boron [email protected] composites for thermal energy storage

Journal Pre-proofs High thermal conductive shape-stabilized phase change materials of polyethylene glycol/boron nitride@chitosan composites for thermal energy storage Xiwen Jia, Qingye Li, Chenghong Ao, Rui Hu, Tian Xia, Zhouhang Xue, Qunhao Wang, Xueyong Deng, Wei Zhang, Canhui Lu PII: DOI: Reference:

S1359-835X(19)30459-2 https://doi.org/10.1016/j.compositesa.2019.105710 JCOMA 105710

To appear in:

Composites: Part A

Received Date: Revised Date: Accepted Date:

3 August 2019 24 October 2019 20 November 2019

Please cite this article as: Jia, X., Li, Q., Ao, C., Hu, R., Xia, T., Xue, Z., Wang, Q., Deng, X., Zhang, W., Lu, C., High thermal conductive shape-stabilized phase change materials of polyethylene glycol/boron nitride@chitosan composites for thermal energy storage, Composites: Part A (2019), doi: https://doi.org/10.1016/j.compositesa. 2019.105710

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High thermal conductive shape-stabilized phase change materials of polyethylene glycol/boron nitride@chitosan composites for thermal energy storage Xiwen Jia,a Qingye Li,a Chenghong Ao,a Rui Hu,a Tian Xia,a Zhouhang Xue,a Qunhao Wang,a Xueyong Deng,a Wei Zhang,a,b,* Canhui Lu,a,b,*

a. State Key Laboratory of Polymer Materials Engineering, Polymer Research Institute at Sichuan University, Chengdu 610065, China

b. Advanced Polymer Materials Research Center of Sichuan University, Shishi 362700, China

* Authors for correspondence: E-mail: [email protected] (W. Zhang), [email protected] (C. Lu); Phone: 86-28-85460607; Fax: 86-28-85402465

Abstract Phase change materials (PCMs) applied in the energy storage and temperature control system are crucial for energy conservation and environmental protection. In this work, boron nitride (BN)@chitosan (CS) scaffolds with three-dimensional (3D) porous structures were fabricated. And effective thermal conductive pathways could be created in the resultant scaffolds. By introducing polyethylene glycol (PEG) into the BN@CS scaffolds, composite PCMs with large latent heat of fusion and excellent shape-stability were obtained. In particular, a high thermal conductivity up to 2.77 W m−1 K−1 could 1

be reached at a relatively low content of BN (27 wt%). Moreover, they also exhibited a satisfactory energy storage density of 136 J g−1. This work demonstrated a facile and environmentally friendly strategy to simultaneously achieve enhancement of thermal conductivity, high energy storage density, shape stability and outstanding thermal repeatability for composite PCMs, which held promising potential in waste heat recovery, cooling system and temperature control system.

Keywords A. Energy materials; Thermal conductivity; B. Thermal properties; A. Biocomposite.

1. Introduction Thermal energy storage (TES) of latent heat, sensible heat and reversible thermochemical reaction has proved to be a promising and low-cost technique in terms of energy conservation and environmental protection [1-3]. Latent heat storage, which utilizes the phase change materials (PCMs) to store or release latent heat [4], has a wide range of applications, such as industrial waste heat recovery, comfort applications in buildings, cooling system, and temperature control system [5-9]. Polyethylene glycol (PEG) is one of the most widely studied PCMs [10]. And it possesses high energy storage density, suitable phase change temperature, stable physical and chemical properties, low toxicity and competitive price [11, 12]. However, several insurmountable defects restrict its final applications. For instance, leakage in the molten state, low thermal conductivity and high degree of supercooling. Such drawbacks lead 2

