Cryogenic Impact on Carbon Fiber-Reinforced Epoxy Composites for Hydrogen Storage Vessels

Abstract

Carbon fiber-reinforced epoxy (CF/EP) composites are attractive materials for hydrogen storage tanks due to their high strength-to-weight ratio and outstanding chemical resistance. However, cryogenic temperatures (CTs) have a substantial impact on the tensile strength and interfacial bonding of CF/EP materials, producing problems for their long-term performance and safety in hydrogen storage tank applications. This review paper investigates how low temperatures affect the tensile strength, modulus, and fracture toughness of CF/EP materials, as well as the essential interfacial interactions between carbon fibers (CFs) and the epoxy matrix (EP) in cryogenic environments. Material toughening techniques have evolved significantly, including the incorporation of nano-fillers, hybrid fibers, and enhanced resin formulations, to improve the durability and performance of CF/EP materials in cryogenic conditions. This review also assesses the hydrogen barrier properties of various composites, emphasizing the importance of reducing hydrogen permeability in order to retain material integrity. This review concludes by highlighting the importance of optimizing CF/EP composite design and fabrication for long-term performance and safety in hydrogen storage systems. It examines the prospects for using CF/EP composites in hydrogen storage tanks, as well as future research directions.

1. Introduction

The use of carbon fiber-wrapped tanks has replaced all-metal tanks for hydrogen storage. Liquid hydrogen may be held in Type I and V tanks. In contrast, Type I tanks were made of stainless steel, which made them heavier and had a poorer mass density when it came to storing hydrogen. All modified CF/EP composites were used in the construction of a Type V vessel without inner liners [1]. Outstanding mechanical properties of CF/EP composites, such as a high strength-to-weight ratio, superior stiffness, and fatigue resistance, are widely recognized [2]. These characteristics make them very desirable for a range of cutting-edge technical applications, such as renewable energy, automotive, along with aerospace [3]. However, the performance of these composites in extreme environments, notably at cryogenic temperatures, remains a crucial topic of research. CF/EP composites present special challenges in cryogenic conditions, as materials can change significantly in their mechanical and interfacial properties at extremely low temperatures. Understanding these changes is critical for the design and implementation of CF/EP composites in cryogenic environments such as space exploration, superconducting magnets, and hydrogen storage tanks [4]. At cryogenic temperatures, epoxy matrices can become brittle, and the interfacial bonding between the CFs and the epoxy matrix may be weakened, resulting in structural integrity failures [5].

Fiber-resin delamination, resin cracking, as well as fiber fracture are possible failure mechanisms for CF/EP composites. Only around 60% of the carbon fibers’ strength is utilized in the cryo-compressed hydrogen (CcH₂) vessel composite layer, and under typical circumstances, carbon fiber strength failure is rare. Moreover, in cryogenic temperatures, the carbon fiber surface roughens, strengthening the fiber-resin interface’s connection [6]. Consequently, at cryogenic temperature as opposed to room temperature (RT), carbon fiber (CF) along with resin delamination happen less frequently. Resin, the principal opposite stress unit in composite layers, degrades the entire cryogenic mechanical characteristics of CF/EP composites by increasing their brittleness in cryogenic environments alongside having a low susceptibility to crack formation and development [7]. Important steps to take in order to enhance the cryogenic mechanical characteristics of CF/EP composites are to make the resin more cryotough and to stop fractures from spreading throughout the matrix [8]. Toughening resins is usually achieved using thermoplastics, elastic rubbers, and nanomaterials. Plastic materials’ high toughness allows them to withstand thermal stress-induced deformation, which has a substantial impact on thermosetting resin toughness. Similarly, nanoparticles can toughen resin while also preventing microcrack growth. A substitute for rubber toughening is the application of thermoplastics as the second stage of a stiff component to improve the toughness of cured EP composites. Thermoplastics have the potential to improve toughness without sacrificing modulus or heat resistance. Thermoplastics eventually split from epoxy resins during the curing reaction to form the co-continuous phase or phase inversion microstructure, which has a stronger toughening effect than the spherical or nearly spherical rubber toughened phase [9]. Research into the evolution pattern of CFRP cryogenic mechanical properties may begin with its fundamental components, namely carbon fiber, resin matrix, and interface. The results show that the cryogenic state influences the microstructure and properties of carbon fibre. Under cryogenic conditions, the surface of carbon fibre becomes rougher, resulting in improved bonding strength at the fibre-resin interface as well as interlaminar shear strength. At −196 °C, carbon fiber and resin have different linear expansion coefficients (CTE), resulting in larger shrinkage for resin. Thus, thermal stress is primarily focused on two distinct regions. One is the interface between the fiber and the matrix, which increases interfacial shear strength. The matrix between two fibers is a primary cause of microcracks in CFRP. Extending a microcrack can lead to cryogenic propellant leakage, tank material breakdown, and failure in serious situations [8]. A liner Type V CcH₂ storage tank may fail due to microcracks that penetrate the composite as well as vacuum insulation layers [10]. It has been shown that increasing resin toughness is a helpful strategy for lowering cracking at cryogenic temperature. Elastic rubber along with thermoplastics are two common materials used for toughening. The resin along with thermoplastic polymers combine and cure to form a stable, homogeneous phase. The characteristics of thermoplastics remain unaffected by cryogenic temperatures in CcH₂ tanks due to their wide temperature tolerance range. Several thermoplastics are often used for resin modification, including polyethylene glycol (PEG) [11], polysulfone [12], polyethersulfone [13], polysiloxane [14], polycarbonate (PC) [15], and polyetherimide (PEI) [16]. Most of the toughening effect of thermoplastic materials on resins is attributed to the co-continuous phase structure that is created during mixing along with curing. The rate of migration of the molecules is constrained by the chemical bonds as well as molecular shrinkage of the resin in cold settings, increasing bonding strength while lowering resin hardness. Conversely, the dispersed phase of thermoplastic toughening agents consists of co-continuous microstructural domains that are related. This phase boosts the cryogenic effect strength of the resins by absorbing thermal stress energy [17]. Additionally, the flexible thermoplastic polymer chains’ soft ether bond segments might lessen internal stress in the brittle resin network. Many of the flexible and soft ether-bonded segments incorporated into the resin network stay partly frozen even in cryogenic settings, which lessens internal stress brought on by bond motions. This inhibits the development of thermal stress microcracks and enhances the cryogenic mechanical characteristics of the resin. Reduction of resin’s coefficient of thermal expansion (CTE) can also be aided by a thermoplastic toughening agent. However, adding thermoplastic modifiers does not always lead to improved resin properties since the unbound volume of the thermoplastic molecules occupies up space in the cured, improved resin network. The redesigned network’s free volume addition lowers the crosslinking density of the polymer. The resin’s ability to cure is reduced when the modifier content is very high. The tensile strength of the modified resin is negatively impacted by the flexible segments that emerge after curing. Rigid rubber fragments can also be used to toughen resin. Rubber vacancies, which encourage shear plastic deformation of the matrix and hence raise yield strength, are the main factor behind the toughening process of rubber-added resin. The energy required to expand the rubber phase, the pore effect brought on by detachment, the extension of the shear band, particularly the influence of anchoring the rubber particles on the cracking process are all connected to resin toughness [18]. The degree of rubber elastic particle size-to-resin toughening is directly correlated. In the toughening process, larger (microscale) particles perform a number of roles, such as fracture healing, deviation, detachment, as well as void creation, which causes void expansion and shear yield. Conversely, fracture fixing and deviation do not happen in tiny (nanoscale) particles considering these mechanisms need a particle diameter greater than the size of the crack [19]. Nanosized particles hence increase shear yield, create voids, and detach from resin to improve toughness [20]. While rubber elastic particles may greatly increase the toughness along with fracture energy of epoxy resins, a number of problems could occur when they are being used at cryogenic temperature. Thermoplastic materials exhibit superior cryogenic toughening of resins compared to rubber particles in linerless cryo-compressed hydrogen storage tanks. The cryogenic mechanical characteristics of resins can be enhanced by nanoparticles. Enhanced interfacial contact might arise from the interaction between active groups on the surface of the nanomaterial and the resin. Stress from the nanoparticles may be transmitted to the epoxy resin regions around them, absorbing fracture energy while increasing matrix ductility [21]. By acting as physical crosslinking sites for the molecules of resin, nanoparticles (NPs) efficiently prevent the development of fractures. As a microcrack expands, it comes into contact with a cluster of nanosheets oriented vertically in the direction of the crack [22]. The propagation channel of the crack is now blocked, resulting in several branching fractures. Popular cryogenic modifiers in resins include boron nitride nanotubes (BNNTs) [23], silica nanotubes (SNTs) [24], CNTs [25], among others. This review study explores the cryogenic impacts on the tensile strength and interfacial bonding of CF/EP composites, with an emphasis on advances and applications in hydrogen storage tanks. We investigate the fundamental mechanisms underlying the cryogenic behavior of CF/EP composites, including the impact of temperature on the matrix, fibers, and fiber–matrix interface. We evaluate the experimental and theoretical investigations undertaken to elucidate cryogenic mechanical features such as tensile strength, modulus, and fracture toughness. We also investigate the effect of cryogenic temperatures on interfacial properties such as adhesion, debonding, and interlaminar shear strength. Furthermore, we highlight advances in cryogenic testing procedures and analytical models for forecasting CF/EP composites cryogenic behavior. We discuss recent advances in tailoring the microstructure and composition of CF/EP composites for improved cryogenic performance. Finally, we investigate the existing and potential uses of CF/EP composites in hydrogen storage vessels focusing on the challenges and opportunities related with cryogenic implementation.

2. Cryogenic Temperature Effects

2.1. The Matrix Influence

Studies into the evolution pattern of CF/EP composites cryogenic mechanical properties may begin with its fundamental constituents, namely carbon fiber, the resin matrix, and the interface. The findings show that the cryogenic state influences the microstructure and properties of carbon fiber. Under cryogenic conditions, the surface of CF becomes rougher, resulting in improved bonding strength at the fibre-resin interface as well as interlaminar shear strength [6]. At −196 °C, carbon fiber and resin have different linear expansion coefficients (CTEs), resulting in larger shrinkage for resin. Thus, thermal stress is primarily focused on two areas. One is the fiber–matrix interface, which increases interfacial shear strength [26]. The other is the matrix between two fibres, which is a primary source of microcrack formation in CF/EP composites. Extending a microcrack can result in cryogenic propellant loss, tank material damage, and failure in an extreme situation [27]. Meiling Yan and her team [8] delve into CF/EP composites unidirectional laminate, from a microstructural standpoint, examine the cryogenic evolution process of the matrix that influences the cryogenic mechanical properties of CF/EP composites. These studies offer light on the behavior of EPs in harsh environments by revealing how cryogenic temperatures affect their mechanical properties. Nowadays, epoxy resin is the most often utilized material for the resin matrix of CFRP tanks [28,29]. The resin samples’ composition and content are described in Figure 1. The contents of the four types of epoxy resin samples mainly include epoxy resin, curing agent and diluent (Figure 1). Among them, the epoxy resin is glycidyl ether resin, include bisphenol A and bisphenol F. According to Figure 1, the major structural distinction is the two methyl groups connected to the C in between the two benzene rings. This may cause steric hindrance effect during the rotation of the molecular chain, resulting in high molecular stiffness and correspondingly poor toughness. The curing agent is mainly aliphatic amine and aromatic amine. The benzene rings of aromatic amines have high steric hindrance, consequently high stiffness. Aliphatic amines are flexible long-chain alkanes. The active diluent containing epoxy groups participates in the curing reaction, resulting in excellent toughness of epoxy resin. The methyl groups of bisphenol A located within benzene rings induce steric hindrance. The stiffness of epoxy resin is directly impacted by the quantity of benzene rings. Compared to DER354, WSR615 and CYD-128 have larger cryogenic moduli because to greater steric hindrance and the quantity of benzene rings. The cryogenic elongation at break of DER354 reduces with increasing concentration of the soft segment (reactive aliphatic diluent). Of the four resin types, cryogenic temperatures have the least effect on EL-2203 (Figure 2). Steric hindrance is caused by the methyl groups of bisphenol A between the benzene rings.

The rigidity of an epoxy resin is directly proportional to the number of benzene rings. The mechanical property growth rates and reduction rates of different types of epoxy resins is presented in Figure 2. Among them, WSR615 and DER354 contain the same curing agent, whereas the epoxy resin contents are bisphenol A and bisphenol F, respectively. The strength and modulus of WSR615 is affected more by cryogenic condition, while the elongation at break and impact strength of DER354 are largely dependent on temperature. With higher steric hindrance, the intermolecular forces in bisphenol A increased during cryogenic shrinkage, therefore cryogenic temperature can significantly promote the strength and modulus of WSR615. On the other hand, with the low steric hindrance in bisphenol F and the application of reactive diluent for flexible segments, DER354 exhibits poor flexibility and lowered elongation at break under cryogenic condition. Comparing to WSR615, CYD-128 has the same epoxy resin structure, but different curing agent. In the same amount of curing agent content, CYD-128 contains more benzene rings. The modulus of CYD-128 varies the most under cryogenic condition, whereas its strength and toughness are less affected. Out of the four resin types, EL-2203 is less impacted by cryogenic temperature. The components of EL-2203 are complicated, with two epoxy resins, three curing agents, and one reactive diluent. Since the proportion and molecular structure of each component remain unknown, it can be difficult to be analyzed in detail. To sum up, it can be concluded that the stiffness of EP is dependent on the number of benzene rings; the more benzene rings are contained, the more stiffness it increases under cryogenic condition. The elongation at break of those with higher portion of flexible segments is relatively affected more by cryogenic condition; their flexibility reduces and toughness is lowered more. Nevertheless, given the complex of epoxy resin, the molecular structural impact on their cryogenic mechanical properties requires further study.

The mechanical properties of several epoxy resins at room temperature and cryogenic temperature (−196 °C) are shown in Figure 3. Significantly impacted by the cryogenic environment are the modulus of CYD-128 and the tensile strength of WSR615, which increased by 286.0% and 70.8%, respectively. Impact strength of WSR615 and failure strain of DER354 reduced by 72.5% and 37.8%, respectively. The mechanical properties of several epoxy resins at room temperature and cryogenic temperature (−196 °C) showed several interesting findings:

• -. The tensile strength (δ) of epoxy resin types was shown to increase in cryogenic temperature (−196 °C) compared to room temperature (Figure 3a).

• -. This increase in tensile strength at cryogenic temperatures is due to the freezing of the main-chain segment of epoxy polymers and some micromolecular structure units, making them immobile.

• -. The presence of -CH₃ (methyl groups) between benzene rings in bisphenol-A was found to increase steric hindrance, which directly influences the stiffness of the epoxy resin.

• -. Under cryogenic conditions, the tensile modulus (E) of all epoxy types increased (Figure 3b).

• -. Certain epoxy types have a higher cryogenic modulus than others as steric hindrance and benzene ring content increased.

• -. The modulus of epoxy is significantly associated with molecular strand segment packing density and free volume fraction, which influence its mechanical properties at cryogenic temperatures.

• -. The failure strain (ε) of the epoxy types was found to decrease at freezing conditions, indicating a decrease in toughness (Figure 3c).

• -. Epoxy becomes brittle and loses hardness at cryogenic temperatures due to the freezing of resin molecules and micromolecular structure units, rendering them immobile.

• -. The cryogenic elongation at break of some epoxy types decreased as the percentage of soft segments (reactive aliphatic diluent) increased.

• -. The impact strength (d) of epoxy was shown to be reduced in cryogenic conditions, indicating a decrease in toughness (Figure 3d).

• -. At cryogenic temperatures, the microstructure of epoxy is critical in determining its impact strength. Factors such as molecular strand segment packing density and free volume fraction influence this structure.

• -. This study emphasizes the relevance of understanding the microstructure of epoxy in determining its cryogenic mechanical properties.

2.2. Mechanical Properties of CFRP at Cryogenic Temperatures

Figure 4a illustrates the analytical results for the cryogenic mechanical properties of CFRP unidirectional laminate. The study on the mechanical properties of CFRP at cryogenic temperature showed several interesting findings:

• -. The 0° and 90° tensile strength of CFRP increased by 0.2% at −196 °C.

• -. The increase in tensile strength at cryogenic temperature can be attributed to changes in the microstructure of the resin matrix and carbon fibers, which result in stronger bonding at the fiber-resin interfaces.

• -. The in-plane shear strength of CFRP increased by 27.5% at −196 °C.

• -. The roughening of the carbon fiber surface under cryogenic temperatures boosted bonding strength at the fiber-resin interface, resulting in higher interlaminar shear strength.

• -. CFRP’s compressive strength increased by 8.2% at −196 °C.

• -. Thermal stress concentration at the fiber–matrix contact and within the matrix between fibers assisted increase compressive strength at cryogenic temperatures.

• -. The failure strain of CFRP decreased from 1.6% at room temperature to 0.9% at −196 °C.

• -. The drop in failure strain at cryogenic temperatures was attributed to resin embrittlement, which resulted in decreased toughness and increased brittleness of the composite.

• -. The fracture surface morphology of CFRP at cryogenic temperatures was smoother and cleaner, indicating a different fracture behavior than at RT (Figure 4b).

• -. Freezing resin molecular segments and microstructural units at cryogenic temperatures resulted in more brittle fracture behavior, which influenced CFRP’s overall mechanical properties.

• -. The resin free volume influenced molecular mobility, stress distribution, fracture behavior, and intermolecular forces, all of which contributed to the composite’s toughness at cryogenic temperatures.

These studies provide insight on the mechanical behavior of CFRP at cryogenic temperatures, including changes in tensile strength, shear strength, compressive strength, failure strain, and fracture behavior. Understanding these changes is critical for improving CFRP performance in cryogenic conditions.

2.3. Epoxy Resin’s Effect on CFRP

Storage modulus G′ and loss modulus G″ can be used to study the movement of an epoxy molecular chain segment at cryogenic temperatures (Figure 5). The G′ of epoxy decreases with temperature, suggesting a considerable decrease in the deformability of epoxy resin. Epoxy becomes more rigid as a result of decreased molecular mobility. The big peak at −30 °C is indicative of the epoxy structural unit’s low amplitude vibration.

Figure 6a displays a tan δ versus temperature curve, indicating that the main-chain segment of resin molecules gradually defrosts as temperature increases. Tangent loss causes a significant increase in frictional force between molecules. Based on these findings, the following explanations and interpretations are given for the influence of epoxy on CFRP at cryogenic temperature:

• -. The peak at 147 °C on the tan δ vs. temperature curve corresponds to the epoxy resin’s glass transition temperature. This temperature marks the point at which the epoxy resin shifts from a glassy to a rubbery state, resulting in changes in its mechanical properties.

• -. The peak at −25 °C in the tan δ vs. temperature curve reflects the resin’s secondary transition temperature. This temperature point implies another phase transition in the epoxy resin, possibly due to changes in molecular mobility or structural rearrangements.

• -. Because −196 °C is substantially lower than the T_g of epoxy and the secondary transition temperature, the main-chain segment of resin molecules and a portion of the micro-molecular structure units freeze and become immobile. Epoxy freezes at temperatures below this range. As a result, epoxy becomes brittle and loses its hardiness. At room temperature, the failure strain is 3.7%, but at −196 °C, it is only 1.1%. CFRP’s cryogenic elongation during break is therefore reduced.

Figure 6b in this study depicts the molecular strand segments of epoxy after reaction with a curing agent. Based on the observations, the molecular structure of epoxy after the reaction with the curing agent is described below:

• -. The molecular structure of epoxy following the curing agent reaction appears amorphous, with molecular strand segments randomly interlocking and tangling. This pattern resembles a reel of thread, implying a complex and interconnected structure of resin molecules. Epoxy resin molecular strand segments are bound together by a variety of intermolecular forces such as physical entanglement, covalent bonds, H-bonds, and van der Waals forces. These forces help to ensure the overall stability and integrity of the resin structure [34].

• -. Epoxy resin’s molecular strand segment packing contains holes and flaws, which produce free volume within the substance. Adequate free volume allows molecular strand segments to move, whereas excessive free volume can result in stress concentration and structural defects. EP’s molecular structure has free volume, which allows for the mobility of molecular strand segments, which is critical for the material’s deformation and mechanical behavior. Proper molecular mobility enhances the material’s overall flexibility and resilience.

• -. The mechanical properties of epoxy resin are heavily influenced by its molecular structure, which is defined by intermolecular forces and free volume. The arrangement and interaction of molecular strand segments affect the resin’s strength, stiffness, and durability. The molecular structure of epoxy resin’s post-curing agent response influences the material’s homogeneity and stability. Proper intermolecular interactions and molecular mobility help to create a more homogenous and strong resin matrix in composite materials like CFRP.

• -. The combination of the EP and curing agent causes the formation of a cross-linked network structure, which improves the material’s mechanical properties and thermal stability. The curing process affects the molecular arrangement and bonding of the resin matrix.

Figure 6c shows that epoxy resin has a CTE of 21.5 × 10⁻⁶/°C at −196 °C and 50.7 × 10⁻⁶/°C at 25 °C, which decreases as temperature decreases. Resin molecules contract under CT. as a result, the polymer’s molecular strand segment uniformity may decrease, causing it to expand locally and reduce free volume. Stress concentrates on areas where free volume increases. Once damaged, fissures expand quickly and destroy the covalent and secondary bonds.

To summarize, epoxy resin behavior at cryogenic temperatures has a substantial impact on CFRP mechanical properties. Changes in resin stiffness, fracture behavior, and intermolecular forces all contribute to increased CFRP strength and modulus at low temperatures. Understanding how epoxy resin affects CFRP mechanical properties at cryogenic temperature is critical for improving the performance of composite materials in cryogenic applications. Figure 7 illustrates the tensile failure morphology of epoxy at room temperature and −196 °C, as well as the corresponding stress–strain curves. The study on the tensile failure morphology of epoxy at room temperature and −196 °C showed several interesting findings:

• -. At room temperature, the tensile failure morphology of epoxy resin exhibits distinctive features such as river patterns that diverge from the breaking point. This pattern represents the direction of crack propagation and the material’s fracture behavior under tensile loading (Figure 7a).

• -. At −196 °C, epoxy resin has a smoother and neater fracture surface than room temperature. Cryogenic conditions impact the material’s fracture behavior, resulting in a distinct surface appearance (Figure 7b).

• -. The stress–strain curve of epoxy at room temperature and cryogenic temperature offers information on the material’s mechanical behavior under various temperature settings. A comparison of the curves reveals how temperature affects the material’s response to applied stress (Figure 7c).

In conclusion, this study gave important insights into the microstructural effects, tensile fracture behavior, and material performance of epoxy resin-based CFRP composites under cryogenic settings. The findings emphasized the importance of resin free volume, intermolecular forces, and microstructural design in enhancing composite mechanical characteristics and toughness for use in extreme temperature conditions.

3. Interfacial Properties

The study’s goal is to better understand how mechanical properties and interfacial bonding strength vary when composite materials are exposed to extremely low temperatures. This evaluation is critical for applications in which carbon fiber/epoxy (CF/EP) composites are exposed to a variety of temperature conditions, ensuring their reliability and performance. Kwon et al. [35] evaluated the interfacial properties of CF/EP composites using electrical resistance measurements. This study investigated the relationship between the tensile properties of individual carbon fiber samples and the electrical resistance ratio (ERR), which revealed a linear relationship. The interfacial evaluation of CF/EP composites using electrical resistance measurements was conducted through the following methodologies:

• -. This study used a single fiber tension test to examine the link between tensile stress as well as the ERR of carbon fiber (ASTM D-3379) [35]. A multimeter was used to measure the electrical signals for the ERR tests, as well as electrical resistance variations were noted at pre-arranged intervals of tensile strain. To lower the resistance of contact among the copper wire as well as the carbon fiber, silver paste was used.

• -. The electrical resistance (ER) fragmentation test sample consisted of a carbon fiber embedded in a semi-transparent EP composite (Figure 8). ER was measured during the carbon fiber fracture in this test. A proportionate relationship between tensile stress and ERR was identified, leading the building of an empirical formula. Rapid increase in electrical resistance indicated carbon fiber fracture, while prolonged fracture durations suggested greater interfacial adhesion between carbon fiber and epoxy matrix. These tests tested several epoxy types and temperature tests, including cryogenic temperatures (−150 °C).

• -. This study used the static contact angle method to determine the wettability of the materials. Contact angles with different epoxy kinds and solvents were measured in order to determine surface energy. The Young equation was used to compute the static contact angle in order to assess the wettability of the materials.

