Effect of Graphene Oxide-Modified CaAl-Layered Double Hydroxides on the Carbon Dioxide Permeation Properties of Fluoroelastomers

Abstract

This work aimed to investigate the CO₂ gas barrier and mechanical properties of fluorine rubber nanocomposites filled with Ca/Al layered hydroxide (graphene oxide [GO]/LDH-Ca₂Al) modified by GO. GO/LDH-Ca₂Al nanocomposite fillers were prepared by depositing Ca/Al layered hydroxide (LDH-Ca₂Al) into the surface of alkalized GO (Al-GO). The prepared GO/LDH-Ca₂Al nanocomposite fillers and complexes were characterized by Fourier infrared spectroscopy (FTIR), X-ray diffraction (XRD), scanning electron microscopy (SEM), and transmission electron microscopy (TEM) for structural and micromorphological characterization. The results showed that GO/LDH-Ca₂Al was successfully prepared with strong interactions between Al-GO and LDH, and the compatibility of GO/LDH-Ca₂Al nanocomposite fillers with the polymer was significantly improved compared with that of LDH-Ca₂Al. Consequently, both the fracture strength (σ_b) and strain (ε_b) of GO/LDH-Ca₂Al nanocomplexes remarkably increased, and they exhibited excellent mechanical properties. Differential scanning calorimetry and thermogravimetric analysis were used to characterize the thermal stability of GO/LDH-Ca₂Al nanocomposite fillers, and GO/LDH-Ca₂Al nanocomposite fillers have better thermal stability than LDH-Ca₂Al. The reaction products (S-LDH-Ca₂Al and S-GO-Ca₂Al) of LDH-Ca₂Al and GO/LDH-Ca₂Al with CO₂ were characterized using XRD and TGA, respectively, and the results show that LDH-Ca₂Al reacts readily and chemically with CO₂, resulting in a lower diffusion coefficient of CO₂ in the LDH-Ca₂Al nanocomplexes than that of the GO/LDH-Ca₂Al nanocomplexes and leading to the destruction of the laminar structure of LDH-Ca₂Al, while GO/LDH-Ca₂Al has better CO₂ resistance stability. GO/LDH-Ca₂Al nanocomplexes exhibited a reduced content of hydroxyl groups with pro-CO₂ nature exposed on the surface of LDH-Ca₂Al, improving the interfacial interaction between the nanofillers and the rubber matrix and enhancing the dispersion of GO/LDH-Ca₂Al in the polymers. Moreover, CO₂ in the soluble GO/LDH-Ca₂Al nanocomposites was significantly reduced, while the diffusion properties demonstrated weak temperature dependence on solubility. The mechanism of the CO₂ gas barrier of polymers filled with GO/LDH-Ca₂Al was proposed on the basis of the Arrhenius equation.

1. Introduction

Climate change is one of the most serious problems facing mankind, and anthropogenic air pollutant emissions and climate change have a direct relationship in which carbon dioxide (CO₂) is the main air pollutant [1]. In the field of oilfield development, CO₂ is an attractive repellant for enhanced oil recovery, but some problems, such as CO₂ leakage during CO₂ transportation, are encountered; hence, a sealing material with excellent barrier performance needs to be developed [2].

Rubber is a highly elastic and amorphous polymer with good toughness, elasticity, and elongation, and they are also known as elastomers; for this class of material, a very small external force can produce a large deformation, which can be restored after removing the external force [3,4]. Rubber materials are widely used in tire liners, chemical protection products, medical packaging, automotive tanks and natural gas storage tanks, and natural long-distance pipeline valve sealing materials, and their gas barrier properties are very important [5]. The gas barrier properties of rubber materials are very important. Given the high free volume fraction between rubber molecular chains, gas molecules can easily diffuse from one end of the rubber material to the other, and most diene-based rubbers such as natural rubber (NR), ethylene–propylene rubber (EPDM), and styrene–butadiene rubber (SBR) have high permeability to gases [6].

The gas barrier properties can be effectively improved by filling a certain amount of nanoparticles into the rubber to form a filling network, thus creating distorted paths and inhibiting gas molecules from penetrating the rubber matrix. More importantly, the strong interface between the nanoparticles and the rubber molecules is significant in limiting chain migration and further reducing the free volume between the nanoparticles and the rubber molecules. Among a range of nanoparticles, layered particles such as layered double hydroxides (LDH) [7,8,9] and graphene (GE) or graphene oxide (GO) [10,11] have higher aspect ratios than spherical and fibrous particles, making them more conducive to enhancing the gas barrier properties of rubber composites [12]. In particular, LDH is considered to be a promising CO₂ adsorbent because of its controllable layer spacing and pro-CO₂ properties [13]. The chemical formula of LDH can be represented by the general expression ${\mathrm{M}}_{1-x}^{2+}{\mathrm{M}}_x^{3+}{\left(\mathrm{O}\mathrm{H}\right)}_2{\left({\mathrm{A}}^{\mathrm{n}-}\right)}_{x/n}·\mathrm{y}{\mathrm{H}}_2\mathrm{O}$ [14,15]. LDH consists of a hydromagnesite-like layer in which a small portion of octahedrally coordinated divalent metal cations is replaced by trivalent metal cations, resulting in a positively charged host layer [16]. Exchangeable inorganic or organic anions are accommodated in the interlayer channels to compensate for the positive charge. In addition, the hydroxyl groups of the host layer are connected to anions or water molecules via hydrogen bonding. The X value is equal to the molar ratio of M³⁺/(M²⁺+M³⁺), and it ranges from 0.17 to 0.33 [16]. Given the controllable M²⁺/M³⁺ molar ratio, the tunability of the metal cations, and the exchangeable charge compensating anions [17], the layer spacing of the LDH can be optimized, which can aid in the design of the lamellar packing that facilitates CO₂ intercalation, further forming a more complex pathway for CO₂ diffusion.

