Attapulgite-reinforced polyimide hybrid aerogels with high dimensional stability and excellent thermal insulation property

Polymer 176 (2019) 196–205

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Attapulgite-reinforced polyimide hybrid aerogels with high dimensional stability and excellent thermal insulation property

T

Tingting Wu, Jie Dong∗∗, Guofen Xu, Feng Gan, Xin Zhao, Qinghua Zhang∗ State Key Laboratory for Modification of Chemical Fibers and Polymer Materials, College of Materials Science and Engineering, Donghua University, Shanghai 201620, PR China

H I GH L IG H T S

nanorods are used as reinforcing nanofiller for PI aerogels via H-bonding interaction. • AT of composite aerogels increases as the density decreases. • Modulus mechanical properties of aerogels is due to AT supporting skeleton. • Improved • AT-reinforced aerogels exhibit an excellent thermal dimensional stability.

A R T I C LE I N FO

A B S T R A C T

Keywords: Polyimide Attapulgite Anti-shrinkage

Strong low-density polymer aerogels have received intensive attention as thermal or sound insulator, filtration media and sensors, etc. However, dimensional instability has been regarded as one of issues that limits the wide application of polymer aerogels in a harsh temperature. In present work, polyimide aerogels were reinforced with attapulgite (AT) nanorods, one of natural fibrillar minerals, through the strong hydrogen-bonding interaction. In hybrid aerogels, AT nanorods act as the rigid skeleton that supporting the framework of aerogels. The resultant hybrid aerogels exhibit improved mechanical properties. At the 80% strain, the compressive stresses for the aerogels containing 10 wt% AT are 100% as high as those of the pure polyimide aerogels. Interestingly, modulus of hybrid aerogels increases as the density decreases. For example, the compression Young's and specific modulus increase by 100% and 105% for the hybrid ODA-based aerogel compared to the pristine sample, meanwhile, the density decreases by 10%. Furthermore, the AT nanorods show strong supporting effects in maintaining the structural integrity of aerogels, which produce a significant effect on reducing the shrinkage of hybrid aerogels at high temperatures and retaining their excellent thermal insulation performance. The detailed investigation reveals that the AT is an effective and low-cost additive for preparing high performance polymer nanocomposites aerogels with improved mechanical property and thermal dimensional stability.

1. Introduction Aerogels have attracted a great attention due to their large internal surface areas, low densities, excellent adsorption and insulation behaviors imparted from their inherent characters of high continuous porosity [1]. Thus, aerogels have been regarded as ideal potential substitutes for traditional foams in adsorption, insulation and mechanical cushioning applications. For aerogels, with the chemical compositions ranging from inorganic oxides (clay [2], silica [3], metal [4], carbon [5]) to polymers (polybenzoxazine [6], cellulose [7], polyurethane [8] and polyimide [9]) that have been intensively investigated up to now,

∗

silica aerogels have played a dominating role in academia and industry. For example, silica aerogels have been utilized as insulating materials in PATHFINDER MARS mission to protect the electric devices in the alternating cold and hot environment [10]. However, a major drawback is that native silica aerogel is brittle without plasticizer due to their intrinsic pearl-necklace-like skeletal framework. To alleviate this limitation, numerous works [11–13] on modifying silica aerogels by crosslinked polymers between the necks of particles have been explored and the strength of silica aerogels could be improved with different degrees. Nevertheless, the relatively poor thermal stability of common polymer reinforcements is the fatal defect for the reinforced silica aerogels,

Corresponding author. Corresponding author. E-mail addresses: [email protected] (J. Dong), [email protected] (Q. Zhang).

∗∗

https://doi.org/10.1016/j.polymer.2019.05.007 Received 18 March 2019; Received in revised form 24 April 2019; Accepted 3 May 2019 Available online 06 May 2019 0032-3861/ © 2019 Published by Elsevier Ltd.

