Composite phase change materials embedded into cellulose/polyacrylamide/graphene nanosheets/silver nanowire hybrid aerogels simultaneously with effective thermal management and anisotropic electromagnetic interference shielding

Exploiting an advanced material simultaneously with effective thermal management (TM) and electromagnetic interference (EMI) shielding capacity is ungently demanded yet challenging for the miniaturized and integrated electronics. Anisotropic networks can be impregnated with phase change materials (PCMs) to fabricate multifunctional shape-stable PCMs (ss-CPCMs) simultaneously with excellent TM and EMI shielding, which is rarely reported. Herein, the anisotropic cellulose/polyacrylamide/graphene nanosheet/silver nanowire (CPGxAy) hybrid aerogels were successfully prepared using directional freeze-drying method, and then utilized as supporting skeletons to embed polyethylene glycol (PEG) via vacuum-assistant impregnation. Profited by the synergistic effect of graphene nanosheets (GNPs) and silver nanowires (AgNWs), the resultant polyethylene glycol@cellulose/polyacrylamide/graphene nanosheet/silver nanowire hybrid aerogel (PEG@CPGxAy) ss-CPCMs exhibit fascinating thermal conductivity (TC) of 0.84 W/m·K (200% increase in comparison with that of pure PEG) and anisotropic average EMI shielding effectiveness (SE) of 71.08 dB along the transverse direction and 35.21 dB along the longitudinal direction, while remaining high melting and crystallization enthalpy efficiency of 93.47% and 93.08%, respectively. In addition, PEG@CPGxAy ss-CPCMs also display great shape stability, thermal stability, and cyclic reusability in the storing/releasing latent heat processes. This investigation sheds new light on designing and fabricating ss-CPCMs with pretty comprehensive properties for TM and EMI shielding of modern electronics.


Introduction
Over the past decades, the ever-increasing human needs promote the rapid iteration of science and technology, which forces the high-frequency and high-power modern electronics to tend toward miniaturization and integration. Thus, the heat accumulation would lead to possible electronic invalids or even damage of the advanced electronics [1]. Traditional TM technologies based on air and liquid forcing convection cooling are not applicable for modern electronics, which is attributed to the requirement of additional power consumption or equipment. Although the highly thermally conductive composites have been widely used to dissipate the heat generated by electronics, only keeping electronics at a relatively constant temperature can ensure their normal operation and service life [2]. Hence, it is essential to develop an TM material with high thermal conductivity (TC) and temperature-control capacity. Specially, phase change materials (PCMs) could control the operating temperature of electronics within a safe range under the condition of no extra power consumption or equipment, because they can store/release a mass of thermal energy at fairly stationary temperature during the phase change period [3]. Currently, many organic PCMs (e.g., polyalcohol [4], paraffin wax [5], and fatty acid [6]) have already been deemed as the promising substitutes for TM materials, arising from the large latent heat enthalpy, admirable chemical reliability, environmental friendliness, recyclability, nontoxicity, suitable phase-change temperature, and so on. Unfortunately, most of organic PCMs suffer from the inherent drawbacks of inferior TC and easy leakage during the phase transition, which seriously impedes their practical applications in TM systems [7]. Therefore, it is essential to explore shape-stable composite PCMs (ss-CPCMs) with high TC for TM of electronics.
Although various strategies, mainly including microcapsule encapsulating [8], polymer shaping [9], coaxial electrospinning [10], chemical crosslinking [11], and nanomaterials adsorbing [12], are put forward to fabricate ss-CPCMs, some exiting problems such as technical difficulty, large sacrifice of enthalpy, and difficulty in promoting TC are detrimental to large-scale production or material performance. Especially, impregnating organic PCMs into three-dimensional (3D) porous networks hybridized with nano-filler (for instance, graphene [13], MXene [14], carbon nanotubes [15], silver nanowires (AgNWs) [16], etc.) can overcome the aforementioned deficit, due to its simple process, easy large-scale production, high adsorption efficiency, and effective highway for phonon conduction. For instance, Wu et al. [17] fabricated multi-responsive CG@MF/PEG ss-CPCMs through vacuum adsorbing PEG into melamine foam-templated cellulose nanofiber/graphene nanoplatelet (CG@MF) 3D skeletons. The CG@MF showed a high loading fraction (up to 96%) for PEG and endowed CG@ MF/PEG with excellent thermal and electrical conductivity, which imparted CG@MF/PEG ss-CPCMs with considerable light/electro-to-thermal conversion capacity. Noteworthily, freeze-drying strategy has been increasingly applied to prepare porous aerogels for encapsulating organic PCMs in recent years, due to its green processes, simple steps, and inexpensive cost [18]. In addition, the generated aerogels with low density, high porosity, and large specific surface area can effectively adsorb PCMs and ensure high melting/ crystallization enthalpy efficiency. Normally, the freezedrying strategy can impart the obtained aerogels with controllable pore structures by regulating the growth direction of ice crystals. Typically, the isotropic and anisotropic pore structures are formed by random (non-directional) freezing and directional freezing, respectively, and subsequent lowtemperature vacuum drying. Moreover, directional freezing includes ice crystals formation and growth along the direction of temperature gradient when the bottom of samples contacts with copper or aluminum plate cooled by liquid nitrogen. Correspondingly, applying them to encapsulate PCMs will lead to anisotropic properties of ss-CPCMs. Lv et al. [19] successfully constructed reduced graphene oxide (rGO)/AgNW aerogels (rGAA) and vacuum-impregnated them into melted lauric acid (LA) to achieve LA@rGAA ss-CPCMs. Besides an ultrahigh enthalpy efficiency of 90% and TC of 0.903 W/m·K, LA@rGAA could effectively maintain a stable temperature by absorbing/dissipating a large amount of heat during heating/cooling process, and exhibited outstanding solar-thermal conversion efficiency of 94.54%, which had potential application in waste heat recovery, intelligent TM, and solar energy utilization.
