The rheometric study was performed on the Ektrontek EKT2005 Rheometer at 180°C for 6 min, and Mooney viscosity was tested at 125°C for 5 min in the MV2000 Mooney Viscometer by Alpha Technologies. The compounded rubber sample (weight of 25–30 g) was kept in the cavity of the Mooney viscometer. It was allowed to preheat for 1 min, and for the remaining 4 min, it was allowed to shear. The Mooney value is denoted as ML (1+4) @ 125°C (López Manchado and Arroyo, 2002).

The mixed sheets were molded at their respective conditions in a compression molding machine (Santosh Industries, India) at 165°C for 30 min at 1,000 psi pressure (For all samples, the same conditions were followed) and tested as per the respective ASTM standards.

To characterize the mechanical behavior of the sample materials, including tensile strength, elongation at break, and modulus of elasticity, standardized testing methodologies were employed. Molded sheets with a controlled thickness of 2.5 ± 0.5 mm were fabricated. From these sheets, dumbbell-shaped specimens were prepared in accordance with the ASTM D-412 standard for tensile property assessment. Similarly, samples for tear strength evaluation were cut following the guidelines outlined in ASTM D-624. All mechanical tests were conducted using an Instron 3366 Universal Testing Machine. Physical properties testing was done before and after aging, and the aging condition was 125°C for 70 h in a hot air oven.

Sample hardness was quantified using a Shore-A durometer (Wallace Cogenix Hardness Tester), adhering to the testing protocol specified in ASTM D 2240-15. Reported hardness values represent the average of three independent measurements.

The compression set of the EPDM rubber samples was determined according to the ISO 815-1972 standard. Testing was conducted for a period of 72 h at a temperature of 125°C, with a constant compression of 25%.

Abrasion resistance index (ARI) was measured in the rotary drum abrader (ASTM D5963), and ASTM Compound B was taken as a reference sample to calculate ARI and volume loss of samples.

Heat Build-up of the samples is studied as per ASTM D-623. This is done by using the Goodrich Flexometer, clamping the sample inside the chamber, and monitoring the increase in the temperature of the sample when it is operated under a constant 30 Hz frequency, 2 KPa stress, and stroke of 1,800 cycles/min conditions.

Resistance to flex cracking was evaluated using the DeMattia test, following the procedures outlined in ASTM D813. Specifically, Molded DeMattia specimens were subjected to repeated flexing and deflexing cycles using a DeMattia Flexometer. The length of crack propagation with the number of cycles completed was monitored and recorded at 30-min intervals. Reported values represent the average of measurements obtained from three independent samples (Muthukannan et al., 2024).

The described methodology aims to assess the ability of a polymer to withstand the development and propagation of cuts or cracks under conditions of cyclic flexing. In this study, the chamber temperature of 70°C was maintained for the testing, and then the crack length was recorded at predetermined cycles of 10 k and 25 k. The standard specimen dimensions are 152 mm in length, 25 mm in width, and 6.4 mm in thickness, as specified by ASTM D1052. A notch approximately 2.5 mm deep is typically introduced on one side of the sample to initiate cracking and evaluate its resistance to crack growth under repeated flexing.

Aramid short fibers were incorporated into EPDM rubber formulations at varying concentrations up to 20 phr, as indicated in Table 1. The tensile strength was assessed and illustrated in Figure 4. A significant reduction in tensile strength for the gum compound occurs up to 12 phr of fiber loading. The tensile strength in both longitudinal and transverse directions drops up to 12 phr fiber loading and subsequently increases gradually in the case of T7 and T8 composites. At 12 phr, the reinforcing effect is inadequate to offset the dilution effect caused by the fibers, resulting in a loss in tensile strength. The limited quantity of fibers prevents uniform stress transfer throughout the matrix, causing localized stress at the fiber-matrix interface and resulting in inefficient reinforcement. It is essential to understand the intrinsic restrictions related to the short fiber length. Short fibers possess a limited aspect ratio (length-to-diameter ratio), which limits their capacity to effectively withstand load and transfer stress across extended distances within the matrix. The reinforcing efficiency at lower fiber loadings is reduced. Increased fiber concentrations result in a greater number of fibers obstructing the fracture front, leading to a more uniform distribution of stress; hence, tensile strength rises with fiber concentration beyond 12 phr. The tensile strength in the transverse direction decreases with fiber content up to 12 phr, the same as in the longitudinal direction of the rubber matrices, after which it remains relatively stable. This results from the weakening of the rubber matrix, which results from transversely oriented fibers. The comparatively stable tensile strength in the transverse direction beyond 12 phr indicates the limited ability of the short fibers to successfully bridge cracks, as we anticipated that over 60% of the short fibers would be oriented in the longitudinal direction following the milling and calendering processes (Pittayavinai et al., 2017; Correia et al., 2017).