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which can weaken the excellent thermal stability and fire-retardant characteristics of original silica aerogels. Recently, a different class of rigid-rod high-performance polymer aerogels, polyimide (PI) aerogels, have intrigued numerous scientists and engineers attributed to the combination of extraordinary thermal, mechanical, and electrical properties. Polyimide aerogels have been fabricated as flexible thin films [9] or monoliths [14] with excellent moisture resistance, and used as mechanically strong particulate air filters [15] and a low dielectric constant substrate for lightweight antennas [16]. In addition, polyimides have a flexible molecular designation, and the specific tailoring of their chemical structure can be regarded as a feasible way in obtaining the desired properties of resultant aerogels. Over the past few years, Meador's group from NASA have done extensive and fruitful fundamental researches to tune the polyimide chemical structures for tailoring the properties of resultant polyimide aerogels. Typically, they found that some readily available compounds, including 1,3,5-benzenetricarbonyl trichloride (BTC) [17] and amine functionalized polyoligomeric silsesquioxane (OAPS) [10], could promote strong polyimide aerogels when used as the cross-linking agent and were regarded as great substitutes for some expensive crosslinkers. Meanwhile, they introduced 2,2-bis(3,4-dicarboxyphenyl)hexafluoropropane dianhydride (6FDA) blocks to the polyimide structure and resulted in a novel fluorinated aerogel with a relatively low dielectric constant (1.084) by increasing the free volume of polymer chains [18], which still had comprehensive moduli of 4–8 MPa (40 times higher than the silica aerogel at the same density). Generally, relative to other polymeric aerogels, polyimides had a higher thermal resistance and could be able to keep their performances at high temperatures. However, Meador's results indicated that PI aerogels always exhibited a shrinkage of 30% at 150 °C due to the collapsing of the aerogel infrastructure [17], which was regarded as one of the critical issues that limited their utilization in some rigorous conditions. By incorporating carboxylic acid-functionalized cellulose nanocrystals (CNC–COOH), the shrinkage of aerogels could be significantly reduced during aging even treated at 200 °C [19], illustrating the effective reinforcement of nanofillers in retaining the structural integrity of aerogels. As previously described, dimensional stability of polymer aerogels is a vital property for rigorous conditions, especially for them applied in the space exploration or a long-term high temperature environment. Over the past few years, the use of nanofillers, including carbon nanotubes [20], carbon nanofibers [21], graphene oxide [22], cellulose nanofibrils (CNF) [7], clay [23] as well as polyacrylonitrile fibers [24,25], to improve the properties of composite aerogels were well investigated. Feng et al. [24] reported that by impregnating the resorcinol-formaldehyde aerogels with oxidized PAN fibers, a great reduction in the shrinkage of the resulted carbon aerogels (from 11% for the pristine aerogel to 5.2% for the hybrid aerogel) was observed in the carbonization process. Guo et al. [26] illustrated that the addition of graphene oxide (GO) into resorcinol-formaldehyde (RF) matrix could effectively reinforce the aerogels' framework, thus suppressing the collapse of nanopores in RF aerogels. They illustrated that by incorporating 2.0 wt% GO, the linear shrinkage of the resultant hybrid aerogels decreased from 28.3% to 2.0%. In other works, aerogels reinforced by PU/PEO electrospun nanofibers [27] or natural CNF [28] exhibited an obvious improvement in flexibility and yield strength in relative to their nonreinforced counterparts. Berglund et al. [29] pioneered aerogels combining poly (vinyl alcohol) (PVA) with clay, and they also added the rod-like CNF as the reinforcement. The yield strength of the CNF corbelled PVA/clay aerogels was nearly 15 times higher than the neat PVA/clay aerogel, and more importantly, the hybrid aerogel showed an excellent dimensional stability even soaked into water for 24 h mainly attributed to the formation of foam-like cellular morphologies of CNF in aerogels. Nevertheless, an obvious deficiency for these polymeric nano-reinforcement is that they are flammable and the thermal stabilities are not sufficiently high, easily resulting in the