Particularly, anisotropic 3D aerogels have been extensively utilized to achieve functional ss-CPCMs with high TM and other capability due to its unique structure. In the previous report of Min et al. [5], polyimide (PI)/rGO aerogels with anisotropic structure were utilized to impregnate paraffin wax (PW). The acquired PI/rGO/PW ss-CPCMs exhibited a higher longitudinal TC of 8.87 W/m·K than transversal TC of 2.68 W/m·K, while keeping with a distinguishing latent heat retention efficiency of 98.7%, which imparted PI/rGO/PW with real-time and fast-charging solar energy conversation along longitudinal direction. Wei et al. [20] manufactured microcrystalline cellulose (MCC)/graphene nanoplatelet (GNP) aerogels with aligned stacking and 3D segregated structures. After being impregnated with PEG, the acquired MCC/GNP/PEG ss-CPCMs exhibited relatively high TC enhancement (232%) at the small GNP loading (1.51 wt%), melting enthalpy efficiency (99.84%), and excellent solar/electrical-thermal transition ability.
In addition to TM problem, the electromagnetic interference (EMI) pollution emerging with the prosperity of modern electronics would downgrade electronic operating precision and jeopardize operators' health, which has gradually become a severe social issue [21]. To this end, it is of great significance to explore an advanced material combining effective TM and superb EMI shielding capability. Generally, EMI shielding mechanism can be explained by the eddying effect theory, electromagnetic field theory, transmission line theory, and so on. Among them, the transmission line theory is the most widely recognized, thanks to its advantages of easy-to-understand, easy calculation, and high accuracy. The theory considers that the partial incident electromagnetic waves (EMWs) will be reflected back into the air by the surface of shielding materials due to a wave impedance mismatch between shielding materials and air interface. Then, the remaining EMWs penetrate into shielding materials and propagate forward, while the electric dipole or magnetic dipole of shielding materials interacts with electromagnetic field to dissipate EMWs into heat. Additionally, EMWs experience multiple reflections and transmissions at the interface between shielding materials. Usually, EMI shielding effectiveness (SE) is used to evaluate the EMI shielding ability, and Schelkunoff formulas (Equations S2-S5) are extensively applied to calculate EMI SE. Particularly, anisotropic aerogels are also considered a unique structure for EMI shielding owing to internal multilayer reflective interface. Du et al. [22] fabricated ultralight 3D MXene/ aramid nanofiber (ANF) aerogels through freeze-drying technology. The results indicated that the anisotropic aerogels were more beneficial to EMI shielding than isotropic aerogels under the same conditions. For example, the average EMI SE T of anisotropic MXene/ANF aerogels could reach 65 dB, yet the one of the isotropic MXene/ ANF aerogels were 55 dB. Chen et al. [23] prepared the anisotropic cellulose nanofibrils (CNFs)/AgNW sponges by directional freeze-drying. The obtained CNF/AgNW sponges possessed anisotropic mechanical property and electrical conductivity, and EMI shielding performances. The average EMI SE T of transversal direction was 9 dB higher than that of longitudinal direction when the AgNW content was 0.3 vol% and the thickness of sample was 3 mm. The above reports suggest that anisotropic structure is conducive to shielding EMWs, while lacking research simultaneously on their TM ability. Obviously, anisotropic structure has contributed greatly to improve either TC or EMI shielding properties. To the best of our knowledge, few relevant researches have been tried to integrate anisotropic structure with PCMs for multifunctional ss-CPCMs simultaneously with excellent TM and anisotropic EMI shielding.
Inspired by the above researches, developing an anisotropic 3D porous skeleton for adsorbing organic PCMs to attain ss-CPCMs with the integrated latent heat enthalpy, TC, and anisotropic EMI shielding properties is greatly significant but challenging. Herein, we made full use of the synergistic effect of AgNWs and GNPs to prepare the anisotropic CNF/polyacrylamide (PAM)/GNP/AgNW (CPGxAy) hybrid aerogels by directional freeze-drying method (x/y represented the ratio of GNPs to AgNWs). After encapsulating PEG, the acquired PEG@CPGxAy ss-CPCMs exhibit remarkable TC and anisotropic EMI shielding performances, while maintaining high enthalpy efficiency. In PEG@ CPGxAy ss-CPCMs, some GNPs contact with each other to provide the connected thermally conductive pathways for phonon conduction. Highly conductive 1D AgNWs can bridge the gaps between GNPs to further aggrandize electrical/thermal conduction and EMI shielding ability [24]. This design strategy takes a tremendous step towards the fabrication and tailoring of ss-CPCMs with both efficient TM and surprising anisotropic EMI SE for prospective applications in electronic devices.