Figure 6 depicts the correlation between tear strength and the fiber content within the composite materials. Generally, an increase in fiber content corresponded to an increase in tear strength, observed in both the longitudinal and transverse directions (Kumar and Thomas, 1995). This enhancement is likely due to the increased physical impedance to crack propagation provided by the higher fiber concentration. Furthermore, at a given fiber loading, tear strength was consistently higher in the longitudinal direction compared to the transverse direction. This anisotropic behavior in tear resistance is attributable to the alignment and orientation of the reinforcing fibers. When tested in the transverse direction, the fibers tend to align parallel to the advancing crack front, offering reduced resistance to tear propagation. The crack can more easily propagate along the fiber-matrix interface when the fibers are parallel to it. This leads to reduced tear values compared to the longitudinal direction, where the fibers are oriented perpendicularly to the crack path, requiring more energy to break or pull them out of the matrix. The variation in tear strength highlights the necessity of considering fiber orientation in the development of composite materials for applications where tear resistance is a crucial performance parameter.

Table 4 presents the retention of tensile properties of the composites before and after aging. The tensile strength retention percentages of all compounds are greater in the transverse direction than in the longitudinal direction. Retention is notably inadequate for T7 and T8, particularly in the longitudinal direction. This indicates that the aging process adversely affects longitudinal tensile strength more than transverse tensile strength. A plausible reason for the variation could originate in the orientation of the aramid fibers and their interaction with the EPDM matrix under aging circumstances. Prolonged exposure to 125°C can accelerate differential thermal expansion between the aramid fibers and the EPDM matrix, resulting in increased stress concentrations at the fiber-matrix interface in the longitudinal direction. This may lead to micro-cracking or debonding at the interface, weakening the material’s tensile strength along the fiber direction. Moreover, the elevated temperature may promote oxidative degradation of the EPDM matrix preferentially in the longitudinal direction due to greater exposure along the fiber length, thereby reducing the load transfer efficiency between the fibers and the matrix. Meanwhile, elongation at break also exhibits poor retention in the longitudinal direction, as illustrated in Table 5. The decrease in elongation at break values upon aging indicates that the material becomes increasingly brittle, especially along the longitudinal axis, likely due to the aforementioned reasons associated with fiber-matrix interface degradation. This behavior is essential to evaluate for applications where the composite has to withstand prolonged exposure to high temperatures, as in power transmission belts, which highlights the necessity for deeper investigation into aging mechanisms and potential strategies to enhance the long-term durability of these composites under such thermal aging conditions. Additionally, investigating the efficiency of specific stabilizing techniques, such as adding antioxidants or changing the fiber surface treatment, could be quite crucial in raising the aging resistance of these composites.

As the concentration of short aramid fibers in EPDM rubber composites increases, the Shore A hardness correspondingly increases. The rigid, high-modulus fibers reduce the flexibility of the rubber matrix, resulting in a material that is less responsive to surface indentation. The higher fiber loading improves the composite’s resistance to deformation under localized pressure, yielding a stiffer surface and enhanced hardness measurements. Moreover, better stress transfer between the rubber and the fibers leads to the noted increase in surface hardness. Based on the data shown in Figure 7, it can be concluded that T1 has the lowest shore A hardness value of 75.8, whereas T8 has the maximum hardness value of 89.8.

