3.1. Microstructural Characterization of CPW 800 Steel: Base Metal and Welded Region
In this stage, the qualitative aspects of the microstructure obtained from two variations of chemical attacks, nital 2% and LePera, are analyzed and discussed. The results are presented and discussed through qualitative analysis by observing the structure to identify the phases present using the chemical reagents.
Figure 3a shows the metallography of the CPW 800 complex phase steel sample in the longitudinal orientation, subjected to a 2% nital etch for 8 s, at 1000× magnification.
Figure 3b displays the metallography of the CPW 800 steel in the transverse orientation, etched with 2% nital for 9 s, at 1000× magnification. In this image, the following phases are identified: Ferrite (F), Martensite (M), Retained Austenite (RA), Pearlite (P), and Bainite (B).
Using the 2% nital chemical reagent, it is possible to outline the grain contours of the structure, revealing two shades of gray: a lighter shade corresponding to ferrite, and a darker shade corresponding to either retained martensite or retained austenite, bainite, and/or pearlite. This makes it challenging to precisely identify the outlined structure. The observed structure is complex, with phases appearing simultaneously, which is characteristic of typical multiphase structures. The most effective method for this chemical etching involved dipping a portion of cotton in the 2% nital solution and rubbing it against the sample surface at room temperature, specifically 23 °C, in the laboratory where the analysis was conducted.
In
Figure 4a,b, the CPW 800 complex phase steel sample, after being subjected to chemical etching with LePera reagent for 25 s, is viewed in a longitudinal orientation at a magnification of 1000×. The LePera reagent, based on sodium metabisulfite and picric acid, is ideal for highlighting the bainitic structure (B), which appears light brown or dark brown [
36,
37,
38]. The phases containing martensite and retained austenite (RA) cannot be easily differentiated, as these substances appear in light blue or white shades, while ferrite (F) appears in light green to light blue tones. The application of this reagent must be preceded by a pre-etching with 2% nital to delineate the grain contours of the structure, enabling the subsequent etching with LePera reagent to highlight the microconstituents with greater contrast [
39]. The average pre-etching time with 2% nital varied between 20 and 25 s.
At this stage of analyzing the CPW 800 complex phase steel sample (subjected to chemical etching with 2% nital) using the Scanning Electron Microscopy (SEM) technique, as shown in
Figure 5, it is possible to identify the phases present in the material.
Figure 5a,b illustrate the following phases: Ferrite (F), Martensite (M), Retained Austenite (RA), and Bainite (B). The resolution of the BSE (backscattered electrons) image is lower than that of SE (secondary electrons), as the regions cover a larger area than those associated with secondary electron release on the analyzed surfaces [
40,
41].
The phase mapping in the CPW 800 complex phase steel sample with 2% nital chemical etching, as carried out with the aid of the SEM technique, clearly shows the phases present and provides a good definition of the grain contours. As evidenced in the literature review, the presence of ferritic, bainitic, and martensitic and/or austenitic phases can be noted, each with a different morphology, which facilitates their identification.
Figure 6 shows the metallography of the welded region of the CPW 800 material. In this image, the small area affected by the welding process and the changes in the microstructure are clear. In the Molten Zone (MZ) region, slightly elongated grain boundaries characteristic of fusion structures were observed. The entire region affected by the welding process has a length of approximately 1300 µm, indicating the minimal influence of the process on the sheets of CPW 800 material. It was observed that the Thermally Affected Zone (TAZ) formed is very small in laser welding. The CO
2 laser used for this welding process heats and cools the joining points of the sheets simultaneously, thus avoiding excessive heating, which generally occurs in other welding processes. In the base metal (BM) region, more homogeneous grain boundaries can be observed, due to the material’s multi-constituted structure, combined with fine precipitates.
3.2. Tensile Test
Table 1 presents the average tensile test results for eight specimens of non-welded CPW 800 material and laser-welded CPW 800 material, following the specified methodology.