to hysteresis of thermal response and degrade the performance in energy storage and thermal regulation [13]. The above problems can be possibly solved by employing the packaging technology and adding fillers of high thermal conductivity [12]. Shape-stabilized PCMs (ss-PCMs) are generally synthesized by blending PCMs with supporting materials. They can maintain a solid shape even at a temperature higher than the melting temperature of PCMs [14, 15]. Metal foams (copper foam [16], nickel foam [17]), porous carbons (carbon nanotubes [18] and graphene aerogel [19]), porous oxides (expanded perlite [20], diatomite [21]) and polymers (poly(vinyl alcohol) [22], polystyrene [23]) have been extensively explored as the supporting materials to fabricate ss-PCMs. However, most of those PCMs are fabricated via complex procedures with corrosive or toxic chemicals, leading to high cost and environmental concerns. Biopolymers have also been used as the supporting materials, but most of them are cellulose. And the actual enthalpies of the composite PCMs combined with cellulose are usually lower than their theoretical values [24, 25]. In addition, crosslinking is always required to enhance their stability and strength. Chitosan (CS) is a renewable bio-based polysaccharide abundant in nature with distinct advantages, such as low cost, good bio-compatibility, non-toxicity, environmentally friendliness [26]. CS can be easily and environment-friendly converted into a three-dimensional (3D) hierarchical porous scaffold without crosslinking [27, 28]. It is suitable for multiple applications, including drug delivery [29], supercapacitors [28], wastewater treatment [30], thermal insulators [31], organic catalysts [32]. Indeed, CS contains numerous 3

amino and hydroxyl groups, which can form hydrogen bonding interactions with PEG and effectively prevent the leakage of PEG melt. Nevertheless, it only has negligible effect on the enthalpy of PEG [24]. By combining thermal conductive fillers with supporting materials, the leakage and low thermal conductivity problems of PEG-based PCMs can be simultaneously solved [11]. The conductive fillers stacked alone with the supporting materials can offer a thermal conductive path for the composite PCMs. In contrast to composites with randomly dispersed fillers, this method realizes high thermal conductivity with less loading of fillers [33]. Otherwise, the overloaded fillers will lead to the deterioration of mechanical properties and the reduction of enthalpy of the composites PCMs. Carbon and metal materials have been extensively studied as thermal conductive fillers [34]. However, the electrical conductivity may seriously limit their applications in some special fields where electrical insulation is compulsive. Boron nitride (BN) has been widely applied in high thermal conductive composites [35, 36]. It has a crystal structure similar to graphene, and possesses high thermal conductivity coefficient, thermal and chemical stability and electrical insulation resistance. In this work, we constructed 3D scaffolds via the ice-templated self-assembly strategy, in which CS served as the supporting material while BN functioned as the high thermal conductive filler. Next, PEG was impregnated into the BN@CS scaffolds to produce composite PCMs with a huge energy storage density. The BN@CS scaffold was particularly important for a high performance PCM as it provided a thermal 4

conductive pathway and an appropriate encapsulation to prevent the leakage of melted PEG owing to the intermolecular interactions between PEG and CS [24]. By changing the freezing temperature during casting, the porous microstructures could be well regulated, which was proved to have direct impact on the thermal conductivity, enthalpy and supercooling temperature of the composite PCMs. Compared to pure PEG, the composite PCMs (27 wt% BN) exhibited a thermal conductivity up to 2.77 W m−1 K−1 corresponding to 804% improvement, and 73% latent heat could be retained. Interestingly, the composite PCMs still kept their initial shape above the melting temperature of PEG.

2. Experimental section 2.1. Materials Boron nitride powder (BN, ~10 µm, 99.9%) was purchased from Shanghai St-nano Science & Technology CO., LTD. Chitosan (CS), acetic acid and PEG (Mn = 6000) were obtained from Chengdu Cologne Chemical CO., LTD. 2.2. Preparation of BN@CS slurries 20 mg mL−1 CS solution was prepared by dissolving CS in 2 wt% an acetic acid solution under magnetic stirring for 12 h. A certain amount of BN was mixed with the CS solution to form a uniform aqueous suspension with different BN concentrations (0, 5, 10, 20 and 30 wt%). Then the suspension was sonicated for 10 min in a noise isolating chamber (JY 99-IIDN, Ningbo Scientz Biotechnology Co., Ltd., China) at 30% power. 5

2.3. Preparation of BN@CS scaffolds The hybrid slurries were poured into square moulds and frozen in a refrigerator (−20 oC)