Synthesis of AgNWs
As shown in Fig. S1, AgNWs were fabricated on the grounds of a previously declared procedure with certain significant revisions [25]. Typically, 350 mg PVP was added into FeCl 3 ·6H 2 O solution (0.34 mg FeCl 3 ·6H 2 O in 20 mL EG), and mixed at room temperature until a uniform EG solution was formed. Sequentially, the freshly prepared AgNO 3 solution (0.34 g AgNO 3 in 20 mL EG) was poured into the obtained solution with vigorous stirring. Sequentially, the mixture solution was transferred into autoclave for hydrothermal reaction at 180 °C for 2.5 h. Finally, the achieved solution was fully washed by C 3 H 6 O and C 2 H 6 O for several times, and then deionized water was employed to regulate the obtained AgNW dispersion concentration to 7.5 wt% for further application.

Preparation of anisotropic CPGxAy hybrid aerogels
According to Fig. 1, hybrid aerogels were prepared using simple solution blending and directional freeze-drying methods [26]. The theoretical concentration of each component in the mixed dispersion was listed in Table S1. First, CNF/ PAM hybrid suspension was prepared by adding 0.25 g PAM to CNF dispersions under magnetic stirring for 4 h. Then, a designed amount of GNP slurry was tardily poured into the above suspension to prepare CNF/PAM/GNP mixture, followed by ultrasonic bath treatment at 240 W for 1 h and subsequent intensive magnetic agitation for 4 h. Then, the synthesized AgNW dispersion was slowly dropped into CNF/ PAM/GNP mixture. The obtained CNF/PAM/GNP/AgNW mixed liquid was allowed to magnetic stirring for 8 h and then degassed by vacuum for 25 min at room temperature. After being poured into cylinder mold and standing at 4 °C for 12 h, the acquired CNF/PAM/GNP/AgNW hybrid hydrogels were placed on a copper block which was permeated in liquid nitrogen for freezing, and then lyophilized at −72 °C for 60 h to fabricate CNF/PAM/GNP/AgNW hybrid aerogels. For comparative study, pure CNF, CNF/PAM, CNF/ GNP, CNF/PAM/AgNW, and CNF/PAM/GNP hybrid aerogels were also fabricated using similar procedures. Based on the AgNW concentration of 1.92 mg/mL, the concentration of GNPs in CNF/PAM/GNP/AgNW mixed dispersion is fixed as 0.32, 0.64, 1.28, and 1.92 mg/mL, respectively. Thus, the GNP/AgNW weight ratios in the corresponding CNF/PAM/GNP/AgNW aerogels were 1:6, 2:6, 4:6, and 6:6, respectively. For brevity, the prepared hybrid aerogels were denoted as CPGxAy, where x/y represented the GNP/AgNW weight ratios.

Fabrication of PEG@CPGxAy ss-CPCMs
PEG@CPGxAy ss-CPCMs were fabricated through a vacuum-assisted melting infiltration approach. Initially, PEG was put into a vacuum oven and fully fused at 80 °C. Subsequently, the prepared CPGxAy hybrid aerogels were impregnated into the melting PEG and subjected to decompression (about 80 Pa) for 24 h, in which the bubbles in CPGxAy hybrid aerogels and molten PEG were removed and PEG was full impregnated into CPGxAy hybrid aerogel skeletons. Eventually, PEG embedded CPGxAy hybrid aerogels were taken out and wiped using soft tissue to eliminate the residual PEG. After cooling to room temperature, PEG@ CPGxAy ss-CPCMs were thus yielded. Similarly, PEG@ CNF, PEG@CNF/PAM, PEG@CNF/GNP, PEG@CPA6, and PEG@CPGx were obtained through the above method. The loading fraction (W f ) of PEG in ss-CPCMs was obtained using Eq. (1).
where m 0 and m 1 represented the weight of hybrid aerogels and ss-CPCMs, respectively.

Characterizations
The morphology and microstructures of AgNWs, GNPs, hybrid aerogels, and ss-CPCMs were observed using scanning electron microscopy (SEM, Quanta 250FEG, America) and transmission electron microscopy (TEM, Talos F200X G2, America). The energy dispersive X-ray spectroscopy (EDS) mappings were utilized to explore the distribution of Ag, O, and C elements in AgNWs and CPGxAy hybrid aerogels. The crystalline structures of hybrid aerogels and ss-CPCMs were investigated by an X-ray diffractometer (XRD, PANalytical, Netherlands), and the samples were scanned in the region of diffraction angles 5 ~ 90°. Fourier transform infrared spectroscopy (FTIR, Nicolet iS50, USA) was employed to estimate the chemical structures of the samples. The thermal properties of PEG and ss-CPCMs were further tested via thermogravimetric analyzer (TG, Schematic roadmap of the fabrication of CPGxAy hybrid aerogels and PEG@CPGxAy ss-CPCMs TG209F1, Germany) and differential scanning calorimeter (DSC, DSC-2500, America). The supercooling degree (ΔT) and melting/crystallization enthalpy efficiency (η) were calculated by Eqs. (2) and (3), respectively.
