Microstructure and thermal characteristics of Mg–Sn alloys as phase change materials for thermal energy storage

Abstract

Latent heat storage proves to be one of the most efficient ways of storing thermal energy. The selection of phase change materials is the key factor in storing thermal energy. In this study, the microstructure and thermal characteristics of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn (wt%) alloys as high temperature phase change materials for thermal energy storage were investigated. The microstructure of Mg–24% Sn alloy mainly consisted of α-Mg matrix and α-Mg + Mg₂Sn eutectic phases. The microstructure of Mg–37% Sn alloy mainly consisted of α-Mg + Mg₂Sn eutectic phases. The microstructure of Mg–37% Sn alloy mainly consisted of the primary Mg₂Sn phase and α-Mg + Mg₂Sn eutectic phases. The melting enthalpies of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn alloys are 105.3, 217.8 and 118.8 J g⁻¹, with the phase change temperatures of 557.6, 554.4 and 557.1 °C, respectively. The melting enthalpy of Mg–37% Sn alloy is the highest among these three alloys due to its higher proportion of α-Mg + Mg₂Sn eutectic phases, comparing with that of Mg–24% Sn and Mg–50% Sn alloys. Besides, the thermal conductivity of alloys decreases with increasing Sn content.

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1. Introduction

Solar thermal power generation could be feasible as a source of power in arid countries, and energy storage systems are widely used due to the intermittent and variable nature of solar resources. Thermal energy storage (TES) proves to be an attractive and economical alternative for large-scale use.¹ Phase change materials (PCMs) have some advantages such as high energy storage density, storing and releasing thermal energy at nearly constant temperature, which are getting more and more attention in TES systems.²

Several the thermal properties are considered when we choose PCMs for TES applications, such as working temperature, melting enthalpy, thermal conductivity, specific heat and thermal stability. The operating temperature of PCMs for TES applications in solar thermal power generation are in the range of 120–1000 °C.³ In this temperature range, the candidates of PCMs generally include molten salts, salt eutectics, metals and alloys. However, one of most disadvantages of inorganic molten salts is low thermal conductivity, which means that it needs a more complex heat exchanger for charging and releasing processes in the TES system. Comparing with inorganic molten salts, metallic PCMs have more prominent advantages, such as high thermal conductivity, high heat storage capacity, stable performance, long life and so on, so it has extensive application prospect in the medium–high temperature thermal storage.^4,5

More than thirty years ago, metallic PCMs were considered as TES materials,^6,7 and have been partly researched for TES applications.^8–10 At present, a lot of research on the aluminum-based alloy thermal storage materials has been conducted at home and abroad. For instance, a waste heat storage device was developed by He and Zhang.¹¹ The device utilized an eutectic Al–Si alloy (AlSi₁₂), whose heat of fusion of alloy was 560 J g⁻¹ and the melting temperature was 576 °C. The reliability of Al–Si alloys as a latent heat energy storage material was investigated by Li et al.¹² They concluded that Al–Si alloys were relatively stable after lots of melting and solidification cycles, and the stability may be improved through having an accurate eutectic composition and controlling cooling rates. The compatibility of Al (60 wt%)–34Mg–6Zn alloy with containment materials such as stainless steel SS304L and carbonaceous steel C20 was studied by Sun et al.⁸ Besides, the behavior of thermal properties of the Al–Mg–Zn alloy according to number of thermal cycles was investigated. The latent heat of melting and the specific heat (the liquid and solid forms) of Al–Si, Al–Si–Mg and Al–Si–Cu alloys were determined by Huang et al.¹³

