of the semi-solid experimental alloy (at.%) Elements Mg Al Zn Si Total (%)
Fig.7a-A
Fig.7a-B
Fig.7a-C
Fig.7a-D
Fig.7a-E
Fig.7b-A
Fig.7b-B
Fig.7b-C
Fig.7b-D
Fig.7b-E
93.70 5.14 1.04 0.12
67.14 30.04 2.73 0.09
60.11 0.63 – 39.26
58.64 0.74 – 40.62
61.43 34.30 4.27 –
94.06 5.81 – 0.13
63.12 32.85 3.92 0.11
62.63 0.67 – 36.70
61.53 0.70 – 37.77
63.56 31.22 5.22 –
100
100
100
100
100
100
100
100
100
100
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379
Fig. 7 – SEM images of the secondary microstructure for experimental alloy for (a) treated at 580 ◦ C for 120min; (b) treated at 585 ◦ C for 120min; (c) local magnification in (a); (d) local magnification in (b).
sible explanation for the modification mechanism of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment, could be obtained by using the Gibbs–Thomson effect (Nishioka and Maksimov, 1996). According to the Gibbs–Thomson formula, the Si concentration in the matrix corresponded to the site where the Mg2 Si phase has larger curvature, can be expressed as (Nishioka and Maksimov, 1996): C˛ (r) = C˛ (∞)exp
2v B
kB Tr
(3)
Where C␣ (r) is the Si concentration at the position with a curvature radius, r, C␣ (∞) is the Si concentration at flat interface, the surface tension, vB the volume of Si atom, T the temperature, kB is the coefficient related to the shape. According to the Eq. (3), smaller the curvature radius, higher the Si concentration is. Since the curvature radius of different positions for a Chinese script shaped Mg2 Si particle
might exist difference, then a gradients of Si concentration could be created between these positions. Therefore, during the semi-solid isothermal heat treatment, the Si atoms would diffuse from the position where the curvature and Si concentration are respectively large and high to flat interface where the Si concentration is lower, and then the balance of local Si concentration between these positions could be broken. Furthermore, in order to keep the balance of Si concentration, these positions with larger curvature could be dissolved. Oppositely, due to the supersaturation of Si concentration, the Mg2 Si phases could form in the ␣-Mg matrix corresponded to the flat interface. As a result, these positions with larger curvature could break, and then the granule and/or polygon shaped Mg2 Si phase whose different positions have close curvature radius, would form. The process can be illustrated in Fig. 8. As shown in Fig. 8, the curvatures of these positions such as ‘A’ and ‘B’ in Fig. 8a are larger, then these positions could be first dissolved (Fig. 8b). Furthermore, with the gradual
Fig. 8 – Sketch of the dissolving and breaking process of the Chinese script shaped Mg2 Si phase in larger curvature site.
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diffusion of Si atoms, the Mg2 Si particle could break in these positions such as ‘A’ and ‘B’ in Fig. 8a, as shown in Fig. 8c and d. Finally, with the prolonging of the semi-solid isothermal holding time, the granule and/or polygon shaped Mg2 Si particles whose different positions have close curvature radius would form (Fig. 8d and e). In spite of the above, the modification mechanism of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment is not completely clear, further investigation needs to be carried out. For example, as shown in Fig. 4b, the majority of Mg17 Al12 and Mg2 Si phases in the as-cast experimental alloy connects with each other. In addition, the Mg2 Si phases in the semi-solid experimental alloy tightly contact with the partial remelted liquid at the grain boundaries, as shown in Fig. 7. Therefore, one question whether the modification of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment was effected by the penetration of remelted liquid and the diffusion of Al atoms in the remelted liquid, still remains. The question is the subject for further study in our group.
4.
Conclusion
(1) Results show that it is possible to produce the Mg–6A1–1Zn–0.7Si alloy with non-dendritic microstructure by the semi-solid isothermal heat treatment. After treated at 575–585 ◦ C for 120 min, the experimental alloy can obtain a non-dendritic microstructure with a 12–21% liquid content and an average size range of 48–67 m of the unmelted primary solid particles. (2) The semi-solid isothermal heat treatment can modify the Chinese script shaped Mg2 Si phases in Mg–6A1–1Zn–0.7Si alloy. After treated at 580 ◦ C or 585 ◦ C for 120 min, the morphology of Mg2 Si phases in the experimental alloy changes from the initial Chinese script shape to granule and/or polygon shapes.