However, the thermal stability of LDH is poor, and it starts to undergo layer structure destruction at approximately 250 °C [18]. Most importantly, LDH is a polar inorganic material with limited compatibility with organic macromolecules such as natural rubber (NR), fluorine rubber (FKM), and nitrile rubber (NBR), leading to agglomeration in the polymer and the formation of weak phase interfaces [8,9]. Therefore, the dispersion and thermal stability of LDH in polymers can be improved by forming a compound with LDH along with new components. GO has high electronic conductivity, good mechanical strength, excellent thermal stability, large specific area, and abundant oxygen-containing groups on its surface, such as epoxide, hydroxyl, and carboxyl groups, which can form strong interactions with polymer molecules via hydrogen or ionic bonding [13,14,15]. It also has good descriptive properties with polymers. Moreover, GO is an almost monolayer layered structure with a large surface area [19]. GO is an ideal carrier for LDH because of its negatively charged layer surface and oxygen-containing groups, which can form hydrogen bonding and strong electrostatic interactions with the positively charged LDH on the surface of the host layer [20,21,22]. Yang et al. prepared FeNi-LDH/GO hybrid nanosheets by alternately stacking GO layers and FeNi double hydroxide ion layers [20].

To meet the application requirements for rubber of the tire, aerospace, and military fields, nanofillers such as graphene oxide, carbon nanotubes, carbon black, and montmorillonite have been widely used to improve the mechanical properties of rubber. The effects of various modifying additives and fillers on the physicochemical and mechanical properties of polymer composites are determined by many factors as follows: (1) filler–polymer interfacial interaction [23,24,25], (2) content of fillers [26], and (3) cross-linking density of polymers [27]. Among these, filler–polymer interfacial interaction exerts an extremely significant effect on the mechanical properties of polymer composites.

In this study, GO/LDH-Ca₂Al nanolamellar fillers with better thermal stability and better compatibility with polymers were prepared via a simple synthesis method (Scheme 1). The layer spacing of this nanofiller increased by approximately 0.04 nm compared with that of pristine LDH-Ca₂Al, thus increasing the chances of interlayer insertion of CO₂ gas molecules. At the same time, the resistance to CO₂ stability was significantly improved, and a stable layer structure can be maintained in a CO₂ environment. The GO/LDH-Ca₂Al nanofiller reduced the content of exposed hydroxyl groups of a pro-CO₂ nature on the surface of LDH-Ca₂Al, increased the interfacial interactions between LDH-Ca₂Al and the rubber matrix, and improved the dispersion of LDH-Ca₂Al in the polymer. Consequently, the CO₂ solubility of the nanocomposites was significantly reduced.

2. Experimental

2.1. Materials

FKM 246, a terpolymer of vinylidene fluoride (VF2); hexafluoropropylene (HFP); and tetrafluoroethylene (TFE), with 68.5% fluorine content, density of 1.86 g/cm³, Mooney viscosity, ML 1 + 10 at 121 °C = 25, and solubility parameter, δ = 19.7 MPa^1/2 were supplied by Shanghai Huayi 3F New Materials Co., Ltd. (Shanghai, China). 2,5-Dimethyl-2,5-di(tert-butylperoxy)hexane (Macklin), triallyl isocyanurate (TAIC, ≥98%, Aladdin), and calcium chloride dihydrate (CaCl₂·2H₂O, ≥98%) were purchased from Beijing Yongchang Haoran Bio-technology Co., Ltd. (Beijing, China). Aluminum chloride hexahydrate (AlCl₃·6H₂O, ≥98%) was obtained from Shanghai Taitan Technology Co., Ltd. (Shanghai, China). Graphene oxide (GO, >99%), sodium hydroxide (NaOH, ≥98%), and deionized water were purchased from Shanghai Boer Chemical Reagent Co. (Shanghai, China). All reagents were of analytical grade, and the purity of the CO₂ test gas used for the experiments was 99.9% (Beijing Chengweixin Industrial Gas Sales Center, Beijing, China).

2.2. Preparation of Layered LDH-Ca₂Al and GO/LDH-Ca₂Al

LDH-Ca₂Al was prepared via co-precipitation [28,29,30]. Approximately 15 g CaCl₂·2H₂O and 7.606 g AlCl₃·6H₂O were dissolved in 300 mL of deionized water to obtain a colorless and transparent Ca₂Al solution. Approximately 7.36 g NaOH was dissolved in 30 mL of deionized water to obtain a colorless and transparent NaOH solution, and the NaOH solution was added to the vigorously stirred (800 rpm) Ca₂Al solution at room temperature. The pH of the suspension was maintained at approximately 11, and stirring was maintained for 1 h. The sample was then filtered, washed with deionized water until the pH was approximately 10, and finally vacuum-dried at 60 °C to obtain the white powder LDH-Ca₂Al.

Approximately 0.5 g GO was added to 1050 mL of deionized water to obtain the GO suspension [31]. Then, an appropriate amount of NaOH was added to adjust the pH of GO suspension at approximately 11 (Al-GO), and 8.15 g LDH-Ca₂Al was added to the solution. The sample was sonicated for 10 min and then placed in a vacuum oven at 60 °C for 24 h to obtain GO/LDH-Ca₂Al suspension. Finally, it was filtered, washed with deionized water, and freeze-dried.

2.3. CO₂ Resistance Stability Testing of LDH-Ca₂Al and GO/LDH-Ca₂Al

S-LDH-Ca₂Al and S-GO/LDH-Ca₂Al, which were intercalated and reacted by CO₂, were obtained by simultaneously placing 1.5 g of LDH-Ca₂Al and GO/LDH-Ca₂Al in an autoclave of CO₂ gas as the ambient medium at 80 °C and 4.3 MPa, and the sample was taken out after 1 day.

2.4. Preparation of FKM/GO/LDH-Ca₂Al and FKM/LDH-Ca₂Al Composites

The raw rubber was sheared on the rolls for 2 min at room temperature by using an open double-roller mill (friction ratio: 1:1.4). Then, different proportions of GO/LDH-Ca₂Al were added for 3–4 min. Finally, 5 phr LUPEROX 101XL-50 and 5 phr TAIC were added to the sample, and the mixture was sheared and agitated for 3–4 min to obtain the unvulcanized FKM/GO/Ca₂Al composite. A rotorless vulcanometer was used to determine the optimal vulcanization time (tc90) of the unvulcanized FKM/GO/LDH-Ca₂Al composites at 160 °C, and was vulcanized to the optimal vulcanization time (tc90) by using a plate vulcanizer at 160 °C and 10 MPa. FKMs unfilled and filled with the same amount of LDH-Ca₂Al were prepared using the same method as the blank control group.