where T c and T m represented the crystallization and melting temperature, respectively. ∆H ss-CPCMs and ∆H PEG were the latent heat (melting/crystallization) enthalpy of ss-CPCMs and PEG, respectively. The thermal reusability was evaluated by thermal cycle testing using a desktop type high-low temperature (wet and hot) test chamber (CEEC-M64H-65, China) in temperature range of 30-80 °C according to ASTM 2423.22-2012, and the samples were hold at 30 °C and 80 °C for 5 min to ensure that PEG was fully solidified/melted. So as to illustrate the reinforcement effect of PAM on pure CNF aerogels, compressive test was implemented on a universal material testing machine (Instron-Model 5567, America) at a compression speed of 1 mm/min, according to ASTM D638-1999. To demonstrate the geometrical stability of ss-CPCMs, the specimens settled on a hot plate were heated up to 80 °C and photographed at 25 °C, 60 °C, and 80 °C by an optical camera. PEG@CPG6A6 ss-CPCMs were imposed by a constant force of 5 N for 30 min, when the temperature reached to 80 °C. The TC was detected by the portable TC instrument (TC3000E, XIATECH, China) at a voltage of 1.5 V. Furthermore, an IR thermal camera (FLIR ONE Pro, America) was applied to achieve the real-time temperature response, as the specimens were set on a hot plate at 80 °C.
The volume electrical conductivity (σ) of the samples with high conductivity was detected using serial resistivity tester (FT-330, Ningbo Rooko, China), and the volume resistivity (ρ) of the samples with low conductivity (< 10 -6 S/m) was measured using super megohmmeter (SM7110, Hioki, Japan) with a precision of 1.67%. The corresponding electrical conductivity was calculated via Eq. (4).

Microstructures of GNPs and AgNWs
As depicted in Fig. S2, GNPs present a plate-like morphology with some creases, which is quite close to the previously reported results [12]. The digital photo of AgNW dispersion (Fig. S3) exhibits a light gray color. The SEM images ( Fig. 2a and b) show that uniform AgNWs with few particles are interrelated with each other to form a network frame, and the TEM image ( Fig. 2c)

Structures and mechanical properties of CPGxAy hybrid aerogels
The inner structures and morphology of the representative hybrid aerogels are exhibited in Fig. 3. Figure 3a1 ~ d1 display digital images of CNF/PAM, CPA6, CPG2A6, and CPG6A6 hybrid aerogels, respectively. In Fig. 3a2 ~ d2 Fig. 3a3 ~ d3, it can be clearly found that the prepared hybrid aerogels exhibit tubelike and honeycomb-like pore structures in the longitudinal (x-o-z cross-section) and transverse (x-o-y cross-section) direction, respectively. This indicates that anisotropic 3D porous skeletons can be constructed by unidirectional freezing, which is attributed to the mechanism of oriented growth of ice crystals. When the bottom of CPGxAy hybrid hydrogels contact with copper plate cooled by liquid nitrogen, two-dimensional ice crystals on transverse direction are formed inside CPGxAy hybrid hydrogels. Afterwards, the two-dimensional ice crystals grow from bottom to top along longitudinal direction under temperature gradient, and CPGxAy compounds are excluded from ice fronts through the driving force derived from the growth process of ice crystals. Finally, anisotropic 3D CPGxAy "cell walls" are generated in ice block. After the subsequent freeze-drying process, anisotropic porous CPGxAy skeletons are finally acquired after ice crystal sublimation. Obviously, the addition of AgNWs and GNPs results in the more regular, thicker, and smoother "cell walls" due to the more chances for them to tightly interconnect with each other, which improves the strength and stability of hybrid aerogels. Furthermore, with the addition of AgNWs and GNPs, the pore numbers are higher, and the average internal pore size of CPA6, CPG2A6, and CPG6A6 hybrid aerogels are ~ 47.23 μm, ~ 46.28 μm, and ~ 44.16 μm, respectively, which are slightly smaller than that of CNF/ PAM hybrid aerogels (~ 48.71 μm) (Fig. 3a4-d4 and Fig. S4), ascribing to the increase of ice crystal nucleation density caused by the incremental GNP and AgNW content during freeze casting. This is favorable to effective PEG adsorption for remaining enthalpy efficiency. Apparently, EDS mapping images (Fig. 3e)   is conducive to high TC and anisotropic EMI SE for PEG@ CPGxAy ss-CPCMs. The XRD and FTIR tests are utilized to determine, respectively, the crystalline structures and chemical structure of CPGxAy hybrid aerogels [28]. According to Fig. S5a-d and the attribution of diffraction peaks (S3.1 and S3.2.), the as-prepared CPG6A6 hybrid aerogels present no new characteristic peak, and only some absorption peaks slightly shift, illuminating the strong physical interaction among the components instead of chemical interaction, such as hydrogen bond, Van der Walls' force.
For systematically analyzing the reinforcement of PAM on CNF aerogels, the typical compressive test is conducted to determine the mechanical properties of CNF and CNF/ PAM hybrid aerogels. The corresponding stress-strain courses are described in Fig. S6a. The results reveal that the compressive stress of CNF/PAM hybrid aerogels is much larger than that of virgin CNF at the same strain. Specifically, CNF/PAM hybrid aerogels present a compressive stress of 12.23 kPa at 50% strain, which is 1.63 times as high as that of CNF aerogels (7.5 kPa). This phenomenon demonstrates that PAM has a significant enhancement on the mechanical properties of CNF/PAM hybrid aerogels, which could be by reason of hydrogen bonding between CNFs and PAM. From Fig. S6b and c, it is visibly observed that the surface of the fabricated PEG@CNF/GNP ss-CPCMs has obvious wrinkles and collapses, while the surface of PEG@ CPG6 ss-CPCMs is relatively intact. In addition, much more black sediment is remained in the residual molten PEG after impregnating CNF/GNP hybrid aerogels, in comparison with CPG6 hybrid aerogels (Fig. S6d and e), which suggests that the strength of CPG6 hybrid aerogels is greatly improved by the incorporation of PAM. All the above results illuminate the enhancement of PAM on the mechanical properties of the aerogels.