However, the long-time and large-scale applications of Al-based PCMs are limited due to their highly corrosive to iron-based containment materials at high temperature.¹⁴ According to the phase diagram, Mg–Fe system has properties of thermodynamic stability and immiscible in temperature interval 400–600 °C.¹⁵ Metal Mg has advantages of suitable melting temperature, high melting enthalpy and thermal conductivity. Hence, Mg-based materials can be used as PCMs. Rodríguez-Aseguinolaza et al.^16–18 mainly focus on the studies about magnesium based used as PCMs. They investigated thermophysical properties of Mg–51% Zn and Mg–24.9% Zn–5.1% Al eutectic metal alloys as phase change materials for thermal energy storage during 300–400 °C temperature range. Comparing with Al-based phase change material, the studies about magnesium used as PCMs whose operating temperature reaches above 400 °C are relatively few. In the previous study, we studied the microstructure and thermal characteristics of a series of Mg–Bi alloys and found the Mg–54% Bi has good melting enthalpy and thermal conductivity.¹⁹ Comparing with metal Bi, the melting enthalpy and thermal conductivity of metal Sn are higher. The aim of this paper is to investigate the microstructure and the thermal characteristics of three compositions of Mg–Sn (wt%) alloys as PCM in the 400–600 °C temperature range so that the reference could be provided for the application of metal alloys as a TES system in solar thermal power generation.

2. Experimental procedure

A series of Mg–Sn alloys were prepared by resistance furnace with pure magnesium ingot (99.98 wt%) and tin ingot (99.99 wt%) in the graphite crucible under an argon gas atmosphere. Table 1 shows the nominal composition of Mg–Sn alloys. The RJ-2 flux refining agent (produced by Liaoning Yangchen Metallurgical Material Co, Ltd, China) was used in the melting processing. Table 2 shows the composition of RJ-2 flux and coating agent. The main function of RJ-2 flux refining agent is removing the inclusion of magnesium alloy during melting. The cavity dimension of iron mold was measuring φ 30 mm × 100 mm, which was preheated up to 200 °C. Then, the samples were sectioned from the bottom center of casting, which were etched by an etchant of 4 vol% nitric acid + ethylalcohol.

Table 1 The nominal compositions of Mg–Sn alloys

Samples	Alloy	Chemical composition (wt%)
		Mg	Sn
a	Mg–24% Sn	76	24
b	Mg–37% Sn	63	37
c	Mg–50% Sn	50	50

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Table 2 The compositions of RJ-2 flux

	Composition (wt%)
	MgCl2	KCl	NaCl	CaCl2	CaF2	BaCl2
RJ-2 flux	43–55	20–30	20–30	3–5	10–15	3–5
Coating agent	65–74	10–20	15	3–5	4–7	—

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Microstructure and phase analysis were carried out by X-ray diffraction (XRD, D8 Advance, BRUKER AXS, Germany) and electron probe micro-analysis (EPMA, JXA-8230, JEOL, Japan) with an energy-dispersive X-ray spectrometer (EDS, NCAX-ACT, JEOL, Japan).

Differential scanning calorimeter (DSC, STA449C/3/G, Netzsch, Germany) was used to measure the melting temperatures and enthalpies, as well as the specific heat of the Mg–Sn alloys. Measurements with heating rate (5 K min⁻¹) were operated between room temperature and 600 °C under argon gas controlled atmosphere.

The thermal diffusivity of samples was measured by Nanoflash device (LFA457, Netzsch, Germany) between 40 °C and 400 °C basing on laser-flash method. The dimension of block sample was 10 mm × 10 mm × 2.5 mm. The testing temperature interval was set to 50 °C, and each testing point was measured at least three times to insure that the result was robust. The relative elongation of samples were tested using pushrod type dilatometer (DIL 402C, Netzsch, Germany) from 30 °C to 450 °C.

The density at high temperature was calculated from the relation:²⁰

ρ = ρ₀ (1 + ΔL/L₀)⁻³ (1)

where ρ₀ is the density of the alloy at 20 °C obtained by the Archimedes method, and ΔL/L₀ is the relative elongation.

The thermal conductivity can be calculated from the equation below:

λ = a × ρ × c_p (2)

where a is the thermal diffusivity, ρ is the density and c_p is the specific heat at constant pressure.

3. Results and discussion

3.1. Structural analysis

The X-ray diffraction patterns of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn alloys are depicted in Fig. 1, from which the expected α-Mg and Mg₂Sn can be clearly distinguished, and the diffraction intensities of α-Mg gradually are weakened with increasing Sn content. It is possible that the proportion of α-Mg and Mg₂Sn in alloys has changed as the increase of Sn content.

Fig. 1 XRD patterns of Mg–Sn alloys.