Table 1. FKM nanocomposites formulation.

Samples	FKM 246	LUPEROX 101XL-50	TAIC	LDH-Ca2Al a	GO/LDH-Ca2Al a
FKM	100	5	5
FKM/LDH-Ca2Al-5	100	5	5	5	5
FKM/LDH-Ca2Al-8	100	5	5	8
FKM/LDH-Ca2Al-10	100	5	5	10
FKM/GO/LDH-Ca2Al-5	100	5	5		5
FKM/GO/LDH-Ca2Al-8	100	5	5		8
FKM/GO/LDH-Ca2Al-10	100	5	5		10

^a The densities of the materials FKM 246, LDH-Ca₂Al, and GO/LDH-Ca₂Al are 1.86, 1.65, and 2.01 g/cm³, respectively.

provides the formulation of FKM/GO/LDH-Ca₂Al composites. Different portions of GO/LDH-Ca₂Al-filled fluoroelastomers were denoted as FKM/GO/LDH-Ca₂Al-x, where x refers to the number of portions of nanofillers in 100 phr of fluoroelastomers.

2.5. Characterization

Fourier transform infrared spectroscopy (FT-IR) measurements were carried out on a TENSOR II spectrometer (Bruker, MA, USA) at the wave number range of 4000–400 cm⁻¹ with a resolution of 4 cm⁻¹ and a 64-scan signal via the KBr pellet technique. An X-ray diffractometer (D8 Focus, Bruker) was used to characterize LDH-Ca₂Al and GO/LDH-Ca₂Al before and after CO₂ immersion and GO, and the measurements were carried out at a wavelength of λ = 0.154056 nm with a step size of 0.01° under Cu/Kα radiation. LDH-Ca₂Al and GO/LDH-Ca₂Al before and after CO₂ immersion, FKM/LDH-Ca₂Al, and FKM/GO/LDH-Ca₂Al were subjected to morphological observation by using a SU8010 cold field emission scanning electron microscope (SU8010, Hitachi, Japan). The CO₂ gas permeability of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al composites were measured via differential pressure by using a homemade gas permeability tester according to ISO 2782-1:2022 [32]. The samples to be tested were cut into dumbbell-shaped specimens. The tensile test was carried out by using a universal testing machine at room temperature with a tensile rate of 500 mm/min, following the GB/T 528–2009 standard [33] and three parallel specimens were used for each group of materials. According to the ISO 7619-1:2010 standard [34], the Shore A hardness of the composites was tested using an LX A-type hardness tester. The samples were subjected to DMA testing by using a DMA (NETZSCH 242 C) in tensile mode at the frequency of 10 Hz. Each sample was scanned in the range of −50 °C to 60 °C at a heating rate of 3 °C/min⁻¹. The samples were subjected to thermogravimetric analysis (TGA) by using a Shimadzu DTG-60 thermogravimetric analyzer, and the weight of the test samples ranged from 8 mg to 10 mg. The samples were heated from 25 °C to 800 °C at a heating rate of 10 °C/min under a nitrogen atmosphere purge of 100 mL/min. A differential scanning calorimeter (DSC, NETZSCH-5) was used to characterize the LDH-Ca₂Al and GO/LDH-Ca₂Al thermal transition. The weight of the test samples ranged from 6 mg to 10 mg and was heated from 25 °C to 250 °C at a heating rate of 10 °C/min under a nitrogen atmosphere of 100 mL/min.

The crosslink density of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al was measured using the equilibrium dissolution method [27,35] according to the standard ASTM D6814-02 [36]. Approximately 0.2 g of the sample was placed in a sealed container containing 25 mL of acetone and held at 25 °C for 3 days, during which the acetone solvent was changed every 24 h. The mass of the swelling equilibrium was measured as $m_S$ and dried in a blower oven at 70 °C for 16 h. Then, the weight of the dried rubber $m_d$ was measured, and the density ${\rho }_d$ was calculated. The crosslink density $v_e$ was calculated as follows:

$$V_r=(m_d/{\rho }_d)/\left[m_d/{\rho }_d+\left(m_S-m_d\right)/{\rho }_S\right]$$

(1)

$$v_e=\frac{-\left[ln\left(1-V_r\right)+V_r+{\chi }_1{V_r}^2\right]}{\left[V_1({V_r}^{1/3}-V_r)/2\right]}$$

(2)

where $V_r$ and ${\rho }_S$ are the volume fraction of rubber and the density of the solvent (the density of acetone is 0.79 g/cm³), respectively, $V_1$ is the molar volume of acetone solvent (73.53 cm³/mol), and ${\chi }_1$ is the interaction parameter of the polymer with the solvent, which is 0.358 in this case. According to the Bristow–Watson equation, the following expression can be obtained [37]:

$${\chi }_1={\beta }_1+\left(\frac{V_1}{RT}\right){({\delta }_s-{\delta }_p)}^2$$

(3)

where ${\beta }_1$ is the lattice constant (typically 0.34), $R$ is the gas constant, $T$ is the absolute temperature, and ${\delta }_p$ and ${\delta }_s$ are the solubility parameters of FKM and solvent acetone, respectively.

