Mechanism of the anisotropic nitroguanidine crystal arrangement on triple-base propellant failure by impact and strategy of structural enhancement

High-performance propellants are crucial components for the powerful firing capabilities of future artillery. Currently, the triple-base propellants in active service can be utilized to charge field artillery, such as large-caliber cannon howitzers. To meet the demands of modern weapon development, specifically large-caliber and long-range artillery, and to enhance the power of high-bore pressure artillery in use, large-size launch charges with high energy, stable, and safe mechanical properties and improved environmental adaptability are essential.^1–6 

Decades of persistent research on the safety of propellant charge firing has revealed that the rupture of propellant grains is primarily caused by the unstable mechanical response due to their anisotropic structure. Many triple-base propellant materials exhibit discontinuous and heterogeneous microstructures, leading to issues if treated as isotropic materials. The high number of unstable propellant grains within a propellant charge inevitably results in uncontrollable energy release patterns under complex impacts. Therefore, it is crucial to differentiate between isotropic and anisotropic materials, studying their physical and mechanical properties separately. Anisotropy refers to the variations in physical and mechanical properties, structural characteristics, and damage modes of triple-base propellants with respect to thickness layer, action surface, and angle. As a result of this anisotropy, triple-base propellants may display distinct deformation characteristics, elastic moduli, and mechanical strengths under different loading directions.^7–10 

Based on internal tissue morphology analysis results, the influence of microstructural features on the mechanical properties of solvent-based triple-base propellants (STPs) was thoroughly examined. Additionally, the radial tensile strength of STP is considerably smaller at low temperatures, explaining the anomalous ballistic performance at low temperatures in actual applications.^11–14 The primary reason for this is the significant difference between axial and radial mechanical properties at low temperatures, making the propellant prone to cracking and fragmentation along the axial direction under high-pressure and high-speed impacts in the chamber. Consequently, triple-base propellants have a limited range of applications, with the relatively poor radial mechanical strength becoming a technical challenge in developing high-performance propellants. Therefore, it is vital to focus on damage modes and characteristics, crack initiation location, expansion, and penetration evolution of propellant system structures under complex shock wave action.

Triple-base propellants are typical highly particle-filled viscoelastic materials,^15–17 exhibiting complex mechanical behavior under various environmental conditions and load categories. Numerous experiments have been conducted recently to investigate the mechanical properties of novel triple-base propellants, primarily focusing on mechanical structural properties under a single loading direction. However, these technical parameters are not suitable for evaluating the architectural integrity of propellant grains from ignition to peak pressure in a large-caliber cannon howitzer as they do not accurately reflect the complex mechanical behavior of propellant grains under high-pressure and high-speed shock waves.^18–20 During the firing process, propellant grains are subjected to triaxial boosted impingement, with their surfaces and internal structures experiencing high pressure and high-speed overloads in multiple directions.

In this study, we initially compared the molding processes of solvent type and solventless type and then obtained STP and solventless triple-base propellants (SLTP) samples. The crystal orientation and matrix-filler interface structure were analyzed. In light of the significant differences in the mechanical structure of novel triple-base propellants, the progressive damage performance in three coordinate directions, such as quasi-static tensile properties, quasi-static compressive properties, and dynamic impact resistance, has been investigated in depth. Finally, the failure mechanism and damage model were proposed.

The nitrocellulose (NC) paste employed in this study, which absorbed nitroglycerine (NG) and diethyleneglycol dinitrate (DEGDN), was procured from Sichuan Nitrocell Co., Ltd. (Sichuan, China). The composition of this paste included 60.0 wt % NC (12.5% nitration level), 28.0 wt % NG, 9.5 wt % DEGDN, 2.0 wt % dimethyl phenyl urea (C₂), and 0.5 wt % TiO₂. Nitroguanidine (NQ, purity 99%), characterized by its elongated needle-shaped structure, was selected for this study and obtained from Liaoning Qingyang Specialty Chemical Industry Group Co., Ltd. (Liaoning, China). Figure 1 displays the crystal sizes, morphologies, and micrographs of the micro/nanoscale pure needle-shaped NQ, revealing the polygonal cross sections of the elongated needle NQ crystals and the angular details along the needle-shaped structures. Analytical reagent grade acetone, ethanol (1:0.95 v/v), and surfactant (emulsifier OP-10) were purchased from Aladdin Chemical Reagent Co., Ltd. (Shanghai, China) and utilized as cosolvents in the propellant preparation process. All chemicals were employed without any additional treatment.