A significant consideration for assessing the interfacial characteristics of CF/EP composites is the link between carbon fiber tensile stress as well as the ERR, as follows:

• -. The findings showed a linear relationship between the ERR and the tensile characteristics of individual carbon fiber samples. In particular, it was demonstrated that the carbon fiber’s ERR and tensile stress were exactly correlated. Single fiber tensile tests were used in experimental research to demonstrate this link.

• -. An empirical formula was created to measure the connection between both of these parameters based on the proportional correlation that was discovered between tensile stress along with ERR. This method yields a valuable tool for analyzing the behavior of CF/EP composites by offering a means of evaluating the interfacial characteristics of ER fragmentation samples.

• -. The relationship between ERR and tensile characteristics worked well for calculating the embedded carbon fiber fracture strength in EP. Researchers’ comprehension of the interfacial characteristics of these composite materials increased when they were able to calculate and predict carbon fiber fracture tensile stress utilizing an empirical formula created from experimental data.

The experimental data and the empirical formulae for ERR vs. strain along with carbon fiber tensile stress are displayed in Figure 9a. Figure 10 illustrates this using drawings. Because of their consecutive C-C bonds and mismatched microfibril structure, carbon fibers often exhibit electrical conductivity. Stress may cause some C-C bonding to break, raising the CF’s electrical resistance.

The following important conclusions were obtained from the study’s interfacial property prediction utilizing the ER fragmentation approach:

• -. The bisphenol A type epoxy ER fragmentation tests conducted at different temperatures provide important insights into the interfacial characteristics of the composite material. In the ER fragmentation investigations, the ERR exhibited distinct behaviors at RT and CT, indicating changes in interfacial adhesion at varying temperatures.

• -. The bisphenol A-type EP ER fragmentation results at two distinct temperatures are shown in Figure 9b and Figure 9c, respectively: RT (Figure 9b) and CT (Figure 9c). The ERR grew more quickly at RT than at CT, according to the ER fragmentation data. Additionally, during CT compared to RT, the tensile stress of the ER fragmentation samples increased more quickly. Compared to CT, the CF’s calculated fracture tensile stress was greater at RT. The interfacial properties of CF and bisphenol A declined at CT. Figure 9d,e show the ER fragmentation of bisphenol F type epoxy at RT and CT. Compared to bisphenol A type, bisphenol F type ER fragmentation generated greater electrical signal noise at room temperature. This results from tensile tension during the ER fragmentation test, which causes fiber sliding effects. Compared to bisphenol A-type EP, bisphenol F-type EP exhibited a reduced ERR at fracture time. This implies that the interfacial characteristics of bisphenol A-type EP are better than those of bisphenol F-type EP. The ERR of bisphenol F-type EP and bisphenol A-type EP varied greatly at different temperatures. ERR values for bisphenol F-type EP were similar at RT and CT. These experiments provided insightful data on the behavior and interfacial characteristics of the composite material at different temperatures. The findings for the two types of EPs at different temperatures are displayed in Figure 9, together with the calculated tensile stress of the CF at the fracture site during the ER fragmentation investigations. CF fracture tensile stress of 2742 MPa was anticipated for bisphenol A-type EP, which had an ERR of 0.023 at RT. A little less than the RT value, 1928 MPa, was the projected CF fracture tensile stress at CT. Additionally, the estimated strain of the CF in fragmentation samples was examined in a number of ways. Due to fiber sliding effects, the projected CF strain raised by 0.07% at CT compared to RT. For specimens with a bisphenol F-type EP matrix, the calculated CF fracture tensile stress was 2146 MPa at RT and only marginally dropped (to 2027 MPa) at CT. These findings suggest that bisphenol F-type EP is a better option for CT applications. Figure 11 depicts a tensile loading model for fragmentation specimens with varying interfacial adhesions between fiber and material. The model illustrated the behavior of fragmentation specimens under tensile loading conditions, with a focus on the varying levels of interfacial adhesion between the fiber and material. For specimens with good interfacial adhesion, the model would show minimal fiber slipping and effective stress transfer between the fiber and material during tensile loading. This scenario would result in high predicted CF fracture tensile stress values, indicating strong interfacial properties between the fiber and material. In contrast, for specimens with poor interfacial adhesion, the model would demonstrate significant fiber slipping under tensile loading conditions. This would lead to poor stress transfer between the carbon fiber and the material, resulting in lower predicted carbon fiber fracture tensile stress values compared to specimens with good interfacial adhesion. In summary, the model of tensile loading of fragmentation specimens with different interfacial adhesion levels offers a visual representation of how the interfacial properties between the fiber and material impact the mechanical behavior of composite materials under tensile stress, providing valuable insights for material characterization and optimization.

• -. Static contact angle tests were used to determine the amount of adhesion between the CF and material. This method involved measuring contact angles for different types of EP with various solvents to assess the wettability and adhesion properties of the composite material. The work of adhesion measurements provided quantitative data on the interfacial interactions between the fiber and the material. Figure 12a shows the surface temperature of EP samples as a function of time after immersion in liquid nitrogen. Temperature affected the static contact angle of a water droplet on the surface of EP samples. EP samples were immersed in a liquid N₂ container for 30 min during this study. The specimen’s surface temperature rapidly increased within 15 s of being extracted from the liquid N₂ container. After 30 s, vapor condensation began, and one minute later, a droplet of water was frozen. As a result, surface energy measurements on EP specimens should be performed 15–30 s after being removed from the liquid N₂ container. Figure 12b depicts the static contact angle data for several EP samples at various temperatures. Figure 12c illustrates the surface energy of various solvents on EP specimens at varying temperatures. Bisphenol F-type EP has a 10° greater static contact angle than bisphenol A-type EP at RT. At CT, the static contact angle of bisphenol F-type EP dropped, whereas that of bisphenol A-type EP rose. The ER fragmentation test method involved conducting tests on fragmentation specimens to evaluate the interfacial properties between the fiber and material. By measuring the electrical resistance changes during CF fracture in the ER fragmentation tests, researchers could assess the stress transfer and interfacial adhesion in the composite material. The method offered a means of quantifying the interfacial properties based on electrical resistance behavior.

In conclusion, this research on evaluating interfacial properties in CF/EP composites using a combination of work of adhesion measurements and ER fragmentation tests provided valuable insights and established a reliable method for assessing the fiber–matrix interface in composite materials. The consistent results and comprehensive evaluation approach contribute to advancing the understanding of interfacial properties and optimizing the performance of composite materials in practical applications.

Shao et al. [36] investigated the effects of cryogenic treatment on the mechanical and interfacial characteristics of CNT fiber/bisphenol-F epoxy matrix, including interfacial shear strength and fragmentation test results. The specimen preparation process in the study involved several key steps:

• -. The CNT fiber for the fragmentation test samples was fixed in the center of a silicone mold, as shown in Figure 13a.

• -. Conductive silver paste was used to connect two copper wires to the CNT fiber.

• -. Following 30 min of full vacuum at RT, the EP mixed with curing agent was poured into the mold containing the CNT fiber.

• -. The matrix was allowed to cure for 10 h at 50 °C.

• -. Cryogenic treatment was conducted at −196 °C.

• -. The treatment involved decreasing the RT from RT to −196 °C at a cooling rate of 2 °C per minute.

• -. The samples were kept at −196 °C for 12 h to achieve thermal equilibrium before gradually increasing the temperature back to RT, as shown in Figure 13b.

The findings on the mechanical and interfacial properties of the CNT fiber/bisphenol-F EP matrix, before after cryogenic treatment are as follows:

• -. The cryogenic-treated epoxy had a 51% higher Young’s modulus compared to the untreated EP.

• -. The tensile strength of the cryogenic-treated EP increased by 27% compared to the untreated epoxy.

• -. After cryogenic treatment, the epoxy’s failure strain remained constant.

• -. The electrical resistances of the CNT fibers before and after the cryogenic process were evaluated to determine the interfacial stress transfer between the CNT fiber and EP. The cryogenic-treated sample’s electrical resistance changed rate with strain, indicating a better fiber/matrix interaction (Figure 14a).

• -. The composite’s interfacial shearing strength increased by 31% after cryogenic treatment compared to the pristine CNT fiber composite. These numerical values provide specific information about the interfacial properties improved by cryogenic treatment in the CNT fiber composite. The increase in interfacial shearing strength and the enhanced fiber/matrix interaction suggest a more robust and efficient load transfer mechanism within the composite material, which is essential for applications requiring high mechanical performance and reliability.

• -. After cryogenic treatment, there was a small decrease in the static resistance of the embedded CNT fiber compared to the untreated sample (Figure 14b).

• -. The gage factor value of the cryogenically treated embedded CNT fiber was significantly greater than that of the untreated sample. The specific numerical values for the gage factor before and after cryogenic treatment were reported as 1.6 for the treated sample and 0.6 for the untreated sample (Figure 14b). These values provide specific insights into the electrical properties of the embedded CNT fiber before and after cryogenic treatment. The decrease in static resistance and the increase in gage factor suggest enhanced interfacial bonding and stress transfer within the composite material, which are crucial for applications requiring reliable load transfer and improved mechanical performance.

Yao et al. [21] looked into the possibility of grafting carbon nanotubes (CNTs) onto fibers in a single process at very low temperatures to enhance the interfacial characteristics of CF/EP composites. At extremely low temperatures, a unique one-step chemical vapor deposition (CVD) technique was used to deposit CNTs on CF surfaces, and molding technology was employed to generate hybrid composites. Studies in dielectric and mechanical aspects were carried out to ascertain the modifications in out-of-plane characteristics subsequent to CNT grafting. The CVD process for CNTs on CF consists of the following steps:

• -. CF fabrics are desized and oxidized to improve the surface for CNT growth.

• -. The oxidized CF fabrics are immersed in a catalyst solution containing Fe(NO₃)₃·9H₂O and Co(NO₃)₂·6H₂O dissolved in C₂H₅OH (ethanol).

• -. Catalyst-coated CF fabrics are placed into a CVD furnace, and the air pressure is reduced to a vacuum before introducing N₂ gas.

• -. The temperature is raised gradually, and a mixture of C₂H₂ and H₂ gases is injected into the furnace at controlled flow rates.

• -. CNT growth occurs on the surface of the CFs at specific temperatures (e.g., 400–485 °C) for a defined duration.

• -. During the CNT growth process, the furnace pressure is maintained at 0.01 MPa.

• -. After the growth period, the injection of gases is stopped, and the remaining carbon source is removed by reducing the air pressure in the furnace.

Mechanical and dielectric tests were conducted to assess the changes in out-of-plane properties after grafting CNTs on the CF surface. The following results were obtained:

• -. The CNTs-coated CF had higher strength and Weibull modulus than desized fibers due to less damage caused by catalysts at low temperatures and the repairing effect of CNTs.

• -. The ILSS of CNTs-coated CF/EP composites increased by approximately 32.29% compared to desized fiber-based composites (Figure 15a).

• -. The IFSS of CNTs-coated CF/EP composites increased by approximately 30.73% compared to desized fiber-based composites (Figure 15b).

• -. The dielectric properties of the materials were improved after the addition of CNTs.

• -. Three forms of carbon nanotube (CNT) fracture in materials have been identified: breaking, debonding with resin, along with fracture of the CNT/fiber junction. For effective load transfer, a new interphase was formed among CF and the EP matrix as a result of the CNT layer’s improved wettability.

Overall, the results indicated significant enhancements in mechanical properties, including tensile strength, IFSS, and ILSS, as well as improvements in dielectric properties after grafting CNTs on the CF surface. The CNT layer-induced interphase played an important role in improving the composites’ out-of-plane properties.

Surface characterizations of the CF and CNT-coated carbon fibers were conducted using various techniques as follows:

• -. The radial fracture surface of the composites is shown in Figure 16. The desized CF-reinforced materials’ shearing fracture demonstrates that the fiber-EP matrix interface was clearly separated, with a substantial portion of the fibers broken and dragged out. Simultaneously, a significant amount of holes developed and remained in the matrix, suggesting a poor fiber-EP matrix interaction and the absence of an effective interphase. The formation of CNTs on the surface of CFs resulted in a neat and flat fracture surface for the materials, and the huge fissures between the CF along with the EP matrix disappeared, indicating a considerable improvement in the materials’ interface properties. Moreover, there was a protracted transitional interphase between the CF along with EP due to their close relationship.

• -. The interface modulus was calculated using the force modulation mode of the AFM in order to verify the existence of the interphase. The circular form of the CF in the material is seen in excellent detail in Figure 17a. The hardness modulus levels in the two locations differ significantly, as seen by the modulus line distribution. The sudden shift in modulus suggested that the material’s interface bonding was inadequate since there was no transitional interphase connecting the two. A fuzzy transition layer that links the CF as well as the EP formed when CNTs were grafted onto CFs (Figure 17b). There was a clear step-like transition in the modulus line distribution, which is known as an interphase region. There was a modulus between the matrix and the reinforcing fiber, and the interphase thickness was 0.41 μm. This demonstrated that the carbon fiber material treated with CNTs created a uniform transition interphase at the interface that was able to distribute, disperse, including absorb stress, enhancing the material’s mechanical characteristics and interface bonding capacities [37].

• -. One side of the CNT was fully immersed in the EP matrix at the material interface, while the other side was attached to the CF. According to the CNTs’ binding force to the matrix or fiber, Figure 18 depicts three different forms of CNT fracture in materials: (i) A break at the CNT/CF interface. There were still a few raw areas among the mostly resin-coated fiber surfaces (Figure 18a). Because some CNTs made poor contact with the fibers, CNTs were easily removed. Generally speaking, there was no gain in composite bearing capacity as a result of these poor connection positions. (ii) CNTs separate from EP. CNTs that were removed with a greater length and micro-scale perforations (shown by the arrows) on their surface demonstrated this kind. This form of the fracture indicates a poor connection among the CNTs and EP. (iii) CNT fracture. Even with their own fracture around the root, some CNTs that have a close link between them and the EP or fibers are challenging to remove. Usually, a smaller portion of these CNTs stayed on the fiber. When composites broke, types (ii) and (iii) were essential in absorbing more energy, which improved the mechanical properties. The enhancing mechanism of the CNTs may become clearer with more precise pictures of the fracture surface. CNT pull-out and crack bridging are the two toughening mechanisms on the fracture surface, as seen in Figure 18. According to Zhang et al. [38], one or more CNTs are most likely to blame for the filamentous bridging among fibers (Figure 18c). The energy required for this bridging step widens the fracture propagation, increasing the damage resistance of the composites [39,40]. In addition, a single CNT pull-out is shown in Figure 18b,d; it is circled in yellow. As the single CNT pull-out process removes the CNT from the resin until it breaks, it could also take in more energy. According to Yao et al. [41], the particular pull-out energy rose as the length of the CNT inserted grew, suggesting that the absorbed energy increased with the pull-out length. Most CNTs have an embedded length longer than 500 nm, as seen by the length of a single CNT that remains on the carbon fiber surface (Figure 18b,d).

Figure 15

(a) ILSS and (b) IFSS of carbon fiber composites [41]. Copyright 2020 Elsevier.

[Figure omitted. See PDF]

Figure 16

The fracture surface in the radial direction of the composites reinforced by (a) desized, (b) CVD-485, (c) CVD-465, (d) CVD-450, (e) CVD-430, and (f) CVD-400 carbon fibers [41]. Copyright 2020 Elsevier.

[Figure omitted. See PDF]

Figure 17

AFM force modulation images and section analysis of the interphase (along the white arrow) in (a,c) desized carbon fiber/epoxy and (b,d) CVD-400 carbon fiber/epoxy composites [41]. Copyright 2020 Elsevier.

[Figure omitted. See PDF]

Figure 18

The fracture surface in the weft direction of CNT-coated carbon fiber/epoxy composites: (a) bare surfaces; (c) filamentous bridging between fibers; (b,d) length of a single CNT left on the surface of carbon fiber [41]. Copyright 2020 Elsevier.

[Figure omitted. See PDF]

Patnaik et al. [42] studied the effect of CT on the mechanical behavior of graphene carboxyl (GCOOH)-grafted CFRP. This work investigates how nanofiller content affects the mechanical characteristics of these matrices at both RT and CT. Several crucial steps were taken to prepare samples for this investigation on the effect of cryogenic temperature on GCOOH-grafted CFRP composites:

• -. CFs were desized to remove the epoxide coating or sizing by heating them to 120 °C for 2 h in an oven.

• -. A cathodic electrophoretic deposition (EPD) unit was built up to deposit GCOOH onto a negatively charged electrode.

• -. In an aqueous solution of GCOOH and Methyl Violet (MV), CF mats were mounted to hardwood frames as cathodes and stainless-steel plates as anodes.

• -. Sonication in distilled water produced the GCOOH suspension, which was then charged and dispersed with MV.

• -. The EPD technique used a constant current of 5A for 30 min.

• -. After deposition, the GCOOH-coated CF mats were air-dried for 24 h before immersing in acetone for 2 h to remove excess MV.

• -. Thermal annealing at 140 °C for 2 h was used to increase the adhesion between the nanofillers and fibers.

• -. The hand lay-up technique was used to create 12-layered laminates with CF reinforcement and epoxy resin matrix.

• -. The EP and curing agent were mixed in a particular ratio, and the stack of 12 layers was crushed in a hot press at controlled temperatures.

• -. Laminates containing CFs modified by electrophoretic deposition (EPD) are named after the bath concentration used during the process.

• -. The laminates were cured at RT for 24 h before being cut to appropriate dimensions for testing.

Following these processes, the researchers were able to manufacture GCOOH-grafted CFRP composites for mechanical testing at RT and cryogenic temperatures. The findings regarding the influence of CT on the mechanical behavior of the matrices are:

• -. During testing at cryogenic temperatures, all of the composites were brittle and failed catastrophically.

• -. The neat flexural sample showed linear elastic behavior before catastrophic failure, but the changed flexural samples showed pseudo-ductile phenomena before failure at RT.

• -. The pseudo-ductile behavior seen in modified samples at RT was not identified at CT.

• -. Experimental results demonstrated that composites with GCOOH-grafted CFs improved ILSS and FS at both RT and CT.

• -. The composites with the largest improvement in ILSS and FS in the EPD bath contained 1.5 g/L of GCOOH.

• -. The presence of GCOOH at the fiber–matrix interface increased the composite’s interface/interphase performance.

• -. The active carboxyl group in GCOOH established bonds with EP and matrix molecules, strengthening the interaction between the fiber and material.

In conclusion, the study examined the behavior of GCOOH-grafted CFRP composites under cryogenic circumstances, revealing a change toward brittle failure as opposed to the pseudo-ductile behavior reported at RT. The inclusion of GCOOH improved mechanical properties and interface quality, which had a substantial impact on composite performance at cryogenic temperatures (Figure 19). Graphene-based nanofillers offer various advantages when employed in multi-scale fiber-reinforced composites for cryogenic applications:

• -. Graphene-based nanofillers improve the interface between reinforcing CFs and the polymer matrix, resulting in increased interfacial adhesion. This enhanced interface is critical for keeping mechanical properties stable at cryogenic temperatures.

• -. The incorporation of graphene-based nanofillers can improve the interlaminar shear strength and flexural strength of composites. This improvement in mechanical properties is crucial for applications that demand high specific strength, such as aerospace and cryogenic storage tanks.

• -. G-COOH nanofillers have been shown to provide homogenous deposition on CFs, resulting in consistent reinforcement throughout the composite.

• -. Graphene-based nanofillers have shown promise in keeping their properties and performance at low temperatures, making them appropriate for cryogenic applications.

• -. The use of graphene-based nanofillers in composites can cause pseudo-ductile behavior, which can help minimize catastrophic failures and improve overall toughness.

In summary, the addition of graphene-based nanofillers to multi-scale fiber-reinforced composites improves the mechanical characteristics, interfacial adhesion, and performance of the materials, making them suitable for demanding cryogenic applications.

4. Toughening Methods of Epoxy Resins

4.1. Plastic Materials

4.1.1. Biscitraconimide Resin (BCI)

Liu et al. [30] investigate how adding biscitraconimide resin (BCI) improves the mechanical properties of epoxy resins, particularly in cryogenic conditions. The study’s preparation technique for epoxy composites toughened BCI included many important procedures to ensure proper sample mixing and curing. The specific steps in the synthesis process are as follows (Figure 20):

• -. Different formulations of DGEBA/BCI composites with varying BCI content (0%, 10%, 15%, 20%, 25%, and 30%) were produced using the processes described in Figure 21.

• -. DGEBA was warmed at 80 °C for about 30 min with vigorous mechanical stirring.

• -. Addition of BCI: The determined amount of BCI was progressively added to the preheated EP at 80 °C while stirring to scatter it and create a homogenous mixture.

• -. The stoichiometric amount of 4,4ʹ-diaminodiphenylmethane (DDM) was added to the DGEBA/BCI mixture using magnetic stirring until it was homogenous.

• -. The clear homogenous solution was degassed with a vacuum pump and an ultrasonic device to remove any air bubbles that could interfere with the specimen’s properties.

• -. To prevent sticking, the degassed a mixture was put into hot stainless-steel molds that had been coated with a release agent.

• -. The blend-containing molds were cured at temperatures ranging from 100 °C to 200 °C for 1 h each to speed up the curing reaction and solidification of the composites.

• -. The mechanical and thermal properties of the cured samples were evaluated using a variety of characterization techniques, including Fourier transform infrared spectroscopy (FTIR), Thermogravimetric analysis (TGA), Dynamic mechanical analysis (DMA), Scanning electron microscopy (SEM), tensile testing, impact testing, and bending testing.

Following this extensive preparation technique, this work effectively synthesized and characterized EP/BCI composites, resulting in improved cryogenic mechanical properties compared to pure EP. This study examined the reactions involved in the modification of epoxy resin with BCI and the curing agent (DDM). The detailed explanation of the reactions observed in this study are shown in the Figure 22:

• -. The epoxy group in the EP reacts with the amino group in the DDM during the curing process. This reaction results in the creation of chemical bonds between the epoxy and amino groups, which cross-links the polymer network.

• -. In the FTIR spectra, typical peaks of the EP group at specific wavenumbers disappear, showing that epoxy groups are consumed as a result of the reaction with amino groups (Figure 22a).

• -. This study also found a Michael addition reaction between the group the double bond in the BCI and curing agent’s amino (DDM).

• -. This reaction involves adding an amino group to the BCI double bond, which results in the formation of additional chemical bonds and contributes to the modified epoxy system’s toughening mechanism.

• -. The FTIR spectra revealed characteristic peaks corresponding to the reaction between the double bond in BCI and the amino group, suggesting the presence of the Michael addition reaction (Figure 22b).

• -. This study discovered a thorough response of the two intermediate products, the EP/DDM and BCI/DDM systems, in epoxy/BCI composites.

• -. The FTIR spectra of the EP/BCI/DDM materials showed characteristic peaks caused by the reactions between EP and DDM, BCI and DDM, and interactions between the intermediate products.

• -. The shifts in typical peaks in the FTIR spectra showed complicated chemical interactions and reactions occurring in the modified epoxy system, resulting in the development of a new composite material with improved properties (Figure 22c).

Based on these findings, the incorporation of BCI to epoxy resins has been shown to dramatically improve the cryogenic mechanical properties of materials. According to the study, the primary methods via which BCI improves the cryogenic mechanical properties of epoxy resins are as follows:

• -. During the curing process, chemical reactions occur in the EP/BCI materials, resulting in the creation of stronger chemical bonds.

• -. The reactions involving the epoxy groups, amino groups from the curing agent, and double bonds in BCI produce a well-crosslinked network with improved intermolecular interaction.

• -. The addition of BCI to the epoxy resin significantly reduces internal stress in the composite.

• -. BCI increases the cross-linked network’s free volume, allowing for increased molecular chain movement.

• -. This reduction in internal stress and greater free volume improves the composite’s flexibility and toughness, particularly at cryogenic temperatures.

• -. BCI, with a unique molecular structure and properties, serves as a toughening agent in the EP system.

• -. The methyl group in BCI changes the symmetry of the chemical structure, increasing solubility and processing capacities.

• -. The methyl group’s steric hindrance improves BCI’s impact strength and toughness when compared to standard toughening agents like BMI.

• -. BCI’s mechanism for improving tensile characteristics was mostly based on energy dissipation [43]. The formation of microcracks and voids induced by BCI debonding was the primary source of energy loss. Nonetheless, unlike previous research, the current study’s BCI was covalently attached to epoxy cross-linked networks rather than simply blending, resulting in stronger interfacial adhesion and easier motility (Figure 23). As a result, applying an adequate amount of BCI increased EP’s tensile qualities significantly.

• -. The addition of BCI significantly increases the cryogenic tensile strength, fracture toughness, impact strength, and bending strength of EP materials.

• -. At 77 K, the EP/BCI materials shown significant gains in tensile strength, bending strength, fracture toughness, and impact strength over pure EP, making them appropriate for cryogenic applications.