3. Results and Discussion

3.1. Fourier Transform Infrared (FTIR) Spectroscopy

The IR spectra of GO, Al-GO, LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al are shown in Figure 1. The broad and strong absorption peaks of the infrared spectra of GO near 3381 cm⁻¹ represent the O–H stretching vibrations in carboxyl and hydroxyl groups, and the peaks at 1223 and 1733 cm⁻¹ represent the out-of-plane bending vibrations of O–H in COOH and its C=O stretching vibrations [38,39]. Alkoxy (C–OH) stretching vibrations and hydroxyl deformation vibrations occurred at 1052 and 1390 cm⁻¹, respectively [40], and while that at 1620 cm⁻¹ can be attributed to the conjugate bond of aromatic C=C-C=C [39]. The peaks of the C=O stretching vibration of COOH and O–H out-of-plane bending vibration of COOH on Al-GO disappeared, and the symmetric and antisymmetric peaks of COO⁻ appeared near 1636 and 1458 cm⁻¹ [41]. The O–H stretching vibration peaks of Al–OH and Ca–OH of LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al appeared at 3634 and 3474 cm⁻¹, respectively, the bending vibration of the interlayer H₂O molecule was observed at approximately 1621 cm⁻¹, and the absorption peaks at 787, 535, and 434 cm⁻¹, which are below 800 cm⁻¹, represent M-O vibrational peaks (M is Ca or Al) [18,41,42]. In the IR spectra of GO/LDH-Ca₂Al, Al-GO peaks appeared at 3741 and 1637 cm⁻¹, but these peaks appear to be blue-shifted, possibly because of the positively charged nature on the LDH-Ca₂Al layer and the presence of the electron-absorbing induced effect. The peaks at 1733 and 1223 cm⁻¹, which represent COOH, disappeared from the infrared spectra of GO/LDH-Ca₂Al and Al-GO, while symmetric and antisymmetric telescopic vibrational peaks near 1458 and 1636 cm⁻¹ for COO⁻ appeared [41]. The symmetric and antisymmetric telescopic vibrational peaks of COO⁻ in the GO/LDH-Ca₂Al spectra also shifted to the high-frequency region, indicating the presence of a strong interaction between GO and LDH-Ca₂Al.

The CO₃²⁻ peaks at 1428 and 875 cm⁻¹ represent small amounts of CaCO₃ impurities in GO/LDH-Ca₂Al with LDH-Ca₂Al [43]. Both LDH-Ca₂Al and GO/LDH-Ca₂Al showed weak interlayer CO₃²⁻ absorption peaks at 2974 and 1486 cm⁻¹, but the interlayer CO₃²⁻ absorption peaks of LDH-Ca₂Al of S-LDH-Ca₂Al and S-GO/LDH-Ca₂Al were significantly enhanced, suggesting that CO₂ can be intercalated in the interlayer of LDH-Ca₂Al mainly in the form of CO₃²⁻ [18].

3.2. XRD Analysis of LDH-Ca₂Al and GO/LDH-Ca₂Al

The XRD results of GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al are shown in Figure 2, where the appearance of CaCO₃ impurity (PDF #97-001-6710 and PDF #97-018-1959) peaks proves the accuracy of the FITR results (Figure 1). The X-ray diffraction pattern of GO/LDH-Ca₂Al (PDF#04-010-4677) does not show GO diffraction peaks, suggesting that the periodic stacking arrangement was not formed by the composite of GO and LDH-Ca₂Al [13]. The GO/LDH-Ca₂Al shows the typical Bragg reflection of LDH-Ca₂Al as the pristine LDH-Ca₂Al, but the diffraction peaks of GO/LDH-Ca₂Al at the (003) crystal plane shifted to a lower angle; according to Bragg equation ($2\mathrm{d}\mathrm{s}\mathrm{i}\mathrm{n}\left(2\mathsf{\theta }\right)=\mathrm{n}\mathsf{\lambda }$) the calculated LDH-Ca₂Al and GO/LDH-Ca₂Al basal spacings were approximately 2.587 and 2.623 nm, respectively, and the thickness of the LDH-like hydromagnesite layer was 0.48 nm [41]. The LDH-Ca₂Al and GO/LDH-Ca₂Al interlayer spacings were 2.107 and 2.143 nm, respectively. Consequently, Al-GO formed a complex with LDH-Ca₂Al, resulting in an increase in interlayer spacing CO₂ by approximately 0.036 nm, as shown in Figure 2b. This phenomenon can be attributed to the strong interactions between the Al-GO and the LDH-Ca₂Al layers, which weakened the electrostatic interactions of the anionic in the interlayer channel with the cationic on the layer, consistent with the FTIR results. Given the increase in GO/LDH-Ca₂Al layer spacing, the specific surface area also increased, which provided more active sites for CO₂ molecular adsorption and increased the possibility of CO₂ gas adsorption and intercalation.

3.3. Structural Characteristics of LDH-Ca₂Al and GO/LDH-Ca₂Al

The SEM of images of LDH-Ca₂Al and GO/LDH-Ca₂Al are shown in Figure 3. LDH-Ca₂Al and GO/LDH-Ca₂Al have obvious lamellar structures, and the diameter of LDH-Ca₂Al flakes is in the range of 0.3–0.5 µm with a thickness of approximately 0.03 µm. The XRD of LDH-Ca₂Al showed a layer spacing of approximately 2.107 nm (Figure 2), indicating that LDH-Ca₂Al is composed of multilayered LDH-Ca₂Al lamellae stacked along the (003) direction. In GO/LDH-Ca₂Al, LDH-Ca₂Al lamellae are much smaller than Al-GO (S1), which is consistent with the TEM results of GO/LDH-Ca₂Al in Figure 4e. A closer look reveals that a large amount of LDH-Ca₂Al is loaded onto the surface of Al-GO because of the existence of a strong interaction force between LDH-Ca₂Al and Al-GO, which is consistent with the FITR results of GO/LDH-Ca₂Al (Figure 1).

The TEM of GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al is shown in Figure S1 and Figure 4, and they show a lamellar structure, which is consistent with the SEM results. The HR-TEM of LDH-Ca₂Al in Figure 4d shows the appearance of obvious (110) and (104) lattice stripes, in which the spacing of the stripes are 0.286 and 0.378 nm, consistent with the results of (110) and (104) diffraction peaks in the XRD diffraction pattern. The SAED plot along the crystal axis (16, −7, 1) is shown in the lower right corner of Figure 4d, consistent with PDF#04-010-4677. A large amount of two-dimensional LDH-Ca₂Al grows on GO/LDH-Ca₂Al, and the HR-TEM of GO/LDH-Ca₂Al clearly demonstrates that the two-dimensional GO in GO/LDH-Ca₂Al grows LDH-Ca₂Al (Figure 4e); a (110) lattice face of LDH-Ca₂Al was observed in the GO/LDH-Ca₂Al lattice streak, and the SAED image in this region shows LDH-Ca₂Al diffraction spots along the (17,16,2) directions [44], indicating that LDH-Ca₂Al is deposited on the GO surface and stacked along the (003) direction on the GO surface.