The macroscopic appearance of raw NQ presents as a white, smooth, low-density powder. The scanning electron micrograph image (Fig. 1) allowed a more detailed view on the microscopic morphology of the industrial grade nitroguanidine crystals. It can be seen that the industrial grade nitroguanidine crystals had a needle-shaped longitudinal appearance with a length of about 20–80 µm and a diameter of about 2–8 µm.

The schematic diagram of the fabrication process for the STP and SLTP strands is shown in Fig. 2. Experimental parameters and the compositions of STP and SLTP with a different gelating production method are shown in Table I.

The aforementioned two samples were sectioned into various solid geometries, such as cylinders, cubes, and dumbbell-shaped pillars, for the evaluation of mechanical properties, including tensile strength, compression strength, and impact properties, respectively. The preparation of these propellants followed the methodology delineated in the General Specification for propellant GJB 1529A-2001.

In this study, the specimens were systematically characterized and evaluated using a variety of techniques, including scanning electron microscopy (SEM), micro x-ray computed tomography (micro-CT), quasi-static tensile resistance testing, quasi-static compression resistance testing, and rapid impact resistance testing. Comprehensive descriptions of these characterization techniques and test methodologies can be found in the supplementary material.

The 3D visualizations of the microstructure for both STP and SLTP samples are presented in Fig. 3, which are assembled by combining orthogonal sections of tomographs from longitudinal specimens. These visualizations offer a full-field, high-resolution 3D view of deformation, in addition to the information regarding the long needle-shaped NQ orientation within the volume element, as depicted in Fig. 3 (the standard micro-CT microtomography results). The segmentation into fibers and matrix, based on their relative x-ray attenuation (i.e., gray levels), is shown in Figs. 3(a) and 3(c), where white represents the long needle-shaped NQ and orange represents the composite matrix. As seen in Figs. 3(a) and 3(c), the areas between the long needle-shaped NQ are filled with the composite matrix. Visualizations of the long needle-shaped NQ [Figs. 3(b) and 3(d)] display two distinct arrangements of NQ and the continuous structure of the matrix, also highlighting the primary differences between the two molding processes concerning NQ crystal orientation. The green color signifies a tendency toward the vertical direction, i.e., the extrusion molding direction, while blue and red colors indicate a tendency toward the horizontal direction, i.e., perpendicular to the extrusion molding direction. It is evident that the majority of NQ in STP samples are arranged in a nearly unidirectional and highly oriented manner, while the NQ in SLTP samples exhibit a more random and disordered orientation. In summary, the highly alignment-oriented NQ in STP samples are oriented parallel to the direction of extrusion driving force load, while the randomly disordered NQ are distributed throughout the matrix of SLTP samples.^21–23 

Further quantitative measurements were conducted to assess the alignment state of rigid NQ particles in STP and SLTP samples. In the STP samples, the average alignment of rigid NQ particles with respect to the molding axial angle Theta ranged from 72.6° to 105.5°. Simultaneously, the average alignment of rigid NQ particles with respect to the molding axial angle Theta in the STP samples spanned from 2.9° to 167.7°. The arrangement of the rigid NQ particles in these two samples exhibits a significant difference in orientation distribution at the macroscopic level. The distribution of NQ crystals in STP samples was highly aligned, ∼80% better than that of the SLTP samples. Consequently, the NQ particles in the SLTP propellant column are arranged in a more disordered orientation compared to the STP propellant column.