There are two aspects to the toughening of BCI on EP in terms of results. The EP/BCI materials underwent a number of chemical reactions throughout the curing period. Secondly, the inclusion of BCI altered the EP cross-linked network. On the one hand, the curing procedure resulted in the following responses (Figure 22): However, by incorporating BCI into the EP system, internal stress in the brittle EP’s cross-linked networks was greatly decreased, improving the material’s mechanical qualities. Moreover, the introduction of BCI resulted in a decrease in the cross-linked network’s volume and an increase in the system’s free volume (Figure 24), facilitating greater molecular chain mobility. Ultra-low temperatures cannot freeze a molecular chain completely, but they can lessen its internal tension. Lastly, the use of BCI can enhance EP’s mechanical qualities and reinforce it [44].

Overall, this study found that including BCI into the epoxy resin system resulted in considerable increases in fracture toughness, impact strength, bending strength, and tensile strength at ambient and cryogenic temperatures. These advancements demonstrate BCI’s efficiency as a toughening agent for EPs, making the improved composites acceptable for demanding applications in cryogenic conditions and industries that require high mechanical performance.

4.1.2. Polyethylene Glycol (PEG)

Feng et al. [5] employed PEG-4000 as a modifier in the DGEBA/MeTHPA epoxy system to improve its mechanical properties, especially at low temperatures. The sample preparation process for this study are as follows:

• -. To create a homogenous mixture, 60 g of DGEBA was initially combined with different quantities of PEG-4000 (3.2, 6.7, 10.6, as well as 15.0 g), and the mixture was agitated at 60 °C for 30 min.

• -. After adding 1.5 g of accelerator benzyl dimethylamine (BDMA) along with 60 g of methyltetrahydrophthalic anhydride (MeTHPA) to the PEG-4000/EP mixtures at 60 °C, the mixtures were vacuum-degassed for 30 min.

• -. Homogeneous solutions were made and poured into an 80 °C steel mold that had been warmed. After that, they were cured for two hours at 130 °C, ten hours at 150 °C, along with five hours at 170 °C.

• -. ASTM D638-96 along with ASTM D-256 standards [5] were followed in the preparation of the dumbbell-shaped tensile as well as impact specimens, respectively.

• -. To conduct impact testing at 77 K, samples were dipped in liquid nitrogen for more than 10 min before testing.

• -. The samples were characterized using positron annihilation lifetime spectroscopy (PALS) and SEM to evaluate free volume and fracture surfaces, respectively.

The results of PALS provide valuable insights into the enhanced mechanical properties observed in the modified EP sample with PEG-4000. PALS is used to study free volume and molecular packing in polymers. The results of PALS for the unmodified and PEG-4000 modified EPs, including τ₃, I₃, and T_g, are presented in Figure 25. The results of PALS are as follows:

• -. The positronium lifetime (τ₃) values increase with increasing PEG-4000 content in the modified EP materials.

• -. This suggests that the average size of free volumes in PEG-4000 modified epoxy materials improves as the PEG-4000 content increases.

• -. The increase in τ₃ values suggests a correlation with the enhanced mechanical properties, such as ductility and influence resistance, observed in the modified epoxy system.

• -. Modified EP materials have significantly higher intensity (I₃) values than unmodified EP materials. This means that the PEG-4000 modified EP networks have a higher concentration of free volumes than the unmodified system.

• -. The higher concentration of free volumes in the modified EP sample may contribute to improved energy absorption and impact resistance at cryogenic temperatures.

Figure 25a–c depict the sizes of simulated structures of a single PEG-4000, DGEBA, and MeTHPA compound. It is obvious that the flexible PEG-4000 chain utilized in this study, which is significantly greater than the minimal single EP compound and the cure agent MeTHPA compound. Figure 26 depicts cured systems in unmodified EP and free volumes in PEG-4000-modified EP networks. It demonstrates that when PEG-4000 chains are inserted into EP samples, the distances between the networks grow. As a result, adding PEG-4000 increases the size and concentration of free volumes in the modified EP network. The incorporation of PEG-4000 influences the cryogenic tensile strength of EPs in the following ways:

• -. The addition of PEG-4000 to EPs creates flexible segments that remain active at low temperatures (77 K). These flexible portions serve to minimize internal stress in the material, allowing for greater deformation before fracture. This greater ductility contributes to higher cryogenic tensile strengths.

• -. PEG-4000 modification permits free volumes to exist within the cured epoxy system, even at cryogenic temperatures. These free volumes enable the material to absorb energy by allowing for molecular motion during low-temperature impact events. This energy absorption process increases the cryogenic tensile strength of the modified EP sample.

• -. The modified EP sample with a specified PEG-4000 content exhibits the highest impact strength at 77 K. This implies that there is an ideal PEG-4000 concentration for maximizing the epoxy resin’s cryogenic tensile strength. Beyond this ideal percentage, any increases in PEG-4000 may not give significant benefits and could potentially the material’s mechanical performance.

In conclusion, incorporating PEG-4000 into EPs proved to be a successful technique for increasing cryogenic mechanical properties, with the modified epoxy system demonstrating increased tensile strength, ductility, and impact resistance at low temperatures. The study’s findings emphasize the use of PEG-4000 as a modifier to adjust the properties of EPs for cryogenic engineering applications.

4.1.3. Polybutylene Terephthalate, Polyetherimide, and Polycarbonate

In this study, He et al. [45] modified epoxy with three distinct types of thermoplastics to improve performance in cryogenic applications. CF reinforced thermoplastic modified EP matrices were also prepared using VARTM. Figure 27 depicts the chemical formulas of the different thermoplastics. The following processes were followed to modify the EP and prepare the CF/modified EP matrices in this study are:

• -. The polyetherimide (PEI) or polycarbonate (PC) modification process involves dissolving PEI or PC in dichloromethane, mixing the solution into the epoxy YD-128 (DGEBA) at RT, degassing the mixture in a vacuum oven at 80 °C, adding the curing agent (polyoxypropylenediamine) while stirring, and curing the blend at 80 °C for 4 h.

• -. The polybutylene terephthalate (PBT) modification process involves adding PBT directly into the epoxy, heating the mixture to 170 °C, adding the curing agent while stirring, and curing the blend at 80 °C for 4 h.

• -. The preparation of CF/modified epoxy composites involves preparing symmetric laminates with nine plies of carbon fabric and fabricating the composite using vacuum-assisted resin transfer molding (VARTM).

Epoxies modified with various thermoplastics have the following impact strength at CT and RT:

• -. The impact strengths of thermoplastic-modified epoxies (PC, PBT, and PEI) at 77 K were higher than those of the pristine EP.

• -. The modified epoxy’s cryogenic impact strength increased as the thermoplastic percentage increased up to 1.5 wt%.

• -. At CT, the PEI-modified EP demonstrated the maximum impact strength value, which was 45% higher than the pristine EP.

• -. At RT, the thermoplastics modified epoxies outperformed the neat epoxy in terms of impact strength.

• -. The impact strength at RT increased as the thermoplastic component increased to 1.5 wt%.

• -. At RT, the impact strength of the PEI-modified EP increased by approximately 59% over that of the pristine EP.

These results indicate that the addition of thermoplastics, especially PEI, can enhance the impact resistance of the EP composite at both CT and RT.

The results of the thermomechanical analysis of the epoxies modified with different thermoplastics are as follows:

• -. The glass transition temperatures (T_g) (°C) for neat epoxy, PEI modified epoxy, PC modified epoxy, and PBT modified epoxy are 77.98 °C, 82.13 °C, 80.65 °C, and 81.77 °C, respectively (Figure 28a).

• -. The coefficient of thermal expansion (CTE) reduction (%) for PBT modified epoxy, PEI modified epoxy, and PC modified epoxy is 14%, 17%, and 23%, respectively (Figure 28b).

• -. The storage modulus (MPa) at −150 °C for neat epoxy, PEI modified epoxy, PC modified epoxy, and PBT modified epoxy is 4000 MPa, 5025 MPa (30% higher), 4733 MPa (21% higher), and 4539 MPa (17% higher), respectively (Figure 28c).

These findings show that the thermoplastic modified epoxies exhibited higher storage moduli, higher T_g, and reduced CTE than the pristine epoxy. The addition of PEI, PC, and PBT improved thermomechanical properties, making the modified epoxies better suited for cryogenic applications.

Figure 28d depicts variations in microcrack density in CF reinforced epoxy laminates. The results of the laminate micro-cracking in CF reinforced EP laminates modified with different thermoplastics are as follows:

• -. Microcracks propagate normal to the carbon fibers along their length, beginning at the laminate’s outside border and ending at the 0°/90° ply interface. Micro-cracks appear to be wider in the neat EP laminate than in the thermoplastic modified epoxy laminates.

• -. PBT-induced toughening improved the laminates’ microcrack resistance slightly.

• -. The PC modified epoxy demonstrated excellent bonding with the EP composite, resulting in increased energy absorption during crack propagation. The dispersed PC reduced the crack density by approximately 37%.

• -. A phase-separated morphology is visible in the PEI modified epoxy, indicating effective toughening and reduced crack density by about 37%.

These observations suggest that PEI and PC are effective toughening system for EP, enhancing micro-crack resistance in CF reinforced EP laminates during cryogenic thermal cycling (Figure 29). In conclusion, this study demonstrated that thermoplastic modification of epoxy at CT enhanced the impact strength, reduced CTE, and improved micro-crack resistance in CF reinforced epoxy laminates, with PEI and PC identified as effective toughening materials for epoxy in cryogenic applications.

4.1.4. Polyurethanes

Qu et al. [46] investigate the use of hydroxyl-terminated polyurethane to toughen epoxy resin, resulting in improved mechanical strength at both room and cryogenic temperatures. The findings emphasize the importance of interfacial adhesion between the matrix and fibers in achieving high tensile properties. Figure 30 depicted the manufacturing procedure of CFRP composites. Several important stages are involved in the manufacture of CFRP composites to guarantee that the CFs interface properly with the EP matrix. The following steps were taken to fabrication process for the study:

• -. The epoxy matrix is formed by mixing the EP and the curing agent.

• -. To toughen the matrix, the epoxy resin is mixed with hydroxyl-terminated polyurethane (HTPU1 or HTPU2).

• -. The EP matrix is applied to the CFs using a hand paste method.

• -. CFRP composites are produced in unidirectional [0°] and [90°/0°/90°] stacking sequence designs.

• -. The matrices are cured using a hot-press method.

• -. The curing process begins with 90 min of curing at 80 °C, followed by another 90 min at 140 °C.

• -. CFRP matrices are post-cured at 140 degrees Celsius for 90 min under a pressure of 5 MPa.

The contact angle between CFs and modified EP can provide insight on the matrix’s wettability on the fibers’ surfaces. Sample microscopy images of neat EP, EP/HTPU1, and EP/HTPU2 droplets produced on the surface of a single CF were presented in Figure 31. The contact angle values obtained from the measurement process are:

• -. The contact angles were determined using the micro-Wilhelmy method.

• -. The average contact angles were 34.3 ± 2.1° for CF-EP, 31.6 ± 1.8° for CF-EP/HTPU1, and 33.9 ± 1.9° for CF-EP/HTPU2.

• -. The results showed that the presence of HTPU1 and HTPU2 slightly decreased the average contact angle compared to neat EP.

• -. The decreased contact angle could be due to the lower viscosity of the modified epoxy with HTPU, as well as the chemical similarity between EP/HTPU and the sizing agent on CF surfaces, which improves wettability.

The study tested the tensile properties of CFRP materials, both transversely and longitudinally. The important findings of the tensile properties of CFRP composites are as follows:

• -. The transverse (90°) tensile stress–strain graphs of CFRP and CF/EP/HTPU materials were studied at RT and 77 K.

• -. The stress–strain curves showed linear elastic behavior until failure due to interfacial debonding between CFs and matrix.

• -. The transverse tensile strength and fracture strain of CF/EP/HTPU materials were compared to pure CFRP composites.

• -. The results showed that the CF/EP/HTPU composites had higher tensile strength and fracture strain at both RT and 77 K compared to the neat composites.

• -. The longitudinal tensile properties of the CF/EP/HTPU materials were also assessed.

• -. The addition of HTPU1 and HTPU2 improved the composites’ tensile strength and fracture strain at both RT and 77 K.

• -. The CF/EP/HTPU2 composite has the highest tensile strength of any composite at both temperatures due to its unique failure process involving rubber-like particles and cavitations.

• -. In contrast to the explosive fracture mode at RT (Figure 32a), the CFRP materials’ cryogenic fracture modes (Figure 32b) included splitting, partial CF fracture, and edge delamination. As a result, CFs’ tensile strength was not fully utilized at CT. As a result, the tensile strength and fracture strain of CF/EP/HTPU materials will significantly decrease at 77 K [47]. CF/EP/HTPU2 also had a higher longitudinal tensile strength than CF/EP/HTPU1 at both RT and 77 K.

• -. The tensile modulus of CF/EP/HTPU1 was slightly greater than that for standard CFRP composites, whereas CF/EP/HTPU2 had a minor decrease.

The study assessed the composites’ flexural properties, such as flexural strengths, and flexural moduli. The important findings of the flexural properties of CFRP composites are as follows:

• -. The CF/EP/HTPU1 system‘s flexural strength rose approximately 15% at RT in comparison to the pristine system.

• -. The CF/EP/HTPU2 system‘s flexural strength dropped by around 6% at RT.

• -. The CF/EP/HTPU1 and CF/EP/HTPU2 composites’ cryogenic flexural strengths were marginally higher than those of pristine CFRP systems, which enhanced the adherence of the matrix to the CF interfacial.

• -. Compressive fracture in the top part of the area was the main cause of the virgin CFRP and CF/EP/HTPU systems’ failure at RT, as seen in Figure 32c. The tighter CF/EP interface resulting from the thermal contraction of the EP matrix led both tensile and compressive fractures to develop simultaneously inside the upper and lower half-regions at 77 K, as seen in Figure 32d. The two main failure types found were delamination and fiber fracture, indicating that delamination needed more energy [48,49].

• -. The flexural moduli of the CF/EP/HTPU1 and CF/EP/HTPU2 matrices increased slightly when compared to neat CFRP matrices.

• -. The improvement in flexural modulus was minimal, demonstrating that toughening with HTPU had little to no effect on the composites’ flexural stiffness.

The main changes in tensile strength between EP/HTPU1 and CF/EP/HTPU2 composites are as follows:

• -. The EP/HTPU1 composite has a better tensile strength than the CF/EP/HTPU2 system at both RT and 77 K.

• -. The tensile strength improvement in EP/HTPU1 composite was attributed to HTPU1′s toughening impact on the EP matrix, which resulted in improved interfacial bonding with CFs.

• -. The CF/EP/HTPU2 system has a lower tensile strength than the EP/HTPU1 composite at both RT and 77 K.

• -. The increased tensile strength in the CF/EP/HTPU2 composite was attributed to a unique failure mode involving rubber-like particles and cavitations, which contributed to the composite’s overall mechanical properties.

In conclusion, including HTPU1 and HTPU2 into the epoxy matrix of CFRP composites shows promise for increasing mechanical properties, particularly tensile and flexural behavior, making them ideal for a variety of high-performance applications.

4.1.5. Polysiloxanes

Li et al. [50] investigate the effect of a flexible polymer containing EP groups (EPSE) on the mechanical properties of EP and laminate at cryogenic temperatures. The flexible polymer with EP groups was synthesized using a co-hydrolytic condensation reaction of dimethyldimethoxysilane (DMDMS) and KH560 (Figure 33a). The specific steps in the synthesis process are as follows:

• -. DMDMS (0.18 mol) was added to a round-bottomed flask and swirled at 60 °C.

• -. Dropwise additions of dilute HCl (2 mL) were made, and the reaction was held at 60 °C for 3.0 h.

• -. After adding KH560 (0.18 mol), diluted HCl (22 mL) was added dropwise.

• -. At 60 °C, the reaction continued for approximately ten hours.

• -. To obtain the product EPSE, H₂O and CH₃OH (a by-product) were removed for approximately two hours at 80 °C under decreased pressure.

To fabricate the CF/EPSE-EP matrix laminate, the VARTM method required several processes (Figure 33b). The following are the specific steps:

• -. EPSE-epoxy resin (8 wt%) was introduced into the vacuum bag via the inlet.

• -. The resin was applied uniformly throughout the bag to impregnate the carbon fabrics.

• -. The composite laminate was cured at prescribed temperatures.

• -. Temperatures ranged from 100 °C for 2 h to 160 °C for 4 h.

The main findings concerning the impact of the flexible molecular chain on the cryogenic mechanical properties of EP and laminate are as follows:

• -. The addition of EPSE significantly increased the EP matrix’s failure strain at both RT and 77 K.

• -. At RT, an EPSE concentration of 8 wt% improved the epoxy matrix’s failure strain by 69.4%.

• -. The EP matrix’s fracture toughness (K_IC) value at RT and 77 K increased by 10.4% with the appropriate amount of EPSE.

• -. The flexible EPSE increased crack resistance and released more fracture energy in the EP matrix.

• -. The CF/EPSE-EP laminate has a higher flexural strength than the CF/neat EP laminate at 77 K.

• -. EPSE’s flexible molecular chains reduced thermal stress at the CF/matrix contact, resulting in increased laminate flexural strength at cryogenic temperatures.

• -. The flexural strength of the CF/EPSE-EP matrix laminate reached its peak as the number of thermal cycles increased.

• -. Thermal cycling caused interface debonding and microcracks, which contributed to changes in the composite laminate’s flexural strength.

These findings show that the flexible molecular chain in EPSE improves the mechanical properties of both the EP matrix and the CF/EP composite laminate, especially under cryogenic conditions, as evidenced by specific numerical improvements in flexural strength, fracture toughness, and failure strain.

The SEM images of the fracture surfaces of neat epoxy resin and EPSE-epoxy resin at different temperatures, as well as the optical micrographs of CF/pristine EP laminate and CF/EPSE-EP laminate after flexural tests at RT and 77 K are shown in Figure 34, provide valuable insights into the materials’ microstructural changes and failure modes:

• -. (a) RT: The SEM image at RT shows a comparatively smooth fracture surface for the neat epoxy resin, showing a low barrier to crack progression.

• -. (b) 77 K: At 77 K, the SEM picture shows a few crack propagation streaks, indicating some resistance to crack formation at CT.

• -. (c) RT: At RT, the SEM picture displays a rough fracture surface with denser crack propagation streaks at 45°, indicating increased crack development resistance and a more ductile failure mode.

• -. (d) 77 K: The SEM image at 77 K shows a relatively rough fracture surface with denser crack propagation streaks, indicating enhanced crack propagation resistance and higher energy dissipation than neat epoxy at CT.

• -. (e) RT: The optical micrograph following the flexural test at RT reveals matrix debonding and fiber shear breaking, showing the laminate’s failure mode under flexural load at RT.

• -. (f) At 77 K, the optical micrograph shows delamination, indicating a layer-by-layer breakdown pattern with higher thermal stress at CT.

• -. (g) RT: The optical micrograph after the flexural test at RT reveals a failure pattern comparable to the CF/neat epoxy laminate, with matrix debonding and fiber shear breaking.

• -. (h) 77 K: At 77 K, delamination is also seen in the optical micrograph of the CF/EPSE-EP laminate, indicating a layer-by-layer failure mode similar to the neat epoxy laminate at CT.

In conclusion, this study highlighted the potential of EPSE as a modifier to enhance the mechanical performance of CF/epoxy composite laminates, including increased flexibility, fracture toughness, and resistance to thermal cycling effects at both RT and CT.

4.1.6. Hyperbranched Polymers

Zhao et al. [44] investigate an efficient method for hydroxy-terminating hyperbranched polyester (HPB)-modified EP at both RT and CT. The effect of free volume on the cryogenic mechanical properties of EP reinforced with hyperbranched polymers was investigated. Figure 35 depicts the manufacturing process for HPB/EP matrices. The preparation process of HPB/EP matrices involved several steps, as follows:

• -. The DGEBA (DER332)/methyl nadic anhydride (MNA) (DER332/MNA) mixture was treated with hyperbranched polymer (H204).

• -. The mixture was swirled for two hours at 120 °C.

• -. The content of polyester H204 in the mixture varied from zero to 10% in weight.

• -. The H204/EP mixtures were treated with accelerator (K54) (1 wt%), an accelerator.

• -. After removing bubbles with a vacuum rotary pump, the mixture was poured into a preheated mold set to 85 °C.

• -. The samples were cured for 16 h at 80 °C, followed by 4 h at 100 °C and 6 h at 125 °C.

The mechanical properties of the HPB/EP matrices are summarized below:

• -. At RT, tensile strength reduced with the incorporation of 2.5 wt% HPB but remained constant as HPB content increased.

• -. At 77 K, tensile strength increased with HPB content, peaking at 74.37 MPa with 7.5 wt% HPB, which is 29.9% higher than that of neat EP. At 77 K, thermal shrinkage decreases the free volume of hyperbranched HPB in the matrix, but still remains (Figure 36).

• -. The use of HPB decreased the crosslinking density of the matrices, influencing their tensile strength at RT.

• -. At RT, the tensile modulus did not change significantly with HPB content, whereas the cryogenic modulus increased at 77 K due to increased intermolecular adhesion.

• -. The composites’ impact strength increased by 41.8% at RT and 59.4% at 77 K as compared to pure epoxy resin.

• -. The cryogenic impact strength of the composites was greater than that of epoxy resin, showing increased toughness at low temperatures.

• -. At 77 K, the flexible ester group in HPB is not fully frozen [51,52]. The impact energy can still be used by segmental movement, and the smaller free volume at 77 K [51,52] can still absorb a large amount of energy. Both of these properties will improve the cryogenic impact strength of HPB/epoxy composites. At 77 K, the free volume remains, albeit smaller (see Figure 36).

• -. The compression strength correlated negatively with the size of the free volume hole, demonstrating that larger free volume holes were associated to decreased compression strength.

• -. The composites’ mechanical properties were linked to their relative fractional free volume.

• -. At 77 K, the impact strength, tensile strength, and tensile modulus all showed a positive correlation with the relative fraction of free volume.

• -. Increasing the relative fraction of free volume improved the HPB/EP composites’ impact and tensile properties.

These results show that HPB modification has a considerable impact on the mechanical properties of EP matrices, with improvements in tensile strength, impact strength, and fracture behavior reported at both CT and CT. Overall, this research demonstrates the effectiveness of HPB modification in enhancing the mechanical properties of EP matrices at CT and highlights the importance of microstructural considerations in optimizing material performance for low-temperature applications.

4.1.7. Block Copolymers (BCP)

Wang et al. [53] explored innovative ways to enhance the performance of CFRP tanks for storing cryogenic liquids like hydrogen. The study explored the effectiveness of low-modulus soft NPs in solving the microcracking challenges inherent in CFRP at cryogenic temperatures. By integrating a BCP into an EP, nano-structured fillers with an average diameter of about 100 nm can be produced. Figure 37 represents the schematic diagram for the preparation of BCP/epoxy nanocomposites and manufacturing of CFRP composites. The preparation of BCP/epoxy nanocomposites involves several steps to ensure the uniform dispersion of BCP within the epoxy resin, as follows:

• -. The BCP is first dried in an oven at a specific temperature (e.g., 110 °C) for a certain duration (e.g., 3 h) to remove any moisture and ensure optimal performance during blending.

• -. The dried BCP is blended into the epoxy resin at different concentrations (e.g., 1.0 wt%, 2.5 wt%, 5.0 wt%). The blending process involves constant stirring of the mixture at an elevated temperature (e.g., 90 °C) for a specified duration (e.g., 12 h) to achieve complete dissolution and uniform dispersion of the BCP within the epoxy matrix.

• -. After the stirring process, the mixture may undergo multiple passes of calendaring to further enhance the dispersion of BCP within the epoxy resin. Calendaring helps improve the distribution of BCP particles and promotes better interfacial interaction between the BCP and EP components.

• -. Once the BCP is well dispersed in the epoxy resin, a curing agent is added to the mixture at the recommended resin-to-hardener ratio (e.g., 100:30). The curing agent initiates the crosslinking reaction in the epoxy resin, leading to the formation of a cured network structure.

• -. The BCP/EP blend is subjected to vacuum degassing at RT for a specific duration (e.g., 30 min) to remove any entrapped air bubbles. Vacuum degassing helps improve the quality of the nanocomposite by eliminating voids and ensuring proper impregnation of BCP within the epoxy matrix.

• -. The BCP/EP blend is cured following a specific curing process recommended by the supplier. This curing process typically involves an initial curing stage at a certain temperature (e.g., 25 °C) for a specified duration, followed by post-curing at a higher temperature (e.g., 120 °C) for a defined period (e.g., 14 h) to achieve the desired mechanical properties and performance of the nanocomposite.