oxidation of polymer matrix when exposed to an environment with sufficient heat and oxygen. Additionally, some other nanofillers, such as GO and carbon nanotubes, are still relatively high-cost and adverse to commercial production. Attapulgite clay (AT) is a typical porous silicate fibrillar mineral containing ribbons. In addition to its natural and unique fibrillar structure and low production cost, these rod-like nanofillers can be efficiently blended with polymer matrix. Furthermore, the abundant silanol moieties (Si–OH) on the surface of AT nanorods provide active sites for forming the hydrogen-bonding interaction with the polymer matrix. Recently, AT nanorods were illustrated to have an obvious reinforcement in a wide range of polymer matrix, including epoxy [30], polyurethane [31], polyamide-6 [32], chitosan/poly (vinyl alcohol) [33] and polylactic acid [34], to name a few. Accordingly, the onedimensional AT nanorods with an ideal geometric structure (length in the micrometer and diameter in only a few nanometers) were always regarded as ideal candidates for nano-building blocks to construct macroscopic structures. Herein, we report a facile approach to prepare ultralight, low shrinkage, high-strength and thermally stable polyimide/AT hybrid aerogels with improved thermal dimensional stabilities by using AT nanorods as the supporting skeleton. Diamines either 2,2′-dimethylbenzidine (DMBZ) or 4,4′-oxydianiline (ODA) were used to synthesize polyimides with different backbone flexibilities, and the reinforcing effects of rigid AT nanorods on the shrinkage behavior, compressive property, thermal stability and thermal dimensional stability as well as the thermal insulation of these two different polyimide matrixes have been investigated in detail. 2. Experimental section Materials. Diamines ODA and DMBZ were both obtained from Changzhou Sunlight Pharmaceutical Co., Ltd. The dianhydride biphenyl-3,3′,4,4′-tetracarboxylic dianydride, BPDA, was provided by Shijiazhuang Haili Chemical Co., Ltd. Cross-linking agent 1,3,5-triaminophenoxybenzene, TAB, was purchased from Haorui Chemicals Co., Ltd. N-methyl-2-pyrrolidone (NMP) was obtained from Adamas Chemicals. Pyridine and hydrochloric acid (HCl) were purchased from Sinopharm Chemical Reagent Co., Ltd. (China). The crude AT was obtained from Jiangsu Jiuchuan Nano-material Technology Co., Ltd. (China). Purification of Attapulgite. The crude AT powder was first dispersed in deionized water with 1 mol/L HCl for 12 h under mechanical stirring, and then was washed using the deionized water until the pH equilibrated to 7. The obtained AT powders were heated in the vacuum oven at 120 °C for 48 h to remove the absorbed and bonded water. 2.0 wt % AT suspension was prepared by dispersing the purified AT in NMP under vigorous stirring and collecting the upper AT/NMP suspensions. Preparation of PI/AT hybrid aerogels. The procedure of preparing PI/AT hybrid aerogels is shown in Scheme 1. The organ-soluble poly (amic acid)s (PAAs) end-capped with anhydride units were firstly synthesized using a molar ratio of diamine: dianhydride of (n+1): (n), where n is the number of repeat units in the oligomers capped with anhydride (n = 30 was designed in this work). For getting a PAA/NMP solution with a proper viscosity, the solid content in the reaction was kept to 10 wt%. Then the cross-linking agent TAB was added into the PAA precursors, and the TAB:BPDA mole ratio was controlled to 1:40. Acetic anhydride (8:1 mol ratio to BPDA) and pyridine (1:1 mol ratio to acetic anhydride) were utilized as the chemical cyclization agent to prepare polyimide wet gels. Finally, the PI wet gels were treated by the supercritical drying process and the PI aerogels were obtained [35]. Various amounts of AT suspensions were introduced to the PAA solutions. Noteworthy, the solution should be poured into a cylindrical mold (with a diameter of 3 cm and a length of 9 cm) immediately as the acetic anhydride and pyridine were added. The gelation would 197

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Scheme 1. Procedure for preparation of AT-supported polyimide hybrid aerogels.

3. Results and discussion

occurred within 20 min and the wet gels were aged for 24 h in the mold. Subsequently, the solvent NMP within the wet gels was then gradually exchanged by acetone using a stepwise replacement method. In detail, the solid wet gels were firstly immersed in a solution of 75 v% NMP in acetone and kept overnight. Then the gels were extracted into a solution of 25% NMP in acetone and soaked for another 24 h. Finally, the solvent was replaced by a pure acetone and samples were kept for one day. The obtained wet gels were treated by the supercritical CO2 extraction under 13 MPa and finally were dried at 70 °C for 12 h. The resulting sample is defined as ODA- or DMBZ-AT-X, where X represents corresponding mass proportions of AT nanorods in the hybrid aerogels. Characterization. FTIR spectra were performed on a Nicolet 8700 spectrometer over the wavenumber range from 500 to 4000 cm−1. The 13 C NMR spectra were obtained on a nuclear magnetic resonance spectrometer (Bruker Avance 400). BET surface areas of the prepared aerogels were tested on a Quantachrome Autosorb 1-MP at 77 K. X-ray spectra were recorded by a Rigaku Dmax-2550 diffractometer with Cu Kα radiation. The nanostructures of AT were characterized by the HRTEM (JEOL 2100F). Morphologies of aerogel nanocomposites were tested by a FESEM (HITACHI S-4800). The skeletal density of aerogels was tested by a Micromeritics Accupyc 1330 helium pycnometer and the porosities of aerogels were obtained by the following equation:

Natural AT is a kind of fibrillar structural hydrated magnesium aluminum silicate mineral. Attributed to their abundant resource, high specific surface areas, good mechanical strength and thermal stability, AT has been widely used as adsorbents for removing the heavy metal [36] and organic pollutants [37] in industrial effluent or as an ideal candidate for reinforcing polymeric materials. As shown in Fig. 1(A), the obtained AT powder showed a muddy color. Typical HRTEM of AT nanorods is shown in Fig. 1(B), and the AT has a one-dimensional fibrous morphology. The diameter of AT is around 10–30 nm, which is comparable to those of MWCNTs and biological nanofibrils of cellulose. When dispersed in NMP by ultrasonication, the dispersion is stable at room temperature, as shown in the inset of Fig. 1(B), and can be used directly in the in-situ preparation of aerogel nanocomposites. The chemical structure of AT nanorods was verified by FTIR. In Fig. 1(C), the appearance of an intense peak at 1030-998 cm−1 is assigned to the asymmetric stretching of Si–O–Si bond of silsesquioxane architecture. The band at 800 cm−1 corresponds to stretching absorption of Si–O–Al. Other strong characteristic absorption bands at 1651 cm−1 and in the region of 3615–3420 cm−1, respectively, are attributed to the stretching resonances of silanol groups (Si–OH) on the surface of AT, which endow that the AT nanorods can strongly interact with the polymer matrices via the hydrogen-bonding interaction. A typical N2 adsorption-desorption isothermal of AT is presented in Fig. 1(D). It corresponds to isotherm Type IV in the IUPAC classification [38], suggesting the predominant mesopores character of the AT. The specific surface area (BET) is calculated to be 153 m2/g, indicating the porous structure of AT nanorods, which is meaningful for preparing ultralight and large specific surface aerogels. Numerous works [39–41] have illustrated that constructing the strong interfacial regions is an essential condition for achieving the excellent mechanical properties within the composites. It can be believed that the presence of Si–OH units on the surface of AT allows the formation of hydrogen-bonding interaction with the polar carbonyl units within polyimide, which benefits the effective strengthening and load transfer from polyimide to AT nanorods. The consistent conclusion can be verified by FTIR results. As shown in Fig. 2(A), the FTIR spectra show that all aerogels show the characteristic imide peaks at

Porosity=(1-ρb/ρs) × 100% where ρs and ρb represent the skeletal and bulk densities of aerogels, respectively. Dimensional change, or the linear shrinkage, of the prepared aerogels was evaluated according to the diameter change of samples. Cylindrical aerogel samples with a height to diameter ratio of 2:1 were used in the compression tests by using a WDW3020 Micro control electronic universal testing machine, and the compression rate was controlled to 2 mm/min. The thermal properties of aerogels were analyzed by Netzsch TG 209 F3 Thermogravimetric Analyzer. Thermographic pictures were obtained using a FLIR A655sc infrared thermography (FLIR Company, USA). During the tests, samples were placed on a hot plate with T = 300 °C and the measurement time was controlled to 300 s.

198

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Fig. 1. (A) The photograph of purified AT; (B) TEM image of AT (inserted picture of NMP-dispersed AT suspension after sonication); (C) FTIR spectra of AT nanorods with main peaks assignment and (D) N2 adsorption-desorption isotherm and pore size distribution curves of the AT powder.

1780 cm−1, 1720 cm−1 and 1380 cm−1. In addition, both ODA-AT-10 and DMBZ-AT-10 samples show broad absorption bands at the range of 3000–3600 cm−1, as shown in the magnification of the spectra, while, which can not be observed for ODA-AT-0 and DMBZ-AT-0. It is well accepted that both Si–OH (ca. 3600-3000 cm−1) and C]O stretching (ca. 1780 and 1720 cm−1) are hydrogen-bonding sensitive, and the –OH and C]O shift to lower (-OH) and higher (C]O) wavenumbers, respectively, upon the formation of hydrogen-bonding interaction [42]. As shown in Fig. 2(B), the C]O asymmetric stretching vibration band at around 1770 cm−1 in the ODA-AT-0 substantially shifts to a higher wavenumber of 1774 cm−1 in ODA-AT-10 at a AT loading of 10 wt%, and correspondingly, the C]O symmetric stretching vibration band shifts from 1714 cm−1 to 1716 cm−1, which are mainly ascribed to the hydrogen bonding between the hydroxyls on the AT surface and the polar C]O groups in the polyimide chains. By deconvoluting the FTIR spectra in the range of 1740–1690 cm−1 based on the identification of different C]O stretching vibrations, we can investigate this strong interfacial interaction more precisely. In Fig. 2(C), four different peaks are separately identified ranging from 1740 to 1790 cm−1 for the ODAAT-0 aerogel. Konieczny and Wunder et al. [43] reported that the band I and band IV were mainly attributed to the carbonyl symmetric and asymmetric stretching from the intermolecular chemical links, and the Band II and III referred to the “free” and hydrogen-bonded carbonyl symmetric stretching from imide rings, respectively. For ODA-derived PI aerogel without incorporating AT, the hydrogen-bonding interaction mainly formed between the N–H of the intermolecular links and the C] O in imide rings, as reported in the Wu's previous work [44]. While, the hydrogen bonding associated C]O symmetric stretching (band III) shifts to the wavenumber of 1716 cm−1 for the ODA-based aerogel containing 10 wt% AT nanorods. Apparently, the band III to band II area ratio increases greatly when incorporating 10 wt% AT nanorods into the PI matrix, indicating the increased percentage of hydrogenbonded C]O in imide rings and illustrating the formation of strong