and liquid nitrogen (−196 oC), respectively. Then BN@CS scaffolds were prepared

by freeze-drying (Freeze-drying chamber, FD-1A-50, Biocool, China) at low temperature (−50 oC) and pressure (1 Pa) for 48 h. 2.4. Preparation of PEG based composite PCMs The composite PCMs were fabricated using vacuum-assisted impregnation method. PEG was put into a vacuum oven at 110 oC. When PEG was fully melted, 1 h degassing was applied to remove the air bubbles. After that, the as-prepared BN@CS scaffolds were immersed into the PEG melt, followed by repeated degassing for 3 h to make sure the complete infiltration of the PEG melt. The obtained PCMs were designated as C0, C5, C10, C20, C30 and N0, N5, N10, N20, N30 according to the preset BN content in the BN@CS slurries and the freezing temperature. Letters C and N represented the freezing temperature of −20 oC and −196 oC, respectively. For comparison, the melt blending composite samples of PEG and BN with accordant BN contents were also prepared, which were marked as S0 (pure PEG), S5, S10, S20, S30. The whole preparation process for the composite PCMs was schematically shown in Fig. 1. 2.5. Characterization The morphologies of the BN@CS scaffolds and the fractured surfaces of PCMs were observed by a field-emission scanning electron microscope (SEM, JEOL JSM-7500F, FESEM, Japan). The thermal conductivity was measured using a Hot Disk Thermal 6

Constant Analyzer (Hot Disk 2500-OT, Sweden) by a transient plane heat source method. The radius of the measuring probe was 3.189 mm. Before measurement, all the square samples were polished with fine sandpaper. To record the thermal response of the composite PCMs, the samples were placed on a hot plate, and a series of infrared thermograph was taken by an infrared thermal imager (TL50sc, FLIR, USA). Thermogravimetric analysis (TGA, TG209 F1, NETZSCH, Germany) was carried out at a heating rate of 10 oC min−1 from 30 to 700 oC in air atmosphere. Phase change temperatures and phase change enthalpies of pure PEG and the composite PCMs were measured by differential scanning calorimetry (DSC, DSC 204 F1, NETZSCH, Germany), which was conducted at the heating and cooling rate of 10 °C min−1 in the range of 0–120 °C under nitrogen atmosphere. X-ray diffraction (XRD, EMPYREAN, Holland) patterns were recorded at a scanning speed of 10o min−1 over the diffraction angles of 5–80° (2θ). The Fourier transform infrared (FTIR) spectra were obtained on the instrument (Nicolet 6700, Thermo Fisher, USA) in the transmission mode, using KBr pellets for all the samples. The leakage test was investigated by observation of the samples placed on a hot stage with increasing temperature. Cycling stability of C30 was explored using DSC (20–100 oC, 10 oC min−1, N2 atmosphere) and FTIR.

3. Results and discussion 3.1. Microstructures of the BN@CS scaffolds and the PCMs The morphologies of the BN@CS scaffolds for C0, C10, C30 and N30 were 7

displayed in Fig. 2. The neat CS scaffolds for C0 exhibited a folded sheet-like structure with a smooth surface, which overlapped with each other and developed into a 3D network (Fig. 2a). The scaffolds of C series had typical porous structures with pore sizes of several hundreds of micrometers. By increasing the content of BN, the wall of the scaffold became thicker (Fig. 2b–c). This should be ascribed to the fact that during the freezing process of BN@CS slurries in molds, the higher the BN concentrations, the more the chances for them to stack with neighboring ones, leading to the thicker walls [37-39]. As shown in Fig. 2d–f, when the freezing temperature decreased to −196 oC,

the micropore became denser and the pore size became smaller. Additionally, the

wall of N30 scaffolds was much thinner than that of C30 scaffolds. These structural differences arose from the ice growth process [40]. The ice solidification comprises two individual but consecutive phases: nucleation and growth of ice crystals, and a competitive balance exists between the two phases [41]. A lower freezing temperature yielded more nuclei of ice crystals and a relatively faster crystallization rate, resulting in the crystals with a small size. In the meantime, the ice-growth with volume expansion would drive CS and BN to assemble into ordered structures, leading to thinner walls and small pore sizes of the resultant porous scaffolds. Note that the BN@CS scaffold played an important role in supplying a supporting structure and generating capillary force for PEG infiltration. The SEM images of the fractured surfaces of pure PEG, C30 and N30 showed that the BN@CS scaffolds well maintained their structures throughout the PEG matrix (Fig. 8

S1). And no obvious interfacial debonding was observed. Such a compact interface was favorable for the reduction of interfacial thermal resistance between BN and PEG matrix. 3.2. Thermal conductivity of the composite PCMs Thermal conductivity (TC) is very important for PCMs as it governs the sensitivity of thermal response and the work efficiency. Fig. 3a illustrated the thermal conductivity of pure PEG and composite PCMs. The pure PEG exhibited an inherently low thermal conductivity of 0.31 W m−1 K−1. The incorporation of BN caused a prominent increase of the TC. Moreover, the C and N series exhibited higher TC than the S series at a similar content of BN. The representative SEM images for different scaffolds at higher magnification (Fig. 4) revealed that the BN plates had a compact stacking of their basal plane parallel to the scaffolds. The as-formed 3D network could enable a more effective thermal conductive pathway inside the PEG matrix. In addition, TC enhancement (η) was used to characterize the TC enhancement efficiency of BN to the composites, which is defined in Eqn. (1) [33], 𝜂=