• -. Multiple layers of unidirectional (UD) dry carbon fiber sheets are stacked onto a mold or substrate.

• -. The BCP-modified epoxy resin is infused into the carbon fiber layers using vacuum assistance for proper impregnation and consolidation.

• -. The composite laminate is cured under specific temperature and pressure conditions to achieve desired mechanical properties and microstructure.

The results regarding the fracture toughness of BCP-modified epoxy are as follows:

• -. According to the experimental findings, adding a BCP to an EP matrix considerably raises the matrix’s fracture energy by 392% at CT and 563% at ambient temperature. Even in configurations that are prone to microcracking, material laminates created with the BCP-modified may survive CT without matrix microcracking. The enhanced fracture resistance of the material is responsible for the toughening effect of the nanostructured block copolymer at CT.

• -. The production of free volume in the EP matrix, which reduces brittleness, is one of the toughening processes of BCP at CT. Higher toughness at CT is exhibited by BCP components such polybutadiene and polymethylmethacrylate (PMMA), which adds to the improved toughening effectiveness of BCP.

• -. Figure 38 shows a graph that contrasts the increase in fracture toughness at CT with RT for both rigid along with soft fillers. In general, rigid particles are proven to have a stronger toughness improvement at CT than at RT. This implies that stiff fillers have a stronger toughening impact in colder climates. Conversely, soft particles perform less well at toughening at CT than at RT. This could be attributed to the reduced effectiveness of the toughening mechanisms associated with soft fillers in extreme cold conditions. While high-stiffness fillers like nanosilica, graphene, and CNTs have been reported to toughen EP at CT, low-stiffness, soft particles may lose their toughening effectiveness at such low temperatures. BCP is highlighted as a promising toughening agent in epoxies, showing effectiveness at RT and demonstrating improved toughening efficiency at CT compared to other soft particles. The discussion highlights the contrasting behavior of rigid and soft fillers in enhancing fracture toughness at different temperatures. While rigid fillers show better performance at CT, soft fillers are more effective at RT. Understanding the temperature-dependent toughening behavior of fillers is crucial for optimizing the mechanical properties of nanocomposites for specific applications that involve varying temperature conditions.

• -. The fracture energy of BCP-modified epoxy shows a significant increase at −196 °C, with the BCP concentration playing a crucial role in enhancing toughness. Figure 39 depicts the fracture surfaces of CFRP with both a pure and BCP-modified composite. The unmodified epoxy and carbon fiber resulted in evident debonding, as illustrated in Figure 39a,d [54]. The smooth EP phase is found in the space between the fibers. In a similar vein, Figure 39b,e’s laminate with a BCP-modified system displayed typical fiber–matrix debonding. The fibers encountered cohesive fracture failure on the surface when they were wrapped in the modified EP system [55]. Figure 39c–f shows how the fracture surfaces of polymers are different from those of plain EP, with BCP-rich phases evidently scattered throughout. The more ductile fracture model of the laminates created with BCP-modified EP is seen in Figure 39c, where BCP micelles debond and develop voids.

• -. The BCP was unable to affect the interlaminar fracture energy of the matrix, but it was possible that they may effectively inhibit through-thickness microcracking at CT because of their ability to significantly increase system fracture energy. This is easily seen in Figure 40, which shows transverse microcracking due to thermally generated stress in the 90° inner layer of [04/904]s laminates with a modified EP system consisting of different BCP loadings, following 45 min of liquid nitrogen cooling. As seen in Figure 40a,b,d, through-thickness microcracks are clearly visible in the unmodified laminate as well as laminates with 1.0 wt% along with 5.0 wt% BCP. As seen in Figure 40c, through-thickness fractures were absent from laminates having an EP matrix containing 2.5 weight percent BCP. Agglomeration happens when the BCP concentration is higher than 2.5 weight percent, such at 5.0 weight percent, because the microcracks diverge. The aggregation process has left some parts of the matrix with either very little or no BCP. These matrix sections therefore continue to be untoughened or just slightly toughened. Consequently, at liquid nitrogen temperatures, the ideal BCP loading is 2.5 weight percent, which successfully reduces matrix microcracking.

4.2. Nanomaterials

4.2.1. CNTs and MWCNTs Nanoparticles

CNTs along with MWCNTs are the most widely utilized nanotubes [56]. Their interactions with the resin are improved by these nanotubes. As the modified resins cool, their surfaces become layered along with rough, indicating that the nanofillers enhance interfacial adhesion as well as resin dispersion. These systems successfully stop micron-sized fractures from spreading. Moreover, the thick and rough fracture surface of the MWCNT-modified resin suggests that the nanotubes influence the direction of crack propagation [57].

Preprocessing the nanoparticle material increases its bonding strength and compatibility with the modified filler and resin. To generate covalent connections with epoxy resin, CNTs and MWCNTs can be treated with amino functionalization on their surfaces. This treatment improves nanoparticle dispersion and interfacial bonding in CFRPs. Furthermore, GO can be modified with Fe₃O₄ magnetic particles to increase the interfacial area. Magnetized GO particles may be effective in preventing microcracks from spreading along the stress direction. Feng et al. [58] investigate the effects of carbon nanotube content on the transverse tensile strength of CF/EP composites at cryogenic temperatures. The sample production procedure is presented in Figure 41. As shown in Figure 41, a Planetary Ball Mill is used to manufacture MWCNT and DGEBA resin mixtures. The following steps were taken to prepare samples for the study:

• -. The preparation of MWCNT/EP blends involved mixing MWCNTs into the DGEBA epoxy resin using a Planetary Ball Mill. Various MWCNT contents (0, 0.2, 0.5, 0.8, and 1.0 wt%) were blended with the EP at 60 °C for 6 h.

• -. The preparation of CF bundles involved treating T-300 carbon fibers (CFs) with acetone to remove the sizing agent before vacuum drying, after which the dry CF bundles were placed in molds to undergo transverse tensile testing.

• -. The sample fabrication process involved making CF-containing molds by fixing the CFs before adding the mixtures (Figure 42). The specimens were then cured at 80 °C for 6 h, followed by 5 h of post-curing at 130 °C, and subsequently cooled to RT.

These stages ensured the appropriate blending of MWCNTs with the epoxy resin, the creation of CF bundles, and the manufacture of samples for further testing to assess the interfacial binding properties between CFs and the EP matrix at cryogenic temperatures.

Figure 43 depicts a schematic of the study’s contact angle measurement technique, which includes both a vertical and a side view. The contact angle measurement procedure is described based on the accompanying diagram as follows:

• -. In the vertical view (Figure 43a), a bundle of CFs is connected to both ends of a slide glass with adhesive to ensure the CFs remain straight. A drop of the MWCNT/EP mixture is carefully squeezed out of a syringe and applied to the CF surface, causing the liquid drop to spread over the CF surface and form a partly spherical droplet due to surface tension.

• -. In the side view (Figure 43b), a different perspective on the contact angle measurement technique is provided. The contact angle is calculated using image analysis software (CAST 2.0), which employs image recognition technology to identify key locations on the droplet, including the right-end point, vertex, and left-end point. The droplet’s width (2r) and height (h) are determined from the coordinates of these points, and the contact angle (θ) is calculated using the formula: θ = 2 arctan (h/r).

Figure 44 of this work shows the contact angles between the MWCNT/EP composite and the T-300 CF at RT for various MWCNT amounts. The contact angle values obtained from the measurement process are:

• -. The contact angle measurements show that the average contact angles for 0 wt%, 0.2 wt%, 0.5 wt%, 0.8 wt%, and 1.0 wt% MWCNTs are approximately 76.40°, 69.29°, 63.34°, 66.45°, and 67.17°, respectively.

• -. The contact angle trend shows an initial decrease with the addition of MWCNTs, reaching a minimum at 0.5 wt% MWCNT content, followed by a gradual increase as the MWCNT content increases. This change in contact angle is attributed to the dispersion and viscosity effects of MWCNTs in the epoxy (EP) matrix.

The cryogenic interfacial normal bond property, as determined by transverse tensile strength measurements at RT and 77 K for various MWCNT contents in the epoxy matrix, are summarized below:

• -. 0 wt% MWCNTs have a transverse tensile strength of 39.3 MPa at RT and 48.9 MPa at 77 K.

• -. Adding 0.5 wt% MWCNTs increases transverse tensile strength by 29.3% to 50.8 MPa at RT and 51.7% to 74.2 MPa at 77 K.

• -. Transverse tensile strength decreases at both RT and 77 K as the MWCNT content beyond 0.5 wt%.

• -. The transverse tensile strength follows a similar trend to the contact angle measurements as the MWCNT content increases.

• -. The greatest improvement in transverse tensile strength is shown at 0.5 wt% MWCNTs, demonstrating better interfacial adhesion at this concentration.

• -. MWCNTs at an appropriate content improve the cryogenic interfacial normal bond property between CFs and the EP matrix.

These findings indicate the remarkable improvement in cryogenic interfacial normal bond properties between CFs and EP matrix produced by adding MWCNTs at the optimum concentration. The increase in transverse tensile strength at both RT and CT demonstrates the usefulness of MWCNTs in improving interfacial adhesion and overall mechanical properties of the composite. The study’s morphological characterization of fracture surfaces provides insight into the adherence of fibers to matrix in fiber-reinforced composites (Figure 45). The morphological analysis of fracture surfaces are as follows:

• -. At CT (77 K), all CF/EP composites exhibit brittle fracture behavior, resulting in smooth fracture surfaces.

• -. Before introducing CNTs into the EP matrix, a little amount of matrix is connected to the CF surfaces.

• -. After introducing MWCNTs to the epoxy resin, matrix fragments were observed to adhere to the fiber surfaces.

• -. At 0.5 wt% MWCNT content, the greatest quantity of matrix is bound to the CF surfaces, resulting in improved transverse tensile strength in the composite.

• -. When the MWCNT content is increased to 1.0 wt%, the viscosity of the CNT/EP mixture significantly increases, making uniform MWCNT dispersion in the EP matrix difficult. This increases CNT aggregates and stress concentrations, resulting in smoother fracture surfaces and decreased interfacial adhesion strength.

Overall, the study shows that including MWCNTs into EP greatly improves the mechanical properties and interfacial bond strength of CF/EP composites, especially at cryogenic temperatures. The findings stress the relevance of adequate MWCNT content in improving composite material performance.

He et al. [59] investigated how functionalized MWCNTs distributed in the material affect the mechanical properties of CF/EP laminate at 77 K. The techniques include specific procedures that ensure proper dispersion and alignment of the carbon nanotubes in the EP matrix, resulting in enhanced mechanical properties of the matrices and laminates. The preparation processes are as follows:

• -. A flask is filled with raw MWCNTs and a specified ratio of H₂SO₄ and HNO₃ acid. The combination is heated and maintained at a specific temperature for a set time. The oxidized MWCNTs are separated, rinsed until neutral pH, and dried. O-MWCNTs are dispersed in deionized water (DI) water and purged with nitrogen. A solution of FeCl₃·6H₂O and FeSO₄·7H₂O in DI water is added to the flask. The mixture is transferred to a water bath for more processing (Figure 46a).

• -. R-MWCNTs (also known as O-MWCNTs or Fe₃O₄/O-MWCNTs) are vacuum-mixed and sonicated with EP. A curing agent is added to produce a homogenous system (Figure 46b). The mixture is cast into a mold and let to cure at RT before being post-cured in an oven.

• -. Laminates are prepared using the VARTM technique. During crosslinking, the sample is placed in a magnetic field that remains constant. The post-curing process is done in an oven (Figure 46c).

The mechanical properties results demonstrate the enhancement in Young’s modulus, tensile strength, and impact strength of the epoxy composites with the addition of different fillers, particularly the Fe₃O₄/O-MWCNTs composite showing the highest mechanical properties among the tested samples. Figure 47 illustrates the crack propagation and suppression behavior of CNT-modified CF/EP matrices:

• -. Describes the behavior of a CF/EP laminate without any CNT modification (Figure 47a).

• -. Demonstrates fracture propagation characteristics in a CF/EP laminate using R-MWCNTs. The weak interaction between R-MWCNTs and the epoxy matrix causes cracks to deflect toward the R-MWCNT interface. Crack deflection improves energy absorption and fracture toughness (Figure 47b).

• -. Shows how cracks propagate in a CF/EP laminate that has been modified using O-MWCNTs. The incorporation of -OH and -COOH groups to the surface of O-MWCNTs promotes covalent interaction with the EP matrix, increasing load transfer capabilities and causing micro-cracks to zigzag in the laminates (Figure 47c).

• -. Illustrates the crack propagation behavior in a CF/EP laminate modified with Fe₃O₄/O-MWCNTs. The oriented Fe₃O₄/O-MWCNTs prevent matrix cracks from extending along the force direction, leading to improved mechanical properties and crack suppression (Figure 47d).

In conclusion, this study demonstrates the significant impact of MWCNT modification on the mechanical properties and micro-crack resistance of EP and CF/EP materials, highlighting the potential for advanced applications in various industries.

4.2.2. GO Nanoparticles

GO is another effective resin modifier because of its greater specific surface area [60]. The GO can efficiently transmit external stress and create a micro deformation zone in the matrix, enabling it to absorb more impact energy when a third-party force damages the matrix and widens cracks. The matrix will get harder as a result. At cryogenic temperatures, this material inhibits the development of cracks by improving interactions at the matrix-fiber interface. The effects of GO on the mechanical properties of CF-reinforced EP composites at both RT as well as CT are examined by Qu et al. [61]. The process of creating CF/GO-modified EP composites is shown in Figure 48. Several stages were involved in the synthesis of CF/GO-modified EP materials, as follows:

• -. The CFs were refluxed in acetone at 80 °C for 72 h in order to eliminate impurities. The CFs were sonicated in deionized water for an hour in order to remove any leftover sizing agent. After that, they were washed multiple times with deionized water and dried in a vacuum for five hours at 90 °C.

• -. Using a hand paste technique, the treated CFs were covered with the GO-modified EP matrix. For further processing, eight impregnated CF layers were placed in a mold in a [0°] configuration.

• -. For 1.5 h, the stacked CF/GO-modified EP composite was allowed to cure at 80 °C. After that, it was pressure-cured for 1.5 h at 140 °C (5 MPa). In order to guarantee a thorough chemical reaction among the EP including the curing agent and the creation of a fully cross-linked matrix, the materials were lastly post-cured for ten hours at 130 °C.

• -. To examine their impacts on mechanical characteristics, different GO amounts (0.0, 0.05, 0.1, 0.2, 0.35, along with 0.5 wt%) were incorporated to the EP matrix of CF/EP composites.

• -. The ASTM D 7264 standards [61] were followed in cutting the samples from matrix plates for the three-point bending mechanical testing such that they were around 80 mm long, 12.5 mm wide, and 2 mm thick.

Figure 49 presents the design schematics of the RT and cryogenic evaluation systems for three-point bending tests, which involve testing at 77 K and sampling in liquid nitrogen. The mechanical properties of CFRP materials, specifically their flexural properties and ILSS, were tested under a variety of conditions. The mechanical property findings are described as follows:

• -. The flexural characteristics of CFRP composites at RT were linearly elastic until the maximum load was reached, at which point they failed with a jagged shape.

• -. The addition of GO to the EP matrix altered the fracture pattern and offered higher energy dissipation than GO-free CFRP materials at RT.

• -. At 77 K, the flexural characteristics of CFRP composites showed a significant decrease in load, followed by catastrophic failure.

• -. The maximum flexural load and displacement at 77 K were much higher than at RT, with CF/0.2 GO-modified EP having a maximum load and displacement that was double that of RT.

• -. Flexural strength and modulus at 77 K increased with GO content, achieving maximum values at 0.2 wt% GO content.

• -. At both RT and 77 K, the ILSS of CFPR materials peaked at 0.2 weight percent GO content, having originally risen with GO concentration.

• -. At RT and 77 K, the inclusion of GO to the EP matrix increased the ILSS by around 17.6% and 8.7%, respectively.

• -. The formation of H-bonds between GO, EP matrix, and CF resulted in a rise in ILSS, which indicated increased interfacial bonding among the EP matrix as well as CFs.

These results demonstrate that the mechanical characteristics of CFRP materials are significantly affected by GO modification, with enhancements observed in the interlaminar shear strength and flexural properties at both RT and CT. After three-point bending tests at RT and 77 K, the macroscopic failure morphologies of the CFRP composites are shown in a side view in Figure 50. The macroscopic morphologies are as follows:

• -. At RT, the side view of the macroscopic failure morphology of the CFRP samples revealed delamination and compressive fracture mostly in the top layers of the CF/EP composite.

• -. The CF/EP matrix showed compressive fracture with some delamination at the top layers, indicating a failure mechanism under bending stress at RT.

• -. At 77 K, all composites tested experienced both compressive (top layer) and tensile (bottom layer) fractures.

• -. Compressive and tensile fractures, as well as delamination and fiber fracturing, were observed in CFRP samples at 77 K testing conditions.

• -. At 77 K, the CF/0.2 GO-modified EP composite showed less delamination and more severe fiber breaking in the tensile fracture region than other composites, demonstrating that GO is beneficial in preventing debonding between CFs and EP matrix under CT.

These findings highlight the differences in failure morphologies between RT and CT testing for CFRP composites, with the presence of GO influencing the composites’ fracture behavior and interfacial adhesion under different temperature conditions. Figure 51 depicts a simplified model of GO dispersion in an epoxy matrix, which helps to explain how GO improves the mechanical properties of CFRP materials. At a GO concentration of 0.05 wt%, the schematic design illustrates nearly flawless GO dispersion in the epoxy matrix. Nonetheless, the effect of GO on the epoxy matrix is restricted due to the low GO content at this concentration (Figure 51a). With a 0.1 wt% GO concentration, the diagram shows excellent GO dispersion in the EP matrix. The addition of GO at this concentration improves the properties of the epoxy matrix (Figure 51b). At a 0.2 wt% GO concentration, the diagram depicts a few small GO aggregates in the EP matrix. Because of the increased GO content, this concentration has the highest enhancement efficiency (Figure 51c). However, at a 0.5 wt% GO concentration, the diagram shows varying sizes of GO aggregates in the EP matrix [60] (Figure 51d). The presence of GO aggregates of varying sizes at this concentration may lead to micro-sized cracks to form at the GO aggregate locations and spread throughout the EP matrix. The aggregation of GO at this concentration may result in poor stress concentration and stress transfer efficiency, ultimately affecting the CFRP composite’s flexural strength. In conclusion, incorporating GO into the epoxy matrix of CFRP composites is a viable strategy for improving their mechanical properties, with implications for applications needing high-performance materials under variable temperature conditions. The study provides insight on how nanofillers like GO can improve the performance of composite materials.

Hung et al. [62] successfully produced GO-containing CFRP composites, with GO contributing to improved mechanical properties, particularly at cryogenic temperatures. Two different methods of using GO into CFRP materials are:

• -. This process involves dispersing GO in acetone and then mixing it with EP to form a GO-reinforced EP system.

• -. The dispersion of GO in the material improves the mechanical properties of the composite by enhancing stress distribution and interfacial adhesion between the matrix and carbon woven fabric (CWF).

• -. This approach is successful at improving the overall mechanical properties of CFRP matrices, particularly their strength and stiffness.

• -. Electrophoretic deposition (EPD) is utilized to apply GO directly to the surface of CWF.

• -. Coating GO onto the CWF surface enhances interfacial adhesion between the reinforcement and matrix, resulting in better transfer of load and mechanical properties for the composites.

• -. This approach guarantees that GO is distributed uniformly across the surface of the CWF, which can help to improve mechanical performance.

• -. GO-CWF method tends to be more in improving the mechanical characteristics of CFRP materials compared to the GO-EP method.

Figure 52 depicts the production process schematically. The steps for preparing GO-containing CFRP composites are as follows:

• -. The CWF is first processed by refluxing it in acetone at 80 °C for 72 h to eliminate any impurities.

• -. It is then rinsed with deionized water and dried in a vacuum oven at 80 °C for 3 h to remove the sizing agent, yielding de-sized CWF.

• -. To ensure uniform dispersion, GO is dispersed in acetone followed by 3 h of sonication.

• -. The GO-acetone mixture is then combined with epoxy resin and agitated at 60 °C for 6 h to eliminate excess acetone.

• -. The mixture is then agitated at 25 °C for 3 h to produce a homogeneous GO-reinforced EP mixture known as GO-EP.

• -. EPD is utilized to coat GO on the surface of CWF.

• -. The cathode is an aluminum plate, while the anode is a piece of CWF.

• -. The EPD procedure is carried out for 15 min at a constant voltage of 10 V with sonication to effectively disperse GO in ethanol.

• -. To produce GO-coated CWF (GO-CWF), the coated CWF is dried in a vacuum oven at 50 °C for 24 h following deposition.

• -. GO-EP is infused with de-sized CWF, while GO-CWF is infused with EP via the VARIM.

• -. To achieve complete cross-linking in the matrix, the composites are pre-solidified at RT in a vacuum for 24 h before being post-cured in a vacuum oven at 60 °C for 6 h.

• -. The composites are subsequently cut into specimens for mechanical testing in accordance with ASTM D790 [62].

The mechanical properties of GO-containing CFRP materials are significantly improved compared to pristine composites. The mechanical property findings are described as follows:

• -. GO-EP/CWF systems exhibit flexural strengths of 523.67 MPa at RT and 508.09 MPa at CT.

• -. GO-CWF/EP composites show even higher flexural strength values (609.54 MPa at RT and 602.66 MPa at CT).

• -. The addition of GO enhances flexural strength, with GO-CWF/EP composites showing superior performance.

• -. Pristine composites exhibit a flexural modulus of 51.73 GPa at RT and 46.23 GPa at CT.

• -. GO-EP/CWF composites exhibit a flexural modulus of 54.4 GPa at RT and 52.16 GPa at CT.

• -. GO-CWF/EP composites exhibit the highest flexural modulus values, measuring 58.21 GPa at RT and 57.00 GPa at CT.

• -. The addition of GO to the composites increases the flexural modulus, with GO-CWF/EP composites showing the greatest improvement.

• -. GO-CWF/EP matrices outperform GO-EP/CWF matrices at RT and CT by 16.40% and 18.61%, respectively.

• -. The presence of GO on the surface of CWF improves the mechanical properties of the matrices by forming a strong bond between CWF and the EP matrix, resulting in better interfacial adhesion and stress transfer.

By examining their fracture surfaces, CFRP composites containing GO’s fracture processes may be comprehended. While Figure 53 shows the fracture surface of three distinct kinds of composite materials at RT, Figure 54 shows schematic depictions of interface failure. Since less matrix sticks to the fiber in the neat matrix, the fiber surface is relatively clean, suggesting that debonding happens fast at the interface between the CWF and matrix because of low interfacial adhesion (Figure 53a and Figure 54a). But in CFRP composites containing GO, more matrix sticks to the fiber’s surface. This phenomenon suggests that adding GO improves the fiber-epoxy interface interaction. Fiber–matrix couplings are made easier by GO, which improves the mechanical and chemical interlocking between CWF as well as matrix [63]. The failure mechanism shifts from interface debonding to matrix failure due to increased stress transfer efficiency from matrix to fiber. In addition, compared to GO-EP/CWF materials, GO-CWF/EP composites show a greater matrix attachment on the fiber surface. In GO-CWF/EP composites, the interfacial adhesion among fiber and matrix improves because GO encircles CWF in a methodical manner. Because of GO settling on the CWF’s surface, only some portions of the fiber on the CWF have strong interfacial bonding with EP in the GO-EP/CWF matrix, whereas other portions of the CWF’s fiber have no interaction with GO. Consequently, the produced composites have pristine matrix characteristics (i.e., a pristine surface devoid of matrix attachment), as seen by the yellow circles in Figure 53b and Figure 54b [64]. However, with GO-CWF/EP composite materials, it is more difficult to obtain a clean surface because more matrix sticks to the fiber’s surface. It results in a larger shift of failure from interface to matrix because of the increased contact between fiber and EP owing to complete GO coverage over the CWF surface (see Figure 53c and Figure 54c). In summary, the study highlights how effective GO is in improving the mechanical characteristics of CFRP materials, with an emphasis on how to apply GO to the surface of CWF to obtain better mechanical performance, especially at cryogenic temperatures.