Morphology and shape stability of PEG@ CPGxAy ss-CPCMs
The cross-sectional SEM images of PEG@CNF/PAM, PEG@CPA6, PEG@CPG2A6, and PEG@CPG6A6 ss-CPCMs are displayed in Fig. S7. It is showed clearly that no legible crevices and holes appear on all the ss-CPCM fractured surfaces, revealing the excellent interface compatibility between the prepared hybrid aerogel skeletons and PEG. Additionally, W f of PEG in PEG@CPGxAy ss-CPCMs is calculated by Eq. (1) and listed in  S8a and b) analytical results suggest that the encapsulation process has nearly no impact on the crystalline structures of PEG and there are hydrogen bonds between PEG and CPGxAy hybrid aerogels rather than the chemical bonds, which guarantees the phase transition characteristics of PEG. PEG, as a typical organic solid-liquid PCMs, is liable to leakage in the phase transition, which restricts the practical thermal-related applications. Encapsulating PEG into the aerogels with 3D interconnected networks has been proved to be one of the valid methods to solve this problem. So as to intuitively observe the form stability of the as-prepared PEG@CPGxAy ss-CPCMs, the representative samples are placed on a heating stage, and their optical photographs at different moments are taken by a digital camera and presented in Fig. S9. It is worth noted that the specimens all kept their integrated shape at 30 °C (below T m of PEG). As the temperature is raised, PEG melts initially at 60 °C (close to T m of PEG) and melts completely into a flowing liquid at 80 °C (above T m of PEG), demonstrating that the original shape of pure PEG cannot be maintained in the phase transition stage. On the contrary, the representative PEG@ CPGxAy ss-CPCMs exhibit no shape changes with little PEG leakage. Noteworthily, PEG leakage phenomenon is obviously reduced as AgNW and GNP contents increase, which can be as a result of the decreasing pore size and the incremental specific surface area of the corresponding hybrid aerogels, providing more hydrogen bonds, surface tension, and strong capillary force to adsorb and encapsulate PEG. Especially, PEG@CPG6A6 ss-CPCMs hold their initial shape well at 80 °C, even pressed by a 100 g weight [29]. These phenomena illustrate that the prepared CPGxAy hybrid aerogels with 3D porous interconnected networks can be applied as prominent supporting materials for encapsulating PEG.

Phase-transition performances of PEG@CPGxAy ss-CPCMs
Both the phase-transition temperature and latent heat enthalpy are important parameters for ss-CPCMs, when they are applied in TM systems. Thus, DSC scans are conducted to investigate the phase-transition properties of PEG and PEG@CPGxAy ss-CPCMs, and the results are displayed in Fig. S10. The specific parameters extracted from the DSC curves, regarding supercooling degree (∆T), melting enthalpy (∆H m ), and crystallization enthalpy (∆H c ), are marked in Fig. S10c and d. Obviously, the lower supercooling degree of PEG@CPGxAy ss-CPCMs is acquired in contrary to that of pure PEG, which is beneficial for outputting higher latent heat in the thermal discharging process. The reason for this phenomenon is that the physical interactions, including capillary force, hydrogen bonding, and surface tension, can promote the crystallization of PEG encapsulated in CPGxAy hybrid aerogels, which leads to the decreased melting temperatures and increased crystallization temperatures.
As depicted in Fig. S10d, PEG exhibits high melting enthalpy of 179.31 J/g and crystallization enthalpy of 175.42 J/g. Both the crystallization and melting enthalpies of PEG@CPGxAy ss-CPCMs tend to reduce with the increase of GNP loading, due to no phase change capacity of hybrid aerogel skeletons. However, it can be found that the latent heat enthalpy efficiency of PEG@CPGxAy ss-CPCMs is quite high (> 93%) and close to the value of the PEG loading fraction. For instance, the ∆H m and ∆H c of PEG@ CPG6A6 ss-CPCMs with the lowest PEG loading fraction (93.67 wt%) are up to 167.6 J/g and 163.28 J/g, respectively, which exhibit 93.47 and 93.08% melting/crystallization enthalpy efficiency. The aforementioned results prove that the obtained PEG@CPG6A6 ss-CPCMs possess excellent heat storage ability with the negligible loss of latent heat enthalpy, which is of great significance for their great applications in thermal energy storage.

Thermal stability and reusability of PEG@ CPGxAy ss-CPCMs
The excellent thermal stability is essential for ss-CPCMs, when they are applied in TM. Thus, TG measurements are performed to evaluate the thermal stability of PEG and PEG@CPGxAy ss-CPCMs (Fig. 4a-b and Table S3).
