3.1 Metallurgical Evaluation
According to the optical image of FeCrNiCuMoSiCx (x = 10 & 15wt%) shown in Fig. 1 (a & b), the system is characterized by three distinct layers namely the friction material, the intermediate transition layer, and the steel backing plate. From Fig. 1 (a & c), the SiC is stable in the microstructure for Mix A. The SiC distribution in the matrix is uniform. The bonding appears to be strong and free from defects such as delamination, cracks, or interfacial products. In addition, it also enhanced the dimensional stability of the brake pad. There are no traces of defects, such as cracks, pores, or undesirable interfacial products, in the composite's bond with Ni-plated steel, as shown in Fig. 1(a). The microstructure of the steel sheet is dominated by the ferrite phase, with some pearlite. The brake pads developed are not a homogenous material in Fig. 1 (b & d). The particle size of each element is not uniform in size and the distribution of the element is not well dispersed in the matrix. The Mix B materials towards either side of the interface were entirely different, crack, and de-bonding occurred frequently during actual applications. In order to overcome the poor joint, future optimization should be carried out in powder processing/ sintering for better usage of Mix-B. As for this paper Mix-A alone is carried out for mechanical and dynamometer tests. The mechanical and physical properties of Mix-A alloy were also determined. According to Ram et al.[18] the appearance of FCC peaks confirms the existence of a solid solution matrix. Intermetallics (FeNi3 and CrNi3) were discovered in FeCrNiCuMoSiCx=10 composite.
The average hardness, porosity, and density of Mix-A after sintering reached 178 ± 4 HV, 5.84% and 6.34 g/cm3 respectively. A decrease in wear loss is associated with an increase in hardness, according to the Archard relation. In addition, the composites have a higher hardness than the individual metallic elements. There is a possibility of strengthening the solidsolution [1]. The effect of alloy structure on the ability to resist wear at different brake energy and sliding speeds is apparent in this study. The HEAs show high wear that may be due to the relatively low hardness of these alloys [18, 24].
3.2 Wear behaviour
The wear response of FeCrNiCuMo (SiCx=10) was tested for three different sliding speeds. Figure 2 presents the coefficient of friction (CoF) evolution for the 100 braking stops with varying sliding speeds. Wear mechanisms are affected by sliding speed, causing material deformation, changing friction interface temperatures, and affecting wear mechanism [25, 26]. An increase in COF is observed with the increasing sliding speed of FeCrNiCuMo SiCX=10. The coefficient of friction curves is typically characterized by two friction regimes as reported in literature [27].
The A1 and A2 alloys exhibit relatively large fluctuation whereas A3 alloy does a small fluctuation. Initially, the A1 shows a rapid and steep running in a stage where the friction coefficients rapidly increase up to a certain value. The second stage, where friction coefficients are observed to fluctuate significantly with increasing stops. There is a fluctuation in CoF, which can be attributed to lose wear debris. Sliding speed is affected by the stuck debris, which causes abrupt fluctuations in CoF when the debris interferes with the sliding. The smallest attainable CoF is due to the wear mechanism of oxidation wear in contrast to the delamination wear. The large fluctuation of the friction coefficient is caused by two reasons. The first is the periodic localized fracture of the surface layer [28, 29]. The second is the periodic accumulation and elimination of debris on the worn surface. The coefficient of friction increases as the large debris accumulates on the worn surface while it decreases as the debris departs from the worn surface [30]. As for the small fluctuation of A3, it is due to its wear mechanism of oxidative wear, which results in small debris.
Figure 3 shows the average coefficient of friction with varying sliding speed. In comparison, A3 and A2 show a higher coefficient of friction than A1. There is a high difference in the wear mass loss between A1 and A3 after friction testing. The results suggest that the A1 has better wear performance than the A3 does. A1 also shows the smallest mass loss, indicating a greater wear resistance. In accordance with the prevailing wear mechanisms at that sliding speed, as the sliding speed increases, the CoF increases. At the higher speed (40ms− 1), the friction between the two surfaces generates more energy, which causes the temperature to rise. In the Fig. 4, the stopping distance and mean torque are shown as a function of sliding speeds. The mean torque increases with increase in the sliding speeds. As the sliding speed increases, more heat is dissipated in the form of frictional energy between the brake pad and rotor, which increases energy dissipation. In addition to dissipating energy, the braking system also experiences an increase in temperature, which may affect wear behavior. Additionally, high sliding speeds (40ms− 1) can create increased abrasive wear as contaminants and debris are present on the surface of the rotor. In consequence, particles are more likely to interact and enter the brake pad material at a faster rate, resulting in a greater loss of material and more rapid wear. Brake mean torque and sliding speed may also be affected by other factors like the brake pad and rotor material composition and properties as well as system design and operating conditions.