Figure 7 compares the behavior of the specimens in the tensile test.
Figure 7a,b,d, respectively, show the behavior graphs for tensile strength, yield limit, and area reduction. Despite differences in values, when compared with
Table 1, the means and standard deviations indicate that the laser welding process had no significant influence on these parameters. This conclusion is supported by the statistical results of the non-parametric Kruskal–Wallis test, as shown in
Table 2. The values obtained for tensile strength, yield limit, and reduction of area can be attributed to the advantages of laser welding. These advantages include high energy density, high welding speed, small heat-affected zones (HAZ), low material loss, precise control of heat input, and high levels of automation. Additionally, compared to traditional welding methods, laser welding improves the microstructure and reduces the tendency for segregation in the weld zone, resulting in high-quality weld joints [
42,
43,
44,
45,
46,
47]. However, the results for elongation (%) presented in
Figure 7c,
Table 1 and
Table 2, show an increase compared to the non-welded CPW 800 material. The obtained values exceeded expectations for CPW 800. A possible explanation could be the phenomenon of softening in the region around the weld, which has been investigated for TRIP, DP, and HSLA steels [
48].
The problem of softening due to the welding process has been studied in recent years; however, for CPW 800 steel, it still needs to be further discussed. Biro et al. [
49] investigated these characteristics for dual-phase steels and reported that the heat input required for softening the HAZ decreased along with an increase in the carbon content in the martensite [
50]. Park et al. [
51] and Kong et al. [
52] evaluated the softening characteristics for CP steel, but as a function of the boron concentration. For nickel superalloys, heating to the tempering temperature during welding provides softening of the HAZ metal several millimeters away from the weld [
53]. Lambert-Perlade et al. [
54] and Lan et al. [
55] investigated the phase transformation of the HAZ and the effect on the toughness of the microstructure in low-alloy structural steels after simulated welding. However, the microstructure and softening behavior of laser-welded 960 MPa grade high-strength steel joints have recently been reported, and the microstructural evolution of HAZ still needs to be further investigated [
56].
To evaluate the softening, a hardness test was performed with a Vickers diamond indenter, square-based pyramidal, with an angle of 136° between the planes, an application of a load of 1.96 N (200 g), and an indentation time of 15 s. Eight indentations were made in the base material and in the welded material to obtain the average hardness of the region. For the CPW 800 steel, 260 ± 0.73 HRV was obtained, and for the welded CPW 800 material, 230 ± 0.94. Therefore, according to the references and the hardness test performed, softening due to welding can be an explanation for the elongation of the CPW 800 steel.
The tensile test results indicate increases in both yield strength and elongation. The statistical test confirms that the welding process influences these parameters, thus suggesting that the material experienced an increase in ductility after undergoing the welding process.
These results are in line with those of Różański et al. (2016) [
57], where the studied CP steel welded joint exhibited tensile resistance properties equal to or greater than those of the base material. Regarding elongation, Sun et al. (2020) [
58] found that annealing CP800 steels increases the bainite and martensite content, thereby improving tensile strength and elongation properties. The temperatures reached in the HAZ and the observed behavior in the test suggest that similar changes may have occurred in the joint, leading to improved elongation properties, as shown in the results.
3.3. Fatigue Test
Axial fatigue tests were conducted on fifteen specimens of CPW 800 steel, standardized according to ASTM E 466, to compare welded and non-welded materials. The specimens ruptured at the weld zone precisely. Stress levels of 98%, 96.5%, 95%, 90%, and 80% of the average yield stress found in the tensile test (as shown in
Table 1) were applied. The results of the axial fatigue test are presented in
Table 3 and graphically plotted as Stress (S) vs. Log (N) in
Figure 8.