Acknowledgements The present work was supported by the National Natural Science Funds for Distinguished Young Scholar (No.50725413), the Natural Science Foundation Project of CQ CSTC (No.2007BB4400), the Major State Basic Research Development Program of China (973)(No.2007CB613704), and Chongqing Science and Technology Commission in China (No.2006AA4012-9-6).
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journal homepage: www.elsevier.com/locate/jmatprotec
Effect of semi-solid isothermal heat treatment on the microstructure of Mg–6A1–1Zn–0.7Si alloy Yang Mingbo a,b,∗ , Pan Fusheng b , Cheng Renju b , Bai Liang b a b
Materials Science & Engineering College, Chongqing Institute of Technology, Chongqing 400050, PR China Materials Science & Engineering College, Chongqing University, Chongqing 400030, PR China
a r t i c l e
i n f o
a b s t r a c t
Article history:
In this paper, the effect of semi-solid isothermal heat treatment on the microstructure
Received 2 August 2007
of Mg–6A1–1Zn–0.7Si alloy, especially on the semi-solid microstructural evolution and the
Received in revised form
modification of Chinese script shaped Mg2 Si phase, are first investigated. The research
29 September 2007
results indicate that it is possible to produce the Mg–6A1–1Zn–0.7Si alloy with non-dendritic
Accepted 14 December 2007
microstructure by the semi-solid isothermal heat treatment. After treated at 575–585 ◦ C for 120 min, the experimental alloy can obtain a non-dendritic microstructure with a 12–21% liquid content and an average size range of 48–67 m of the unmelted primary solid parti-
Keywords:
cles. In addition, the semi-solid isothermal heat treatment can modify the Chinese script
Mg–A1–Si-based alloys
shaped Mg2 Si phases in experimental alloy. After treated at 580 ◦ C or 585 ◦ C for 120 min,
Semi-solid
the morphology of Mg2 Si phases in the experimental alloy changes from the initial Chinese
Semi-solid isothermal heat
script shape to granule and/or polygon shapes.
treatment
© 2007 Elsevier B.V. All rights reserved.
Mg2 Si phase
1.
Introduction
Magnesium alloys are the lightest structural alloys commercially available and have great potential for applications in automotive, aerospace industries and others. However, in recent years, improving the elevated temperature properties has become a critical issue for the further application of magnesium alloys. Previous investigations (Luo and Pekguleryuz, 1994; Dargusch et al., 2004) showed that the Mg–Al–Si-based alloys are a potential elevated temperature magnesium alloys, because the Mg2 Si phase in the Mg–Al–Si-based alloys has high melting point, high hardness, low density, high elastic modulus and low thermal expansion coefficient, and the Mg2 Si phase is very stable and can impede grain boundary sliding at elevated temperature (Lu et al., 2001). However, the challenge of Mg–Al–Si-based alloys is that the Mg2 Si phase is prone
to forming undesirable, coarse Chinese script shape under low solidification rate, which would damage the mechanical properties of the magnesium alloys. Due to the above-mentioned reason, the research about modification of Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys by microalloying, has received much attention all over the world, and consequently many researches have been carried out. It had been reported that the Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys could be modified by Sb (Yuan et al., 2002; Srinivasan et al., 2005; Nam et al., 2006), Ca and P additions (Quimby et al., 2006; Kim et al., 1999; Yoo et al., 2004). However, some investigations also indicated that Sb addition was not an effective modifier of Mg2 Si phase (Quimby et al., 2006), Ca addition could result in cast defects such as a hot-crack (Tang et al., 2005), P addition could produce ignition and the amount of P addition was also diffi-
∗ Corresponding author at: Materials Science & Engineering College, Chongqing Institute of Technology, Chongqing 400050, PR China. Tel.: +86 23 68667455; fax: +86 23 68667714. E-mail address: [email protected] (Y. Mingbo). 0924-0136/$ – see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2007.12.041