hydrogen-bonding interaction between PI and AT nanorods. For the DMBZ-derived hybrid aerogels, in the curve-resolved FTIR spectra as shown in Fig. S1, the DMBZ-AT-10 aerogel also shows an increase in area ratio of band III to band II compared to the DMBZ-AT-0 sample. However, band III to band II area ratio for the DMBZ-AT-10 is much lower than that of the ODA-AT-10 aerogel, meaning that a stronger hydrogen-bonding interaction inclines to be formed between the rigid AT nanorods and the flexible ODA-derived PI polymer chains. Solid state 13C NMR spectra of the prepared aerogels were shown in Fig. 3(A), a carbonyl peak at δ = 165 ppm in imide rings and broad aromatic peaks in 120–145 ppm can be observed in all spectra. For DMBZ-based aerogels, an additional peak at δ = 20.2 ppm corresponding to the pendant methyl groups of DMBZ in relative to the ODAbased samples. Any acid and amide signals from the precursor PAAs intermediate at ca. 175 ppm are undetectable, revealing the full imidization of five-membered polyimide aerogels. Furthermore, the incorporated AT can not be detected by the 13C NMR spectra due to the absence of carbon atoms in the molecular formula. X-ray diffraction measurements were conducted to determine the crystalline structures in variable aerogels. As shown in Fig. 3(B), the patterns of pure ODAbased and DMBZ-based PI aerogels and AT are included for the comparison purpose. The diffraction patterns of ODA-AT-10 and DMBZ-AT10 hybrid aerogels exhibit the intense and characteristic AT crystalline peak at 2θ = 8.1° matching extremely well with the AT and two pure PIs patterns, indicating that AT-nanorods and PI structures are present in the hybrid aerogels. All samples exhibit a broad reflection peak at 2θ = 15.5°, representing their amorphous nature. Additionally, ODAand DMBZ-based aerogels show sub-diffraction peaks at around 2θ = 23.4° and 26.6°, respectively, indicating their π-π stacking order. It is obvious that the addition of AT nano-rods results in the increased intensity of the diffraction peak at 2θ = 26.6° attributed to the (004) plane of AT, which is also located at around 2θ = 26°. The pore size distributions and surface areas of the resulted aerogels 199

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Fig. 2. (A) ATR-FTIR spectra of neat ODA and DMBZ-based PI and PI/AT hybrid aerogel monoliths with 10 wt% AT; (B) ATR-FTIR spectra of ODA based hybrid aerogels with various AT loadings in the 1700-1800 cm−1 region and curve-resolved FTIR spectra in the range of 1690–1740 cm−1 for (C) neat ODA based PI and (D) the ODA based hybrid aerogel with 10 wt% AT nanorods.