𝜆𝑐𝑜𝑚 ― 𝜆𝑝𝑒𝑔 𝜆𝑝𝑒𝑔

(1)

× 100%

where λcom and λpeg represent the thermal conductivity of the composite PCMs and pure PEG, respectively. Fig. 3b showed η as a function of the BN content for all the composite PCMs, and the η increased with the increase of BN content. As depicted in Fig. 4 and Fig. S2, with the increase of BN loading, the thickness of the scaffold walls increased, leading to the gradual enhancement of TC for the composite PCMs [42]. The 9

samples of C series exhibited higher thermal conductivity as compared with the N series samples at a similar content of BN. Although the denser pores in the scaffolds of N series were favorable for building more thermal conductive pathways, the simultaneously increased surface areas in N series would lead to considerable heat loss [43]. Moreover, the BN sheets dispersed in the scaffolds of N series having a larger surface area could not form confluent thermal conductive pathways. As indicated in Fig. S2c, the thin wall of N30 covered by BN appeared to have many defects, giving rise to remarkable phonon scattering thereof [44]. Notably, the thermal conductivity for C30 (27.01 wt% BN) and N30 (28.09 wt% BN) reached 2.78 and 2.45 W m−1 K−1, which were 804% and 698% higher than that of pure PEG, respectively. The higher η for the C series suggested that the freezing temperature played a significant role in the formation of the thermal conductive network. When the freezing temperature increased, the obtained thicker scaffold walls and larger pore size reduced the thermal conductive interface, and made the heat transfer more effective. 3.3. Thermal Management Capability of the composite PCMs To demonstrate the thermal management capability, the surface temperature variations of the composite PCMs with time during heating and cooling processes were recorded by an infrared thermal camera. The samples of C0, C10, C30, N30 were placed on the same hot stage. Fig. 5 showed the temperature distribution images. In the heating process, the color of C30 changed faster than N30, and C30 also exhibited the fastest change among all samples, which was in accordance with the results of thermal 10

conductivity. As expected, the same phenomenon was also verified in the cooling process. Thus, the high thermal conductivity could endow the materials with better thermal response. 3.4. Thermophysical properties of the composite PCMs The thermal stability and the loading percentage of BN in the composite PCMs were evaluated by thermal gravimetric analysis (TGA). There was almost no mass loss below 200 oC, revealing that the composite PCMs had excellent thermal stability within their working temperature range. Pure PEG began to decompose at 225 oC, and a lower percentage of BN in PCMs could improve the thermal stability. However, an excess of BN accelerated the decomposition of PEG possibly due to the enhanced thermal conductivity [45]. Pure PEG and CS were completely burnt out in the air atmosphere when heated to 700 oC. Whereas there was a negligible loss of mass for BN due to its remarkable stability [39, 41]. Hence, the weight percentage of the final residue should account for the content of BN in the composite PCMs. From Fig. 6a, the BN contents in C5, C10, C20 and C30 could be estimated to be 4.1 wt%, 8.15 wt%, 18.99 wt% and 27.01 wt%, respectively. The results agreed well with the compositions as we designed. Phase change enthalpy is always considered as the most reliable indicator for evaluating the thermal energy storage capacity of PCMs. Because the supporting scaffold did not undergo a phase transition process within the work temperature range, only PEG should account for the latent thermal heat [25, 46]. And the latent heat would decrease when the mass fraction of PEG was reduced. Differential scanning calorimetry 11

(DSC) was performed to explore the thermophysical properties of pure PEG and the composite PCMs (Fig. 6b and Fig. S3). Their melting/crystallization temperature (Tm/Tc), melting/crystallization enthalpy (ΔHm/ΔHc), theoretical enthalpy (ΔHt) and supercooling temperature (ΔT) were summarized in Tab. 1. ΔHt is determined by Eqn. (2)[47], (2)