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cult to control (Yuan et al., 2002). Therefore, further research about the modification of Chinese script shaped Mg2 Si phases in Mg–Al–Si-based magnesium alloys needs to be considered. It is well known that the semi-solid metal (SSM) processing which has low processing temperature that can effectively resolve the problem of oxidation and burning of magnesium alloys, now is becoming an important forming process for the fabrication of magnesium alloys parts. In general, the semi-solid metal (SSM) processing is mainly composed of semi-solid material production, partial remelting and thixoforming, thereinto the semi-solid materials production is the most important one. At present, the methods of obtaining semi-solid materials with non-dendritic microstructure include mechanical stirring (MS), electromagnetic stirring (ES), strain-induced melt activation (SIMA), spray deposition (SD), liquidus cast and semi-solid isothermal heat treatment (Chen et al., 2001; Kim et al., 2002). Among these methods, the semi-solid isothermal heat treatment is a new way being found in the 1990s, which omits the special procedure to fabricate the semi-solid materials but fulfils the semisolid non-dendritic microstructure during heating prior to thixoforming. In addition, recent results indicated that the dendtrtic Mg2 Si phases in Al–Si alloys could be modified to globular and/or elliptic shapes by semi-solid isothermal heat treatment (Qin et al., 2005). Therefore, the semi-solid isothermal heat treatment is thought as a potential method for the modification of Chinese script shaped Mg2 Si phase in Mg–Al–Si-based magnesium alloys. But up to now, the investigation about the effect of semi-solid isothermal treatment on the microstructure of Mg–Al–Si-based alloys has not been carried out. In order to provide a theoretical guide for the properties improving of Mg–Al–Si-based magnesium alloys by SSM processing, the present work investigates the effect of semi-solid isothermal heat treatment on the microstructure of Mg–6Al–1Zn–0.7Si alloy, and particular attention is paid to the semi-solid microstructural evolution and the modification of Chinese script shaped Mg2 Si phase.
2.
Experimental procedure
The Mg–6Al–1Zn–0.7Si alloy was prepared by adding the following materials: commercial AM60 alloy, pure Al, Mg and Zn (>99.9 wt%), and Al–30 wt%Si master alloy. The experimental alloy was melted in a crucible resistance furnace and protected by a flux addition. After treated by C2 Cl6 , the melt was held at 740 ◦ C for 20 min and then poured into a permanent mould whose size of mould cavity is ø8 mm × 120 mm. The actual composition of experimental alloy was analyzed and given as following (wt%): 92.37 Mg, 5.92 Al, 0.79 Zn, 0.24 Mn and 0.68 Si. The semi-solid isothermal treatment experiment was carried out in a box-like resistance furnace. Firstly, the differential scanning calorimetry (DSC) was carried out by using a NETZSCH STA 449C system in order to obtain the temperature of semi-solid isothermal heat treatment. Sample weighted around 30 mg was heated in a flowing argon atmosphere from 30 ◦ C to 700 ◦ C for 5 min before being cooled down to 100 ◦ C. The heating curve was recorded at a controlling speed of 15 ◦ C/min. Fig. 1 shows the DSC heating curve of the experimental alloy. As shown in Fig. 1, there are two peaks in
375
Fig. 1 – DSC heating curve of the as-cast experimental alloy.
heating curve, respectively corresponding to the melting of ␣-Mg matrix and second phase transformation, and the onset and peak temperatures of the matrix melting are 518.8 ◦ C and 612.4 ◦ C, respectively. Therefore, in order to obtain relatively high liquid fraction and reduce the oxidation and burning, the temperatures of semi-solid isothermal heat treatment, 575 ◦ C, 580 ◦ C and 585 ◦ C which approximately correspond to a range of 35–43% liquid content (the detailed calculation is given in the following section), were selected for the experimental alloy. Samples with the dimensions of 15 mm long and 8 mm in diameter were cut from the above-mentioned alloy and put into a box-like chamber of resistance furnace. The specimen temperature was measured by means of a thermocouple which was fixed on the specimen surface. The heating rate and temperature fluctuating range of the resistance furnace were 1 ◦ C/min and ±1 ◦ C, respectively. After samples reached the experimental temperature, they were respectively held for 30 min, 60 min, 90 min and 120 min, and then taken out for water quenching quickly. Then the specimens for the following semi-solid microstructural analysis were obtained. Fig. 2 shows the schematic drawing of the sample location for the semi-solid microstructural analysis. After the as-cast and/or semi-solid samples were etched in a 8% nitric acid solution in distilled water, the centeral microstructures of these samples were examined by Olympus Optical microscope and JOEL JSM-6460LV type scanning
Fig. 2 – Schematic drawing of the sample location for semi-solid microstructural analysis.
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Table 1 – EDS results of the as-cast experimental alloy (at.%) Positions
Fig.4b-A Fig.4b-B Fig.4b-C Fig.4b-D
Elements
Total (%)
Mg
Al
Zn
Si
92.26 67.29 64.16 77.00
4.31 2.02 32.50 21.69
3.43 – 3.34 1.31
0.11 30.69 – –
100 100 100 100
3.