He et al. [65] explored the use of Fe₃O₄/GO to enhance the micro-crack resistance of CF/EP laminates at 77 K, offering promising applications in cryogenic environments. Fe₃O₄/GO was prepared using a modified co-precipitation method. This synthesis process involved dissolving FeSO₄·7H₂O and FeCl₃·6H₂O in deionized water, then adding the solution to GO via sonication. Ammonium hydroxide was then used to adjust the pH, resulting in the formation of Fe₃O₄ on the GO surface. The resulting Fe₃O₄/GO composite demonstrated superparamagnetic behavior without remanence or coercivity. Fe₃O₄/GO had a saturation magnetization value of 54.36 emu/g, which was slightly lower than pristine Fe₃O₄ particles, which was attributed to the presence of GO. The Fe₃O₄/GO composite could be distributed in water as a stable colloidal suspension and manipulated with an external magnetic field. Figure 55a represents the schematic diagram for the synthesis of Fe₃O₄/GO. To modify the epoxy with Fe₃O₄/GO for cryogenic applications, a synthesis procedure was followed. Fe₃O₄/GO was prepared by dissolving FeSO₄·7H₂O and FeCl₃·6H₂O in deionized water, adding the solution to GO, adjusting the pH with ammonium hydroxide, and heating the mixture. The resulting Fe₃O₄/GO composite was then magnetically separated, washed, and collected. The EP matrix’s mechanical and thermal properties were improved for cryogenic applications by incorporating this Fe₃O₄/GO composite. The Fe₃O₄/GO modified CF/EP laminates were prepared through VARTM. Figure 55b illustrates the process for producing Fe₃O₄/GO/CF/EP laminate. This study demonstrated that Fe₃O₄/GO effectively improved the mechanical properties of the EP matrix at 77 K while also decreasing the matrix’s CTE. At CT, the Fe₃O₄/GO modification significantly improved the CF/EP composites’ microcrack resistance. The CTE of the Fe₃O₄/GO modified EP material was reduced by 51.6% compared to pristine EP, while the micro-crack density of the modified CF/EP material at 77 K decreased by 60.0%. These results demonstrate the effectiveness of Fe₃O₄/GO in improving the cryogenic performance of CF/EP laminates by enhancing mechanical properties and reducing micro-crack formation.

Fillers added to EP significantly improved the tensile properties of the cured EP at room temperature and 77 K. The tensile strength of GO/EP materials at RT increased from 68.7 MPa to 73.5 MPa, a 6.9% increase over pristine EP and modified EP materials. Tensile strength increased by 12.1% at 77 K, from 81.6 MPa to 91.5 MPa. The Fe₃O₄/GO modification produced even better results, with the tensile strength of Fe₃O₄/GO/EP matrices at RT rising from 68.7 MPa and 73.5 MPa to 78.4 MPa, a 14.1% increase. Tensile strength increased from 81.6 MPa, 91.5 MPa, and 96.4 MPa at 77 K, representing an 18.1% and 5.3% increase, respectively. The enhanced interfacial interaction between Fe₃O₄/GO and the EP matrix aided in stress transfer and avoided the formation of microcracks in the EP matrix, resulting in improved tensile characteristics for the modified composites.

Figure 56d depicts the CTE values for the system below T_g. The morphology of microcracks and crack density were investigated in CF/EP laminates prepared with VARTM. Micro-cracks propagated perpendicular to the CFs along their length, beginning at the laminate’s outer edge and ending at the 0°/90° ply interface, according to reflected optical microscopy images (Figure 56). The microcracks showed no significant variation based on the type of filler utilized. Furthermore, more microcracks were found on the laminate’s outside side, most likely due to the roughness of the vacuum backside, which provided more crack initiation sites. This effect has been noted in previous studies and may influence the overall micro-crack behavior of the laminates. Figure 57 is a schematic illustration of the crack propagation and suppression behavior of filler-modified CF/EP materials, illustrating the mechanism behind the improved micro-crack resistance in the presence of fillers like Fe₃O₄/GO. When CFs are covered with EP containing fillers, micro-cracks propagate through the EP matrix (Figure 57a). The incorporation of -COOH and -OH groups to the surface of Fe₃O₄/GO enhances covalent bonding with the EP matrix, improving load transfer capability and causing micro-cracks to zigzag in the laminate (Figure 57b). The specific surface area of Fe₃O₄/GO further increases with Fe₃O₄ particles, enhancing micro-crack suppression between CFs. Excellent interfacial contact between fillers and EP matrix, along with the low CTE of the fillers, leads to effective micro-crack suppression at the interface between CFs.

4.2.3. SiO₂ Nanoparticles

For EPs, SiO₂ NPs work well as toughening agents [66]. The toughness as well as cyclic fatigue resistance of the matrix are increased when SiO₂ NPs are combined with EP. SiO₂ nanoparticles significantly increase the viscosity of epoxy monomers due to their tiny particle size [67]. Bray et al. found that the toughness of the modified EP increased with the size along with concentration of SiO₂ NPs during fracture tests. As the concentration rose, the SiO₂ NPs in the EP system created local shear bands. The banding that developed in the plastic gaps of the resin strengthened the matrix [67]. Particle size has an effect on fracture toughness in the 2–42 μm microscale range [68]. Furthermore, SiO₂ NPs were added to elastic rubber to modify it [69]. The mixed EP exhibited a high Young’s modulus when the two modifiers’ proportions in the rubber/SiO₂ NPs/EP combination system were suitably matched. Rubber particle shear yielding along with SiO₂ NPs void generation may be responsible for the combination treatment’s increase in EP toughness [70]. SiO₂ NPs also effectively stop cracks from spreading, which allows more energy to drain at the fracture tip region and makes the material tougher than rubber alone. Moreover, NPs can reduce the CTE of the resin.

Li et al. [43] developed an EFPS/nano-SiO₂ composite to improve the mechanical characteristics of EP and CF-reinforced EP laminates under room temperature and cryogenic conditions. Figure 58 depicts a schematic diagram for the preparation of a CF reinforced EFPS/nano-SiO₂ epoxy laminate. Figure 59 depicts a schematic depiction of the co-hydrolytic condensation process for producing EFPS/Nano- SiO₂. The CF reinforced laminate is prepared using the following steps:

• -. By ultrasonically dispersing it in ethanol (C₂H₅OH), the nano-SiO₂ dispersion solution was created. After that, the dispersion solution was added to a 250 mL three-necked round-bottom flask that was equipped with a reflux condenser along with an electromagnetic stirrer. At the same time, 8.15 g of DMDPS, 12.03 g of DMDMS, along with 23.60 g of KH560 were added. 16 mL of a diluted HCl solution (pH ≈ 2) was added dropwise while stirring after the mixture had heated to 60 °C. The co-hydrolytic reaction of condensation was allowed to finish by keeping the resultant mixture at 60 °C for an additional 12 h. Subsequent reactions were conducted for approximately one hour at 80 °C under decreased pressure in order to remove the generated methanol and water.

In order to create EP matrices with different epoxy-functionalized polysiloxane (EFPS)/Nano-SiO₂ mass ratios, the following procedures were followed:

• -. By ultrasonically dispersing EFPS/nano-SiO₂ in C₂H₅OH at RT, a homogeneous EFPS/nano-SiO₂ solution was created.

• -. The NPEF-170 EP was then added to the C₂H₅OH solution containing the specified amounts of EFPS/nano-SiO₂, and it was rapidly stirred for 30 min at 25 °C. For thirty minutes, keep stirring at 80 °C to further disperse the nano-SiO₂ and eliminate the C₂H₅OH. Using a vacuum oven set at 80 °C, the residual ethanol in the EP mixture was extracted.

• -. The EP mixture was then mixed with a determined amount of DDM, the curing agent, and heated to 80 °C until homogeneity was reached. The mixture was put in a Teflon mold along with cured at 100 °C for two hours, 130 °C for two hours, along with 160 °C for four hours after degassing in a vacuum oven.

• -. The VARTM process was utilized to manufacture the CF-reinforced epoxy laminate.

• -. Six plies of the 50 cm by 50 cm CF fabric were laid on a hot plate that had been warmed.

• -. The carbon textiles were covered with a layer of peel ply and distribution mesh, and then the laminate was sealed with a vacuum bag.

• -. Using the vacuum assistance approach, the carbon fabric was impregnated with an EP matrix containing 5.0 weight percent EFPS/Nano-SiO₂. The resulting composite was subsequently dried at 100 °C for two hours, 130 °C for two hours, and 160 °C for four hours.

By following these procedures, the EFPS/nano-SiO₂ composite may be successfully synthesized and its nano-SiO₂ dispersion in the epoxy resin matrix can be enhanced. The mechanical qualities of EP and CFRP are enhanced by this composite, particularly in cold settings. Tensile testing were carried out in an environment chamber sprayed with liquid nitrogen to establish a cryogenic environment (Figure 60). The tensile characteristics of EFPS/Nano-SiO₂ EP matrices yielded the following findings:

• -. The greatest tensile strength of the EFPS/Nano-SiO₂ EP materials was 203.82 MPa at 90 K and 98.80 MPa at RT. These were greater than the pure EP by 9.30% as well as 17.07%, respectively (90.39 MPa at RT as well as 174.10 MPa at 90 K).

• -. The epoxy composites’ elastic modulus rose with the addition of EFPS/nano-SiO₂, increasing their stiffness and structural integrity.

• -. In terms of failure strain, the EFPS/Nano-SiO₂ EP matrices fared better than the clean EP, exhibiting increased ductility as well as deformation capacity.

• -. The creation of plastic voids and nano-SiO₂ debonding, two energy dissipation mechanisms, were responsible for the significant increase in fracture toughness of the EFPS/nano-SiO₂ EP matrices. Below are drawings and microfractured morphologies of the failure mechanisms of CF-reinforced pristine EP laminate at 90 K and EFPS/Nano-SiO₂ EP laminate:

• -. The weak interfacial connection among the CFs along with the pristine EP matrix is demonstrated by the microfracture morphology of the pristine EP laminate reinforced with CF at 90 K, which clearly shows debonding and fiber pull-out processes (Figure 61a).

• -. The EFPS/nano-SiO₂ EP laminate reinforced with CF exhibits distinct microfracture morphologies at 90 K, which may suggest enhanced interfacial bonding among the CFs as well as the EFPS/nano-SiO₂ EP matrix. Improved bonding may cause more energy to be released as a fracture spreads at the contact (Figure 61c).

• -. A weak interfacial bonding state is shown in the failure mechanism diagram for the pristine EP laminate reinforced with CF. This leads to low interfacial shear strength and quick crack propagation at the interface (Figure 61b).

• -. The failure mode schematic of the CF-reinforced EFPS/nano-SiO₂ EP laminate shows a better mechanical situation where the EFPS/nano-SiO₂ EP material resists rapid crack propagation, transmits shear stress effectively, along with enhances the interfacial bonding that exists among the EP material and the CFs (Figure 61d).

In cryogenic settings, the mechanical properties of EP and CF reinforced EP laminates are markedly enhanced by the combination of EFPS and nano-SiO₂. Several noteworthy information about the impact of this modification on the mechanical characteristics include:

• -. The EFPS results in a more uniform distribution of NPs by increasing the dispersion capacity of nano-SiO₂ particles in the EP matrix. This increased dispersion may help the composite material’s mechanical qualities.

• -. The tensile strength of epoxy resin may be improved by the use of EFPS/nano-SiO₂, especially at CT. The epoxy resin’s internal thermal stress may be released by the flexible molecular chain of EFPS, increasing the material’s tensile strength at low temperatures.

• -. The EFPS along with nano-SiO₂ particles work together to strengthen the EP when they are appropriately distributed and bonded. By using this reinforcing approach, mechanical qualities like fracture toughness along with tensile strength may improve.

• -. The epoxy resin’s increased fracture toughness can result from the debonding of nano-SiO₂ particles, which can improve energy dissipation at the fracture front. In cold settings, this process may be essential for maintaining mechanical integrity.

In conclusion, the tensile characteristics of the epoxy resin were enhanced by adding EFPS/nano-SiO₂. These improvements show how effective EFPS/nano-SiO₂ is in enhancing EP materials’ mechanical performance, particularly their tensile properties.

4.2.4. Metal Oxide Nanoparticles

Another cryogenic modifier for resin is metal oxide nanoparticles (NPs), such as Al₂O₃ [71], Fe₂O₃ [72], ZrO₂ [73], TiO₂ [74]. These metal oxides help accelerate CFRP curing because they are uniformly distributed throughout the resin. Moreover, there is now more interaction between the nanofillers and the resin. Stress transmission from the resin to the nanofillers is enhanced as a result. Because metal oxide nanoparticles increase interfacial strength and stop microcrack propagation, they can thereby enhance the cryogenic mechanical characteristics of the resin. Moreover, evenly dispersed metal oxide nanoparticles (NPs) can help stop cracks from propagating through the matrix as a result of route deviations. The resin becomes more resistant to cryogenic fractures as a result. The interfacial force among NPs as well as the matrix and EP curing may both be enhanced by adding NH₂ groups to the surface of ZrO₂. The preparation process for modified nano-ZrO₂/epoxy composites is shown in Figure 62. The tensile stress–strain curves of bisphenol F EP treated with ZrO₂ are shown in Figure 63. The immaculate EP’s tensile strength was 79.6 MPa at room temperature, and its failure strain was 4.57%. The EP matrices’ tensile strength and failure strain reached their maximums of 103.6 MPa and 6.84%, respectively, when the quantity of modified ZrO₂ NPs was increased to 3 weight percent. Compared to the pristine EP, CT enhanced both the tensile strength (166.6 MPa) and the failure strain (1.66%) by up to 26.4% as well as 21.1%, respectively [73].

Islam et al. [54] developed a novel approach for modifying the surface of nano-cupric oxide (nCuO) with polydopamine (PDA), which has a dimension of less than 50 nm and a low thermal expansion coefficient. The efficiency of this innovative approach is evaluated and compared to unmodified nCuO and nanosilica (nSiO₂). The study’s methods of manufacturing include surface functionalization and manufacture of NCs, as well as laminated NCs. The following procedures were followed:

• -. Surface functionalization of nCuO particles with a polydopamine (PDA) coating was carried out through the addition of TRIS buffer solution and dopamine monomer in water at 30 °C, when the nCuO particles were added into the solution, PDA polymerized in situ. Figure 64a depicts the schematic for PDA covering of nCuO NPs.

• -. Rubber molds were used to produce NCs with original nCuO, PDA-coated CuO NPs, and nano-silica. Figure 64b schematically depicts the dispersion of PDA-coated and original CuO NPs in EP via probe sonication. Different weight percentages of nano-silica were employed for manufacturing test specimens.

• -. Vacuum infusion was used to produce carbon-epoxy composite laminates (Figure 64c). To produce the appropriate compositions, nano-silica, original nCuO, and PDA-coated CuO NPs were mixed into the epoxy resin at varying concentrations.

• -. Dry CF plies were chopped into pieces and arranged on a glass surface in a specified fiber architecture. A precrack was produced in the laminate, and the stack was put together with bagging materials for resin infusion.

• -. Resin infusion was performed using several EPs at a certain vacuum pressure. Following curing at RT and post-curing, mechanical testing involved slicing the laminates into double cantilever beam test samples.

This study reported significant improvements in the mechanical properties of the EP NCs toughened with nSiO₂ and nCuO NPs at both RT and CT. The important findings about the mechanical properties obtained from this study are as follows:

• -. The fracture toughness of nCuO-EP matrices with varying concentrations (1, 2, 4, and 8 wt%) increased at RT and CT. CT showed a higher percentage increase than RT.

• -. The addition of PDA-coated nCuO NPs at various concentrations (1, 2, 4, and 8 wt%) increased the fracture toughness of the NCs, with a greater percentage increase at CT than at RT.

• -. PDA-coated nCuO NPs exhibited up to a 260% increase in fracture toughness, surpassing improvements reported in the literature for other NPs.

• -. The nanocomposites’ tensile strength increased at CT due to the significant compressive stress on the nanoparticle surfaces, resulting in better composite strength.

The presence of nano-silica, nCuO, or PDA-treated nCuO NPs increased the Young’s modulus of the epoxy, with the nanoparticle-containing epoxy having a greater modulus than the unmodified epoxy.

These results demonstrate the effectiveness of incorporating low-thermal expansion nanoparticles in improving the mechanical properties, fracture toughness, and strength of CF composites, especially at cryogenic temperatures, for applications such as lightweight fuel storage. Figure 65a of the study depicts the curing network of pristine EP and the improved curing network achieved by the PDA layer on nCuO NPs. The improved crosslinking network is critical for increasing the mechanical and fracture characteristics of epoxy NCs. The PDA coating on nCuO NPs promotes a stronger interaction between the NPs and the EP matrix via covalent and hydrogen bonding interactions. This improved bonding results in better stress transfer and an increase in critical energy release rate, which contributes to the observed enhancements in the mechanical properties and fracture toughness of the NCs. The schematic picture in Figure 65b likely depicts a denser and more interconnected network than the neat epoxy, demonstrating the usefulness of the PDA coating in reinforcing the crosslinking structure and improving the overall performance of the NCs. The figure illustrates three key toughening mechanisms:

• -. Nanoparticle debonding is the separation of NPs from the EP matrix under stress.

• -. This mechanism allows for energy dissipation and toughening of the material by absorbing energy during debonding.

• -. Nanoparticle debonding can result in the creation of voids and microcracks, which contribute to nanocomposites’ toughness.

• -. Plastic void growth involves the expansion of voids within the material due to applied stress.

• -. As the material deforms, voids form plastically, absorbing energy and enhancing toughness.

• -. NPs may influence void growth behavior and improve the toughening impact in NCs.

• -. Matrix shear banding is the creation of localized shear deformation zones inside an EP matrix.

• -. These shear bands help redistribute stress concentrations and prevent catastrophic failure by promoting plastic deformation.

• -. The presence of NPs can influence the initiation and propagation of shear bands, contributing to the overall toughness of the NCs.

Figure 66 shows how various toughening mechanisms, including nanoparticle debonding, plastic void formation, and matrix shear banding, work together to increase the fracture toughness and mechanical characteristics of modified EP NCs at CT. In conclusion, this study showed that employing low-thermal expansion NPs can improve the mechanical characteristics and fracture toughness of EP NCs, particularly CT. The findings have important implications for the design of new composite materials for high-performance applications in extreme temperatures.

The same group [75], has developed a new technique to lower thermal residual stresses and boost the fracture energy of CFRPs, thereby enhancing their performance in cryogenic applications. The manufacturing processes involved in the study include the preparation of bulk epoxy polymers and the fabrication of CFRP laminates are as follows:

• -. ZrW₂O₈ NPs were weighed to the desired concentration.

• -. Nanoparticles were mixed with epoxy resin using a three-roll mill at 100 rpm.

• -. Manual mixing of the epoxy/nanoparticle mixture with the curing agent (ratio of 1:0.3).

• -. Degassing the EP/NP/curing agent mixture under vacuum pressure of 0.95 atm for 30 min.

• -. Preparation of tensile test specimens and fracture energy (G_{IC, bulk}) measurements.

• -. To adhesively bond two material adherends, a 2.5 mm thick layer of unmodified EP or EP matrix containing varying concentrations of uncoated (0.5, 1, 4, and 8 wt%) and PDA-coated (0.5 and 1 wt%) ZrW₂O₈ NPs was used.

• -. CFRP laminates were manufactured using unidirectional CF fabrics.

• -. Formation of a pre-crack using a PTFE film in the laminate stack.

• -. Infusion of unmodified and modified EPs with ZrW₂O₈ nanoparticles into the fiber architecture.

• -. Initial cure at 25 °C for 12 h followed by post-cure at 120 °C for 14 h.

• -. Fabrication of blocked cross-ply laminates for studying matrix microcracking at cryogenic temperatures (−196 °C).

These processes were crucial in evaluating the effectiveness of ZrW₂O₈ nanoparticles in improving the mechanical properties and cryogenic performance of CFRP materials. Zirconium tungstate (ZrW₂O₈) nanoparticles functionalized by polydopamine play a crucial role in mitigating cryogenic microcracking in CFRPs through the following mechanisms:

• -. ZrW₂O₈ nanoparticles exhibit negative thermal expansion properties, which can help reduce the thermal residual stresses that occur due to the mismatch in coefficients of thermal expansion between CFs and the EP matrix at CT. This reduction in thermal stresses can prevent the initiation and propagation of microcracks in the CFRP laminate.

• -. The addition of ZrW₂O₈ NPs functionalized by PDA can significantly enhance the fracture energy of the EP matrix in CFRPs. This enhancement in fracture energy helps in enhancing the interlaminar fracture toughness of the composite, making it more resistant to microcracking under cryogenic conditions.

• -. By incorporating these nanoparticles into the EP matrix, the modified matrix effectively suppresses matrix microcracking at CT. This is achieved by reducing the brittleness of the EP polymer and improving its toughness, thereby enhancing the overall performance of the CFRP laminate in cryogenic conditions.

In summary, the functionalization of ZrW₂O₈ NPs with polydopamine helps in reducing thermal stresses, enhancing fracture energy, and preventing matrix microcracks in CFRPs at CT, ultimately improving the mechanical properties and durability of the matrix. The mechanical properties of epoxy polymers modified with ZrW₂O₈ NPs functionalized by PDA were investigated. The important findings about the mechanical properties obtained from this study are as follows:

• -. The tensile strength of the NCs generally increased as the concentration of nanoparticles was increased. However, at higher amounts (e.g., 8 wt% uncoated ZrW₂O₈ NPs), the tensile strength was found to be roughly equivalent to that of the neat EP. This could be due to poor dispersion and agglomeration of NPs at greater amounts, resulting in stress concentrations and a decrease in tensile strength.

• -. The Young’s modulus of the nanocomposites showed an increasing trend with the concentration of NPs. The modulus values were higher for nanocomposites containing higher amounts of NPs, indicating a stiffening effect due to the presence of ZrW₂O₈ nanoparticles. However, agglomeration of nanoparticles at higher concentrations could affect the modulus values.

• -. A reduction in ductility was observed at CT (approximately −196 °C) compared to RT in the tensile tests. This reduction in ductility at low temperatures can be considered an indication of increased brittleness in the nanocomposites under cryogenic conditions.

• -. SEM micrographs revealed that agglomeration of nanoparticles increased with higher nanoparticle concentrations. Agglomeration of nanoparticles could lead to stress amounts in the NCs, potentially reducing their tensile strength and affecting mechanical properties (Figure 67).

One important metric that was examined in the study was the fracture energy of the EP-ZrW₂O₈ NCs. When employing PDA-coated ZrW₂O₈ NPs as opposed to plain NPs, the fracture energy of the bulk epoxy-nanoparticle systems, more precisely the fracture energy G_IC, _bulk, increased. This increase in fracture energy was ascribed to the more energy needed to induce fiber-bundle debonding and matrix cohesive failure, which enhanced the nanocomposites’ toughness. Micrographs of the fracture surfaces of EP NCs containing 1 weight percent untreated and PDA-coated ZrW₂O₈ NPs at −196 °C are shown in Figure 68, which was obtained using SEM. On the fracture surface of the 1 weight percent untreated ZrW₂O₈-EP NCs, cracked pinning along with void development are visible (Figure 68a); on the fracture surface of the 1 weight percent PDA-coated ZrW₂O₈-EP NCs, on the other hand, more crack-pinning, river-lines in the matrix, larger void development, microscopic cracking, alongside breakage of the NPs are visible. The void size of the ZrW₂O₈ NCs coated with 1 weight percent PDA is notably greater than that of the uncoated nanoparticles, as Figure 68a further illustrates. In fact, coated nanoparticles exhibited 70% wider void widths than uncoated NPs, as seen in Figure 68b. The increased robustness that the PDA coating produces might be explained by this bigger void size. The PDA coating may raise the critical displacement for interfacial failure, leading to increased plastic void formation, in accordance with a cohesive zone idea for the interface. The main strengthening processes NP debonding, plastic void expansion, crack pinning, along with fracture deflections caused by NPs are shown in Figure 69. These mechanisms are depicted in the micrographs taken from Figure 68′s SEM. The percentage increase in G_IC and mass at RT and CT following the addition of untreated as well as PDA-coated ZrW₂O₈ NPs is compared in Figure 70 to findings for other kinds of NPs reported in the literature [43,76,77]. As seen in Figure 70, the addition of 1 weight percent PDA-coated ZrW₂O₈ NPs to the EP resulted in the largest absolute boost in fracture energy at CT. The PDA coating on ZrW₂O₈ NPs significantly enhanced the fracture properties of the EP NCs. he effectiveness of surface functionalization in enhancing mechanical characteristics was demonstrated by the discovery that a 1 weight percent concentration of PDA-coated ZrW₂O₈ NPs may increase fracture energy by up to 140% when compared to unmodified EP matrices.