The compression set of EPDM rubber formulations was evaluated according to ISO 815-1972, employing a 72-h test duration at a temperature of 125°C and 25% compression. Compression set is a critical performance parameter, particularly for V-belt applications where resistance to permanent deformation under sustained load is essential. The results, presented in Figure 8, demonstrate a clear correlation between fiber content within the rubber matrix and compression set. Specifically, an increase in fiber loading leads to a corresponding increase in the observed compression set percentage, indicating a reduced ability of the material to recover its original dimensions after sustained compressive stress over time. The increase in compression set with increased fiber loading results from the stiff nature of the fibers, which restricts the elastic deformation and recovery of the EPDM matrix. As fiber concentration increases, the mobility of polymer chains diminishes, resulting in greater permanent deformation during prolonged compression. This leads to a high compression set value, indicating reduced resilience, which may affect the long-term ability to maintain its shape and contact pressure in V-belt applications. T8 is showing the highest compression set % value of 56.3 among all rubber composites due to having the highest short aramid fibers.

Abrasion resistance refers to the ability of polymeric materials to withstand surface wear caused by mechanical action. This test becomes very important for products that are in continuous friction with another surface, like the base of the power transmission belt, which is in continuous contact with the pulley, and the tire tread, which is in continuous contact with the road. The higher value of ARI indicates better abrasion resistance, while the lower value of ARI indicates poorer abrasion resistance. Abrasion resistance registers an improvement with an increase in fiber content (Corrêa et al., 1998). These are calculated with standard given formulas. Abrasion  Loss  in  mm 3 = Test  rubber  mass  loss  in  mg × Normal  abrasiveness / Density  of  test  rubber × Standard  Rubber  Mass  Loss

Where constant values according to ASTM D5963 are: Normal abrasiveness-200.0, Standard Rubber Mass Loss-215.0.

To calculate the abrasion resistance index (ARI) ARI = Standard  Rubber  Mass  Loss × Density  of  Test  Rubber / Test  rubber  mass  loss × Density  of  standard  rubber × 100

Density of standard Rubber-1.353.

The abrasion resistance was performed using a DIN abrader. The abrasion sheet used was corundum (aluminum oxide). The diameter of the cylinder is 150 mm, and the length is 500 mm, rotating at a speed of 40 rpm. The diameter and thickness of the test specimen were 16 mm and 6 mm, respectively. The density of the samples was measured. The test run was performed, and the results were taken as the average of three test runs. The weight of the specimen before and after the test was measured.

From Figure 9, it is clear that by adding a high concentration of Fibers into the elastomer, the abrasion loss (mm³) gets reduced to a certain level, which means higher fiber concentration leads to better abrasion resistance properties.

Similarly, the Abrasion Resistance index value is High for T8 and low for T1 as shown in Figure 10, which also determines the same. It is necessary for the base rubber of a power transmission belt to have a higher abrasion resistance index. So, depending upon the application, we can select the particular fiber concentration compound. But in case of higher fiber concentration, many factors play a vital role, like High Mooney viscosity, Hardness, and brittle nature, leading to fast rupture of base rubber on flexing, which is not suitable for power transmission belts.

From the DeMattia Fatigue Study, it can be clearly seen from Figure 11 that T7 and T8 with 15 phr and 20 phr fiber concentration, respectively, exhibit significantly reduced fatigue life as recorded at approximately 109 kilocycles and 105 kilocycles, the main cause of this low fatigue life is the composites increasing brittleness at higher fiber loadings, which reduces elastic deformation and promotes early fracture initiation. Particularly, the T8 specimen shown in Figure 12 shows behavior typical of brittle failure, in which stress concentration becomes more dominant than energy dissipation, by showing crack initiation from the center of the test strip. In contrast, T5 and T6 composites with 10 phr and 12 phr aramid fiber exhibit a significant increase in fatigue life; it reaches 377 kilocycles and 420 kilocycles, respectively. This implies that a modest fiber loading level offers enough reinforcement to inhibit the spread of cracks and preserves matrix flexibility to absorb cyclic stress, hence improving the fatigue performance. Furthermore, we can see that fatigue life is improved in the case of T4, and it has been achieved up to 552 Kilocycles, which is the highest among all. Notably, in the case of T3 Compound with 6 phr of fiber concentration, showing the best Fatigue Life, in which the crack initiated after 367 Kilocycles, which is the highest among all, and full rupture of the specimen occurred at 545 kilocycles (Li and Xin, 2018).