During braking, the friction coefficient of the body determines the stopping distance of the body. A higher sliding speed leads to a larger mean break torque and a shorter stopping distance for FeCrNiCuMoSiCx=10. Shorter stopping distances are seen at A1 as compared to A2 and A3. As a result of the higher oxygen content at A1 (20ms− 1) and the lower wear rate, the third body friction is enhanced and the stopping distance decreases. Considering these factors, it can be deduced that A1 sliding speed has a higher wear resistance and mean torque as well as a better braking performance than the A2 and A3.
The higher temperatures cause the constituents to oxidize, which in turn causes mass loss. As a result, contact surfaces exhibit enhanced diffusion activity, which leads to deterioration in wear properties (Fig. 5). High sliding speeds cause delamination of particle/matrix if braking force is applied, resulting in significant mass loss. Beyond this sliding speed, the oxide scale thickness increases, cracking of the scales will accelerate faster. When the friction temperature rises, the material softens, weakening the particle/matrix interface and resulting in mass loss. Up to the critical speed, increasing the sliding speed facilitates the formation of oxides and a tribolayer at the wearing surface, thereby preventing the accumulation of wear per unit sliding distance until the critical speed is reached [31]. This suggests that the A1 is suitable for use as a brake pad wear protector, which is a practical application.
The SEM technique to analyse the morphologies and compositions of the friction surface after braking tests and the corresponding results are shown in Fig. (6–8). Based on the wear pattern shown in Fig. 5, the A1 (20ms− 1) exhibits a higher wear resistance than the other sliding speeds, with smooth surface and peeling area is less. The main wear mechanism is therefore suggested to be abrasive and slight fatigue wear. Fatigue, brittle fracture, or low deformation cause the loss of material in this method. Moreover, the wear debris between the two surfaces is also used as an abrasive, which is referred to as three-body abrasion [32, 33]. According to the EDS results, the A1 surface showed a high oxygen content in the worn surface, indicating that an extent of oxidation wear had occurred on the surface and formation of hard oxide particles which are effective on the abrasive wear mechanism. The formation of a hard oxide layer leads to changes in the wear mechanism and a reduction in the effect of adhesion. However, the brittle characteristic of the oxide layer might be effective on this mechanism, as applied stress can break the oxide layer with a lower amount of deformation.
When the sliding speed was changed to A2 (30ms− 1), the wear mechanism observed in Fig. 7 was found to be a combination of abrasive wear and severe adhesive wear. This led to a significant increase in wear loss and the area of the spalling pit on the surface. As the wear progressed, numerous pits and flakes started to appear on the worn surface, accompanied by noticeable internal cracks that had the potential to expand further [33]. The presence of hard particles within the spall pit was also observed. At this stage, the wear mechanism can be categorized as delamination wear, where significant deformation causes the formation of adhesion points. These adhesion points can detach subsequently due to shearing forces, resulting in the generation of debris. This phenomenon is commonly referred to as adhesive wear. Over time, the adhesion points can grow, leading to the peeling off the material matrix in the form of lamellae, under the influence of high shear forces. This process ultimately gives rise to severe delamination wear. Additionally, as the sliding speed increases, the brittle oxide layer becomes more vulnerable to breaking under the heat generated by friction. This transition from oxidative wear to adhesive wear further exacerbates the severity of the wear observed.
In Fig. 8 it was found that the A3 (40ms− 1) displayed many plastic deformations and spalling which indicate that serious fatigue wear and adhesive wear occurred. Increasing the sliding speed caused the FeCrNiCuMoSiC10 to develop large spalling pits and deep grooves. The repeated action of sliding at high speeds resulted in a gradual softening of the contact point of the friction surface, which in turn led to a gradual weakening of the surface's resistance to plastic deformation as the temperature of the friction surface. Spalling pits and grooves formed on the worn surface because of the peeled-off debris. According to Zhang et al. [34], continuous friction caused the worn surface to form more pits and grooves. Additionally, the EDS analysis revealed a notable presence of oxygen, indicating the formation of an oxide film on the surface during the wear process. Although the quantity of these oxide phases may not be substantial, the localized temperature rises from friction caused the surface to undergo oxidation.
In the sliding speed of A2 and A3 there is a reduction in lamellar shedding, the worn surface becomes shallower, the oxide film deteriorates, and hard particles are evident. Wear occurs through abrasive and adhesive mechanisms currently. It appears the friction pair removed the oxide film from the composites, leaving behind the worn composites. Due to the shear force, the relatively soft matrix was quickly worn away because of the shear force, leaving behind the harder reinforced particles within the matrix and dividing the adhesion points into discontinuous points. The harder particles cause the abrasion of the matrix material, resulting in the formation of pits. These pits can cause an increase in the friction coefficient and further accelerate the wear rate.