5. Applications in Hydrogen Storage Vessels

Increasing the CFRP’s cryomechanical properties to enable the CFs to endure stresses of up to 3000 MPa while being largely unaffected by cryogenic settings is the primary challenge in CcH₂ storage vessels. When CRFP fails, the resin is usually to blame. The material’s elastic modulus increases at low temperatures, but its capacity to withstand deformation under high pressure decreases. As microcracks appear, they spread throughout the matrix [46]. Therefore, it is essential to change CFRP resins in order to increase cryotoughness and stop microcrack development. There aren’t any high-performing, dependable, or safe linerless type V cryo-compressed hydrogen (CcH₂) storage tanks in production at the moment [78]. To solve this problem, Yan et al. [79] provide a unique method for creating a linerless CcH₂ storage tank. First, a modified resin coating is made by mixing thermoplastics and nanoparticles to stop hydrogen leaking in. In addition to removing flaws from the film, hot rolling enhances molecular alignment. Subsequently, the altered resin films undergo surface treatment to enhance their ability to adhere to cold surfaces. The CFRP films are positioned between the modified resin layers to create a structure that has H₂ barrier qualities. Toughening agents, such as compounds or nanoparticles, are used to lessen cryogenic heat stress. In CFRP, the creation and propagation of microcracks are reduced. This approach produces a composite film for CcH₂ storage tanks with good cryogenic mechanical characteristics and H₂-barrier performance, preventing tank failure from flaws in metal or plastic liners.

PEG 600 modified EP was employed by Zhang et al. [80] to enhance the cryogenic mechanical characteristics of a cryocompressed H₂ storage tank. The cryogenic mechanical characteristics of the composite layer are examined in the context of a cryo-compressed hydrogen storage tank using the procedures listed below:

• -. Typically, the CFRP composite layer, the liner layer, the vacuum insulation layer, together the exterior protective layer makes up the composite layer of a CcH₂ storage tank.

• -. The composite film is subjected to two extreme working conditions: high pressure (35 MPa) and CT (20 K).

Figure 71a depicts the production of a CcH₂ storage tank with PEG modified EP. The steps for preparing the CcH₂ storage tank with PEG modified EP are as follows:

• -. The process begins with the infiltration of carbon fibers in an EP solution to create a ply. This step ensures that the carbon fibers are evenly distributed and impregnated with the modified epoxy resin [81].

• -. The impregnated carbon fibers are then used to wind the CcH₂ storage vessel. The winding process involves layering the impregnated fibers to form the composite layer of the vessel. The winding direction can be hoop winding and longitudinal winding [82], with specific angles and ratios determined based on design requirements. Figure 71b shows the design flow chart for the material film [81].

• -. The E51 epoxy resins with specific epoxy value (0.51/100 g) and molecular weight (392.2) and curing agent (Triethylenetetramine) are mixed in a specific ratio, and the modifier, Polyethylene glycol 600 (PEG 600), is added at varying concentrations (1 wt%, 3 wt%, and 5 wt%) (Figure 72).

• -. The mixture is stirred thoroughly to ensure uniform distribution of the modifier.

• -. The modified EP solution is poured into molds and cured at high temperatures for the specified time.

• -. The modified epoxy resins are categorized into different groups based on the concentration of PEG 600 added (e.g., MER-1, MER-3, MER-5).

• -. A control group using neat epoxy resin (EP) is also prepared for comparison.

Measurements of tensile strength and compression are crucial tests used to assess the mechanical characteristics of materials, particularly EP modified with PEG. Tensile stress–strain profiles for the hardened EP at room temperature and CT are displayed in Figure 73a. The sample initially experiences elastic deformation and then fractures at the maximum tensile strength. The specimen’s elongation at break increased and this tendency became more evident as the amount of PEG increased, but its tensile strength and elastic modulus decreased significantly at room temperature following the addition of the PEG modifier. All materials’ tensile strength and elastic modulus are significantly increased by CT. In a CT environment, molecules and chemical bonds contract, strengthening the binding force between them while reducing their mobility. The results of the tensile strength and modulus tests for the improved EP are shown in Figure 73b. The tensile strength of ME-5R, where R denotes room temperature, is only 25.99 MPa when the PEG content is 5 weight percent. MER-5R has a tensile strength that is 53.63% less than ER-R, the control group. The tensile strength of MER-5R is 53.63% less than that of ER-R, the control group. The redesigned EP will no longer fulfill the strength requirements of the CcH₂ storage tank at room temperature if the amount of PEG is raised further. The redesigned EP will no longer fulfill the strength requirements of the CcH₂ storage tank at room temperature if the amount of PEG is raised further. The MER-1C and MER-3C, which are modified EPs with a C indicating that the experiment was carried out at CT, show higher tensile strengths than ER-C. However, MER-3C experiences a 45.31% reduction in tensile modulus, although elongation at break and fracture toughness are concurrently improved. The PEG alteration gives EP its “strength and toughness” characteristics by drastically lowering its tensile modulus while preserving its high strength at CT. The primary reason for the experimental phenomenon is that a high-temperature hardened EP becomes more brittle and generates a considerable amount of internal thermal stress when it is cooled to CT. Internal stress in the brittle EP chain network is decreased by the soft ether bond in the flexible PEG polymer chain [83]. Bond motion can still minimize internal stress and a high number of soft elastic ether bond segments injected into the EP will not completely freeze under cryogenic circumstances, improving the cryogenic mechanical properties of the resin. Nevertheless, because PEG free volume occupies space in the hardened modified EP, its addition lowers the crosslink density of the epoxy system. The tensile strength of the hardened material decreases when the amount of PEG is very large (>5 wt%) because the flexible segment’s negative impact factors take precedence when computing tensile strength. Figure 73c,d show the modified EP’s compressive properties at RT and CT. Compressive strength and modulus, like tensile test results, decrease with increasing PEG modifier level. CT, on the other hand, has a higher compressive strength and modulus than RT, so it increases at first and then decreases as the PEG modifier concentration increases.

In the context of a CcH₂ storage tank, winding plies, stress, and displacement are critical factors that influence the structural integrity and performance of the composite layer. To solve the problem, this study proposes two design strategies for a CcH₂ storage tank, as seen in Figure 74. Here is an overview of winding plies, stress, and displacement considerations in the design and analysis of a CcH₂ storage vessel:

• -. Winding plies refer to the layers of carbon fiber composite material that are wound around the vessel to form the composite layer.

• -. The number of winding plies is determined based on the mechanical properties of the composite, the operating conditions of the vessel, and the desired structural strength.

• -. The goal is to achieve a balance between the number of winding plies, the material properties, and the structural requirements to ensure the vessel’s safety and performance.

• -. The optimal number of winding plies is calculated based on stress analysis, material properties, and design standards to meet the vessel’s mechanical requirements

• -. Stress analysis is conducted to evaluate the distribution of internal stresses within the composite layer of the CcH₂ storage vessel.

• -. The analysis considers factors such as internal pressure, temperature variations, and material properties to determine the stress levels experienced by the composite material.

• -. The stress analysis helps in identifying critical areas of areas requiring reinforcement, potential failure points, and high stress concentration.

• -. Understanding the stress distribution is crucial for optimizing the design, material selection, and manufacturing process of the composite layer.

• -. Displacement refers to the movement or deformation of the composite layer in response to applied loads, such as internal pressure or external forces.

• -. Analyzing displacement helps in predicting the structural behavior of the tank under different loading conditions and ensuring that deformations are within acceptable limits.

• -. Excessive displacement can lead to structural instability, leakage, or failure of the storage vessel, highlighting the importance of accurate displacement analysis.

• -. By considering displacement in the design phase, engineers can optimize the vessel’s geometry, material properties, and reinforcement to minimize deformations and enhance structural performance.

• -. Through advanced modeling methods, such as finite element analysis (FEA), engineers can simulate the behavior of the composite layer under various loading scenarios.

• -. By iteratively adjusting the number of winding plies, material properties, and design parameters, designers can optimize the vessel’s structural performance while minimizing weight and material usage.

• -. The goal of design optimization is to ensure that the CcH₂ storage vessel meets safety standards, performance requirements, and durability criteria while maximizing efficiency and reliability.

Overall, the findings of the work highlight the importance of material optimization, resin modification, winding plies design, stress analysis, and design optimization in enhancing the properties and performance of the composite film in CcH₂ storage tanks.

The same group [84], has developed a novel approach using polyethylene (PE) layers to enhance the performance of CF composites under extreme conditions. To prevent hydrogen permeation, PE films were added to the CF interlayers. Figure 75 depicts two alternative composite preparation techniques. Traditional approaches involve inserting hydrogen-barrier layers between prepreg fibers [85]. Composites having a sandwich structure are formed after heating and curing at 110 °C. The resin anchor effect is used to connect the films and fibers, which can easily lead to delamination failure in cryogenic environments. This group developed a new procedure in which hydrogen-barrier films initially were inserted between films of dry CF to ensure uniform fiber orientation. The specimen was then placed in a furnace for two hours of hot pressing at 210 °C and 0.1 MPa. In this process, the PE films melt and begin to flow outwards. Under pressure, PE traverses the spaces between CF bundles. The hydrogen-barrier layers, which had previously been separated by CF layers, were connected interlayer. Molten PE grids wrapped around the CF bundles, resulting in a unique composite structure known as a cured network. Moreover, the CFs and PE were impregnated using the vacuum-assisted resin converted molding technique [86]. Finally, the specimen was returned to the furnace to cure the resin at 110 °C. The key findings are as follows:

• -. Through a variety of tests and analysis, the cryogenic mechanical characteristics of the combined CF materials in the study were assessed. The semi-product of CFs with PE film following hot pressing treatment is depicted in Figure 76a,b. There is no EP connection between the CFs, which keeps them apart. PE melts penetrate the CF layers and create crosslinks after being hot-pressed. A mixed CF and PE region is created as the melted PE envelops the nearby CFs. As the melted PE flows, three regions form successively inside the semi-product’s cross-section: the CF area (fibers), the PE zone (PE polymer), and the mixed region (fibers wrapped in PE). The mixed area has a diameter of approximately 100 μm. A cross-section micrograph of the matrix exhibiting a cross-linked network structure may be found in Figure 76c. The three zones are firmly linked together by the resin once the resin is impregnated and allowed to set, creating a single substance. The boundaries of the mixed zone are less clear at this time. PE grids surround the CFs and resin. The combination is shown in Figure 76d. The resin will cure first since the curing temperature (110 °C) is much lower than the PE melting temperature (180–210 °C). Consequently, CFs and PE are held together in the sandwich structure by the mechanical anchor effect of the resin. But the anchors will weaken at CT, which will cause a delamination flaw. PE sheets are melted by hot pressing and then placed into dry CFs. The CFs are then encircled by the PE, creating a cross-linked network structure inside the matrix. The cross-linking action of PE preserves the laminate’s superior mechanical characteristics at CT while removing its layered structure. The interlayer adhesion forces of the two structures are tested, as shown in Figure 76e,f. The PE and fibers are firmly bonded at RT using resin. The resin creates mechanical anchors by penetrating polyethylene layers and fiber surfaces [87]. There is very little difference in the adhesion force values between the two materials. Nevertheless, CT distorts and shrinks the resin, lessening the anchor effect. Consequently, the adhesion force of the sandwich structure immediately drops by 28.8%. On the other hand, the cross-linked network structure makes use of PE’s cross-linking influence to integrate the multi-layer CFs into a whole. The adhesion force is still quite powerful even if the resin’s anchor impact is diminishing. At CT, a cross-linked network structure has an adhesion force that is 22.7% more than a sandwich structure. Macroscopic photographs of shear contacts are shown in Figure 76g. A portion of the PE sticks to the laminates along the shear plane of the sandwich structure at room temperature. While the PE in cross-linked network structures is strongly attached to the laminates, the PE in sandwich structures totally separates from the laminates under freezing temperatures. The longitudinal along with transverse tensile strengths of the matrices rose by 23.2% and 21.2%, respectively, at CT. This improvement in tensile strength shows how PE coatings may enhance a material’s mechanical qualities in harsh environments.

• -. One important factor that was assessed to make sure the CF materials in the study were suitable for use in CcH₂ storage containers was their H₂-barrier characteristic. The concept of the differential pressure technique is shown in Figure 77a. The sample is seen being put between two chambers in Figure 77b. Although the bottom chamber is vacuumed at a value below 10 Pa, the upper half (chamber) is connected to 0.1 MPa of high-pressure gas [88]. The bottom chamber’s pressure is continually monitored by a sensor, and the experiment automatically finishes when pressure equilibrium is attained. A cryogenic fatigue test that was conducted on the laminates to replicate the real-world operational circumstances of the CcH₂ storage tank is shown in Figure 77c. Following 10,000 loading cycles, the laminates are cut into circular samples, and the coefficients of hydrogen permeability are determined. The cross sections of the matrix are shown in Figure 77d. Numerous microcracks emerge at cryogenic pressure, facilitating quick H₂ penetration. The method in which the materials’ linked microcracks allow gas to escape is seen in Figure 77e. On the other hand, PE film may keep a considerable amount of its toughness at CT, which stops microcracks from propagating through materials. Diffusion is therefore the most typical method of H₂ permeation. It has been demonstrated that microcracks produce noticeably more leakage than diffusion [89].

In summary, the research successfully addressed the challenges of hydrogen leakage and delamination defects in cryogenic environments by developing composite materials with improved mechanical strength and hydrogen-barrier performance, paving the way for the application of these composites in CcH₂ storage vessels.

Liquid hydrogen can be stored in type V linerless vessels, which is composed entirely of composite materials. However, even at this low temperature, the epoxy resin matrix exhibits some difficulties with gas barrier and mechanical performance. A different study conducted by the same group explores the co-modification of montmorillonite and polyethylene glycol to enhance the properties of EP at low temperatures [90]. The specimen production process diagram is shown in Figure 78. The epoxy resin samples in the study were prepared by following a specific procedure:

• -. The epoxy resin (EP model E51) solution was prepared with the specified components.

• -. Various proportions of PEG and montmorillonite K10 (MMT) were added to the EP solution based on the sample requirements.

• -. To ensure that the modified fillers were evenly dispersed, the solution was stirred at 1000 r/min at RT for 10 min.

• -. With each stirring, any bubbles formed were likely removed to maintain a uniform mixture.

The gas-barrier property results in the study were obtained through specific tests and analyses:

• -. Helium (0.26 nm) gas was used as the testing gas due to its smaller molecular diameter compared to H₂ gas (0.289 nm) [91,92].

• -. Before starting the experiment, the EP samples’ surfaces were cleaned.

• -. The samples were placed in a test instrument with upper and lower chambers, and real-time monitoring was performed using pressure sensors.

• -. The lower chamber (less than 20 Pa) was vacuumed, and He gas was introduced into the upper chamber (0.1 MPa) until pressure balance was achieved [90].

• -. Figure 79 shows how the gas permeability coefficients of different samples were measured and compared. As the MMT content increases, the gas permeability coefficient decreases. This is due to MMT NPs’ ability to make the gas diffusion and leakage channel in the EP matrix more convoluted, thereby increasing the diffusion path of gas molecules, as shown in Figure 80. As MMT content increases, so does the quantity of nanoparticles, and the diffusion channel becomes more convoluted. As MMT content increases, the gas permeability coefficient decreases. The gas permeability coefficient increases as PEG content increases. Figure 81a depicts two types of nanoparticles used to reduce polymer gas permeability coefficients: intercalation and exfoliation. Exfoliated NPs were widely assumed to have increased disorder and improved gas barrier properties [93]. MMT’s hydrophilicity, on the other hand, made uniform dispersion in the polymer difficult [93,94,95]. As a result, some MMT may exist in the form of agglomerations [96,97]. Mixing PEG, a macromolecular chain, with MMT changed its properties. Figure 81b shows that the macromolecular chain had a less significant effect on the intercalation and exfoliation of montmorillonite nanoparticles than the resin matrix. The inclusion of PEG reduced the amount of resin, resulting in an improvement in the gas permeability coefficient. When enhancing the mechanical properties, the amount of PEG used should be kept to a minimum.

The mechanical properties of the EP samples were evaluated at RT and CT in this study:

• -. The tensile strength of the MMT-25 specimen at RT was reported as 23.93 MPa, with a decrease of 57.54% compared to the base EP.

• -. The Young’s modulus at RT increased by 59.56% to 2.17 GPa, indicating a more brittle material.

• -. The addition of PEG was aimed at strengthening and toughening the resin matrix due to the observed decrease in mechanical properties.

• -. The MMT25-PEG10 sample at CT showed a significant improvement in tensile strength, reaching 57.02 MPa, which was a 98.12% increase compared to RT.

• -. The Young’s modulus decreased by 11.22% to 2.69 GPa at CT, indicating improved toughness and strength of the EP matrix.

• -. The epoxy resin’s elongation at break increased at RT and CT after the addition of PEG, increasing the material’s toughness.

In conclusion, the co-modification of EP with PEG and MMT offers a promising approach to enhance the performance of resin matrices for cryogenic applications, such as in Type V linerless vessels, ensuring both gas-barrier integrity and mechanical robustness.

6. Conclusions and Future Perspective

The current review focuses on the main challenges and developments in resin toughening and hydrogen barrier properties for cryo-compressed hydrogen storage tanks. It emphasizes that resin toughening improves the cryogenic mechanical properties of CF/EP, which is critical for preventing cryogenic pressure deformation. The use of thermoplastic toughening compounds is chosen because they can improve strength and toughness at cryogenic temperatures. The work also analyzes the effectiveness of hydrogen-barrier films and modified resins in decreasing hydrogen leakage. These films may effectively prevent microcrack propagation and convert permeation processes, increasing the overall safety and efficiency of hydrogen storage systems. However, problems such as interfacial defects and nanoparticle limitations in preventing microcrack propagation are highlighted. According to the review, addressing these challenges effectively requires a combination of strategies for manipulating hydrogen barrier properties. Overall, the findings highlight the potential for improved hydrogen storage solutions through innovative material design and engineering approaches, paving the way for future advances in sustainable energy technologies.

Author Contributions

All the authors equally contributed. All authors have read and agreed to the published version of the manuscript.

Conflicts of Interest

The authors declare no conflict of interest.

Abbreviations

CF	Carbon fiber
CFs	Carbon fibers
EP	Epoxy resin
EPs	Epoxy resins
VARTM	Vacuum-assisted resin transfer molding
RT	Room temperature
TFBT	Transverse fiber bundle tension testing
CNT	Carbon nanotube
EPD	Electrophoretic deposition
CT	Cryogenic temperature
ILSS	Interlaminar shear strength
MWCNTs	Multi-walled carbon nanotubes
FS	Flexural strength
DDM	4,4′-diaminodiphenylmethane
TGA	Thermogravimetric analysis,
BMI	Bismaleimide resin
CTE	Coefficient of thermal expansion
GO	Graphene oxide
CFRP	Carbon fiber-reinforced polymer
ILSS	Interlaminar shear strength
DETDA	Diethyl toluenediamine
HPB	Hyperbranched polymer
DMDPS	Dimethoxydiphenylsilane
DMDMS	Dimethyldimethoxysilane
EFPS	Epoxy-functionalized polysiloxane
NPs	Nanoparticles
CWF	Carbon woven fabric
EPD	Electrophoretic deposition
PALS	Positron annihilation lifetime spectroscopy
MeTHPA	Methyltetrahydrophthalic anhydride
MW	Molecular weight
NCs	Nanocomposites
SEM	Scanning electron microscopy
DCB	Double cantilever beam
SENB	Single-edge notch bending
PBT	Polybutylene Terephthalate
PEI	Polyetherimide
PC	Polycarbonate
DMA	Dynamic mechanical analysis
TRIS	Tris(hydroxymethyl)aminomethane
ERR	Electrical resistance ratio
ER	Electrical resistance
CVD	Chemical vapor deposition
POSS	Polyhedral oligomeric silsesquioxane
PMMA	Polymethylmethacrylate
PTFE	Polytetrafluoroethylene
TPs	Thermoplastics
FTIR	Fourier transform infrared spectroscopy
KH560	γ-(2,3-epoxypropoxy) propytrimethoxysilane
AFM	Atomic force microscope
DGEBA	Diglycidyl ether of bisphenol A
FEA	Finite element analysis
ERR	Electrical resistance ratio
DI water	Deionized water
CcH2	Cryo-compressed hydrogen

Footnotes

Disclaimer/Publisher’s Note: The statements, opinions and data contained in all publications are solely those of the individual author(s) and contributor(s) and not of MDPI and/or the editor(s). MDPI and/or the editor(s) disclaim responsibility for any injury to people or property resulting from any ideas, methods, instructions or products referred to in the content.

Figures

Figure 1. Molecular structures of epoxy resin, curing agents, and components of four types of epoxy [8]. Copyright 2022 Elsevier.

Figure 2. Relative growth rate of mechanical properties of four type epoxy resin at RT and −196 °C [8]. Copyright 2022 Elsevier.

Figure 3. Mechanical properties of different types of epoxy resin at RT and −196 °C: (a) δ tensile strength, (b) E tensile modulus, (c) ε failure strain, (d) impact strength [8,30,31,32,33]. Copyright 2022 Elsevier.

Figure 4. (a) Mechanical properties of CFRP at RT and −196 °C; tensile failure morphology of CFRP. (b) RT and (c) −196 °C (1: fibre pull-out; 2: interfacial debonding; 3: fibre breakage; 4: matrix crack) [8]. Copyright 2022 Elsevier.

Figure 5. (a) G′ vs. temperature curve, (b) G″ vs. temperature curve [8]. Copyright 2022 Elsevier.

Figure 6. (a) Tan δ vs. temperature curve; (b) crosslinking structure of epoxy; (c) CTE vs. temperature curve [8]. Copyright 2022 Elsevier.

Figure 7. Tensile failure morphology of epoxy at (a) RT and (b) −196 °C; (c) stress–strain curve of epoxy at RT and −196 °C [8]. Copyright 2022 Elsevier.

Figure 8. Electrical resistance (ER) fragmentation tensile test specimen [35]. Copyright 2015 Elsevier.

Figure 9. (a) Data and empirical correlation calculations between △R/R0 and stress to strain; (b) tensile data of the fragmentation sample (bisphenol A) at RT; (c) tensile findings of the fragmentation sample (bisphenol A) at CT; (d) tensile data of the fragmentation sample (bisphenol F) at RT; and (e) tensile findings of the fragmentation sample (bisphenol F) at CT [35]. Copyright 2015 Elsevier.

Figure 10. Schematic model of carbon fiber with tensile loading [35]. Copyright 2015 Elsevier.

Figure 11. Model of tensile loading of fragmentation specimen with different interfacial adhesion between fiber and matrix [35]. Copyright 2015 Elsevier.

Figure 12. (a) Static contact angle test system under the cryogenic condition; (b) static contact angle results of epoxy resin for different temperature conditions; and (c) static contact angle of epoxy resin at RT and CT [35]. Copyright 2015 Elsevier.

Figure 13. (a) Diagrams and dimensions of fragmentation testing specimens; (b) cryogenic treatment process for different samples [36]. Copyright 2017 Elsevier.

Figure 14. (a) Electrical resistance-changing ratio of CNT fiber and fragmentation specimen before and after cryogenic treatment under tensile testing; (b) static resistance and gage factor of the embedded CNT fiber before and after cryogenic treatment [36]. Copyright 2017 Elsevier.

Figure 19. Schematic depicting behavior of neat composite at RT and CT and of modified composite at CT [42]. Copyright 2020 Elsevier.

Figure 20. Molecular structures of (a) DGEBA, (b) BCI, and (c) DDM [30]. Copyright 2022 Elsevier.

Figure 21. Schematic of the preparation process of the EP/BCI composites [30]. Copyright 2022 Elsevier.

Figure 22. (a) The reaction between the epoxy and amino groups; (b) in BCI, the amino group reacts with the double bond via the Michael addition reaction; (c) a complete response to the two intermediate products [30]. Copyright 2022 Elsevier.

Figure 23. Schematic of the deformation of pure EP and the EP/BCI composite molecular curing networks under tensile load [30]. Copyright 2022 Elsevier.

Figure 24. Schematic illustration of the cross-linked network of (a) pristine EP and (b) EP/BCI materials [30]. Copyright 2022 Elsevier.

Figure 25. The τ3, I3, and Tg of unmodified and PEG-4000 modified epoxy resins and schematic depicting: (a) a single PEG-4000, (b) DGEBA, and (c) MeTHPA compounds [5]. Copyright 2014 Elsevier.

Figure 26. Mechanism of epoxy resins toughened with PEG [5]. Copyright 2014 Elsevier.

Figure 27. Chemical structures of epoxy, D-230, PBT, PEI, and PC [45]. Copyright 2013 Elsevier.

Figure 28. (a) thermal expansion, (b) CTE values, (c) DMA curves, and (d) average crack densities of neat epoxy and thermoplastic modified EP [45]. Copyright 2013 Elsevier.