3.4. Thermal Stability of GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al

The DSC and TGA curves of GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al are shown in Figure 5a–c, respectively. Figure 5a shows a physisorbed water removal near 125.4, 136.4, and 117.5 °C for GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al, respectively. Meanwhile, GO shows the oxidation of oxygen-containing functional groups near 198 °C, indicating an exothermic reaction [45]. The removal of interlayer H₂O molecules from LDH-Ca₂Al and GO/LDH-Ca₂Al occurred near 192.6 and 200.3 °C, respectively, suggesting that GO/LDH-Ca₂Al has better thermal stability than LDH-Ca₂Al. Meanwhile, GO/LDH-Ca₂Al did not show the exothermic reaction of GO near 198 °C, indicating that the thermal stability of GO in GO/LDH-Ca₂Al was significantly improved due to the presence of strong interaction between Al-GO and LDH-Ca₂Al.

The thermal stability of GO, LDH-Ca₂Al, and GO/LDH-Ca₂Al is also depicted in Figure S2 and Figure 5b. Similar to the DSC results (Figure 5a), GO underwent three key weight loss steps at <125 °C corresponding to the removal of physically adsorbed water, 130–350 °C corresponding to the removal of oxygen functional groups, and 360–800 °C corresponding to the carbon framework of the oxidative pyrolysis [39]. Based on Figure 5b, three distinct weight loss steps can be observed for LDH-Ca₂Al at 30–100 °C corresponding to the loss of physisorbed water, 100–170 °C corresponding to the process of interlayer water removal, 215–400 °C corresponding to the dehydroxylation of the layers and the removal of interlayer chloride ions, and 440–550 °C corresponding to the further elimination of the hydroxyl groups on the layers [18,41,46]. Compared with LDH-Ca₂Al, the thermal stability of GO/LDH-Ca₂Al was significantly improved between 30 and 200 °C. Although the temperature at which approximately 1.2% of physisorbed water was lost was approximately 7 °C lower than that of LDH-Ca₂Al, the temperature at which the interlayer water was removed increased by nearly 43 °C (ΔT₁). Interestingly, the temperature of releasing interlayer chloride ions decreased by about 20 °C (ΔT₂), which was caused by the enlarged interlayer spacing of LDH-Ca₂Al. At the same time, the thermal stability of the GO layer on GO/LDH-Ca₂Al was improved, in which the carbon framework decomposition temperature increased by approximately 29 °C, respectively. The residual mass of S-LDH-Ca₂Al with a value of approximately 59% at 800 °C (m₂), which is approximately 5% higher than that of the LDH-Ca₂Al, can be attributed to the decomposition of the carbonates that occurred after 450 °C based on the XRD patterns of S-LDH-Ca₂Al; this finding revealed the presence of carbonates (Figure 6a) [18].

3.5. Carbon Dioxide Resistance Stability of LDH-Ca₂Al and GO/LDH-Ca₂Al

The XRD of S-LDH-Ca₂Al and S-GO/LDH-Ca₂Al are shown in Figure 6a,b. LDH-Ca₂Al and GO/LDH-Ca₂Al were placed in a pure CO₂ environment at a pressure of 4 MPa and a temperature of 80 °C. After 1 day, the XRD diffraction peaks of the reaction products S-LDH-Ca₂Al and S-GO/LDH-Ca₂Al were very similar to those of Ca₄Al₂ (OH)₁₂ (OH₀.₄ (CO₃)_0.8 (H₂O)₄ compounds (PDF#97-026-3123). According to the Bragg equation ($2d\mathrm{s}\mathrm{i}\mathrm{n}\left(2\theta \right)=n\lambda$), the calculated layer spacings of LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al were 2.106, 2.143, 2.106, and 2.138 nm, respectively (Figure 2b and Figure 6b) [41], and the kinetic diameter of CO₂ was approximately 0.33 nm [47]. The IR spectra of S-LDH-Ca₂Al and S-GO/LDH-Ca₂Al also showed significant interlayer CO₃²⁻ absorption peaks (PDF #97-018-1959 and PDF #97-000-0150, Figure 1), suggesting that CO₂ was intercalated into the interlayer of LDH and appeared in the interlayer in the form of CO₃²⁻. The intensity of the CaCO₃ diffraction peaks increased in the S-LDH-Ca₂Al, and the micro-morphology of the S-LDH-Ca₂Al showed a large number of massive structures (Figure 6c), indicating that CO₂ reacted with LDH-Ca₂Al, which led to an increase in CO₃²⁻ content in S-LDH-Ca₂Al. By contrast, the intensity of the CaCO₃ diffraction peaks in the XRD pattern of S-GO/LDH-Ca₂Al decreased, no significant change was observed in the microscopic morphology (Figure 6d), and GO/LDH-Ca₂Al hardly reacted with CO₂, indicating that GO/LDH-Ca₂Al has a very good CO₂-resistant chemical stability, which is consistent with the TGA results (Figure 5).

Table 2. Contents of CaCO₃ in LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al ^a.

Sample	Content (CaCO3) (%)
LDH-Ca2Al	0.23
GO/LDH-Ca2Al	5.97
S-LDH-Ca2Al	59.94
S-GO/LDH-Ca2Al	2.37

^a XRD data of LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al were obtained via fine-tuning by using FullProf 2023.2 software.

also provides information on the change of CaCO₃ content in LDH-Ca₂Al, GO/LDH-Ca₂Al, S-LDH-Ca₂Al, and S-GO/LDH-Ca₂Al.