nanorods show a higher average surface area than the CNC–COOH as illustrated in Fig. 1(D). The BJH pore size distributions of neat aerogels and hybrid aerogels are shown in Fig. 4(B) and (D). The pore size distribution peaks are in the range of 30–40 nm, and the incorporated AT nanorods have little effect on the pore size distribution. The visual microstructures of the AT-reinforced PI hybrid aerogel monoliths were examined by SEM (Fig. 4(C) and (F)). Aerogels resemble porous and nanofiber-like morphologies tangle together with fiber diameters in the range of 15–40 nm, which are different from those of silica aerogels with the typical clusters or strings of particles. The AT nanorods, which diameters are about an order of magnitude larger than the fine structure

were analyzed by BET and BJH model. As shown in Fig. 4(A), surface areas range from 288 to 329 m2/g for ODA-based aerogels. For comparison sake, in Fig. 4(D), surface areas on the average for DMBZ-based aerogels are in the range of 394–424 m2/g, which are significantly larger (∼100 m2/g) than those made from ODA. The DMBZ-AT-7.5 aerogel has the largest surface area of 424 m2/g, which is around 30 m2/g higher than that of the DMBZ-AT-0. Nguyen et al. [19] have illustrated that incorporating carboxylic acid-functionalized CNC (CNC–COOH) into PI matrix resulted in the significantly decreased surface areas of hybrid aerogel monoliths due to the low surface area of the CNC–COOH (84 m2/g) and their aggregation. Comparatively, AT

Fig. 3. (A) 13C NMR spectra of neat PI and PI/AT aerogel hybrid monoliths with 5 wt% and 10 wt% AT nanorods; (B) X-ray diffraction patterns of AT nanorods, neat PI and PI/AT hybrid aerogels. 200

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Fig. 4. Typical N2 adsorption-desorption isothermals (A, D) and BJH pore size distributions (B, E) of PI/AT aerogel hybrid monoliths with various AT loadings; SEM micrographs of the fracture surface of ODA-AT-10 (C) and DMBZ-AT-10 aerogel hybrid monoliths (F).

density of AT (∼2.1 g/cm3) in comparison to the neat PI aerogels. Correspondingly, a decrease of shrinkage contributes to the increase in porosity of the aerogels with the AT content (Fig. 5(C)). Apparently, AT nanorods have a larger effect on improving porosities of ODA-based aerogels than those of the DMBZ sets. Nguyen et al. [19] reported a similar result when investigating CNC–COOH reinforced polyimide aerogels, and they explained it that high modulus nanofillers have a more obvious effect on increasing the aerogel porosity with a flexible polymer backbone than that of the rigid one. Thus, for AT nanorods in the present work, it would be expected to not have an obvious effect on the porosity of the DMBZ-based aerogel. Stress-strain curves and the detailed compression modulus from compression testing of the prepared aerogel monoliths are shown in Fig. 6. These comprehensive curves for both sets exhibit a typical deformation behavior of open honeycomb-like foams: aerogels can tolerate compression strains in a linear elastic region of an about 0–1% strain (as shown in the inset of Fig. 6(A) and (B)) followed by a plastic yielding plateau. Finally, aerogels undergoes a densification region exhibiting a yield up to about 90% strain, which are much better than the brittle silica aerogels. The AT nanorods braced architecture in the hybrid aerogels results in unique mechanical properties; aerogels exhibit obvious increase in yield stress and Yang's modulus with the increased AT content. For example, at 80% strain, the compressive stress increases from 11.5 MPa for DMBZ-AT-0 to 20.1 MPa for DMBZ-AT-10 containing 10 wt% AT, and a similar increment has been found in ODAAT-10 aerogel at the same strain, illustrating that these rigid AT

of the aerogel, stand out randomly (marked as red arrows) throughout the aerogel while the PI nano-fibers are intimately wrapped around these AT nanorods. Seen from the strong interfacial adhesion between the aerogel matrix and AT nanorods from the FTIR analysis, the two phases combine tightly, which likely favor efficient stress transfer from the polymer matrix to the AT nanorods and thus enhance mechanical properties of the nanocomposites. Nevertheless, the strong Van der Waals attraction associated to the large aspect ratio of AT nanorods promotes the aggregation at higher concentrations. The effects of AT nanorods on the liner shrinkage, density and porosity of the resulted aerogels in the supercritical drying process are shown in Fig. 5. As the concentration of AT nano-rods increases, the liner shrinkage decreases by 14% for ODA-AT-10 compared to the ODAAT-0, and by 13.5% for DMBZ-AT-7.5 compared to the DMBZ-AT-0. While, continually increasing the AT loading to 10 wt% in DMBZ-based hybrid aerogel, the shrinkage increases slightly, which can be attributed to inhomogeneous dispersion of AT and the formation of AT bundles in the aerogels as shown in the SEM micrographs. The DMBZ-based aerogels shrink less than those ODA-based during the supercritical drying process, which can be accounted for the higher rigidity of biphenyl structure contained in DMBZ backbones. Simultaneously, as shown in Fig. 5(B), the densities of these two sets of hybrid aerogels exhibit different trends. ODA-based aerogels show continually decreased densities with increasing AT content, while, DMBZ-based aerogels exhibit a minimum density at the AT loading of 5 wt%, and more AT loadings result in the increase of densities due to the higher