Results and discussion
fore, under the permanent moulding casting of this paper, the Mg2 Si phases in the experimental alloy exhibit typical Chinese script shaped morphology. It is well known that the presence of fine and uniform phases distributed along the grain boundaries in microstructure is easier to act as an effective straddle to the dislocation motion thus improving the properties of engineering alloys (Balasubramani et al., in press). Apparently, the coarse Chinese script shaped Mg2 Si phases in the as-cast experimental alloy would give a detrimental effect on the mechanical properties of the alloy since the long cracks can easily nucleate along the interface between Chinese script shaped Mg2 Si particles and ␣-Mg matrix (Kim et al., 1999).
3.1.
As-cast microstructure
3.2.
Fig. 3 – XRD result of the as-cast experimental alloy.
electron microscope (SEM) equipped with Oxford energy dispersive X-ray spectrometer (EDS). The grain size was analyzed by the standard linear intercept method using a Olympus stereomicroscope. The phases in the experimental alloy were also analyzed by D/Max-1200X type analyzer operated at 40 kV and 30 mA.
Fig. 3 shows the XRD result of the as-cast experimental alloy. It is found from Fig. 3 that, similar to the common Mg–Al–Sibased magnesium alloys (Dargusch et al., 2004; Srinivasan et al., 2005; Nam et al., 2006; Quimby et al., 2006; Kim et al., 1999), the main phases in as-cast experimental alloy are ␣Mg, Mg17 Al12 and Mg2 Si. Fig. 4 shows the optical and SEM images of the as-cast experimental alloy. As shown in Fig. 4, the experimental alloy exhibits typical dendritic microstructure (Fig. 4a), and in Fig. 4b the coarse Chinese script shaped phases and white second phases are also observed. According to the EDS result of the as-cast experimental alloy (Table 1), the Chinese script shaped phases and white second phases in Fig. 4b are Mg2 Si and Mg17 Al12 , respectively. In general, the Mg2 Si phases in Mg–Al–Si-based alloys are prone to forming coarse Chinese script shape under low solidification rate (Luo and Pekguleryuz, 1994; Dargusch et al., 2004). There-
Semi-solid microstructural evolution
Fig. 5 shows the semi-solid microstructural evolution of the experimental alloy treated at 580 ◦ C for different holding times from 30 min to 120 min. It is found from Fig. 5 that during the initial period of holding (30 min), the liquid is dispersed discontinuously, and the “entrapped liquid” pool is also present inside the primary solids. With the holding time increasing from 60 to 120 min, the amount of the liquid phase both inside the grains and between the grains increases, and the liquid is dispersed continuously along grain boundaries. The mechanism of partial melting in the initial stage of the semi-solid isothermal holding, could be related to the dissolutions of the last solidified phases with low melting temperatures (Wang et al., 1997; Zoqui and Robert, 2001). Fig. 6 shows the semi-solid microstructures of the experimental alloy treated at different temperatures for 120 min. Comparing Figs. 4–6, it is found that the semi-solid microstruc-
Fig. 4 – As-cast microstructures of the experimental alloy for (a) optical image; (b) SEM image.
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377
Fig. 5 – Semi-solid microstructures of the experimental alloy treated at 580 ◦ C for different times for: (a) 30 min; (b) 60 min; (c) 90 min; (d) 120 min.
ture, which comprises of unmelted primary solid particles surrounded by a transformation product of the remelted liquid, is essentially different with the as-cast microstructure. As shown in Figs. 5 and 6, the morphology of the unmelted primary solid particles is approximately globular shape, and with the increasing of the holding time and semi-solid isothermal temperature, the size and globular trendance of unmelted primary solid particles decrease and become more obvious, respectively. For example, after treated at 575 ◦ C, 580 ◦ C and 585 ◦ C for 120 min, the measured average size of unmelted primary solid particles in semi-solid experimental alloy is 67 m, 53 m and 48 m, respectively. In addition, it is surprising that the coarsening of unmelted primary solid particles by the mechanisms of coalescence and Ostwald ripening (Czerwinski, 2005; Wang et al., in press; Chen et al., 2003), is not observed under the experimental condition of this paper. At the same time, another surprising result is found from Figs. 5 and 6, namely the volume fraction of liquid phase in the semi-solid microstructures of experimental alloy is relatively low. According to the Scheil equation (Eq. (1)) (Zhang et al., 2007), the volume fraction of liquid fL in the semi-solid experimental alloy should be a constant value at a given temperature of semi-solid isothermal heat treatment, with the assumption that the homogenization of the liquid is complete and no diffusion is in the solid.