Overall, the data clearly show that aramid fibers facilitate mechanical reinforcement, but excessive loading causes brittleness, which could cause quick fracture failure. Maximizing fatigue life in fiber-reinforced elastomer composites depends thus upon optimizing fiber content. The numerical values in Figure 12 are in Kilo Cycles units.

During dynamic applications, excessive heat generation in rubber products promotes further curing, hence increasing hardness and brittleness. Dynamic stress-strain cycles, therefore, easily cause premature fractures in engineered or consumer products. It directly impacts the service life, durability, and performance under dynamic conditions, especially in applications like tires, V-belts, and seals, where continuous flexing occurs. The use of short aramid fibers in rubber composites usually causes an increase in heat buildup, resulting from increased internal friction and limited mobility of polymer chains. With an increase in fiber content, the interactions at the fiber–matrix interface become more prominent, resulting in greater energy dissipation under cyclic loading. Although aramid fibers provide mechanical reinforcement, their low thermal conductivity and rigidity could block heat dissipation, leading to localized temperature increases. The cumulative effect results in increased heat buildup in the composite as the percentage of short aramid fibers rises.

From Table 6, it is clear from T1 to T8 that the temperature of rubber composites increases as the fiber content increases after predetermined cycles, which clearly indicates that the high fiber content in the elastomeric composite leads to the generation of higher heat.

Traditional V-belts, typically fabricated from natural or common synthetic rubber formulations, are generally limited to operating temperatures below 90°C. Furthermore, at even lower temperatures around 60°C, accelerated degradation of the rubber material becomes a concern. Maintaining equivalent drive durability and performance under elevated temperature conditions necessitates making it larger and stronger than usual to withstand the increased thermal stress and prevent premature failure. It is estimated that a 10°C increase in operating temperature above 60°C reduces V-belt lifespan by approximately 50%. Consequently, achieving equivalent service life at a working temperature of 90°C would necessitate utilizing 8 times more belts, a solution that is both economically impractical and mechanically inefficient. In contrast, EPDM belts offer superior thermal stability, maintaining consistent performance across a temperature range of −40°C–110°C and are capable of operating at temperatures exceeding 130°C. This makes EPDM an ideal candidate for high-temperature belt applications where long-term durability is essential. So, in this study, to simulate realistic service conditions and assess the crack growth behavior under dynamic flexing, the Ross Flex testing was conducted at a chamber temperature of 70°C. As shown in Figure 13. Then, the Crack length was recorded at predetermined cycles of 10 k and 25 k, which are given in Table 7, given below. It is very clear that T3 shows the best results among all tested specimens. After 10 kilocycles, the crack length was around 4 mm, and after 25 kilocycles, it was just 13 mm, while in the case of T6, T7, and T8, all specimens were fully ruptured after 25 Kilocycles. But the initial crack length after 10 kilocycles in T7 and T8 was very low, 3 mm and 5 mm, respectively. This clearly shows that a higher content of aramid fibers in the compound resists the crack growth, but up to a certain limit. If the content is very high, like 15 and 20 phr, then other factors like high viscosity, high hardness, and brittle nature weaken the compound and cause rapid rupture of the compound on flexing and deflexing.

In the case of T2 and T1, crack growth values were slightly increased in comparison to T3 due to the lower content of aramid fibers. This clearly shows that aramid fibers play a vital role in crack growth resistance. So, from the data analysis, it is clear that the aramid fiber content of 6 phr gave better results in crack growth resistance.