Fig. 5. Graphs of (A) shrinkage, (B) density and (C) porosity of prepared ODA- and DMBZ-based PI/AT aerogel hybrid monoliths with various AT loadings. 201

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Fig. 6. (A, B) Graphs of stress-strain curves and (B, D) detailed Young's and special Young's modulus of PI/AT aerogel hybrid monoliths with various AT loadings.

behaviors of these two sets of aerogels at 550 °C. At 700 °C, ODA-based aerogels show reduced weight losses with the AT concentration, while, no strong trend with composition can be observed for the DMBZ sets, indicating that AT nanorods provide thermal resistance for flexible polymer aerogels at the higher temperature region. As reported previously, most polyimide aerogels still underwent some liner shrinkage (40–50%) occurring at temperatures from 150 to 200 °C at air atmosphere [19]. The application of aerogels at high temperature is always limited by their dimensional instability, usually resulting in a decrease in thermal insulation and an increase in dielectric constant. The shrinkage changes of the prepared aerogels in the isothermal aging as a function of AT loading were measured at temperature range of 100–200 °C, as shown in Fig. 8. Obviously, the shrinkage is reduced with increasing the AT concentration in the aging, emphasizing the effect of the AT reinforcement in retaining the structural integrity of the aerogels. For example, the DMBZ-AT-0 aerogel monolith shows a shrinkage of 26% when thermally treated at 150 °C for 24 h due to the collapse of micropores, while, the shrinkage decreases to 16% as incorporating 10 wt% AT. Another interesting phenomenon is that the introduction of AT nanorods results in an obvious shrinkage reduction for ODA sets than the DMBZ sets, indicating the more obvious reinforcing effects on the ODA-based aerogels. The shrinkage of aerogels inevitably leads to an increase in density. For example, in Fig. 8(F), the density of DMBZ-AT-0 with no AT filler increases from 0.12 to 0.52 g/cm3 when aged at 200 °C for 24 h, while, the density of DMBZ-AT-10 containing 10 wt% AT increases to 0.41 g/ cm3. The original ODA-based aerogels and aged samples at 100 °C show higher densities than the DMBZ-based aerogels with the same AT concentration. However, the reverse tendency takes place for the aerogels aged at 150 and 200 °C, which can be due to the greater shrinkage of DMBZ-based aerogels at higher temperatures. Thus, it can be concluded that AT nanorods have stronger reinforcing effect on the polymer aerogels with flexible backbones. Nevertheless, it should be noticed that

nanorods produce a significant improvement in the mechanical properties of aerogels. In Fig. 6(C) and (D), Yang's modulus (obtained from the slope of the linear elastic region in s-s curves) and specific moduli (defined as a ratio of compressive modulus to density) of the aerogels are both shown to increase with AT content. The Yang's modulus and specific moduli values of DMBZ-AT-10 reach 45 MPa and 373 J/g, respectively, the highest values among all prepared aerogels in the present study. Previous researches [45–47] on aerogels illustrated the strong proportional dependency of modulus with density for similar backbone families, i.e., lower density always meaning lower modulus. Young's modulus (E) of some organic aerogels always obeys the powerlaw scaling as a function of the bulk density (ρ) by E∼ρn (n = 2–4) [48]. As opposed to previous works, we show a decreasing density with increasing AT concentration, accompanied with an increase in modulus, especially for the ODA-based aerogels. This is of great significance for the development of aerogels with both high mechanical properties and ultra-light weights. To have practical potential, particularly for applications in hightemperature circumstance, polymer aerogels must possess excellent thermal stabilities. Fig. 7 shows TGA and resultant characteristic thermal degradation data of the prepared aerogels as well as the AT nanorods. A small weight loss at the beginning of the AT is assigned to the loss of absorbed water in the air. In the high temperature region even above 950 °C, AT shows little weight loss, indicating itself possesses a more desirable thermal stability. Hence, it can be expected that the addition of AT nanorods is beneficial for enhancing the thermal stability of the hybrid aerogels. As seen, there is no weight loss occurring with the neat or their hybrid aerogels below 500 °C. Weight loss for both ODA- and DMBZ-based aerogels at 550 and 700 °C is shown in Fig. 7(B). At 550 °C, all DMBZ-based aerogels show about 10% weight loss, higher than those of the ODA containing aerogels (∼2.5%), attributing to the degradation of the pendent methyl groups in DMBZ. Meanwhile, AT nanorods appear to have little effect on the thermal 202