From the analysis of the S-logN curve (
Figure 8a) for CPW 800 complex phase steel, it can be concluded that this material exhibits a higher fatigue resistance limit compared to other advanced high-strength steels such as Dual Phase (DP) steel and Transformation-Induced Plasticity (TRIP) steel, which are also used for structural purposes in vehicles. CPW 800 steel demonstrated a fatigue resistance limit of approximately 650 MPa, compared to 570 MPa for TRIP 780 and 580 MPa for DP 780. This indicates that complex phase steel can be effectively used in vehicle components that require good fatigue resistance, thereby increasing its usability. The superior fatigue limit of CPW 800 steel, compared to DP 780 and TRIP 780, is attributed to its microstructure, which exhibits a homogeneous hardness of its micro-constituents. This homogeneity leads to better performance under fatigue demands [
59,
60,
61,
62].
As shown in
Figure 8b, laser-welded specimens exhibited lower levels of fatigue life compared to the non-welded specimens, even when subjected to lower stress levels. Even under lower voltage conditions, fatigue life never reached the infinite life threshold of (>10
6) cycles. Moreover, fatigue life remained relatively constant at approximately 10
4.4 cycles, regardless of stress variations. This reduction in fatigue life can be attributed to the microstructural changes induced by the welding process. As will be shown, the weld bead forms a martensitic structure. Although this microstructure is harder and more resistant, it weakens the region, leading to a decrease in the material’s fatigue life.
Table 4 presents the non-parametric test results for the axial fatigue test (number of cycles). At a significance level of 5%, the results for specimens with and without laser welding were statistically different. Therefore, it can be concluded that the laser welding process has a significant influence. The pronounced influence of laser welding is evident, as the specimens ruptured in the weld bead precisely. This may be attributed to welding failures, such as low energy efficiency (~10%), material roughness, and mechanical defects.
As seen in
Section 3.1, CP800 steel presents a refined microstructure (
Figure 3,
Figure 4 and
Figure 5). After the laser welding process, the microstructure is altered, even with the laser beam acting on a small thickness, as shown in
Figure 6. This alteration causes a rearrangement of the structure, implying a decrease in resistance and fatigue life. This conclusion is in line with the results of Zhou et al. (2024) [
63], who observed that the strength and fatigue lifetime of complex phase steel increase with the decrease in grain size.
3.4. Impact Test
The results from the Charpy impact test, conducted on ASTM E 23 (sub-sized) specimens, are presented in
Table 5 as absorbed energy (J). The graphical representation of these results is shown in
Figure 9.
Comparing the results obtained, it can be observed that the energy absorbed between temperatures of −40 °C to 60 °C is practically the same (within the dispersion range of the results). This indicates that complex phase steel does not show significant differences in impact energy absorption within the evaluated range, ensuring the reliability of structural calculations used during vehicle design, regardless of temperature variations in potential collision scenarios. When compared to other classes of advanced high-strength steel, such as DP steels and TRIP steels, complex phase steel demonstrates superior impact resistance [
29,
64,
65]. Therefore, complex phase steel can be effectively used in vehicles, particularly in components requiring high impact resistance, such as bumper supports, which are currently produced using DP or TRIP steels.
This increases the usability percentage of complex phase steel in automotive applications.
Table 6 presents the non-parametric test results for the impact test (absorbed energy). At a significance level of 5%, the results are statistically the same, indicating that there is no difference between using and not using the laser welding process. This suggests that laser welding had no influence on the parameter analyzed, likely due to the high-quality weld beads, high welding speed, and high flexibility [
66].
3.5. Fractographic Analysis of Fractured Samples from the Fatigue Test by Scanning Electron Microscopy
Figure 10 shows images obtained from the fractured region during axial fatigue tests, using Scanning Electron Microscopy (SEM), for the stress level of 95% of the yield stress of CPW 800 steel, with laser welding.