fL =
T − T (1/1−k) M L TM − T
(1)
where TM is the melting point of pure metal, TL the liquidus temperature of the alloy, T the temperature of semi-solid isothermal heat treatment, k is the equilibrium distribution coefficient. Under the assumption that the interactions among various alloying elements are neglected, the k of the multicomponent is treated using the equivalent pseudo-binary method (Li et al., 1998), which is expressed as
k=
mck ii i mi ci
(2)
where ci is the initial composition of the i alloying element, mi and ki are the liquidus slope and partition coefficient of i element in the Mg-i binary alloys, respectively. Under the experimental condition of this paper, i is Zn and Al elements, respectively. Based on the DSC result (Fig. 1), the value of TL is 612.4 ◦ C. In the present paper, the values of TM is 650 ◦ C (Wang et al., in press; Zhang et al., 2007). According to the reference (Zeng et al., 2006), mAl = −6.87; mZn = −6.04; kAl = 0.37; kZn = 0.12. cAl and cZn are 5.92 and 0.79 wt%, respectively. So according to this Eqs. (1) and (2), the liquid fractions of semi-solid experimental alloy at 575 ◦ C, 580 ◦ C and 585 ◦ C are 0.35, 0.39 and 0.43, respectively. While in the present experiment, the practical values of fL measured by the auto image analysis system respectively are 0.12, 0.16 and 0.21 for the samples treated at 575 ◦ C, 580 ◦ C and 585 ◦ C for 120 min. The possible reason is that the liquid–solid diffusion equilibrium is not achieved due to the short isothermal holding time in the present experiment. However, the exact reason is not completely clear. It is
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Fig. 6 – Semi-solid microstructures of the experimental alloy treated at different temperatures for 120 min for: (a) 575 ◦ C; (b) 580 ◦ C; (c) 585 ◦ C.
that little Si dissolves in the matrix after semi-solid isothermal heat treatment, and Si is still in the form of Mg2 Si in the alloy. However, it is surprisingly observed from the Fig. 7 that the Mg2 Si phases in the semi-solid experimental alloy exhibit granule and/or polygon shapes, indicating that the semi-solid isothermal heat treatment could modify the Chinese script shaped Mg2 Si phases in the as-cast alloy.
the subject for further study in our group. In addition, considering that a 40–60% liquid content is generally defined as a working range of SSM in order to obtain adequate flow ability for a semi-solid slurry (Czerwinski, 2005), then in further study a longer isothermal holding time than 120 min and more appropriate temperature range of semi-solid isothermal treatment should be considered for the experimental alloy. Fig. 7 shows the SEM images of the experimental alloy treated at 580 ◦ C and 585 ◦ C for 120 min. In Fig. 7, the morphology of precipitates whose EDS results are listed in Table 2, is clearly exhibited, and in Fig. 7c and d the fine secondary ␣Mg phases surrounded by eutectic Mg17 Al12 phases, are also observed. At the same time, the eutectic Mg17 Al12 phases both inside the grains and between the grains are found to respectively exhibit granule and sawtooth shapes, which is consistent with the results of other Mg–Al-based alloys (Czerwinski and Zielinska-Lipiec, 2005; Chen et al., 2007; Zhang et al., 2008). In addition, it is proved from Tables 1 and 2
3.3.
Modification mechanism of Mg2 Si phase
Since the melting temperature of Mg2 Si phase is 1085 ◦ C (Czerwinski, 2005), the modification of Chinese script shaped Mg2 Si phase during the semi-solid isothermal heat treatment could not be achieved by melted mode. It is well known that the temperature and solute concentration exist fluctuation during the solidification process of as-cast experimental alloy. Therefore, the surface of Mg2 Si phase should exist curvature fluctuation. Based on the above analysis, one pos-
Table 2 – EDS results of the semi-solid experimental alloy (at.%) Elements Mg Al Zn Si Total (%)
Fig.7a-A
Fig.7a-B
Fig.7a-C
Fig.7a-D
Fig.7a-E
Fig.7b-A
Fig.7b-B
Fig.7b-C
Fig.7b-D
Fig.7b-E
93.70 5.14 1.04 0.12
67.14 30.04 2.73 0.09
60.11 0.63 – 39.26
58.64 0.74 – 40.62
61.43 34.30 4.27 –
94.06 5.81 – 0.13
63.12 32.85 3.92 0.11
62.63 0.67 – 36.70
61.53 0.70 – 37.77
63.56 31.22 5.22 –
100
100
100
100
100
100
100
100
100
100
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379
Fig. 7 – SEM images of the secondary microstructure for experimental alloy for (a) treated at 580 ◦ C for 120min; (b) treated at 585 ◦ C for 120min; (c) local magnification in (a); (d) local magnification in (b).