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Fig. 7. (A) TGA curves in air and (B) percent weight loss of ODA- and DMBZ-based PI/AT hybrid aerogels.

and DMBZ-AT-10 aerogels, respectively. This result illustrates that the ODA-AT-10 aerogel has a much better thermal insulation property, which can be ascribed to the lower shrinkage percent and better dimensional stability in the high temperature condition supported by the AT nanorods skeleton within the ODA-AT-10 aerogel. As shown in Fig. 9(B-D), the pristine DMBZ-based aerogel shrinks dramatically in volume, however, the size of the ODA-AT-10 aerogel is fairly stable with a shrinkage of ca. 10%, which further highlights the impact of AT nanorods on the shrinking behavior of the hybrid aerogels. The lower shrinkage is beneficial for maintaining the high porosity and lower thermal conductivity of resultant aerogels.

even combining rigid AT nanofillers into the polyimide matrix as the supporting skeleton via the strong interfacial interaction, the prepared aerogels still exhibit a high degree of shrinkage at the relatively high temperature condition. The microstructure evolution of polyimide aerogels in the thermal aging will be probed in future studies for better understanding the phenomenon and further reducing the shrinkage. As we know, aerogels are one kind of well-known thermal insulation materials with low thermal conductivities due to their highly porous structure. An investigation about the heat-insulating property of the AT reinforced PI aerogels was performed by placing the aerogels on a hot plate with the temperature of approximately 300 °C. The dynamic temperature distribution on the aerogels was monitored by infrared camera. All aerogel monoliths were of approximately the same size (cylindrical samples with a height of 2 cm). As shown in Fig. 9(A), a gradient distribution of temperature from the hot plate through the sample is observed. For DMBZ-based aerogel with no filler, the temperature at the top surface gradually increases from 37.7 °C to 58.8 °C as the heating time is prolonged from 10 s to 300 s. It can be seen that, during the running process, the top surface temperature for the DMBZAT-10 hybrid aerogel drops by 10 °C in comparison with the pristine DMBZ-based aerogel. Comparatively speaking, the ODA-AT-10 aerogel exhibits the lowest top surface temperature with a detailed value of 37.6 °C, which is 21.2 and 10.5 °C lower than that of the DMBZ-AT-0

4. Conclusions The present study established a facile method for the preparation of polyimide aerogels supported by the AT nanorods. Two different diamines, either the flexible ODA or the rigid DMBZ, were chosen to create the backbone of polyimide hybrid aerogels. The structure characterization, compressive property, dimensional stability and thermal insulation behavior of the aerogels were investigated in detail. Because of the hydrogen-bonding interaction between the AT and polyimide matrix, and the rigid AT supporting skeleton, the compressive stress and modulus were greatly enhanced as the AT content increased for both

Fig. 8. Shrinkage (A–C) and density (D–F) of ODA-based and DMBZ-based PI/AT aerogels with different contents of AT nanorods after different aging condition. 203

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Fig. 9. (A) Infrared images of neat ODA-based and DMBZ-based PI aerogels and PI/AT-10 aerogel hybrid monoliths on a 280 °C heating stage for 10, 150 and 300 s; (B–D) optical images of the original aerogel monoliths and the corresponding samples treated at 280 °C for 300 s: (B) DMBZ-AT-0, (C) DMBZ-AT-10 and (D) ODA-AT-10.

ODA- and DMBZ-based aerogels. Another effect of incorporating AT nanorods into the aerogels is their reinforcement on the dimensional stability of aerogels on the harsh temperatures. The resultant hybrid aerogels, especially those synthesized by the flexible ODA, shrunk much less even treated at 200 °C for 24 h and resulted in less dense aerogels, endowing them with an excellent thermal insulation property. Accordingly, such multiscale reinforced polyimide hybrid aerogels showed potential applications as the thermal insulation devices or high temperature filtration media in some rigorous conditions.

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Acknowledgements [13]

This work was supported by Special Fund of China Postdoctoral Science Foundation (2018T110323), China Postdoctoral Science Foundation (2017M611418), National Natural Science Foundation of China (No. 21774019) and the Program of Shanghai Academic Research Leader (18XD1400100).

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Appendix A. Supplementary data [16]

Supplementary data to this article can be found online at https:// doi.org/10.1016/j.polymer.2019.05.007.

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