∆𝐻𝑡 = 𝑥∆𝐻𝑃𝐸𝐺

where ΔHt is the theoretical enthalpy of the composite PCMs; ΔHPEG is the enthalpy of pure PEG; x is the weight fraction of PEG in the composite PCMs. The theoretical and actual enthalpies of the composite PCMs were shown in Tab. 1 and Tab. S1. The actual enthalpies for the melting process of the C and S series approached to their theoretical values, but the values for N series were lower than the theoretical ones (Fig. 6c). The crystallizing process demonstrated similar behaviors. Another critical parameter to compare the thermal properties is the heat storage efficiency (γ), which was calculated through the loss percentage. γ is determined by Eqn. (3) [48],

(

∆𝐻𝑐

)

(3)

𝛾 = 1 ― ∆𝐻𝑚 × 100%

where ΔHm is the melting enthalpy, ΔHc is the crystallization enthalpy. As shown in Fig. 7a, heat loss of N series was higher than C and S series. The difference between C, N and S series might be interpreted as follows: the pore size of N series was extremely small, which restricted the crystal arrangement and orientation of PEG molecular chains, leading to a decline of the regularities of crystal regions and an increase in lattice defects [11, 49]. This inference was well supported by the crystallinity from the XRD results 12

(Fig. 6c). As is well recognized, it is rather difficult to achieve high thermal conductivity and large energy storage density of PCMs simultaneously. In order to reveal the overall performance of the as-prepared PCMs, the thermal conductivity and the relative melting enthalpy (R) for C series and some previously reported organic PCMs were compared in Fig. 6d [25, 41, 50-60]. Relative melting enthalpy (R) is defined as the ratio of melting enthalpies between the composites and pure PEG [61]. Considering the fact that only PEG should account for the melting enthalpy, the larger the R value, the higher the PEG content and the lower the filler content in the composite PCMs. The results indicated that the composite PCMs in C series exhibited superior thermal conductivity over many early reported ones at a similar loading of thermal conductive fillers, while could still maintain a high energy storage density. The supercooling temperature (ΔT) [62] is determined as the difference between Tm and Tc. The corresponding evaluation results for the composite PCMs were summarized in Tab. S1. The introduction of BN improved the thermal conductivity of the composite PCMs and consequently accelerated the phase change speed of PEG [63]. In addition, BN played a role of heterogeneous nucleating seeds, which promoted the crystallization rate of PEG [63-65], and also affected the phase change temperature. Thus, the introduction of BN suppressed the supercooling to some extent. The crystallization temperature of C and N series was obviously higher than that of S series. It might be caused by the hydrogen bonding interactions between CS and PEG, which gave rise to 13

a higher crystallization temperature [46]. The higher melting temperature of N series should be ascribed to the small pore size that restricted the movement of PEG molecular chain [41, 44]. Fig. 7b showed that the ΔT of the C series were 13.3%, 13.7%, 14.4%, 17.2% and 19.6% lower than that of pure PEG. Whereas the ΔT for the N30 and S30 decreased only 11.2% and 9.8%, respectively. These results suggested that the supercooling extent of PEG could be favorably reduced by introducing CS and BN, especially for the BN@CS scaffolds frozen at a higher temperature. Moreover, the XRD patterns in Fig. 7c and Fig. S4 were used to figure out the crystallization behaviors of pure PEG and the composite PCMs. Pure PEG exhibited strong reflections at 19.4o and 23.5o and weak reflections at 26.3o, 31.1o, 36.3o and 39.8o [25, 66]. Compared with pure PEG, the peak positions of the composite PCMs did not change. And the two main diffraction peaks at 19.4o and 23.5o corresponding to the (120) and (032) planes for PEG crystal became wider and flatter with the incorporation of BN@CS scaffold in the composite PCMs. It indicated that CS and BN did not affect the crystal structure of PEG except for a reduced spherulite size. This was in accordance with the results of DSC and FTIR that no chemical reactions took place between PEG and CS or BN. The obvious peak at 26.7

o

was attributed to the BN. Its intensity

increased with the increase of the BN content, overlapping the week peak of PEG at 26.3 o. Estimation of crystallinity by comparing the intensity count values suggested almost no reduction in crystallinity of C and S series, but a little reduction in N series as compared with pure PEG. 14