In the STP samples, the alignment tendency of needle-shaped NQ crystals is strictly oriented, with crystals connected back and forth and parallel to each other. In contrast, the alignment tendency of needle-shaped NQ crystals in SLTP samples is undirected, with crystals interspersed horizontally and vertically to provide mutual support. The combination of NQ rods is separated and surrounded by an NC/NG/DEGDN matrix. Additionally, the SLTP samples possess a mutually supported three-dimensional spatial structure, forming numerous stable “triangular” and “fence” structures for well-rounded force transfer. This spatial structure contributes to the dispersion of external impact, leading to excellent mechanical properties. In comparison, STP specimens display a significant planar laminar structure, which does not disperse and buffer stresses effectively. It is evident that the crystal morphology and spatial arrangement of NQ play a crucial role in the mechanical structure of propellants.^24–26 

In this paper, α is set to be 3452, T_α is set to be 100 °C, η_∞ is 4004 Pa s, η₀ is set to be 8.7 × 10⁵ Pa s, λ is set to be −0.25, and n is set to be 0.14. The density of the propellant material was 1650 kg/m³, the specific heat capacity was 1000 J/(kg K), and the thermal conductivity was 0.0454 W/(m K). The extruded fluid is an incompressible viscous laminar fluid with an inlet velocity of 0.09 mm/s and an outlet pressure of 0 (relative to atmospheric pressure).

The STP and SLTP sample grains are extruded using an annular capillary die where the material flow exhibits non-uniform shear rate and velocity distributions (Poiseuille flow), with a higher shear rate near the die walls, as shown in Fig. 4. The primary reason for the differences in internal organizational structure is that the needle-shaped NQ crystals in the SLTP samples are unable to reorganize sufficiently in an oriented manner due to strong surface friction and viscous flow restriction during the extrusion process.

The microstructural analysis of cold brittle fracture surfaces in both STP and SLTP samples revealed a distinct difference in the interfacial structure between the two sample types. Additionally, the test was also utilized to analyze the interaction between the NQ and the matrix, as well as the impact of high temperature and pressure on the bonded surface.^27,28

As depicted in Fig. 5, the SEM image of the STP samples’ fracture surface exhibits a smooth texture, characterized by large and shallow fracture areas. Notably, evident debonding of NQ and cavities following peeling is observed in the dimple regions. In Fig. 5(c), NQ-matrix bridging and breakage are identified on both sides of the crack, with the bridging fibers’ length and fracture width extending ∼300–400 nm. This suggests that the bridging crack and NQ-matrix breakage stem from post-bonding separation. In contrast to STP samples, the NQ surface of SLTP samples exhibits a more abundant matrix coating, as shown in Fig. 5(e), implying enhanced adhesion. This improvement in the SLTP samples can be attributed to the distinct bonding interface between the filler and composite matrix. In the SLTP samples, no NQ crystal debonding interface is apparent. Moreover, a more compact structure may be achieved by uniformly dispersing such NQ particles within the propellant containing an NC/NG/DEGDN matrix.^29,30

The observed debonding patterns suggest that the composite’s concave sizes correspond to the NQ particle profile. Due to the incomplete bonded interface, the interfacial adhesion between NQ particles and the matrix is weak, creating a plane of vulnerability. As the bond strength is lower than the tensile strength of NQ crystals, debonding and fractures initiate at the interface between NQ crystals and the matrix when local stress at the interface surpasses interfacial strength. Ductile crack propagation leads to the growth and eventual coalescence of the debonding gap, culminating in composite fracture and cavity formation. The presence of numerous craters (highlighted with circles) following the removal of NQ from the fracture surface indicates a significant amount of poorly bonded interfacial areas between the NQ and matrix, consequently reducing the composite’s structural strength.^31–33 

The primary factor contributing to debonding formation is the volatilization of substantial quantities of Volatile Organic Compounds (VOCs), which induces shrinkage and fluffing in the NC/NG/DEGDN matrix, thereby degrading the interface bonding performance between NQ crystals and the matrix in STP samples. Consequently, the formation of a stable NQ-matrix bonding interface becomes challenging. The stress transfer at the NQ-matrix interface grows less efficient with an increase in bond defects. Additionally, the NC/NG/DEGDN binder on the exposed NQ crystal surface is insufficient to enhance the bonding property at the NQ-matrix interface. Another contributing factor is the excellent thermal conductivity of the NC/NG/DEGDN matrix, which causes the matrix to re-soften at elevated temperatures and pressures, thereby improving adhesion between the NQ-matrix. This behavior results in fewer or smaller voids by effectively filling the gaps in the NQ-matrix, ultimately strengthening the microstructure and enhancing the material’s overall strength. In summary, high temperature and pressure processes can effectively inhibit the debonding and stripping of NQ-matrix interfaces, particularly when NQ particles are well bonded to the binder, bearing the primary load.