Figure 29. Reflected optical microscopy images of cryogenically cycled CF/EP laminates: (a) CF/neat epoxy, (b) CF/PBT modified epoxy, (c) CF/PC modified epoxy, and (d) CF/PEI modified epoxy laminates [45]. Copyright 2013 Elsevier.

Figure 30. The schematic graphic depicts the CFRP composite fabrication process [46]. Copyright 2021 Elsevier.

Figure 31. Optical pictures of contact angles with microdroplets on CF for (a) neat EP, (b) EP/HTPU1, and (c) EP/HTPU2 [46]. Copyright 2021 Elsevier.

Figure 32. Digital pictures of longitudinal tensile (a,b) and digital images of the bending after failure test samples (c,d) of CFRP and CF/EP/HTPU materials following tensile failure: (a,c) RT; (b,d) 77 K [46]. Copyright 2021 Elsevier.

Figure 33. (a) The synthetic pathway of flexible polymer containing epoxy groups (EPSE), and (b) illustration diagram of the CF/EP composite laminate [50]. Copyright 2020 Elsevier.

Figure 34. SEM images of the fracture surfaces of neat epoxy at (a) RT and (b) 77 K, EPSE-EP with 8 wt% EPSE content at (c) RT and (d) 77 K, optical images of CF/neat epoxy laminate after flexural test at (e) RT and (f) 77 K, and CF/EPSE-EP laminate after flexural evaluate at (g) RT and (h) 77 K [50]. Copyright 2020 Elsevier.

Figure 35. The procedure for producing HPB/epoxy composites [44]. Copyright 2021 Elsevier.

Figure 36. Schematic of hyperbranched polymer toughened epoxy at CT [44]. Copyright 2021 Elsevier.

Figure 37. Manufacturing procedures for producing nanocomposite and CF composites with the nanocomposite as the sample: (a) BCP-EP nanocomposites and (b) vacuum-assisted infusion of CF composites with BCP/epoxy matrix [53]. Copyright 2024 Elsevier.

Figure 38. Rigid and soft fillers improve nanocomposites’ fracture toughness at (a) RT and (b) −196 °C; (c) comparison of fracture toughness enhancement at CT vs. RT [53]. Copyright 2024 Elsevier.

Figure 39. Fractographs of the fracture surfaces of DCB samples evaluated at RT and CT ((a–c), (d–f), respectively). (a) neat epoxy; (b) polymer with 2.5 wt% BCP; (c) high resolution image of polymer in (b); (d) neat epoxy; (e) polymer with 2.5 wt% BCP; and (f) high resolution image of polymer in (e) [53]. Copyright 2024 Elsevier.

Figure 40. Optical microscopy images of cracks in cross-ply [04/904]s laminates after quenching in liquid nitrogen (−196 °C): (a) neat sample; (b) matrix with 1.0 wt% BCP; (c) sample with 2.5 wt% BCP; (d) sample with 5.0 wt% BCP [53]. Copyright 2024 Elsevier.

Figure 41. Schematic depicting the sample production procedure for contact angle testing, as well as the immersion of a CNT/EP mixture into CF bundles to create transverse fiber bundle tension testing (TFBT) samples [58]. Copyright 2014 Elsevier.

Figure 42. (a) Transverse tensile specimen geometry (in mm), (b) TFBT specimen shot, and (c) cryogenic tensile testing schematic [58]. Copyright 2014 Elsevier.

Figure 43. The contact angle measurement process is represented schematically in two views: (a) vertical and (b) [58]. Copyright 2014 Elsevier.

Figure 44. Optical micrographs of epoxy matrix droplets with various MWCNT percentages on CF bundles [58]. Copyright 2014 Elsevier.

Figure 45. SEM photographs of the matrix samples’ fracture surfaces during transverse tensile testing at 77 K and RT [58]. Copyright 2014 Elsevier.

Figure 46. (a) Scheme to preparing Fe3O4/O-MWCNTs and O-MWCNTs; (b) diagram of the design for preparing ordered Fe3O4/O-MWCNTs modified EP systems; (c) scheme for preparing Fe3O4/O-MWCNTs/CF/EP laminate (left), and picture of Fe3O4/O-MWCNTs/CF/EP laminate (right) [59]. Copyright 2018 Elsevier.

Figure 47. Propagating cracks and suppression characteristics of CNT-modified CF/EP composite materials. CF/EP laminate (a), R-MWCNT-modified CF/EP laminate (b), O-MWCNT-modified CF/EP laminate (c), and Fe3O4/O-MWCNT-modified CF/EP laminate (d) [59]. Copyright 2018 Elsevier.

Figure 48. Schematic representation of the preparation of CF/GO-modified epoxy systems [61]. Copyright 2020 Elsevier.

Figure 49. Schematic representation of the RT and cryogenic evaluation for three-point bending tests, (a) RT, (b) 77K and (c) sample immersed in liquid nitrogen [61]. Copyright 2020 Elsevier.

Figure 50. Macroscopic failure morphologies of CFRP samples after three-point bending tests at (a) RT and (b) 77 K, shown from the side [61]. Copyright 2020 Elsevier.

Figure 51. GO dispersion in the epoxy matrix with various GO concentrations: (a) 0.05 wt%, (b) 0.1 wt%, (c) 0.2 wt%, and (d) 0.5 wt% [61]. Copyright 2020 Elsevier.

Figure 52. Schematic representation of the production process of GO-containing CFRP matrices: (a) GO-EP/CWF matrices and (b) GO-CWF/EP matrices [62]. Copyright 2019 Elsevier.

Figure 53. SEM morphological characteristics of the fracture surface of (a) pristine, (b) GO-EP/CWF, and (c) GO-CWF/EP matrices at RT [62]. Copyright 2019 Elsevier.

Figure 54. Schematic diagrams of failure at the interface for (a) pristine, (b) GO-EP/CWF, and (c) GO-CWF/EP matrices [62]. Copyright 2019 Elsevier.

Figure 55. (a) Scheme for the manufacturing of Fe3O4/GO, (b) Scheme for manufacture of Fe3O4/GO/CF/EP laminate and image of Fe3O4/GO/CF/EP [65]. Copyright 2018 Elsevier.

Figure 56. Reflected optical microscope pictures of cryogenically cycled CF/EP laminates: (a) CF/pristine EP laminate; (b) CF/EP laminate with 0.5 wt% GO; and (c) CF/EP laminate with 0.5 wt% Fe3O4/GO. (d) average crack density for the different systems [65]. Copyright 2018 Elsevier.

Figure 57. A diagram depicting the crack growth and suppression actions of CF/EP materials adjusted with additives: (a) CF/EP laminate, (b) GO or Fe3O4/GO modified CF/EP laminate [65]. Copyright 2018 Elsevier.

Figure 58. Schematic representation of the procedure used to prepare a EFPS/Nano-SiO2 EP laminate with CF reinforcement [43]. Copyright 2020 Elsevier.

Figure 59. Schematic representation of the EFPS/Nano-SiO2 fabrication process [43]. Copyright 2020 Elsevier.

Figure 60. (a) Schematic diagram and (b) test picture of the cryogenic tensile testing, (c) schematic diagram and (d) a test picture of the three-point bending testing a cryogenic environment [43]. Copyright 2020 Elsevier.

Figure 61. At 90 K, CF-reinforced (a) pristine epoxy laminate and (c) EFPS/Nano-SiO2 epoxy laminate exhibit microfractured morphologies. A diagram depicting the failure mode of CF-reinforced (b) pristine epoxy laminate and (d) EFPS/Nano-SiO2 epoxy laminate [43]. Copyright 2020 Elsevier.

Figure 62. Diagram of the preparation process of the modified nano-ZrO2/EP matrices [73]. Copyright 2016 RSC.

Figure 63. Tensile stress–strain diagrams for ZrO2 treated resin at RT and CT [73]. Copyright 2016 RSC.

Figure 64. (a) Method for PDA coating CuO NPs, (b) Illustration of PDA coated CuO dispersed in PDA coated CuO-EP NCs, and (c) Vacuum assisted infusion technique for angle-ply laminates [54]. Copyright 2021 Elsevier.

Figure 65. (a) Curing network of pristine EP; (b) improved curing network by PDA-coating on nCuO [54]. Copyright 2021 Elsevier.

Figure 66. Mechanisms of NP debonding, matrix shear banding, and plastic void formation, in modified EP at CT [54]. Copyright 2021 Elsevier.

Figure 67. The arrangement of ZrW2O8 NPs in EP-ZrW2O8 NCs. SEM pictures of fracture surfaces of tensile samples at CT with amounts of (a) 1 wt% uncoated, (b) 1 wt% PDA-coated, (c) 4 wt% uncoated, and (d) 8 wt% uncoated ZrW2O8 NPs [75]. Copyright 2023 Elsevier.

Figure 68. Fracture surfaces of the epoxy-ZrW2O8 nanocomposites containing (a) 1 wt.% ZrW2O8 and (b) 1 wt.% PDA-coated ZrW2O8 tested at −196 °C. The blue arrows show the direction of crack growth. The blue arrows show the direction of crack propagation [75]. Copyright 2023 Elsevier.

Figure 69. An illustration of the primary toughening mechanisms in nanocomposites, which include debonding, crack deflection, fracture pinning, and plastic void expansion [75]. Copyright 2023 Elsevier.

Figure 70. The effectiveness of ZrW2O8 NPs in toughening EP systems was compared to SiO2 NPs, POSS and rubber NPs at (a) RT and (b) CT [75]. Copyright 2023 Elsevier.

Figure 71. (a) Preparation procedure and important parameters for CcH2 storage tank; (b) composite layer design [80]. Copyright 2023 Elsevier.

Figure 72. Preparation and testing of PEG-modified epoxy [80]. Copyright 2023 Elsevier.

Figure 73. Mechanical properties of PEG modified epoxy at both RT and CT: (a) Stress-strain curves; (b) Tensile strength and tensile modulus; (c) stress-strain curves; and (d) compressive strength and compressive modulus [80]. Copyright 2023 Elsevier.

Figure 74. Two possibilities for design for a CcH2 storage vessel [80]. Copyright 2023 Elsevier.

Figure 75. Preparation technique for composites with hydrogen barrier properties [84]. Copyright 2023 Elsevier.

Figure 76. Adhesion testing and structural characterization. (a) A composite of CFs and polyethylene (PE) film. (b) Distinct zones inside the semi-product. (c) A cross-linked network framework. (d) Sandwich construction. (e) Curves for force displacement during adhesion tests. (f) Adhesive force values. (g) Shear interface images at the macroscopic photographs [84]. Copyright 2023 Elsevier.

Figure 77. Hydrogen-barrier performance testing. (a) Differential pressure technique. (b) Specimen. (c) Cyclical loading curve for cryogenic fatigue. (d) Microcracks. (e) Diffusion and penetration via microcracks [84]. Copyright 2023 Elsevier.

Figure 78. Sample production process diagram [90]. Copyright 2024 Elsevier.

Figure 79. Change in the epoxy matrix gas permeability coefficient [90]. Copyright 2024 Elsevier.

Figure 80. Nanocomposites-modified polymers make gas diffusion more tortuous [90]. Copyright 2024 Elsevier.

Figure 81. The types of MMT in epoxy resin and PEG are: (a) in epoxy resin, in the type of agglomeration, intercalation, and exfoliation; (b) in PEG, it is mainly in the type of agglomeration [90]. Copyright 2024 Elsevier.

References

1. Air, A.; Shamsuddoha, M.; Gangadhara Prusty, B. A Review of Type V Composite Pressure Vessels and Automated Fibre Placement Based Manufacturing. Compos. Part B Eng.; 2023; 253, 110573. [DOI: https://dx.doi.org/10.1016/j.compositesb.2023.110573]

2. Liu, H.; Chen, Y.; Chai, M.; Wu, Y.; Xue, K.; Liu, L.; Huang, Y. High-Performance Carbon Fiber-Reinforced Epoxy Resin Composites Based on Novel, Environmentally Friendly, Stability, and Heat-Moisture Resistant Nano Emulsion Encapsulated POSS Sizing Agent Interface Modification. Compos. Part B Eng.; 2024; 284, 111697. [DOI: https://dx.doi.org/10.1016/j.compositesb.2024.111697]

3. Zheng, H.; Zhang, W.; Li, B.; Zhu, J.; Wang, C.; Song, G.; Wu, G.; Yang, X.; Huang, Y.; Ma, L. Recent Advances of Interphases in Carbon Fiber-Reinforced Polymer Composites: A Review. Compos. Part B Eng.; 2022; 233, 109639. [DOI: https://dx.doi.org/10.1016/j.compositesb.2022.109639]

4. Wetzel, B.; Rosso, P.; Haupert, F.; Friedrich, K. Epoxy Nanocomposites—Fracture and Toughening Mechanisms. Eng. Fract. Mech.; 2006; 73, pp. 2375-2398. [DOI: https://dx.doi.org/10.1016/j.engfracmech.2006.05.018]

5. Feng, Q.; Yang, J.; Liu, Y.; Xiao, H.; Fu, S. Simultaneously Enhanced Cryogenic Tensile Strength, Ductility and Impact Resistance of Epoxy Resins by Polyethylene Glycol. J. Mater. Sci. Technol.; 2014; 30, pp. 90-96. [DOI: https://dx.doi.org/10.1016/j.jmst.2013.08.016]

6. Zhang, Y.; Xu, F.; Zhang, C.; Wang, J.; Jia, Z.; Hui, D.; Qiu, Y. Tensile and Interfacial Properties of Polyacrylonitrile-Based Carbon Fiber after Different Cryogenic Treated Condition. Compos. Part B Eng.; 2016; 99, pp. 358-365. [DOI: https://dx.doi.org/10.1016/j.compositesb.2016.05.056]

7. Unnikrishnan, K.P.; Thachil, E.T. Toughening of Epoxy Resins. Des. Monomers Polym.; 2006; 9, pp. 129-152. [DOI: https://dx.doi.org/10.1163/156855506776382664]

8. Yan, M.; Liu, Y.; Jiang, W.; Qin, W.; Yan, Y.; Wan, L.; Jiao, W.; Wang, R. Mechanism of Matrix Influencing the Cryogenic Mechanical Property of Carbon Fibre Reinforced Epoxy Resin Composite. Compos. Commun.; 2022; 33, 101220. [DOI: https://dx.doi.org/10.1016/j.coco.2022.101220]

9. Mimura, K.; Ito, H.; Fujioka, H. Improvement of Thermal and Mechanical Properties by Control of Morphologies in PES-Modified Epoxy Resins. Polymer; 2000; 41, pp. 4451-4459. [DOI: https://dx.doi.org/10.1016/S0032-3861(99)00700-4]

10. Qu, C.-B.; Xiao, H.-M.; Huang, G.-W.; Li, N.; Li, M.; Li, F.; Li, Y.-Q.; Liu, Y.; Fu, S.-Y. Effects of Cryo-Thermal Cycling on Interlaminar Shear Strength and Thermal Expansion Coefficient of Carbon Fiber/Graphene Oxide-Modified Epoxy Composites. Compos. Commun.; 2022; 32, 101180. [DOI: https://dx.doi.org/10.1016/j.coco.2022.101180]

11. Zavareh, S.; Samandari, G. Polyethylene Glycol as an Epoxy Modifier with Extremely High Toughening Effect: Formation of Nanoblend Morphology. Polym. Eng. Sci.; 2014; 54, pp. 1833-1838. [DOI: https://dx.doi.org/10.1002/pen.23733]

12. Sun, Z.; Xu, L.; Chen, Z.; Wang, Y.; Tusiime, R.; Cheng, C.; Zhou, S.; Liu, Y.; Yu, M.; Zhang, H. Enhancing the Mechanical and Thermal Properties of Epoxy Resin via Blending with Thermoplastic Polysulfone. Polymers; 2019; 11, 461. [DOI: https://dx.doi.org/10.3390/polym11030461] [PubMed: https://www.ncbi.nlm.nih.gov/pubmed/30960445]

13. Jung, H.-S.; Park, Y.; Nah, C.-W.; Lee, J.-C.; Kim, K.-Y.; Lee, C.S. Evaluation of the Mechanical Properties of Polyether Sulfone-Toughened Epoxy Resin for Carbon Fiber Composites. Fibers Polym.; 2021; 22, pp. 184-195. [DOI: https://dx.doi.org/10.1007/s12221-021-9261-4]

14. Sun, X.; Chen, R.; Gao, X.; Liu, Q.; Liu, J.; Zhang, H.; Yu, J.; Liu, P.; Takahashi, K.; Wang, J. Fabrication of Epoxy Modified Polysiloxane with Enhanced Mechanical Properties for Marine Antifouling Application. Eur. Polym. J.; 2019; 117, pp. 77-85. [DOI: https://dx.doi.org/10.1016/j.eurpolymj.2019.05.002]

15. Bakar, M.; Kobusińska, J.; Szczerba, J. Mechanical Properties of Epoxy Resin Modified with Polycarbonate and Reactive Polybutadiene. J. Appl. Polym. Sci.; 2007; 106, pp. 2892-2897. [DOI: https://dx.doi.org/10.1002/app.26898]

16. Deng, S.; Djukic, L.; Paton, R.; Ye, L. Thermoplastic–Epoxy Interactions and Their Potential Applications in Joining Composite Structures—A Review. Compos. Part A Appl. Sci. Manuf.; 2015; 68, pp. 121-132. [DOI: https://dx.doi.org/10.1016/j.compositesa.2014.09.027]

17. Yang, G.; Zheng, B.; Yang, J.-P.; Xu, G.-S.; Fu, S.-Y. Preparation and Cryogenic Mechanical Properties of Epoxy Resins Modified by Poly(Ethersulfone). J. Polym. Sci. Part A Polym. Chem.; 2008; 46, pp. 612-624. [DOI: https://dx.doi.org/10.1002/pola.22409]

18. Bagheri, R.; Marouf, B.T.; Pearson, R.A. Rubber-Toughened Epoxies: A Critical Review. Polym. Rev.; 2009; 49, pp. 201-225. [DOI: https://dx.doi.org/10.1080/15583720903048227]

19. Johnsen, B.B.; Kinloch, A.J.; Mohammed, R.D.; Taylor, A.C.; Sprenger, S. Toughening Mechanisms of Nanoparticle-Modified Epoxy Polymers. Polymer; 2007; 48, pp. 530-541. [DOI: https://dx.doi.org/10.1016/j.polymer.2006.11.038]

20. Tsang, W.L.; Taylor, A.C. Fracture and Toughening Mechanisms of Silica- and Core–Shell Rubber-Toughened Epoxy at Ambient and Low Temperature. J. Mater. Sci.; 2019; 54, pp. 13938-13958. [DOI: https://dx.doi.org/10.1007/s10853-019-03893-y]

21. Mi, X.; Liang, N.; Xu, H.; Wu, J.; Jiang, Y.; Nie, B.; Zhang, D. Toughness and Its Mechanisms in Epoxy Resins. Prog. Mater. Sci.; 2022; 130, 100977. [DOI: https://dx.doi.org/10.1016/j.pmatsci.2022.100977]

22. Guo, L.; Chen, Z.; Han, H.; Liu, G.; Luo, M.; Cui, N.; Dong, H.; Li, M.-Z. Advances and Outlook in Modified Graphene Oxide (GO)/Epoxy Composites for Mechanical Applications. Appl. Nanosci.; 2023; 13, pp. 3273-3287. [DOI: https://dx.doi.org/10.1007/s13204-022-02653-w]

23. Huang, X.; Zhi, C.; Jiang, P.; Golberg, D.; Bando, Y.; Tanaka, T. Polyhedral Oligosilsesquioxane-Modified Boron Nitride Nanotube Based Epoxy Nanocomposites: An Ideal Dielectric Material with High Thermal Conductivity. Adv. Funct. Mater.; 2013; 23, pp. 1824-1831. [DOI: https://dx.doi.org/10.1002/adfm.201201824]

24. Shi, X.; Wu, J.; Wang, X.; Zhou, X.; Xie, X.; Xue, Z. Novel Sound Insulation Materials Based on Epoxy/Hollow Silica Nanotubes Composites. Compos. Part B Eng.; 2017; 131, pp. 125-133. [DOI: https://dx.doi.org/10.1016/j.compositesb.2017.07.055]

25. Kim, M.-G.; Moon, J.-B.; Kim, C.-G. Effect of CNT Functionalization on Crack Resistance of a Carbon/Epoxy Composite at a Cryogenic Temperature. Compos. Part A Appl. Sci. Manuf.; 2012; 43, pp. 1620-1627. [DOI: https://dx.doi.org/10.1016/j.compositesa.2012.04.001]

26. Yan, M.; Jiao, W.; Li, J.; Huang, Y.; Chu, Z.; Chen, X.; Shen, F.; Wang, Y.; Wang, R.; He, X. Enhancement of the Cryogenic-Interfacial-Strength of Carbon Fiber Composites by Chemical Grafting of Graphene Oxide/Attapulgite on T300. Polym. Compos.; 2020; 41, pp. 5072-5081. [DOI: https://dx.doi.org/10.1002/pc.25775]

27. Meng, J.; Zheng, H.; Wei, Y.; Liu, D.; Shi, H.; Wang, P.; Lei, H.; Fang, D. Leakage Performance of CFRP Laminate under Cryogenic Temperature: Experimental and Simulation Study. Compos. Sci. Technol.; 2022; 226, 109550. [DOI: https://dx.doi.org/10.1016/j.compscitech.2022.109550]

28. Guo, F.-L.; Tan, D.; Wu, T.; Huang, P.; Li, Y.-Q.; Hu, N.; Fu, S.-Y. Experimental Characterization and Molecular Dynamics Simulation of Thermal Stability, Mechanical Properties and Liquid Oxygen Compatibility of Multiple Epoxy Systems for Cryotank Applications. Extrem. Mech. Lett.; 2021; 44, 101227. [DOI: https://dx.doi.org/10.1016/j.eml.2021.101227]

29. Zhou, Z.; Qian, J.; Zhang, J.; Cui, Y.; Li, Y.; Tan, D.; Long, J.; Wu, H.; Liu, Q.; Guo, F. et al. Phosphorus and Bromine Modified Epoxy Resin with Enhanced Cryogenic Mechanical Properties and Liquid Oxygen Compatibility Simultaneously. Polym. Test.; 2021; 94, 107051. [DOI: https://dx.doi.org/10.1016/j.polymertesting.2021.107051]

30. Liu, N.; Wang, H.; Ma, B.; Xu, B.; Qu, L.; Fang, D.; Yang, Y. Enhancing Cryogenic Mechanical Properties of Epoxy Resins Toughened by Biscitraconimide Resin. Compos. Sci. Technol.; 2022; 220, 109252. [DOI: https://dx.doi.org/10.1016/j.compscitech.2021.109252]

31. Wu, T.; Liu, Y.; Li, N.; Huang, G.-W.; Qu, C.-B.; Xiao, H.-M. Cryogenic Mechanical Properties of Epoxy Resin Toughened by Hydroxyl-Terminated Polyurethane. Polym. Test.; 2019; 74, pp. 45-56. [DOI: https://dx.doi.org/10.1016/j.polymertesting.2018.11.048]

32. Chen, Z.-K.; Yang, J.-P.; Ni, Q.-Q.; Fu, S.-Y.; Huang, Y.-G. Reinforcement of Epoxy Resins with Multi-Walled Carbon Nanotubes for Enhancing Cryogenic Mechanical Properties. Polymer; 2009; 50, pp. 4753-4759. [DOI: https://dx.doi.org/10.1016/j.polymer.2009.08.001]

33. Chang, W.; Rose, L.R.F.; Islam, M.S.; Wu, S.; Peng, S.; Huang, F.; Kinloch, A.J.; Wang, C.H. Strengthening and Toughening Epoxy Polymer at Cryogenic Temperature Using Cupric Oxide Nanorods. Compos. Sci. Technol.; 2021; 208, 108762. [DOI: https://dx.doi.org/10.1016/j.compscitech.2021.108762]

34. Wang, X.Q.; Jian, W.; Buyukozturk, O.; Leung, C.K.Y.; Lau, D. Degradation of Epoxy/Glass Interface in Hygrothermal Environment: An Atomistic Investigation. Compos. Part B Eng.; 2021; 206, 108534. [DOI: https://dx.doi.org/10.1016/j.compositesb.2020.108534]

35. Kwon, D.-J.; Wang, Z.-J.; Choi, J.-Y.; Shin, P.-S.; DeVries, K.L.; Park, J.-M. Interfacial Evaluation of Carbon Fiber/Epoxy Composites Using Electrical Resistance Measurements at Room and a Cryogenic Temperature. Compos. Part A Appl. Sci. Manuf.; 2015; 72, pp. 160-166. [DOI: https://dx.doi.org/10.1016/j.compositesa.2015.02.007]