In the FTIR spectra (Fig 7d), the pure chitosan had distinct amide band and hydroxy band at 1560 and 1380 cm−1, respectively [67]. The amino group had a characteristic absorption band in the region of 3400–3500 cm−1, which was overlapped by the broad absorption band of the OH group at 3446 cm−1. The spectrum of pure PEG was characterized by the stretching vibration of O–H at 3450 cm−1, the C–H stretching vibration at 2886 cm−1, the C=O stretching vibration at 1633 cm−1 and the C–O stretching vibration at 1105 cm−1 [10]. In the spectra of composite PCMs, the peaks for O–H and C–O groups shifted to lower wavenumbers as compared with those for pure PEG and CS, indicating that there existed intermolecular hydrogen bonding interactions between PEG and CS [14]. 3.5. Thermal reliability of the composite PCMs The thermal stability of C30 was examined against 50 heating and cooling cycles. As shown in Fig. 8a, the phase change temperatures and enthalpies had little change, indicating the excellent thermal stability of the composite PCMs. Also, no change of the shape or the position for the characteristic peaks was observed in the FTIR spectra (Fig. 8b), which disclosed that no chemical reactions took place and the chemical structure was stable. These results consistently implied that the composite PCMs possessed excellent thermal reliability and stability for a long service life. 3.6. Shape-stability of the composite PCMs Shape stability is another crucial challenge for the practical application of organic PCMs. The leakage test for pure PEG, C0, C30 and N30 were conducted on a hot stage. 15

Fig. 9 showed that when heated to 70 °C (above the melting temperature of PEG), pure PEG gradually melted with obvious liquid leakage. In contrast, no obvious leakage was observed for the composite PCMs. When heated to 90 °C, PEG completely melted into liquid, while the composite PCMs remained intact with negligible leakage. Furthermore, unlike the sample S30, the composite PCMs with BN@CS scaffolds could well retain their shapes even after compressed by a weight of 500 g at 90 °C and only slight liquid leakage could be observed (Fig. S5). This should be mainly ascribed to the strong capillary force and intermolecular hydrogen bonding interactions between PEG and CS in the composite PCMs. The PEG molecules were tied to the surface of CS by the confinement effect of strong intermolecular hydrogen bonding and lost their freedom of motion. It is noteworthy that the leakage was slightly visible for the composite PCMs produced at a higher freezing temperature, which was consistent with the previous results. 3.7 Potential applications of the composite PCMs in thermal management To demonstrate the potential applications of the as-prepared composite PCMs in thermal management, a simple temperature control system was designed. As schemed in Fig. 10a, a hot stage was utilized to simulate the heat source at 90 oC. Then the hot stage was removed to simulate the cooling process. The dimeticone’s temperature gradually increased at different rates during the heating process owing to the different thermal conductivity of the PCMs (Fig. 10b). For the cooling process, the curve of the dimeticone with C30 presented a distinct platform, suggesting that C30 was capable to 16

maintain the temperature of dimeticone at about 37.5 oC for a longer time due to the release of latent heat. In addition, the temperature response of dimeticone with C30 was more sensitive than that with C0, indicating the higher efficiency of C30 in energy storage and thermal regulation in the presence of BN [53, 68].

4. Conclusions Novel PEG based PCMs with high thermal conductivity, large energy storage density, remarkable thermal reliability and excellent shape-stabilization were successfully prepared by incorporation of BN@CS scaffolds which were obtained via the icetemplated self-assembly strategy. The composite PCMs, especially with BN@CS scaffolds obtained at a higher freezing temperature, possessed more effective 3D thermal conductive pathways. And their actual thermal energy storage density approached to the theoretical value. With 27.01 wt% of BN, the composite PCMs exhibited a high thermal conductivity of 2.78 W m−1 K−1, 804% higher than that of pure PEG. In addition, they also demonstrated a satisfactory energy storage density of 136 J g−1. Importantly, the composite PCMs could maintain thermal reliability and reusability even after 50 melting and cooling cycles. The incorporation of BN@CS scaffolds could also prevent the leakage of PEG owing to the combined actions of hydrogen bonding and capillary force. Possessing such outstanding properties, the present composite PCMs were envisaged to have promising applications in waste heat recovery, comfort applications in buildings, cooling system, temperature control system, and so on. 17

Acknowledgements This work was supported by the National Natural Science Foundation of China (51433006 and 51861165203), Sichuan Science and Technology Program (2019YJ0125), State Key Laboratory of Polymer Materials Engineering (sklpme20192-19), and the Fundamental Research Funds for the Central Universities.