3.6. Interfacial Interactions of FKM Matrix with LDH-Ca₂Al and GO/LDH-Ca₂Al

The interaction between the nanofillers and the polymer matrix influences the dispersion state of the nanofillers in the polymer, thus remarkably affecting the gas barrier properties and mechanical properties of the nanocomplexes [48,49,50]. GO/LDH-Ca₂Al nanofillers had an excellent dispersion state in the FKM/GO/LDH-Ca₂Al nanocomposites, and the energy storage modulus of the FKM/GO/LDH-Ca₂Al complexes substantially increased because of the existence of a good phase interface between the GO/LDH-Ca₂Al nanofillers and the rubber matrix compared with the LDH-Ca₂Al nanofillers (Figure 7a,b). This interfacial interaction will be quantitatively measured next using two-phase modeling [51]. The loss tangent relationship between the filler and the unfilled polymer matrix satisfies the following equation [27,51,52]:

$$tan\delta =\frac{tan{\delta }_m}{1+1.5B\phi }$$

(4)

where $tan{\delta }_m$ and $tan\delta$ represent the tangent values of the filler and the unfilled polymer matrix, $\phi$ is the volume fraction of the filler, and $B$ is a phenomenological interaction parameter, which can represent the strength of the interface interaction between the filler and the matrix; the larger the value of $B$, the stronger the interface interaction between the filler and the polymer matrix [27]. The $tan{\delta }_m$, $tan\delta$, T_g, and $B$ values of FKM, FKM/LDH-Ca₂Al-8, and FKM/GO/LDH-Ca₂Al-8 are shown in

Table 3. Phenomenological interaction parameter ($B$), $tan{\delta }_m$,$tan\delta$, T_g, and $\phi$ of FKM, FKM/LDH-Ca₂Al-8, and FKM/GO/LDH-Ca₂Al-8.

Sample	$$\mathit{t}\mathit{a}\mathit{n}{\mathit{\delta }}_{\mathit{m}}$$	$$\mathit{t}\mathit{a}\mathit{n}\mathit{\delta }$$	φ (%)	$$\mathit{B}$$	Tg (°C)
FKM	1.480	-	-	-	3.3
FKM/LDH-Ca2Al-8	-	1.396	8.254	0.487	4.9
FKM/GO/LDH-Ca2Al-8	-	1.404	6.88	0.528	5.6

. The loss tangent $tan\delta$ and T_g of FKM/GO/LDH-Ca₂Al shifted toward high temperature, and the $B$ value of FKM/GO/LDH-Ca₂Al-8 is approximately 0.528, which is slightly higher than that of FKM/LDH-Ca₂Al-8 by 0.041. This finding indicates that the interfacial interactions between GO/LDH-Ca₂Al and the FKM matrix were enhanced compared with that of LDH-Ca₂Al, thus hindering the movement of the molecular chain segments. this finding explains that the crosslink density of FKM/GO/LDH-Ca₂Al is always higher than that of FKM/LDH-Ca₂Al for the fluoroelastomer composites filled with the same content of fillers, as shown in Figure 8. The SEM images of FKM/LDH-Ca₂Al-8 and FKM/GO/LDH-Ca₂Al-8 are shown in Figure 7a,b. The results indicate a good phase interface between GO/LDH-Ca₂Al and the rubber matrix, which is caused by the strong interfacial interaction between GO/LDH-Ca₂Al and the fluoroelastomer matrix than that of LDH-Ca₂Al.

3.7. Gas Barrier Properties of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al Composites

3.7.1. Effect of Filler Content on the Gas Barrier Properties of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al Composites

The gas barrier properties of rubber composites, which were measured in terms of the gas permeability coefficient ($Q=DS$), are related to the shape, content, and degree of dispersion of the filler [5,12,53], and lamellar fillers with large aspect ratios have good gas barrier properties [54]. LDH-Ca₂Al and GO/LDH-Ca₂Al are lamellar fillers with large aspect ratios (Figure 3a,b). The aspect ratios of LDH-Ca₂Al are in the range of 10–16, and the $Q$ and $D$ of the rubber filled with 8 phr LDH-Ca₂Al fillers were reduced by nearly 9% and 33%, respectively. The relationships of permeability coefficient ($Q$), diffusion coefficient ($D$), and solubility ($S$) with the contents of LDH-Ca₂Al and GO/LDH-Ca₂Al are illustrated in Figure 8a–c. The results show that the permeability coefficients and diffusion coefficients of the rubber composites decrease significantly with the increase in filler content, which is consistent with the change in the crosslinking density of the rubber composites (Figure 7e). This phenomenon can be explained using the free volume theory [55,56]. The nanofillers can be approximated as physical cross-linking points and the cross-linking density increases with the increase in nanofiller content. The increase in cross-linking density shortens the distance between the rubber chains, reduces the mobility of the rubber chains, and leads to a decrease in the free volume fraction of the permeation path of the gas molecules [6]. The free volume fraction of the gas molecules’ permeation path decreased.

The variation of solubility (S) with filler content is given in Figure 8a. The solubility of FKM/LDH-Ca₂Al increased significantly with the increase in LDH-Ca₂Al content, and the solubility of the rubber composite increased by 108% when filled with 10 phr of LDH-Ca₂Al. By contrast, the solubility of FKM/GO/LDH-Ca₂Al gradually decreased with the increase in GO/LDH-Ca₂Al content, and the solubility of the rubber composite decreased by about 19% when filled with 10 phr of GO/LDH-Ca₂Al. This finding was obtained because LDH-Ca₂Al is a nanofiller with high CO₂ absorption [48]. Moreover, the loading of LDH-Ca₂Al onto the GO surface reduces the content of exposed polar hydroxyl groups on the LDH-Ca₂Al layer, thus forming strong hydrogen bonding interactions with O in the CO₂ molecule and increasing the content of CO₂ adsorbed on the surface of the filler. In addition, compared with the LDH-Ca₂Al, the GO/LDH-Ca₂Al in the composites has good compatibility (Figure 7a,b), which consistently supports the interfacial interactions between GO/LDH-Ca₂Al and the rubber matrix (Table 3). Therefore, more constrained polymer regions are present near the interface between the rubber matrix and the filler [5,57], thus reducing the adsorption capacity of CO₂ molecules at the interface.