Similar to PEG, all the PEG@CPGxAy ss-CPCMs undergo a typical one-step thermal decomposition behavior from 350 to 420 °C, which is mainly resulting from the volatilization of the linear PEG molecules. Both PEG and PEG@CPGxAy ss-CPCMs barely fluctuate until 200 °C, which indicates that PEG and PEG@CPGxAy ss-CPCMs possess ideal thermal stability at normal practical conditions. Reusability is a critical parameter to assess the service life of ss-CPCMs in practical applications. Hence, the 50 and 100 heating-cooling cycle tests over 30-80 °C are performed to measure the reusability of PEG@CPG6A6 ss-CPCMs. Then, DSC curves, latent heat enthalpy, FTIR spectra, and XRD patterns of PEG@CPG6A6 ss-CPCMs after 50/100 heating cycles are given in Fig. 4c-f. Apparently, the DSC curves of PEG@CPG6A6 ss-CPCMs after 50 and 100 thermal cycles are both in accordance with those of the original PEG@CPG6A6 ss-CPCMs (Fig. 4c). Moreover, the calculated loss percentage of melting and crystallization enthalpy after 50 thermal cycles are only 0.46% and 0.42%, respectively, and the ones after 100 thermal cycles are merely increased to 1.28% and 1.32%, respectively, indicating that PEG@CPG6A6 ss-CPCMs maintain their high latent heat enthalpy during such a long cyclic test (Fig. 4d). After 50 and 100 thermal cycle experiments, the characteristic peaks of XRD patterns and FTIR spectra are almost unchanged, signifying the favorable stability of chemical structures and crystalline structures for PEG@CPG6A6 ss-CPCMs during the heating/cooling cycles (Fig. 4e and f). Moreover, the TG thermograms (Fig. S11) after 50 and 100 thermal cycles do not obviously change in contrast to those of PEG@CPG6A6 before cyclic test, which reveals that cyclic thermal tests have nearly no influence on the thermal stability of PEG@CPG6A6. In brief, the fabricated PEG@ CPG6A6 ss-CPCMs with great reusability have immense application potential in the TM area.

Thermal conductivity of PEG@CPGxAy ss-CPCMs
TC is one of the important factors for ss-CPCMs, which determines the heat transfer efficiency in the phase-transition procedure. The higher TC corresponds to the less time needed for thermal energy transfer. PEG corresponds an inherent small TC of 0.28 W/m·K, which is unfavorable for heat transfer. Thus, it is necessary to elevate TC by introducing thermal conduction networks. TC of PEG@CPGx ss-CPCMs is gradually augmented with the incremental GNP content (Fig. S12). Detailedly, PEG@CPG6 ss-CPCMs display a rather high TC of 0.72 W/m·K at the GNP content of 2.19 wt%. After incorporating AgNWs of 2.19 wt%, TC of PEG@ CPGxAy ss-CPCMs is further improved in comparison with that of PEG@CPGx ss-CPCMs at the same GNP loading fraction (Fig. 5a). For example, PEG@CPG6A6 ss-CPCMs filled with 2.19 wt% AgNWs and 2.19 wt% GNPs exhibit the maximal TC of 0.84 W/m·K, enhanced by 200% in contrary to that of pure PEG, suggesting that CPGxAy hybrid aerogels combined with AgNWs and GNPs tend to construct more perfect thermally conductive networks. This phenomenon is on account of the synergetic effect of 1D AgNWs and 2D Fig. 4 a TG and b DTG thermograms of PEG and PEG@CPGxAy ss-CPCMs as indicated in the graphs. c DSC curves, d latent heat enthalpy, e XRD patterns, and f FTIR spectra of PEG@CPG6A6 ss-CPCMs before and after 50/100 thermal cycles GNPs; that is, 1D AgNWs can bridge the gaps between 2D GNPs and consequently promote the free pathways for phonons conduction [30].
In addition, the comparison of the filler loading fraction, TC enhancement, and latent heat enthalpy efficiency for the prepared PEG/CPG4A6 and PEG/CPG6A6 with those of some other reported ss-CPCMs are enumerated in Fig. 5b and Table S4. It is rather difficult to achieve simultaneously fascinating TC and latent heat enthalpy efficiency in the case of ss-CPCMs. Generally, high filler loading is usually desired to reach the satisfactory TC, but accompanied with the sacrifice of latent heat enthalpy because of no phase change capacity of the thermally conductive filler. In this study, it is obvious that PEG/CPG6A6 ss-CPCMs exhibit a competitive TC and latent heat enthalpy efficiency at a rather low filler content (2.19 wt% AgNWs and 2.19 wt% GNPs for PEG/CPG6A6 ss-CPCMs), which prevail over most the reported ss-CPCMs. The comparison results demonstrate that the acquired PEG/CPG6A6 ss-CPCMs have enormous potential applications in TM.