These three samples were characterized by EPMA and representative images are shown in Fig. 2. The compositions of the intermetallic phases obtained from energy spectrum analysis (EDS) are given in Table 3. The region (A, B, D–F) and point (C) are chosen for analysis of EDS. In the low magnification EPMA micrographs of these three alloys, it can be observed that the microstructure of Mg–24% Sn and Mg–50% Sn alloys mainly contains two distinct phases, which are distinguished by the black and grey color. Combined with the high magnification EPMA micrographs and the results of XRD and EDS, the black dendrite shape phase is primary α-Mg phase (A), and the grey punctate shape phase is α-Mg + Mg₂Sn eutectic phase (B) in Mg–24% Sn alloy. The black punctate shape is α-Mg + Mg₂Sn eutectic phase (E), and grey flake shape is primary Mg₂Sn phase (F) in Mg–50% Sn alloy. However, the microstructure of Mg–37% Sn mainly contains one distinct phase, which is distinguished by grey color. The grey phase takes punctate shape. According to the results of XRD and EDS analysis, it is observed that the grey phase is α-Mg + Mg₂Sn eutectic phase (D).

Fig. 2 EPMA micrographs of Mg–Sn alloys.

Table 3 The compositions of intermetallic phases determined by EDS in Fig. 2 (in at%)

Phase	A	B	C	D	E	F
Mg	97.15	85.51	94.65	84.52	87.53	65.52
Sn	2.85	14.49	5.35	15.48	12.47	34.48
Closest phase	α-Mg	α-Mg + Mg2Sn	α-Mg	α-Mg + Mg2Sn	α-Mg + Mg2Sn	Mg2Sn

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The Mg–Sn binary phase diagram is shown in Fig. 3, and the eutectic point of Mg-rich side in Mg–Sn alloy is at 63.1% Mg and 36.9% Sn. In this point, the eutectic transformation L → α-Mg + Mg₂Sn takes place. In the position of Mg–24% Sn (point a), a transformation L → α-Mg takes place firstly in the process of solidification. After that the eutectic transformation L → α-Mg + Mg₂Sn takes place in the remaining liquid metal. Consequently, the primary α-Mg is formed in this alloy. In the position of Mg–37% Sn (point b), the eutectic transformation L → α-Mg + Mg₂Sn takes place. Accordingly, the Mg–37% Sn alloy almost contains α-Mg + Mg₂Sn eutectic phase. While in the position of Mg–50% Sn (point c), the primary Mg₂Sn phase is formed due to the transformation L → Mg₂Sn which takes place before eutectic transformation.

Fig. 3 Binary phase diagram of Mg and Sn.²¹

3.2. Phase transition properties

The measured melting enthalpies and phase change temperatures of samples are depicted in Fig. 4. The data in Table 4 show that the melting enthalpies of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn alloys are 105.3, 217.8 and 118.8 J g⁻¹, with the melting temperature of 557.6, 554.4 and 557.1 °C, respectively. The melting enthalpies of Mg–37% Sn alloy are 106.8% and 83.3% higher than those of Mg–24% Sn and Mg–50% Sn, respectively. The reasons may be as follows: the three compositions of Mg–Sn are prepared in the same condition. These three Mg–Sn alloys all contain α-Mg + Mg₂Sn eutectic phases (Fig. 2), and the main difference is the proportion of this phase in alloy. The proportions of α-Mg + Mg₂Sn eutectic phases in EPMA micrographs are measured, which account for 45–48%, 96–99% and 64–67% in Mg–24% Sn, Mg–37% Sn and Mg–50% Sn, respectively. Combined with the results of melting enthalpies in Table 4, it is interpreted that Mg–37% Sn alloy has higher melting enthalpy due to it has more proportion of α-Mg + Mg₂Sn eutectic phases. Because melting is the inverse process of solidification. According to Mg–Sn binary phase diagram in Fig. 3, the transformation of α-Mg + Mg₂Sn → L takes place when temperature reaches 554–557 °C. Then the transformations of α-Mg → L and Mg₂Sn → L take place when the temperature rise to a certain temperature. However, the melting point of Mg and Mg₂Sn are 651 and 778 °C,²² respectively. They should not melt in the temperature range of DSC melting peak (554–580 °C). Combined with DSC curves in Fig. 4, the melting of α-Mg + Mg₂Sn eutectic occurs mainly in the range of 500–600 °C. It needs to absorb heat during the melting of metal. The heat is the melting latent heat, which is generally represented by melting enthalpy. The more proportion of α-Mg + Mg₂Sn eutectic phases it has, the more heat needs to be absorbed during the melting. Hence, the melting enthalpy is higher. Compared to our previous studies about the phase transition properties of Mg–Bi alloys, the melting enthalpy of Mg–37 Sn% is 20.7% higher than that of Mg–54% Bi alloy.