sible explanation for the modification mechanism of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment, could be obtained by using the Gibbs–Thomson effect (Nishioka and Maksimov, 1996). According to the Gibbs–Thomson formula, the Si concentration in the matrix corresponded to the site where the Mg2 Si phase has larger curvature, can be expressed as (Nishioka and Maksimov, 1996): C˛ (r) = C˛ (∞)exp
2v B
kB Tr
(3)
Where C␣ (r) is the Si concentration at the position with a curvature radius, r, C␣ (∞) is the Si concentration at flat interface, the surface tension, vB the volume of Si atom, T the temperature, kB is the coefficient related to the shape. According to the Eq. (3), smaller the curvature radius, higher the Si concentration is. Since the curvature radius of different positions for a Chinese script shaped Mg2 Si particle
might exist difference, then a gradients of Si concentration could be created between these positions. Therefore, during the semi-solid isothermal heat treatment, the Si atoms would diffuse from the position where the curvature and Si concentration are respectively large and high to flat interface where the Si concentration is lower, and then the balance of local Si concentration between these positions could be broken. Furthermore, in order to keep the balance of Si concentration, these positions with larger curvature could be dissolved. Oppositely, due to the supersaturation of Si concentration, the Mg2 Si phases could form in the ␣-Mg matrix corresponded to the flat interface. As a result, these positions with larger curvature could break, and then the granule and/or polygon shaped Mg2 Si phase whose different positions have close curvature radius, would form. The process can be illustrated in Fig. 8. As shown in Fig. 8, the curvatures of these positions such as ‘A’ and ‘B’ in Fig. 8a are larger, then these positions could be first dissolved (Fig. 8b). Furthermore, with the gradual
Fig. 8 – Sketch of the dissolving and breaking process of the Chinese script shaped Mg2 Si phase in larger curvature site.
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diffusion of Si atoms, the Mg2 Si particle could break in these positions such as ‘A’ and ‘B’ in Fig. 8a, as shown in Fig. 8c and d. Finally, with the prolonging of the semi-solid isothermal holding time, the granule and/or polygon shaped Mg2 Si particles whose different positions have close curvature radius would form (Fig. 8d and e). In spite of the above, the modification mechanism of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment is not completely clear, further investigation needs to be carried out. For example, as shown in Fig. 4b, the majority of Mg17 Al12 and Mg2 Si phases in the as-cast experimental alloy connects with each other. In addition, the Mg2 Si phases in the semi-solid experimental alloy tightly contact with the partial remelted liquid at the grain boundaries, as shown in Fig. 7. Therefore, one question whether the modification of Chinese script shaped Mg2 Si phases during semi-solid isothermal heat treatment was effected by the penetration of remelted liquid and the diffusion of Al atoms in the remelted liquid, still remains. The question is the subject for further study in our group.
4.
Conclusion
(1) Results show that it is possible to produce the Mg–6A1–1Zn–0.7Si alloy with non-dendritic microstructure by the semi-solid isothermal heat treatment. After treated at 575–585 ◦ C for 120 min, the experimental alloy can obtain a non-dendritic microstructure with a 12–21% liquid content and an average size range of 48–67 m of the unmelted primary solid particles. (2) The semi-solid isothermal heat treatment can modify the Chinese script shaped Mg2 Si phases in Mg–6A1–1Zn–0.7Si alloy. After treated at 580 ◦ C or 585 ◦ C for 120 min, the morphology of Mg2 Si phases in the experimental alloy changes from the initial Chinese script shape to granule and/or polygon shapes.
Acknowledgements The present work was supported by the National Natural Science Funds for Distinguished Young Scholar (No.50725413), the Natural Science Foundation Project of CQ CSTC (No.2007BB4400), the Major State Basic Research Development Program of China (973)(No.2007CB613704), and Chongqing Science and Technology Commission in China (No.2006AA4012-9-6).
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