More specifically, the NQ-matrix in SLTP samples forms a more efficient cross-linking system. When the NQ-matrix interfacial adhesion is robust, stress transfer at the NQ-matrix interface is more efficient. It is important to note that the cross-linking system not only functions to connect the NQ-matrix bonding interface and maintain efficient stress transfer but also plays a crucial role in constructing the NQ-matrix 3D chain network itself, thereby preserving the structural integrity of the SLTP samples.

The stress–strain curves of STP samples and SLTP samples at different temperatures (−40, 25 and 50 °C) are shown in Fig. 6.

Based on the stress–strain curves shown in Fig. 6, the mechanical properties (tensile strength, modulus of elasticity, and failure strength) of the STP samples and the SLTP samples are given in Table II.

Table II reveals that the tensile strength, fracture strength, and modulus of elasticity for all samples peak at −40 °C. As temperatures increase, these properties deteriorate. At 25 °C, the heightened kinetic energy of the molecules accelerates their motion and expands the loosely bonded structure, thereby decreasing the aforementioned properties. With further temperature increases, molecular motion intensifies, reaching levels (50 °C) close to the glass transition point; this causes the samples to become supple and rubbery. Consequently, mechanical properties, such as tensile strength, elastic modulus, and fracture strength, are significantly reduced near the glass transition temperature, yielding extremely low values. The material’s softening, however, is responsible for changes in these mechanical properties as a function of temperature.^34,35

As expected, the SLTP samples exhibit greater tensile strength than the STP samples at ambient temperature. The highly aligned orientation in STP energetic composite materials cannot counteract the effects of weakened voids and concentrated stress points. These data suggest that at a specific loading of needle-shaped NQ, SLTP energetic composite materials could be denser and more resistant to tension than conventional STP composites.

In accordance with Fig. 6, the tensile strength of SLTP samples surpasses that of STP samples across all tensile directions. There is virtually no significant variation in the tensile strength results of SLTP samples across the three mutually perpendicular collection planes, with a standard deviation of 1.09. In contrast, the tensile strength results of STP samples in three mutually perpendicular collection planes exhibit considerable dissimilarity, with a standard deviation of 3.93. The tensile strengths in the X and Y collection planes are roughly equal and both lower than the tensile strength in the Z plane.

Figure 6 displays the quasi-static tensile strength in the Z acquisition plane for STP and SLTP samples at various temperatures. In comparison to STP samples, the tensile strength of SLTP samples is stronger at −40, 25, and 50 °C, with respective increases of 9.53, 8.33, and 5.89 MPa. It is noteworthy that this improvement is quite substantial. A significant enhancement in the quasi-static tensile strength of SLTP samples is observed. This can be attributed to the uniform orientation distribution of NQ particles and the effective stress transfer between the NQ and the matrix, which reduces the higher stress concentration within the SLTP samples. Consequently, a relatively large tensile strength is observed in the quasi-static tests. In summary, as the test acquisition temperature increases, the tensile strength values of all tested samples decrease, while the opposite trend is observed for fracture elongation.

As seen in Table II, the elastic modulus of SLTP samples is greater than that of STP samples, whereas the fracture elongation is lower. In addition to providing the force for matrix composites to elongate while the SLTP samples are under tension, the tensile load also induces a shear stress at the NQ and matrix contact. When the interface bonding is excellent, the strain in the SLTP samples will be equivalent to that of the matrix composites. Consequently, the stress transfer efficiency from the matrix to the NQ would decline as the NQ begins to debond, resulting in nonlinear behavior. The matrix structure fractures when the NQ is debonded as the stress cannot be transferred from the matrix to the NQ. Following is an explanation for this: On the one hand, the NQ in SLTP samples has a better interface connection with the matrix, which improves the NQ and matrix’s interfacial interaction and reduces the stress concentration close to the NQ. On the other hand, random distribution of NQ prevents the formation of bundles or “ ropes” that would have effectively transferred and dispersed stresses. Therefore, a randomly oriented NQ and an efficient bonding interface can enhance the tensile strength of the composite more effectively.

This suggests that the NQ in the SLTP sample can accommodate more direction under external force, allowing for better stress transfer and dispersion within the matrix. The morphology of the STP samples all indicate the presence of potential voids near the NQ surface. These voids could negatively affect the mechanical strength of the energetic composite material as the stress transfer from the matrix to the NQ would be suboptimal. Removing these near-surface voids on the NQ would likely further improve the mechanical properties of the SLTP samples.