To observe intuitively the heat transfer property of PEG@ CPGxAy ss-CPCMs, an IR thermal camera is applied to record the surface temperature of the representative PEG@ CPGxAy ss-CPCMs settled on a hot plate at 80 °C for different time. As displayed in Fig. 5c, the surface color of PEG@CPG6A6 ss-CPCMs change more quickly than those of PEG@CNF/PAM, PEG@CPA6, and PEG@CPG2A6 ss-CPCMs, suggesting the thermal response rate is enhanced after filling AgNWs and GNPs. It is obvious that the thermal response rate tendentiously displays a similar growing tendency to TC. Furthermore, the temperature-time plots (Fig. 5d) reveal distinctly that the surface temperature of showing the temperature of the representative PEG@CPGxAy ss-CPCM surface settled on a hot plate at 80 °C. d The temperature changes of the representative PEG@CPGxAy ss-CPCMs with the increase of heating time the representative PEG@CPGxAy ss-CPCM samples at the same heating time is in the following sequence: PEG@ CPG6A6 > PEG@CPG2A6 > PEG@CPA6 > PEG@CNF/ PAM. This result illuminates the heat transfer property of PEG@CPGxAy ss-CPCMs is gradually enhanced with the increase of GNP content, which is ascribed to the more perfect thermally conductive networks at higher GNP content. More importantly, the obvious temperature "platform" between the temperatures of 48 and 57 °C is observed in PEG/CPG6A6 ss-CPCMs during the heating process; thermal energy generated by hot plate is absorbed and stored, which is caused by the solid − liquid phase-transition of PEG/CPG6A6 ss-CPCMs. Therefore, if it is attached on the surface of electronics, the operating temperature will be kept within a certain range close to the phase change temperature of the ss-CPCMs, which can protect electronics from overheating and/or dramatic temperature fluctuation.

Anisotropic EMI shielding performances of PEG@CPGxAy ss-CPCMs
ss-CPCMs are extensively applied in advanced electronics, which can efficiently store/release thermal energy in phase transition period. Apart from pretty TM, EMI shielding property is also requisite for keeping long-term operational life of electronics. It is well known that EMI shielding performance is significantly impacted by the electrical conductivity. According to Fig. 6a, the volume electrical conductivity of CNF/PAM hybrid aerogels in the longitudinal direction (along the freezing direction) is < 10 −5 S/m,  [2,3,14,[39][40][41][42] proving the fact that CNF/PAM hybrid aerogels, like most the polymers, are almost the insulator. After introducing AgNWs, the longitudinal volume electrical conductivity of the obtained CPA6 hybrid aerogels increases sharply to 55.15 S/m, indicating AgNWs construct continuous electrical pathways in CPA6 hybrid aerogels. Noticeably, the volume electrical conductivity of CPG1A6 hybrid aerogels in the longitudinal direction significantly drops to 0.34 S/m in contrast to that of CPA6 hybrid aerogels. The most probable reason is that the dispersing agents in GNP slurry increase contact resistance between AgNWs and GNPs. Moreover, it is worth noting that the volume electrical conductivity of CPG6A6 hybrid aerogels is gradually augmented to 16.25 S/m in the longitudinal direction with the increasing GNP content, demonstrating the effective electrical networks are built at the AgNWs:GNPs of 6:6. Furthermore, the transverse conductivity is always moderately lower than the longitudinal conductivity for the same specimen (Fig. 6d). For instance, CPG6A6 hybrid aerogels exhibit the smaller conductivity of 8.03 S/m in the transverse direction, which illuminates the obvious anisotropic electrical conductivity. However, all the volume electrical conductivity of PEG@ CPGxAy ss-CPCMs is reduced by one order of magnitude than that of the corresponding CPGxAy hybrid aerogels (Fig. S13), which attributes to the physical isolation of insulating PEG.
The EMI shielding performances of PEG@CPGxAy ss-CPCMs with a thickness of 6.0 mm are displayed in Fig. 6b-c, Fig. 6e-f, and Fig. S14. It is observed that PEG@ CNF/PAM ss-CPCMs are almost apparent to electromagnetic waves (EMWs) over x-band with the average EMI SE T of only 2.01 dB in the longitudinal direction and 2.04 dB in the transverse direction, because of their intrinsic nonconductive property. What's more, it can be easily found that the growing trend of EMI SE T of PEG@CPGxAy ss-CPCMs is in line with that of the volume electrical conductivity. Specially, the average EMI SE T of PEG@CPA6, PEG@ CPG1A6, PEG@CPG2A6, PEG@CPG4A6, and PEG@ CPG6A6 ss-CPCMs in the longitudinal direction is 36.56, 21.06, 25.47, 30.89 and 35.21 dB, respectively, satisfying the demands for the commercial EMI shielding materials (> 20 dB) [43]. Noteworthily, EMI SE T of PEG@CPGxA6 is gradually improved with the rising GNP content, which attributes to the enhancement of the corresponding volume electrical conductivity. Appreciably, the EMI SE T along the transverse direction is greatly superior to that along the longitudinal direction for all the PEG@CPGxA6 ss-CPCMs. Particularly, the transverse average EMI SE T of PEG@ CPA6, PEG@CPG1A6, PEG@CPG2A6, PEG@CPG4A6, and PEG@CPG6A6 ss-CPCMs is enhanced to 84.30, 55.52, 61.10, 67.11, and 71.08 dB, respectively, indicating the excellent anisotropy of EMI shielding. All the above results prove that the designed CPGxAy hybrid aerogel skeletons with ideal conductive networks constructed by AgNWs and GNPs endow the PEG@CPGxAy ss-CPCMs with the excellent anisotropic EMI shielding performances. Importantly, PEG@CPG6A6 ss-CPCMs with a relatively small amount of 2.19 wt% AgNWs and 2.19 wt% GNPs exhibit the large average EMI SE T difference (∆SE T ) of 35.87 dB in the different directions. The underlying mechanism is that EMWs need to pass through more "cell walls" of the oriented tubelike pores along the transverse direction, which contributes to the increased multiple reflection and absorption of EMWs. Regardless of GNP content and incident direction, EMI SE A is always higher than EMI SE R , owing to the increasing moving charge carriers together with the formation of conductive networks. For example, EMI SE T , SE A , and SE R of PEG@CPG6A6 ss-CPCMs in the transverse direction are 71.08, 64.54, and 6.54 dB, respectively, revealing 90.80% of SE A in SE T . Natheless, Fig. 6g and h reveal that the coefficient R is always bigger than the coefficient A, which demonstrates that the shielding principle is dominated by microwave reflection rather than microwave absorption.