36. Shao, Y.; Xu, F.; Liu, W.; Zhou, M.; Li, W.; Hui, D.; Qiu, Y. Influence of Cryogenic Treatment on Mechanical and Interfacial Properties of Carbon Nanotube Fiber/Bisphenol-F Epoxy Composite. Compos. Part B Eng.; 2017; 125, pp. 195-202. [DOI: https://dx.doi.org/10.1016/j.compositesb.2017.05.077]

37. Gao, B.; Zhang, J.; Hao, Z.; Huo, L.; Zhang, R.; Shao, L. In-Situ Modification of Carbon Fibers with Hyperbranched Polyglycerol via Anionic Ring-Opening Polymerization for Use in High-Performance Composites. Carbon; 2017; 123, pp. 548-557. [DOI: https://dx.doi.org/10.1016/j.carbon.2017.08.008]

38. Zhang, L.; De Greef, N.; Kalinka, G.; Van Bilzen, B.; Locquet, J.-P.; Verpoest, I.; Seo, J.W. Carbon Nanotube-Grafted Carbon Fiber Polymer Composites: Damage Characterization on the Micro-Scale. Compos. Part B Eng.; 2017; 126, pp. 202-210. [DOI: https://dx.doi.org/10.1016/j.compositesb.2017.06.004]

39. Gojny, F.H.; Wichmann, M.H.G.; Köpke, U.; Fiedler, B.; Schulte, K. Carbon Nanotube-Reinforced Epoxy-Composites: Enhanced Stiffness and Fracture Toughness at Low Nanotube Content. Compos. Sci. Technol.; 2004; 64, pp. 2363-2371. [DOI: https://dx.doi.org/10.1016/j.compscitech.2004.04.002]

40. Garcia, E.J.; Wardle, B.L.; John Hart, A. Joining Prepreg Composite Interfaces with Aligned Carbon Nanotubes. Compos. Part A Appl. Sci. Manuf.; 2008; 39, pp. 1065-1070. [DOI: https://dx.doi.org/10.1016/j.compositesa.2008.03.011]

41. Yao, Z.; Wang, C.; Qin, J.; Su, S.; Wang, Y.; Wang, Q.; Yu, M.; Wei, H. Interfacial Improvement of Carbon Fiber/Epoxy Composites Using One-Step Method for Grafting Carbon Nanotubes on the Fibers at Ultra-Low Temperatures. Carbon; 2020; 164, pp. 133-142. [DOI: https://dx.doi.org/10.1016/j.carbon.2020.03.060]

42. Patnaik, S.; Gangineni, P.K.; Prusty, R.K. Influence of Cryogenic Temperature on Mechanical Behavior of Graphene Carboxyl Grafted Carbon Fiber Reinforced Polymer Composites: An Emphasis on Concentration of Nanofillers. Compos. Commun.; 2020; 20, 100369. [DOI: https://dx.doi.org/10.1016/j.coco.2020.100369]

43. Li, S.; Chen, D.; Gao, C.; Yuan, Y.; Wang, H.; Liu, X.; Hu, B.; Ma, J.; Liu, M.; Wu, Z. Epoxy-Functionalized Polysiloxane/Nano-SiO₂ Synergistic Reinforcement in Cryogenic Mechanical Properties of Epoxy and Carbon Fiber Reinforced Epoxy Laminate. Compos. Sci. Technol.; 2020; 198, 108292. [DOI: https://dx.doi.org/10.1016/j.compscitech.2020.108292]

44. Zhao, Y.; Huang, R.; Wu, Z.; Zhang, H.; Zhou, Z.; Li, L.; Dong, Y.; Luo, M.; Ye, B.; Zhang, H. Effect of Free Volume on Cryogenic Mechanical Properties of Epoxy Resin Reinforced by Hyperbranched Polymers. Mater. Des.; 2021; 202, 109565. [DOI: https://dx.doi.org/10.1016/j.matdes.2021.109565]

45. He, Y.; Li, Q.; Kuila, T.; Kim, N.H.; Jiang, T.; Lau, K.; Lee, J.H. Micro-Crack Behavior of Carbon Fiber Reinforced Thermoplastic Modified Epoxy Composites for Cryogenic Applications. Compos. Part B Eng.; 2013; 44, pp. 533-539. [DOI: https://dx.doi.org/10.1016/j.compositesb.2012.03.014]

46. Qu, C.-B.; Wu, T.; Huang, G.-W.; Li, N.; Li, M.; Ma, J.-L.; Liu, Y.; Xiao, H.-M. Improving Cryogenic Mechanical Properties of Carbon Fiber Reinforced Composites Based on Epoxy Resin Toughened by Hydroxyl-Terminated Polyurethane. Compos. Part B Eng.; 2021; 210, 108569. [DOI: https://dx.doi.org/10.1016/j.compositesb.2020.108569]

47. Meng, J.; Wang, Y.; Yang, H.; Wang, P.; Lei, Q.; Shi, H.; Lei, H.; Fang, D. Mechanical Properties and Internal Microdefects Evolution of Carbon Fiber Reinforced Polymer Composites: Cryogenic Temperature and Thermocycling Effects. Compos. Sci. Technol.; 2020; 191, 108083. [DOI: https://dx.doi.org/10.1016/j.compscitech.2020.108083]

48. Jia, Z.; Li, T.; Chiang, F.; Wang, L. An Experimental Investigation of the Temperature Effect on the Mechanics of Carbon Fiber Reinforced Polymer Composites. Compos. Sci. Technol.; 2018; 154, pp. 53-63. [DOI: https://dx.doi.org/10.1016/j.compscitech.2017.11.015]

49. Cheng, X.; Liu, L.; Feng, X.; Shen, L.; Wu, Z. Low Temperature-Based Flexural Properties of Carbon Fiber/Epoxy Composite Laminates Incorporated with Carbon Nanotube Sheets. Macromol. Mater. Eng.; 2019; 304, 1900247. [DOI: https://dx.doi.org/10.1002/mame.201900247]

50. Li, S.; Chen, D.; Yuan, Y.; Gao, C.; Cui, Y.; Wang, H.; Liu, X.; Liu, M.; Wu, Z. Influence of Flexible Molecular Structure on the Cryogenic Mechanical Properties of Epoxy Matrix and Carbon Fiber/Epoxy Composite Laminate. Mater. Des.; 2020; 195, 109028. [DOI: https://dx.doi.org/10.1016/j.matdes.2020.109028]

51. Nishijima, S.; Honda, Y.; Okada, T. Application of the Positron Annihilation Method for Evaluation of Organic Materials for Cryogenic Use. Cryogenics; 1995; 35, pp. 779-781. [DOI: https://dx.doi.org/10.1016/0011-2275(95)90913-Z]

52. Nishijima, S.; Honda, Y.; Tagawa, S.; Okada, T. Study of Epoxy Resin for Cryogenic Use by Positron Annihilation Method. J. Radioanal. Nucl. Chem.; 2005; 211, pp. 93-101. [DOI: https://dx.doi.org/10.1007/BF02036260]

53. Wang, J.; Chang, W.; Islam, M.S.; Huang, F.; Wu, S.; Rose, L.R.F.; Zhang, J.; Wang, C.H. Toughening Epoxy by Nano-Structured Block Copolymer to Mitigate Matrix Microcracking of Carbon Fibre Composites at Cryogenic Temperatures. Compos. Sci. Technol.; 2024; 251, 110548. [DOI: https://dx.doi.org/10.1016/j.compscitech.2024.110548]

54. Islam, M.S.; Benninger, L.F.; Pearce, G.; Wang, C.-H. Toughening Carbon Fibre Composites at Cryogenic Temperatures Using Low-Thermal Expansion Nanoparticles. Compos. Part A Appl. Sci. Manuf.; 2021; 150, 106613. [DOI: https://dx.doi.org/10.1016/j.compositesa.2021.106613]

55. Carolan, D.; Ivankovic, A.; Kinloch, A.J.; Sprenger, S.; Taylor, A.C. Toughened Carbon Fibre-Reinforced Polymer Composites with Nanoparticle-Modified Epoxy Matrices. J. Mater. Sci.; 2017; 52, pp. 1767-1788. [DOI: https://dx.doi.org/10.1007/s10853-016-0468-5]

56. He, Y.; Zhang, L.; Chen, G.; Li, X.; Yao, D.; Lee, J.H.; Zhang, Y. Surface Functionalized Carbon Nanotubes and Its Effects on the Mechanical Properties of Epoxy Based Composites at Cryogenic Temperature. Polym. Bull.; 2014; 71, pp. 2465-2485. [DOI: https://dx.doi.org/10.1007/s00289-014-1202-6]

57. Quan, D.; Urdániz, J.L.; Ivanković, A. Enhancing Mode-I and Mode-II Fracture Toughness of Epoxy and Carbon Fibre Reinforced Epoxy Composites Using Multi-Walled Carbon Nanotubes. Mater. Des.; 2018; 143, pp. 81-92. [DOI: https://dx.doi.org/10.1016/j.matdes.2018.01.051]

58. Feng, Q.-P.; Deng, Y.-H.; Xiao, H.-M.; Liu, Y.; Qu, C.-B.; Zhao, Y.; Fu, S.-Y. Enhanced Cryogenic Interfacial Normal Bond Property between Carbon Fibers and Epoxy Matrix by Carbon Nanotubes. Compos. Sci. Technol.; 2014; 104, pp. 59-65. [DOI: https://dx.doi.org/10.1016/j.compscitech.2014.09.006]

59. He, Y.; Yang, S.; Liu, H.; Shao, Q.; Chen, Q.; Lu, C.; Jiang, Y.; Liu, C.; Guo, Z. Reinforced Carbon Fiber Laminates with Oriented Carbon Nanotube Epoxy Nanocomposites: Magnetic Field Assisted Alignment and Cryogenic Temperature Mechanical Properties. J. Colloid Interface Sci.; 2018; 517, pp. 40-51. [DOI: https://dx.doi.org/10.1016/j.jcis.2018.01.087]

60. Shen, X.-J.; Liu, Y.; Xiao, H.-M.; Feng, Q.-P.; Yu, Z.-Z.; Fu, S.-Y. The Reinforcing Effect of Graphene Nanosheets on the Cryogenic Mechanical Properties of Epoxy Resins. Compos. Sci. Technol.; 2012; 72, pp. 1581-1587. [DOI: https://dx.doi.org/10.1016/j.compscitech.2012.06.021]

61. Qu, C.-B.; Huang, Y.; Li, F.; Xiao, H.-M.; Liu, Y.; Feng, Q.-P.; Huang, G.-W.; Li, N.; Fu, S.-Y. Enhanced Cryogenic Mechanical Properties of Carbon Fiber Reinforced Epoxy Composites by Introducing Graphene Oxide. Compos. Commun.; 2020; 22, 100480. [DOI: https://dx.doi.org/10.1016/j.coco.2020.100480]

62. Hung, P.; Lau, K.; Qiao, K.; Fox, B.; Hameed, N. Property Enhancement of CFRP Composites with Different Graphene Oxide Employment Methods at a Cryogenic Temperature. Compos. Part A Appl. Sci. Manuf.; 2019; 120, pp. 56-63. [DOI: https://dx.doi.org/10.1016/j.compositesa.2019.02.012]

63. Zhu, J.; Imam, A.; Crane, R.; Lozano, K.; Khabashesku, V.N.; Barrera, E.V. Processing a Glass Fiber Reinforced Vinyl Ester Composite with Nanotube Enhancement of Interlaminar Shear Strength. Compos. Sci. Technol.; 2007; 67, pp. 1509-1517. [DOI: https://dx.doi.org/10.1016/j.compscitech.2006.07.018]

64. Zhao, Z.; Teng, K.; Li, N.; Li, X.; Xu, Z.; Chen, L.; Niu, J.; Fu, H.; Zhao, L.; Liu, Y. Mechanical, Thermal and Interfacial Performances of Carbon Fiber Reinforced Composites Flavored by Carbon Nanotube in Matrix/Interface. Compos. Struct.; 2017; 159, pp. 761-772. [DOI: https://dx.doi.org/10.1016/j.compstruct.2016.10.022]

65. He, Y.; Chen, Q.; Yang, S.; Lu, C.; Feng, M.; Jiang, Y.; Cao, G.; Zhang, J.; Liu, C. Micro-Crack Behavior of Carbon Fiber Reinforced Fe₃O₄/Graphene Oxide Modified Epoxy Composites for Cryogenic Application. Compos. Part A Appl. Sci. Manuf.; 2018; 108, pp. 12-22. [DOI: https://dx.doi.org/10.1016/j.compositesa.2018.02.014]

66. Sprenger, S. Fiber-Reinforced Composites Based on Epoxy Resins Modified with Elastomers and Surface-Modified Silica Nanoparticles. J. Mater. Sci.; 2014; 49, pp. 2391-2402. [DOI: https://dx.doi.org/10.1007/s10853-013-7963-8]

67. Hsieh, T.H.; Kinloch, A.J.; Masania, K.; Taylor, A.C.; Sprenger, S. The Mechanisms and Mechanics of the Toughening of Epoxy Polymers Modified with Silica Nanoparticles. Polymer; 2010; 51, pp. 6284-6294. [DOI: https://dx.doi.org/10.1016/j.polymer.2010.10.048]

68. Rosso, P.; Ye, L.; Friedrich, K.; Sprenger, S. A Toughened Epoxy Resin by Silica Nanoparticle Reinforcement. J. Appl. Polym. Sci.; 2006; 100, pp. 1849-1855. [DOI: https://dx.doi.org/10.1002/app.22805]

69. Bajpai, A.; Carlotti, S. The Effect of Hybridized Carbon Nanotubes, Silica Nanoparticles, and Core-Shell Rubber on Tensile, Fracture Mechanics and Electrical Properties of Epoxy Nanocomposites. Nanomaterials; 2019; 9, 1057. [DOI: https://dx.doi.org/10.3390/nano9071057]

70. Leelachai, K.; Kongkachuichay, P.; Dittanet, P. Toughening of Epoxy Hybrid Nanocomposites Modified with Silica Nanoparticles and Epoxidized Natural Rubber. J. Polym. Res.; 2017; 24, 41. [DOI: https://dx.doi.org/10.1007/s10965-017-1202-y]

71. Liang, W.; Duan, Z.; Sun, K.; Liang, H.; Wang, Z.; Guo, C. Mechanical Performance of Nano-Al₂O₃-Modified Composites at Cryogenic and Room Temperatures. Polym. Compos.; 2018; 39, pp. 3006-3012. [DOI: https://dx.doi.org/10.1002/pc.24301]

72. Sun, T.; Fan, H.; Wang, Z.; Liu, X.; Wu, Z. Modified Nano Fe₂O₃-Epoxy Composite with Enhanced Mechanical Properties. Mater. Des.; 2015; 87, pp. 10-16. [DOI: https://dx.doi.org/10.1016/j.matdes.2015.07.177]

73. Li, J.; Peng, C.; Li, Z.; Wu, Z.; Li, S. The Improvement in Cryogenic Mechanical Properties of Nano-ZrO₂/Epoxy Composites via Surface Modification of Nano-ZrO₂. RSC Adv.; 2016; 6, pp. 61393-61401. [DOI: https://dx.doi.org/10.1039/C6RA08047B]

74. Eker, Y.R.; Özcan, M.; Özkan, A.O.; Kırkıcı, H. The Influence of Al₂O₃ and TiO₂ Additives on the Electrical Resistivity of Epoxy Resin-Based Composites at Low Temperature. Macromol. Mater. Eng.; 2019; 304, 1800670. [DOI: https://dx.doi.org/10.1002/mame.201800670]

75. Islam, M.S.; Chang, W.; Sha, Z.; Wang, J.; Wu, S.; Rose, L.R.F.; Kinloch, A.J.; Wang, C.H. Mitigating Cryogenic Microcracking in Carbon-Fibre Reinforced Polymer Composites Using Negative Thermal-Expansion Nanoparticles Functionalized by a Polydopamine Coating. Compos. Part B Eng.; 2023; 257, 110676. [DOI: https://dx.doi.org/10.1016/j.compositesb.2023.110676]

76. Mishra, K.; Gidley, D.; Singh, R.P. Influence of Self-Assembled Compliant Domains on the Polymer Network and Mechanical Properties of POSS-Epoxy Nanocomposites under Cryogenic Conditions. Eur. Polym. J.; 2019; 116, pp. 283-290. [DOI: https://dx.doi.org/10.1016/j.eurpolymj.2019.02.034]

77. Zhao, Y.; Chen, Z.-K.; Liu, Y.; Xiao, H.-M.; Feng, Q.-P.; Fu, S.-Y. Simultaneously Enhanced Cryogenic Tensile Strength and Fracture Toughness of Epoxy Resins by Carboxylic Nitrile-Butadiene Nano-Rubber. Compos. Part A Appl. Sci. Manuf.; 2013; 55, pp. 178-187. [DOI: https://dx.doi.org/10.1016/j.compositesa.2013.09.005]

78. Andújar, J.M.; Segura, F.; Rey, J.; Vivas, F.J. Batteries and Hydrogen Storage: Technical Analysis and Commercial Revision to Select the Best Option. Energies; 2022; 15, 6196. [DOI: https://dx.doi.org/10.3390/en15176196]

79. Yan, Y.; Zhang, J.; Li, G.; Zhou, W.; Ni, Z. Review on Linerless Type V Cryo-Compressed Hydrogen Storage Vessels: Resin Toughening and Hydrogen-Barrier Properties Control. Renew. Sustain. Energy Rev.; 2024; 189, 114009. [DOI: https://dx.doi.org/10.1016/j.rser.2023.114009]

80. Zhang, J.; Yan, Y.; Zhang, C.; Xu, Z.; Li, X.; Zhao, G.; Ni, Z. Properties Improvement of Composite Layer of Cryo-Compressed Hydrogen Storage Vessel by Polyethylene Glycol Modified Epoxy Resin. Int. J. Hydrogen Energy; 2023; 48, pp. 5576-5594. [DOI: https://dx.doi.org/10.1016/j.ijhydene.2022.11.138]

81. Zhao, X.; Yan, Y.; Zhang, J.; Zhang, F.; Wang, Z.; Ni, Z. Analysis of Multilayered Carbon Fiber Winding of Cryo-Compressed Hydrogen Storage Vessel. Int. J. Hydrogen Energy; 2022; 47, pp. 10934-10946. [DOI: https://dx.doi.org/10.1016/j.ijhydene.2022.01.136]

82. Xing, J.; Geng, P.; Yang, T. Stress and Deformation of Multiple Winding Angle Hybrid Filament-Wound Thick Cylinder under Axial Loading and Internal and External Pressure. Compos. Struct.; 2015; 131, pp. 868-877. [DOI: https://dx.doi.org/10.1016/j.compstruct.2015.05.036]

83. Zhu, T.; Lu, C.; Lu, X.; Zhi, J.; Song, Y. Curing Process Optimization and Mechanical Properties Improvement of Epoxy Resin Copolymer Modified by Epoxy-Terminated Hyperbranched Polyether Sulfone. Polymer; 2022; 241, 124535. [DOI: https://dx.doi.org/10.1016/j.polymer.2022.124535]

84. Zhang, J.; Lei, L.; Zhou, W.; Li, G.; Yan, Y.; Ni, Z. Cryogenic Mechanical and Hydrogen-Barrier Properties of Carbon Fiber Composites for Type V Cryo-Compressed Hydrogen Storage Vessels. Compos. Commun.; 2023; 43, 101733. [DOI: https://dx.doi.org/10.1016/j.coco.2023.101733]

85. Yonemoto, K.; Yamamoto, Y.; Okuyama, K.; Ebina, T. Application of CFRP with High Hydrogen Gas Barrier Characteristics to Fuel Tanks of Space Transportation System. Trans. Jpn. Soc. Aeronaut. Space Sci. Space Technol. Jpn.; 2009; 7, pp. Pc_13-Pc_18. [DOI: https://dx.doi.org/10.2322/tstj.7.Pc_13]

86. Lu, Y.; Liu, T.; Wang, S.; Sun, Y.; Zhang, Y.; Kang, J.; Li, B.; Gao, Y.; Gao, X.; Fan, W. A Mild, Environmental, and Low-Destructive Recycling Strategy for the 3D Multiaxial Braided Composites Cured by the Vacuum-Assisted Resin Transformed Molding Process. Compos. Commun.; 2022; 35, 101354. [DOI: https://dx.doi.org/10.1016/j.coco.2022.101354]

87. Zhang, T.; Li, Y.; Yuan, Y.; Cui, K.; Wei, W.; Wu, J.; Qin, W.; Wu, X. Spatially Confined Bi2O3–Ti3C2Tx Hybrids Reinforced Epoxy Composites for Gamma Radiation Shielding. Compos. Commun.; 2022; 34, 101252. [DOI: https://dx.doi.org/10.1016/j.coco.2022.101252]

88. Chang, Y.; Zhu, Y.; Xu, F.; Zhang, M.; Liu, Y.; Wang, H. A β-Phase Crystal Poly (Vinylidene Fluoride) Incorporated Epoxy-Based Composite Coating with Excellent Oxygen Barrier and Anti-Corrosion. Compos. Commun.; 2023; 37, 101458. [DOI: https://dx.doi.org/10.1016/j.coco.2022.101458]

89. Flanagan, M.; Grogan, D.M.; Goggins, J.; Appel, S.; Doyle, K.; Leen, S.B.; Ó Brádaigh, C.M. Permeability of Carbon Fibre PEEK Composites for Cryogenic Storage Tanks of Future Space Launchers. Compos. Part A Appl. Sci. Manuf.; 2017; 101, pp. 173-184. [DOI: https://dx.doi.org/10.1016/j.compositesa.2017.06.013]

90. Lei, L.; Zhang, J.; Li, G.; Ni, Z.; Yan, Y. Improve the Gas-Barrier and Mechanical Performance of Epoxy Resin by Co-Modification from Montmorillonite and Polyethylene Glycol. Int. J. Hydrogen Energy; 2024; 65, pp. 751-758. [DOI: https://dx.doi.org/10.1016/j.ijhydene.2024.04.037]

91. van Rooyen, L.J.; Karger-Kocsis, J.; David Kock, L. Improving the Helium Gas Barrier Properties of Epoxy Coatings through the Incorporation of Graphene Nanoplatelets and the Influence of Preparation Techniques. J. Appl. Polym. Sci.; 2015; 132, 42584. [DOI: https://dx.doi.org/10.1002/app.42584]

92. Ogasawara, T.; Ishida, Y.; Ishikawa, T.; Aoki, T.; Ogura, T. Helium Gas Permeability of Montmorillonite/Epoxy Nanocomposites. Compos. Part A Appl. Sci. Manuf.; 2006; 37, pp. 2236-2240. [DOI: https://dx.doi.org/10.1016/j.compositesa.2006.02.015]

93. Azeredo, H.M.C. de Nanocomposites for Food Packaging Applications. Food Res. Int.; 2009; 42, pp. 1240-1253. [DOI: https://dx.doi.org/10.1016/j.foodres.2009.03.019]

94. Baghitabar, K.; Jamshidi, M.; Ghamarpoor, R. Exfoliation, Hydroxylation and Silanization of Two-Dimensional (2D) Montmorillonites (MMTs) and Evaluation of the Effects on Tire Rubber Properties. Polym. Test.; 2023; 129, 108265. [DOI: https://dx.doi.org/10.1016/j.polymertesting.2023.108265]

95. Ghamarpoor, R.; Jamshidi, M.; Fallah, A.; Eftekharipour, F. Preparation of Dual-Use GPTES@ZnO Photocatalyst from Waste Warm Filter Cake and Evaluation of Its Synergic Photocatalytic Degradation for Air-Water Purification. J. Environ. Manag.; 2023; 342, 118352. [DOI: https://dx.doi.org/10.1016/j.jenvman.2023.118352] [PubMed: https://www.ncbi.nlm.nih.gov/pubmed/37311344]

96. Gu, H.; Ma, C.; Liang, C.; Meng, X.; Gu, J.; Guo, Z. A Low Loading of Grafted Thermoplastic Polystyrene Strengthens and Toughens Transparent Epoxy Composites. J. Mater. Chem. C; 2017; 5, pp. 4275-4285. [DOI: https://dx.doi.org/10.1039/C7TC00437K]

97. Ghamarpoor, R.; Fallah, A.; Jamshidi, M.; Salehfekr, S. Using Waste Silver Metal in Synthesis of Z-Scheme Ag@WO₃-CeO₂ Heterojunction to Increase Photodegradation and Electrochemical Performances. J. Ind. Eng. Chem.; 2023; 128, pp. 459-471. [DOI: https://dx.doi.org/10.1016/j.jiec.2023.08.010]