When a highly rigid material such as needle-shaped NQ is incorporated into a relatively softer energetic composite materials matrix, it can lead to an increase in the elastic modulus of the energetic composite materials. In other words, it can enhance the material’s deformation resistance. However, it reduces its elongation at break.

The effects of ambient temperature and needle-shaped NQ orientation on the mechanical properties of energetic composite materials were investigated through quasi-static compression tests. Figure 7 shows the stress–strain curves of STP and SLTP samples with different compression test planes as well as the stress–strain curves of Z-compression test planes of STP samples and SLTP samples under different ambient temperatures.

The SLTP samples exhibited consistent mechanical behavior for compression resistance across different compression test planes. In contrast, the STP samples demonstrated varying nonlinear mechanical behaviors depending on the compression test plane. The compressive stress–strain curves of STP samples and SLTP samples under compressive loading can be described through three stages: the linear elastic stage, the high elasticity stage, and the strain hardening stage. Initially, the energetic composite materials enter the linear elastic stage, during which the stress–strain relationship of the cubic block sample approximately satisfies Hooke’s law, and the compressive stress does not cause damage to the matrix. In the high elasticity region, the cubic block samples exhibit nonlinear behavior, with the stress increasing at an accelerated rate as strain increases. Following this region, strain hardening occurs. In this phase, the filler NQ particles play a significant supportive role.

Table III, derived from the stress–strain curves in Fig. 7, illustrates the compressive strengths of the STP samples and SLTP samples at various temperatures.

Figure 7 displays the stress–strain curves of the STP and SLTP samples up to failure. The stress–strain curves of STP samples demonstrate the compressive properties along the x- and y-axes. The curves of the STP samples, tested with compression resistance along the X, Y, and Z axes, reveal that samples with compression resistance along the Z-axis exhibit higher stress intensity than those with compression resistance along the X and Y axes. This is primarily because the NQ particles in the STP samples are highly aligned along the Z-axis. The X and Y axis directions are the two perpendicular directions along the orientation arrangement of the NQ particles. Since the distribution of NQ alignment along the X and Y axes is similar, the compressive stress–strain curves present similar maximum compressive strength.

Figure 7 shows that under low-temperature quasi-static compression conditions (−40 °C), the STP sample displayed the highest value of compressive stress resistance at 65.13 MPa when compressed along the Z-axis. However, the compressive stress values along the X and Y axes are the lowest of them all, at 51.67 and 48.17 MPa, respectively. In comparison, the compressive stress values along the X, Y, and Z axes of the SLTP samples under low-temperature quasi-static compression (−40 °C) were quite uniform and higher than the minimum compressive strength of the STP samples: 58.40, 58.19, and 60.18 MPa. Notably, the standard deviation for the compressive strength of the STP samples in the three mutually perpendicular test planes was 7.28, while the standard deviation for the compressive strength of the SLTP samples in the three mutually perpendicular test planes was 1.10. This suggests that the overall compressive strength of the SLTP samples was more stable and excellent.

The compressive strengths along the Z-axis of STP samples were higher than those along the Z-axis of SLTP samples across the entire temperature range. Specifically, the compressive strength along the Z-axis of the STP sample was 4.49, 21.28, and 26.26 MPa higher than that of the SLTP sample at low temperature (−40 °C), room temperature (25 °C), and high temperature (50 °C), respectively. It is also worth mentioning that the STP and SLTP samples exhibited the same trend of failure strain along the Z-axis as the compressive strength along the Z-axis, that is, the failure strains along the Z-axis of SLTP samples at low temperature (−40 °C), room temperature (25 °C), and high temperature (50 °C) are smaller than the failure strains along the Z-axis of STP samples. This can be attributed to the higher density of the SLTP samples and the smaller interfacial bonding gap of the components, leading to a smaller volume limitation that can be compressed. In other words, the high densification of SLTP samples suppresses the degree of compression softening of the material. This is due to the fact that the experimental test temperature did not bring the energetic composite materials’ molecular chains close to their glass transition temperature, which is between 48 and 57 °C. This is because the motion of the chain segments is limited by van der Waals forces and hydrogen bonds. Consequently, the lower the ambient test temperature in the energetic composite materials system, the higher the compressive strength of the sample to be tested.