So as to evaluate EMI SE of this present work with the latest reported ss-CPCMs, the comparison of filler content and average EMI SE T are summarized in Fig. 6i and Table S5. Conspicuously, PEG@CPG6A6 ss-CPCMs in this work exhibit prominent average EMI SE T value of 71.08 dB at relatively low filler content (2.19 wt% AgNWs and 2.19 wt% GNPs), which exceeds those of the most ss-CPCMs reported in previous literatures. Combining the previously discussed high TC enhancement (200%) and latent heat enthalpy efficiency (93.47%), this study provides valuable guidance for the preparation of ss-CPCMs simultaneously with excellent TC and wonderful anisotropic EMI SE towards potential applications in TM and EMI shielding areas.
To intuitionally illustrate the EMI shielding mechanisms, the process of the EMWs transferring across PEG@CPGxAy ss-CPCMs in longitudinal and transverse directions is displayed in Fig. S15a and b, respectively. First, the partial incident EMWs is immediately reflected when EMWs reach the surface of PEG@CPGxAy ss-CPCMs because of the impedance mismatch between the conductive PEG@CPGxAy ss-CPCMs and the nonconductive air [44]. When EMWs are incident along the transverse direction, EMWs hit the x-o-z plane with the higher longitudinal electrical conductivity and thus is more difficult to enter the sponge shields, directly leading to the higher transverse SE R . In addition, the penetrating EMWs along the transverse direction will meet a large number of "cell walls" for the tubelike pores, which scatters them forth and/or back in the tubelike pores until full dissipation. Thus, such multiple reflections across the tubelike pores can largely augment the absorption of EMWs and overall shielding performance in transverse direction. Moreover, the electrical conduction networks constructed by AgNWs and GNPs provide the conductive pathways for electron migration and hopping, resulting in the conduction losses. Finally, the extremely small part of EMWs transmits across PEG@CPGxAy ss-CPCMs, and high EMI SE T in the transverse direction is achieved. In contrast, when EMWs are incident along the longitudinal direction, a portion of EMWs can stealthily transit along the aligned channels, inducing the poorer EMI shielding property. Thereby, the anisotropic CPGxAy hybrid aerogel skeletons account for the remarkable anisotropy of EMI shielding.
The Bluetooth signal blocking experiment between Apple phone and Android phone is conducted to intuitively assess the EMI shielding performance of PEG@ CPG6A6 ss-CPCMs (Fig. 7). Under naked conditions, the Bluetooth of Apple phone and Android phone can be successfully connected with one another (Fig. 7a). When the android phone is completely wrapped by tinfoil, the Bluetooth fails to connect because of pretty EMI shielding of tinfoil (Fig. 7b). After making a small hole on the tinfoil, the Bluetooth is successfully connected again (Fig. 7c). Interestingly, the Bluetooth connection is subsequently broken down when the tinfoil hole is covered by PEG@CPG6A6 ss-CPCM samples, no matter whether the longitudinal direction or the transverse direction is perpendicular to the cross-section of tinfoil hole (Fig. 7d and  e). The above demonstrations suggest that the designed PEG@CPG6A6 ss-CPCMs have great practical application for EMI shielding in advanced electronics.

Conclusion
In summary, the thermally conductive PEG@CPGxAy ss-CPCMs with superb latent phase change enthalpy and prominent anisotropic EMI shielding property were successfully manufactured by embedding PEG into CPGxAy hybrid aerogels. The anisotropic CPGxAy hybrid aerogels were firstly constructed via directional freeze-drying technology. The porous CPGxAy skeletons with pretty thermal and conductive networks impart PEG@CPGxAy ss-CPCMs with the integrated wonderful form stability, high enthalpy efficiency, outstanding cyclic reliability, outstanding TM, and excellent anisotropic EMI shielding property. Arising from the synergistic effect of GNPs and AgNWs, PEG@ CPG6A6 ss-CPCMs exhibit excellent TC of 0.84 W/m·K (200% increase in contrast with that of neat PEG) and anisotropic average EMI SE T (up to 71.08 dB in the transverse direction and 35.21 dB in the longitudinal direction) at a relatively low content of AgNWs (2.19 wt%) and GNPs (2.19 wt%). Additionally, the melting and crystallization enthalpy are as high as 167.6 and 163.28 J/g, respectively, by virtue of the extremely high PEG loading fraction (93.67%). This effort opens a new route for the construct and modulation of ss-CPCMs simultaneously with efficient TM and wonderful anisotropic EMI shielding performances towards potential applications in modern electronics.