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28

Figure captions

Fig. 1. Schematic route for the preparation of PEG/BN@CS PCMs. Fig. 2. SEM images of the BN@CS scaffolds for (a) C0, (b) C10, (c) C30 and (d–f) N30. Fig. 3. (a) TC and (b) TC enhancement of pure PEG and the composite PCMs. Fig. 4. Schematic illustration for the thermal conductive mechanism of the composite PCMs of (a) C10, (b) C30 and (c) N30. Fig. 5. (a) Infrared thermal images and (b) surface temperature variation with heating time for C0, C10, C30 and N30; (c) surface temperature variation with cooling time and (d) infrared thermal images for C0, C10, C30 and N30. Fig. 6. (a) TGA curves and (b) DSC curves of pure PEG, C0, C10, C30, N30, S30; (c) theoretical and actual melting enthalpies and crystallinity for composite PCMs; (d) comparison of the thermal conductivity for C series and the previously reported organic PCMs at different relative melting enthalpy (R). Fig. 7. (a) Heat loss of the composite PCMs; (b) Phase change temperatures and the variation of supercooling temperatures; (c) XRD patterns of BN, pure PEG and composite PCMs (the inset shows enlarged XRD patterns of PEG, C0, C10 and C30); (d) FTIR spectra of CS, BN, pure PEG and composite PCMs.

29

Fig. 8. (a) The measured latent heat of C30 during 50 melting-freezing cycles (the inset shows DSC curves of C30 tested for 50 cycles) and (b) FTIR spectra for C30 before and after 50 times of thermal cycling. Fig. 9. Photographs for PEG, C0, C30 and N30 during a continuous temperatureincreasing procedure. Fig. 10. (a) Schematic illustration of the temperature control system; (b) the temperature changes of dimeticone during the heating and cooling processes.

30

Table

Tab. 1. Thermophysical properties of pure PEG and composite PCMs. w (wt%)

Tm (oC)

ΔHm (J g-1)

ΔHtm (J g-1)

Tc (oC)

ΔHc (J g-1)

ΔHtc (J g-1)

PEG

0

66.1

189

–

37.6

183.6

–

C0

0

66.5

186.2

187.2

41.8

180.8

181.8

C5

4.1

65.4

180.5

181.3

40.8

175.2

175.2

C10

8.15

65.7

172.6

173.6

41.3

167.6

167.6

C20

18.99

65

151.7

153.1

41.4

147.3

147.3

C30

27.01

64.6

136.9

137.9

41.7

132.9

132.9

N30

28.09

66.2

128.9

135.9

40.9

124.6

124.6

S30

29.01

65.1

133.1

134.2

39.4

129.3

129.3

31

Figure

Fig. 1. Schematic route for the preparation of PEG/BN@CS PCMs.

32

Fig. 2. SEM images of the BN@CS scaffolds for (a) C0, (b) C10, (c) C30 and (d–f) N30.

33

Fig. 3. (a) TC and (b) TC enhancement of pure PEG and the composite PCMs.

34

Fig. 4. Schematic illustration for the thermal conductive mechanism of the composite PCMs of (a) C10, (b) C30 and (c) N30.

35

Fig. 5. (a) Infrared thermal images and (b) surface temperature variation with heating time for C0, C10, C30 and N30; (c) surface temperature variation with cooling time and (d) infrared thermal images for C0, C10, C30 and N30.

36

Fig. 6. (a) TGA curves and (b) DSC curves of pure PEG, C0, C10, C30, N30, S30; (c) theoretical and actual melting enthalpies and crystallinity for composite PCMs; (d) comparison of the thermal conductivity for C series and the previously reported organic PCMs at different relative melting enthalpy (R).

37

Fig. 7. (a) Heat loss of the composite PCMs; (b) Phase change temperatures and the variation of supercooling temperatures; (c) XRD patterns of BN, pure PEG and composite PCMs (the inset shows enlarged XRD patterns of PEG, C0, C10 and C30); (d) FTIR spectra of CS, BN, pure PEG and composite PCMs.

38

Fig. 8. (a) The measured latent heat of C30 during 50 melting-freezing cycles (the inset shows DSC curves of C30 tested for 50 cycles) and (b) FTIR spectra for C30 before and after 50 times of thermal cycling.

39

Fig. 9. Photographs for PEG, C0, C30 and N30 during a continuous temperatureincreasing procedure.

40

Fig. 10. (a) Schematic illustration of the temperature control system; (b) the temperature changes of dimeticone during the heating and cooling processes.

41