As depicted in Fig. 8, the impact strengths of both STP and SLTP samples exhibited a similar trend as their tensile strengths, demonstrating a notable temperature dependence. With increasing test temperature, the effect of temperature on impact strength becomes more pronounced. It was observed that when impact resistance tests were conducted perpendicular to the direction of the long rod-shaped propellant grains, the impact fracture strength of SLTP samples was higher at low (−40 °C) and room (25 °C) temperatures compared to STP samples. However, at high temperature (50 °C), the impact fracture strength of SLTP samples was significantly lower than that of STP samples. This phenomenon can be closely correlated with the glass transition temperature examined in prior tests.

Additionally, matrix shrinkage at low temperatures increases the interfacial space between the binder and filler phases, leading to a weakened intermolecular force between the binder system and NQ. This results in reduced solid surface bonding in the propellant sample, thus increasing stress point formation and decreasing impact strength. As the ambient temperature decreases, the impact strength of energetic composite materials significantly diminishes.

The mechanical response of energetic composite materials is contingent upon the type of loading force, such as tensile and compression resistance. Based on prior multiple impact tests, Fig. 9 presents typical dynamic impact strength data for STP and SLTP samples, which exhibit varying mechanical strengths at different loading temperatures. The impact fracture surface results indicate that the mechanical failure of the stick material predominantly manifests as tensile fracture failure on the reverse side of the impingement under low and ambient temperature impact resistance, suggesting that tensile stress is greater than compressive stress. Conversely, stick material’s mechanical failure primarily exhibits compression collapse failure on the forward side of impingement under high temperature impact resistance, indicating that tensile stress is less than compressive stress. In summary, STP and SLTP samples display distinctly different dynamic mechanical failure mechanisms.³⁶ 

The dynamic impact resistance behavior of energetic composite materials is significantly influenced by the properties of viscoelastic materials, such as tensile and compression resistance. To explore the contributions of rigid NQ particles and the energetic composite materials matrix substructure network to macroscopic dynamic mechanical properties, the impact resistance failure behavior was deconstructed.

The equivalent schematic diagrams of impact stress resistance for STP samples and SLTP samples are presented in Fig. 9. Impact resistance failure is a combination of tensile and compressive failure. According to Sec. III C, the tensile strength of SLTP samples along the Z-axis surpasses that of STP samples throughout the temperature domain. Concurrently, the compressive strength of SLTP samples along the Z-axis is lower than that of STP samples, exhibiting a smaller failure strain along the Z-axis. Using the glass transition temperature as a dividing boundary, the primary factor limiting mechanical structural failure below the glass transition temperature may be attributed to tensile resistance, while above the glass transition temperature, it may be attributed to compression resistance.

Two destructive failure mechanisms can be summarized from the impact strength data and impact failure surface analysis: tensile fracture failure on the reverse side of the impact and compression collapse failure on the forward side of the impact. It can be hypothesized that structural failure of energetic composite materials at low (−40 °C) and room (25 °C) temperatures occurs when internal tensile stress reaches the material’s tensile fracture limit, leading to cracking on the tensile surface and subsequent separation fracture failure. Structural failure at high temperature (50 °C) arises when internal compressive stress reaches the material’s compressive collapse limit, causing the compression surface to bend inward and fail through significant collapse failure. Both failures dissipate mechanical impact kinetic energy by utilizing the fracture and debonding of the three-dimensional network chain structure. The pronounced difference in dynamic impact failure resistance between STP and SLTP samples is determined by the distribution of tensile and compressive stresses within the material and the material’s ability to resist internal stresses. The macroscopic dynamic impact mechanical properties further verify stress transfer and distribution between energetic composite materials’ three-dimensional network chain structures.

By employing an efficient and environmentally friendly propellant manufacturing process, solventless type triple base propellant (SLTP) has been successfully processed and manufactured and compared with conventional solvent type triple base propellant (STP). The high-temperature and high-pressure plasticized manufacturing integration method ensures uniform distribution of needle-shaped nitroguanidine crystal organization within the propellant and tight bonding of each tissue interface. SLTP products demonstrate higher relative density, more accurate molding dimensions, enhanced stability and reliability of overall mechanical properties, and broader environmental adaptability.