Fig. 4 DSC curves of Mg–Sn alloys.

Table 4 Phase transition properties of Mg–Sn alloy

Compounds	Melting temperature/(°C)		Melting enthalpy/(J g^−1)
	Onset	Peak	ΔHm
Mg–24% Sn	557.6	564.2	105.3
Mg–37% Sn	554.4	564.1	217.8
Mg–50% Sn	557.1	564.4	118.8

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3.3. Thermal conductivity

The temperature dependence of the relative elongation of test alloys and error bars are shown in Fig. 5. The experimental uncertainty is estimated to be at most 3.6%. Fig. 5 shows that the relative elongation of alloys enhances with the increase of temperature. The density at 20 °C of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn are 2.142 ± 0.025 g cm⁻³, 2.284 ± 0.018 g cm⁻³ and 2.721 ± 0.049 g cm⁻³, respectively. The experimental uncertainty is estimated to be at most 1.8%. The temperature dependence of the densities used for the eqn (1) and error bars are shown in Fig. 6. It can be seen that the density of test alloys decreases with increasing temperature. The density of test alloys increases with the increase of Sn content. The reason is that the density of pure Sn is four times higher than pure Mg, which are 7.265 g cm⁻³ and 1.738 g cm⁻³,²² respectively.

Fig. 5 Temperature dependence of the relative elongation of Mg–Sn alloys. Error bars indicate standard deviation.

Fig. 6 Temperature dependence of density of Mg–Sn alloys. Error bars indicate standard deviation.

Thermal diffusivity and specific heat of Mg–Sn alloys are depicted in Fig. 7, in which the error bars are also shown. The experimental uncertainties are estimated to be at most 2% and 1.9%, respectively. Fig. 7(a) shows that the thermal diffusivity decreases with increasing Sn content. The thermal diffusivity increases with increasing temperature from 40 °C up to 350 °C (300 °C). After that, it shows a light decrease until 400 °C. The reason for this phenomenon may be as follows: the three alloys all contain α-Mg and Mg₂Sn phases. The difference is the content and morphology of the two phases. The thermal diffusivity of Mg decreases when the temperature increases from 20 °C to 500 °C.²² The thermal diffusivity of Mg₂Sn phase has not been reported, but it can be inferred that the thermal diffusivity of Mg₂Sn phase increases with increasing temperature from 40 °C up to 400 °C according to Fig. 7(a). When the temperature increases from 40 °C to 350 °C (300 °C), the thermal diffusivity of Mg₂Sn phase increases dominantly. So the thermal diffusivity of alloys has a slow increase from 40 °C to 350 °C (300 °C). Then when the temperature exceeds 350 °C (300 °C), the thermal diffusivity of α-Mg decreases dominantly. So the thermal diffusivity of alloys shows a light decrease from 350 °C (300 °C) to 400 °C. As shown in Fig. 7(b), the specific heat of Mg–Sn alloys decreases with increasing Sn content. Then, it can be seen that the specific heat has a slow increase from 40 °C to 300 °C and shows a drastic increasing trend from 300 °C and higher temperatures. The reason may be as follows: the specific heat of Mg are 0.996, 1.07 and 1.16 kJ (kg⁻¹ K⁻¹) at 20 127 and 327 °C,²² respectively. It shows an increasing trend from 20 °C to 327 °C. According to Debye model, it is inferred that the specific heat of Mg₂Sn also increases with increasing temperature. The three alloys all contain the α-Mg and Mg₂Sn. So the specific heat of Mg and Mg₂Sn increase when the temperature rises due to a uniform increasing trend from 300 °C to higher temperatures.