Figure 8b shows the schematic relationship between filler content and diffusion coefficient, in which the diffusion coefficient of the rubber composites decreased significantly with the increase in the filler content of LDH-Ca₂Al and GO/LDH-Ca₂Al, which is consistent with many previous findings [5,58,59,60,61]. This finding was obtained because the filler increases the gas molecule diffusion path and hinders the gas molecule movement [6], while LDH-Ca₂Al and GO/LDH-Ca₂Al are multilayer fillers (Figure 2, Figure 3 and Figure 4) with a layer spacing of more than 2.1 nm, which is significantly higher than the kinetic diameter of the CO₂ molecule (~0.33 nm). Hence, the CO₂ molecule can be interpolated into the layer spacing of the filler, thus significantly improving the gas molecule diffusion path. Accordingly, the diffusion coefficient of the rubber composite material was significantly reduced. Moreover, when the filler content of the filler reached 10 phr, the diffusion coefficient of the rubber composite material was reduced by nearly 57%. However, the diffusion coefficient of FKM/LDH-Ca₂Al is lower than that of FKM/GO/LDH-Ca₂Al at the same filler content, in which even the diffusion coefficient of FKM/LDH-Ca₂Al-10 is approximately 47% lower than that of FKM/GO/LDH-Ca₂Al-10. This finding was obtained because CO₂ entered the spacing of the LDH-Ca₂Al layers and reacted with the LDH-Ca₂Al to generate inorganic compounds containing CO₃²⁻ (Table 2), such as CaCO₃. Consequently, the diffusion activation energy of CO₂ gas remarkably decreased during diffusion in FKM/LDH-Ca₂Al (

Table 4. Activation energy of D and S of FKM, FKM/LDH-Ca₂Al-8, and FKM/GO/LDH-Ca₂Al-8.

Sample	Activation Energy of D(ΔE) a				Activation Energy of S(ΔH) a
	${\mathit{D}}_0$ b	$∆{\mathit{E}}_{\mathit{D}0}$	$∆{\mathit{E}}_{\mathit{D}\mathit{f}}$	$∆{\mathit{E}}_{\mathit{D}\mathit{c}}$	${\mathit{S}}_0$ c	$∆{\mathit{H}}_{\mathit{S}0}$	$∆{\mathit{H}}_{\mathit{s}\mathit{f}}$	$∆{\mathit{H}}_{\mathit{s}\mathit{c}}$
FKM	2.6 × 10−2	−52.8	-	-	1.8 × 10−11	44.5	-	-
FKM/LDH-Ca2Al-8	1.0 × 10−6	−25.2	22.4	5.2	2.2 × 10−6	44.5	−32.2	-
FKM/GO/LDH-Ca2Al-8	7.7 × 10−6	−30.4	22.4	-	4.7 × 10−7	44.5	−32.2	3.8

^a the unit of activation energy is KJ/mol. ^b $D_0$ and ^c $S_0$ are the dimensionless constants.

). For example, the diffusion activation energy of FKM/LDH-Ca₂Al-8 increased by approximately 27.6 J/mol. However, GO/LDH-Ca₂Al has better-stabilizing properties during CO₂ diffusion (Figure 2, Figure 3b and Figure 6), and the CaCO₃ content of S-GO/LDH-Ca₂Al did not change significantly compared with that of GO/LDH-Ca₂Al (Table 2). Hence, although FKM/GO/LDH-Ca₂Al lost some of the ability to reduce the diffusion coefficient brought by LDH-Ca₂Al, it improved the CO₂ resistance stability of FKM/GO/LDH-Ca₂Al.

3.7.2. Effect of Temperature on the Gas Barrier Properties of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al Composites

The barrier mechanism of GO/LDH-Ca₂Al and LDH-Ca₂Al nanofillers in FKM/GO/LDH-Ca₂Al with FKM/LDH-Ca₂Al composites against CO₂ was investigated by conducting a temperature-dependence study on the gas barrier properties of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al composites. The permeability coefficients ($Q$), diffusion coefficients ($D$), and solubility ($S$) of the rubber nanocomposites with temperature are given in Figure 9. The diffusion coefficients of FKM, FKM/GO/LDH-Ca₂Al-8, and FKM/LDH-Ca₂Al-8 remarkably increased with temperature, but the solubility remarkably decreased with the increase in temperature. Based on this phenomenon, the Arrhenius equation can be used to explain the relationship between the ln(S) and ln(D) of rubber nanocomposites and temperature (T), which can be expressed by Equations (5) and (6) as follows:

$$Ln\left(S\right)=Ln\left(S_0\right)+(∆H_{S0}+∆H_{sf}+∆H_{sc})/RT$$

(5)

$$Ln\left(D\right)=Ln\left(D_0\right)+(∆E_{D0}+∆E_{Df}+∆E_{Dc})/RT$$

(6)

where $S_0$ and $D_0$ are pre-finger factors, which are related to the nature of the nanocomposites; $∆H_{S0}$ and $∆E_{D0}$ represent the heat of dissolution and diffusion activation energy of the pure polymer; $∆H_{sf}$ and $∆E_{Df}$ is the packing effect caused by nanofillers, which are related to the interaction of nanofillers with polymers and diffusion gases; and $∆E_{Df}$ is related to factors such as the aspect ratio, size, and shape of the nanofiller. In this experiment, $∆H_{sc}$ refers to the change in the heat of dissolution caused by the compound effect of GO and LDH-Ca₂Al in GO/LDH-Ca₂Al. $∆E_{Dc}$ is the activation energy of the reaction between the nanofiller and the diffusion gas. Table 4 summarizes the relevant parameters of Equations (5) and (6) for the FKM, FKM/GO/LDH-Ca₂Al-8, and FKM/LDH-Ca₂Al-8 samples. The diffusion activation energy (ΔE) of the polymers is less than −25 KJ/mol, and the solubility activation energy (ΔH) is higher than 12.3 KJ/mol, which is consistent with the diffusion coefficient ($D$) and the solubility ($S$) with temperature. The $∆H_{sf}$ from the packing effect of LDH-Ca₂Al is −32.2 KJ/mol, resulting in a significant decrease in the heat of dissolution of CO₂ in the FKM/LDH-Ca₂Al nanocomposites to approximately 12.3 KJ/mol, which explains the strong interaction between CO₂ and LDH-Ca₂Al. This phenomenon led to an increase in the solubility of CO₂ in the FKM/LDH-Ca₂Al nanocomposites. However, the change in the heat of dissolution caused by the compound effect of GO with LDH-Ca₂Al ($∆H_{sc}$) is 3.8 KJ/mol, which explains the existence of a strong interaction of GO/LDH-Ca₂Al with the polymer matrix that reduced the CO₂ adsorption effect on the nanofiller surface. The diffusion activation energy of the nanocomposites remarkably increased, which rendered the activation energy of the reaction of CO₂ with GO/LDH-Ca₂Al negligible, in which the value increased by approximately 42.5%. Hence, the diffusion coefficients of the FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al nanocomposites were insensitive to the temperature variation. In addition to the packing effect caused by the filler ($∆E_{Df}$), a chemical activation energy was observed between CO₂ and LDH-Ca₂Al ($∆E_{Dc}\approx 5.2\mathrm{KJ}/\mathrm{mol}$) in the ΔE of the FKM/LDH-Ca₂Al nanocomposites. Therefore, it has a significantly lower diffusion coefficient than that of FKM/GO/LDH-Ca₂Al. The permeability coefficient of the composites is a combination of the diffusion coefficient and solubility ($Q=DS$) [5,48,62,63]. As shown in Figure 9c, the solubility of the nanocomposites substantially contributes to the permeability coefficient at 60–80 °C, resulting in the low permeability coefficient of FKM/GO/LDH-Ca₂Al. By contrast, the diffusion coefficient of the nanocomposites substantially contributes to the permeability coefficient at 80–100 °C. Therefore, the permeability coefficient of the FKM/LDH-Ca₂Al nanocomposites is lower than that of FKM/GO/LDH-Ca₂Al.