Microstructure morphology observation revealed that SLTP samples possess a mutually supportive three-dimensional spatial structure, and their NQ crystals exhibit tighter and more efficient bonding at the interface with the matrix. Quasi-static structural failure tests, comprising quasi-static tensile resistance and quasi-static compression resistance, determined that SLTP samples have superior tensile strength and stable compressive strength. The quasi-static tensile properties indicated that the tensile strength of SLTP samples was higher than that of STP samples in all three axes, exhibiting even and stable strengths with a standard deviation of 1.09. In contrast, STP samples’ tensile strength in the triaxial axis displayed significant differences and unstable strength with a standard deviation of 3.93. Furthermore, the tensile strengths of SLTP samples exceeded those of STP samples at −40, 25, and 50 °C, with increases of 9.53, 8.33, and 5.89 MPa, respectively.

Quasi-static compression resistance tests demonstrated that SLTP samples exhibit consistent mechanical behavior in different compression test planes. The standard deviation of compressive strength for SLTP samples in the three axes was 1.10, while for STP samples, the value was 7.28, indicating that the combined compressive strength of SLTP samples is more stable and superior. On the other hand, STP samples’ three-axis compressive strength exhibited significant fluctuations.

Dynamic structural damage testing is represented by rapid impact resistance tests. Impact resistance test results revealed that the structural failure of energetic composite materials at low temperature (−40 °C) and room temperature (25 °C) is caused by internal tensile stress reaching the material’s tensile fracture limitation, followed by separation fracture failure after the appearance of cracks on the tensile surface. The structural failure of energetic composite materials at high temperature (50 °C) occurs when internal compressive stress reaches the material’s compression collapse limitation, resulting in the compression surface bending inward after a significant collapse failure.

The supplementary material contains a comprehensive description of the test methods for scanning electron microscopy (SEM), micro-x-ray computed tomography (micro-CT), quasi-static tensile resistance testing, quasi-static compression resistance testing, and rapid impact resistance testing.

This work was supported by the instrument and equipment fund of the Key Laboratory of Special Energy Materials, Ministry of Education, School of Chemistry and Chemical Engineering, Nanjing University of Science and Technology, Nanjing, China. We also gratefully acknowledge Carl Zeiss (Shanghai) Management Co., Ltd. for providing electron microscopic image analysis, micro-CT 3D structure reconstruction, and visualization. We are thankful to Luzhou North Chemical Industry Co., Ltd. and Liaoning Qingyang Special Chemical Co., Ltd. for providing the physicochemical parameters analysis. This work made use of facilities supported, in part, by the NETZSCH Scientific Instruments (Shanghai) Co., Ltd. for providing the testing of thermophysical properties of materials and the School of Mechanical Engineering of Nanjing University of Technology for providing the mechanical strength test of materials. The authors are grateful to Professor Ping Du, Professor Fengqiang Nan, and Professor Xiaoan Wei for their experimental coordination, valuable discussions, and input. We also appreciate the experimental contributions and technical support from engineer Yannian Yu and test assistant Shenglai Yin.

The authors declare that they have no known competing financial interest or personal relationships that could have appeared to influence the work reported in this paper. There is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of the manuscript entitled. All the authors listed have approved the enclosed manuscript for publication.

Yao Zhu: Conceptualization (equal); Data curation (equal); Formal analysis (equal); Methodology (equal); Validation (equal); Writing – original draft (equal). You Fu: Data curation (equal); Investigation (equal); Methodology (equal). Xijin Wang: Methodology (equal); Project administration (equal). Qian Chen: Methodology (equal); Resources (equal); Software (equal). Jing Yang: Validation (equal); Visualization (equal). Bin Xu: Project administration (equal); Writing – review & editing (equal). Zhitao Liu: Formal analysis (equal); Funding acquisition (equal); Methodology (equal). Feiyun Chen: Conceptualization (equal); Methodology (equal). Xiaoan Wei: Conceptualization (equal); Supervision (equal). Xin Liao: Conceptualization (equal); Funding acquisition (equal); Writing – review & editing (equal).

The data that support the findings of this study are available from the corresponding author upon reasonable request.