Fig. 7 Temperature dependence of (a) thermal diffusivity and (b) specific heat of Mg–Sn alloys. Error bars indicate standard deviation.

According to the eqn (2), the value of thermal conductivity is the product of the thermal diffusivity, specific heat and density of a material, which is shown in Fig. 8. The error bars is also depicted in Fig. 8. The experimental error of the thermal conductivity depends on the experimental error of the thermal diffusivity, specific heat and density. The experimental uncertainty of the thermal conductivity is estimated to be at most 5% by calculating. Among all the three Mg–Sn alloys, the thermal conductivity increases with increasing temperature in Mg–24% Sn and Mg–37% Sn alloys, while thermal conductivity of Mg–50% Sn alloy increases with increasing temperature from 40 °C up to 350 °C and has a decrease up to 400 °C. As seen in Fig. 7(a) and 8, the variation of thermal conductivity and thermal diffusivity is very similar from 40 °C to 300 °C. According to eqn (2), the relationship between thermal conductivity and thermal diffusivity is determined by the product of values of density and specific heat. Before 300 °C, the density of test alloys decreases with increasing temperature and the of test alloys increases with increasing temperature, so the product of the value of these two factors did not change much with the variation of temperature. It means that the values of thermal conductivity are mainly influenced by the thermal diffusivity rather than the product of density and specific heat. Hence, the variation of thermal conductivity is similar to that of the thermal diffusivity before 300 °C. Then, when the temperature rises up to 300 °C, the density continues to decrease and the thermal diffusivity has a slight change. However, the specific heat increases dramatically, which play a leading role in calculation of thermal conductivity. Hence, the thermal conductivity of the Mg–Sn alloys increases as the increase of temperature.

Fig. 8 Temperature dependence of thermal conductivity of Mg–Sn alloys. Error bars indicate standard deviation.

Back to Fig. 8, the thermal conductivity of alloys decreases as the increase of Sn content. Compared to Mg–24% Sn and Mg–37% Sn alloys, the value of thermal conductivity of Mg–50% Sn alloy roughly drops by 25 W mK⁻¹ and 20 W mK⁻¹ at 400 °C, respectively. According to kinetic molecular theory, the thermal conductivity of materials increases with the increase of the mean free path of electrons and phonons. In the case of Mg–Sn alloys, Sn in solid solution may act as scattering centers for electrons and phonons, the scattering centers limit the mean free path of electrons and phonons,^23,24 which leads to the decrease of the thermal conductivity.

4. Conclusions

The microstructure and thermal characteristics of Mg–Sn alloys as latent heat energy storage were measured and analyzed. The conclusions are summarized as follows:

(1) The microstructure of Mg–24% Sn alloy mainly consists of α-Mg matrix and α-Mg + Mg₂Sn eutectic phases, the microstructure of Mg–37% Sn alloy α-Mg + Mg₂Sn eutectic phases, and the microstructure of Mg–37% Sn alloy the primary Mg₂Sn phase and α-Mg + Mg₂Sn eutectic phases.

(2) The melting enthalpies of Mg–24% Sn, Mg–37% Sn and Mg–50% Sn alloys are 105.3, 217.8 and 118.8 J g⁻¹, respectively, with the phase change temperatures in the 554–558 °C range. The Mg–37% Sn alloy has a higher value of melting enthalpies. It is possible that there is higher proportion of α-Mg + Mg₂Sn eutectic phases in alloy.

(3) The values of the specific heat and thermal conductivity decrease with increasing Sn content. Compared to Mg–24% Sn and Mg–37% Sn alloys, the value of thermal conductivity of Mg–50% Sn alloy roughly drops by 25 W mK⁻¹ and 20 W mK⁻¹ at 400 °C.

(4) Comparing with the thermal characteristics of Mg–24% Sn and Mg–50% Sn alloys, the Mg–37% Sn alloy has considerable potential to be used as the energy storage material for CSP applications. In addition, further work about the compatibility of Mg-based alloys with containment materials at high temperature will be conducted.

Acknowledgements

This work was supported by the National Key Technology Research & Development Program of China (Grant No. 2012BAA05B05) and the National Natural Science Foundation of China (Grant No. 51206125).

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