3.8. Mechanical Properties of FKM/LDH-Ca₂Al and FKM/GO/LDH-Ca₂Al Nanocomposites

The effect of LDH-Ca₂Al and GO/LDH-Ca₂Al on the mechanical properties of fluoroelastomers was quantified by determining the key mechanical parameters of the nanocomposites with different filler contents (5, 8, and 10 phr), including 100% modulus (E100), Shore A hardness, σ_b, and ε_b (Figure 10 and

Table 5. Mechanical properties of FKM, FKM/LDH-Ca₂Al, and FKM/GO/LDH-Ca₂Al complexes.

Sample	100% Modulus, E100(MPa)	Tensile Strength, σb (MPa)	Strain at Break, εb (%)	Shore A Hardness
FKM	1.5 ± 0.1	9.6 ± 1.1	251 ± 17	50 ± 1
FKM/LDH-Ca2Al-5	2.1 ± 0.2	7.3 ± 0.2	215 ± 18	55 ± 1
FKM/LDH-Ca2Al-8	2.3 ± 0.08	9.9 ± 1.1	260 ± 19	56 ± 1
FKM/LDH-Ca2Al-10	2.4 ± 0.03	10.8 ± 0.6	254 ± 8	58 ± 1
FKM/GO/LDH-Ca2Al-5	2.3 ± 0.01	9.2 ± 0.5	238 ± 10	55 ± 1
FKM/GO/LDH-Ca2Al-8	2.2 ± 0.03	10.6 ± 0.8	270 ± 12	56 ± 1
FKM/GO/LDH-Ca2Al-10	2.4 ± 0.08	11.1 ± 0.6	287 ± 8	57 ± 1

). The mechanical properties of the nanocomposites were significantly improved compared with those of FKM, and interestingly, the ε_b of the nanocomposites gradually increased with the content of nanofillers, in which the largest ε_b (ε_b > 287%, σ_b > 11.1 MPa) value was obtained for the 10 phr filled FKM/GO/LDH-Ca₂Al nanocomposites. This finding can be attributed to the orientation arrangement of the nanofillers with two-dimensional layered structures in the complexes when subjected to force, leading to the stable transfer of stress inside the material [9]. In addition, FKM/GO/LDH-Ca₂Al nanocomposites exhibit superior mechanical properties to FKM/LDH-Ca₂Al nanocomposites, in which σ_b, and ε_b increased by approximately 25% and 12% (Table 5) because of the strong interfacial interaction between GO/LDH-Ca₂Al and the polymer matrix, which promotes the efficient transfer of stress from the polymer matrix to GO/LDH-Ca₂Al, consistent with the results described by the phenomenological interaction parameters (Table 3) [19].

4. Conclusions

In this paper, the CO₂ gas barrier and mechanical properties of nanocomposites filled with Ca/Al layered hydroxide modified by GO (GO/LDH-Ca₂Al) were investigated. The GO/LDH-Ca₂Al nanocomposite filler was prepared by depositing LDH-Ca₂Al on the surface of alkalized graphene oxide (Al-GO). This nanocomposite filler has good thermal stability compared with the pristine LDH-Ca₂Al, and the temperature of the stripped interlayer water increased by nearly 43 °C. At the same time, its CO₂-resistant stability was significantly improved, and almost no carbonate generation occurred at 80 °C and 4.3 MPa pure carbon dioxide environment, resulting in low CO₂ reactivity. The barrier mechanism of modified and unmodified LDH-Ca₂Al-filled nanocomposites to carbon dioxide was also proposed, and the diffusion activation energy and solubility activation energy of the nanocomposites consisted of three components: (1) $∆H_{S0}$ and $∆E_{D0}$, which were determined by the nature of the matrix; (2) $∆H_{sf}$ and $∆E_{Df}$ due to packing effect; and (3) the compound effect $∆H_{sc}$ and reaction activation energy $∆E_{Dc}$. In the nanocomposites, the multilayer stacked LDH-Ca₂Al acts as a CO₂ gas barrier to form a distorted diffusion path, and this effect induces an activation energy of $∆E_{D0}$. Considering the pro-CO₂ multilayer structure and strong chemical reactivity of LDH-Ca₂Al, it provides the opportunity for CO₂ intercalation and reaction, resulting in the reaction activation energy of $∆E_{Dc}$. These two factors led to an increase in the diffusion activation energy of FKM/LDH-Ca₂Al nanocomposites by approximately 42.5%. The solubility activation energy of FKM/GO/LDH-Ca₂Al nanocomposites increased by approximately 3.8 KJ/mol compared with that of FKM/LDH-Ca₂Al nanocomposites, which was caused by the compound effect of GO and LDH-Ca₂Al. This phenomenon explains the existence of strong interfacial interactions between GO/LDH-Ca₂Al and the polymer matrix, thus reducing the CO₂ adsorption effect on the nanofiller surface. The permeability coefficients of the nanocomposites to CO₂ gradually decreased, and the mechanical properties (Shore A, σ_b, and ε_b) were significantly improved with the increase in nanofiller content. The best CO₂ permeation resistance and mechanical properties were obtained for the FKM/GO/LDH-Ca₂Al nanocomposites due to the strong interfacial interactions between GO/LDH-Ca₂Al and FKM.