Figure 10a shows the initial region of the sample fractured by axial fatigue after laser welding, at a magnification of 100×. Regions of nucleation and propagation of the fatigue crack were observed for a short duration due to the high level of tension. In
Figure 10b, the final region of the same sample is shown, also at 100× magnification, where the ductile characteristics of CPW 800 steel were observed. This indicates that significant plastic deformation occurred prior to fracture during the fatigue test. This is usually accompanied by low levels of hardness and resistance, and considerable tolerance to discontinuities in the material. In
Figure 10c, the initial region of the sample fractured by axial fatigue after laser welding is shown at a magnification of 500×. Regions of nucleation and propagation of the fatigue crack were observed due to the high level of tension, i.e., 95% of the yield stress.
Figure 10d shows the final region of the sample fractured by axial fatigue after laser welding, at a magnification of 500×. Ductile fracture was observed, resulting from the application of excessive force to the metal, which has the ability to deform permanently or plastically before fracture. The presence of shallow dimples or alveoli over the entire fracture surface due to overload was also noted [
67].
Figure 10e shows the final region of the sample fractured by axial fatigue after laser welding, at a magnification of 1000×. Ductile fracture was observed, with alveoli present across the entire fracture surface. The material’s ductile characteristics indicate significant plastic deformation in the fracture region, where the applied shear stress must exceed the shear strength before other fracture modes can occur. Ductile fracture originates from microcavities or dimples (close to the center of the necking region) caused by deformation on the fracture surface [
67].
In
Figure 10f, the central region of the sample fractured by axial fatigue with laser welding is shown at a magnification of 1000×. The presence of striations in the CPW 800 steel was observed, indicating stable crack growth. At this stage, the crack advances progressively with each stress cycle. This stage is highly characteristic of the fatigue process, in which specific marks, such as striations, develop. The final fracture occurs during the last stress cycle, when the crack reaches a critical size, leading to unstable propagation and catastrophic failure [
68,
69].
Figure 11 refers to the fatigue-fracture surfaces of the CPW 800 material specimens (without the laser welding process) subjected to a stress level of 95% of the yield stress.
Figure 11a shows the fracture at 95% of the yield stress, showing the ductile rupture mode of the sample (end region of the sample), at a magnification of 100×.
Figure 11b shows the fracture at 95% of the yield stress, showing the ductile rupture mode (end region of the sample), at a magnification of 500×.
Figure 11c shows the fracture at 95% of the yield stress, showing the region with fatigue striations, at a magnification of 1000×. And
Figure 11d shows the fracture at 95% of the yield stress, showing the shear band region, at a magnification of 1000×.
From the images obtained with the aid of the SEM, the ductile appearance of the fracture surface of the material can be observed, along with the presence of microcavities and small dimples, indicating the presence of fine precipitates in the material. Additionally, scattered cracks are noted on the fracture surfaces. In each case, the fracture originated from a different point on the surface of the specimens, with no fractures occurring due to inclusions. The presence of wide fatigue bands was attributed to the high loads applied to the specimen during the test (95% of the yield stress), and characterizes the onset of fatigue cracking.
According to Bathias (2001) [
68], in the giga-cycle regime, internal defects or variations in the grain size of the material compete with surface defects to cause fatigue fractures. From a probabilistic point of view, it is evident that the greatest presence of defects is concentrated inside the material, compared to its surface. However, if the defect density is higher at the surface, competition may occur between the surface and the interior of the material, leading to fracture initiation at the surface. This phenomenon was observed in the present work. None of the analyses carried out using the scanning electron microscope (SEM) revealed fractures starting in inclusions.
Furthermore, numerous models have been proposed for the nucleation of microcracks, such as those observed. Based on the fractographs presented in
Figure 10 and
Figure 11, which refer to the axial fatigue tests of complex phase steel, it is evident that the mechanisms operated at the nucleation sites. The model proposed by Wood (1956) [
69] suggests that the formation of microcracks is identical to the continuous growth of intrusions. As these intrusions increase and facilitate the formation of new intrusions, microcracks will nucleate.
The fatigue behavior of metals is sensitive to their microstructure. The formation of slip bands and the initial propagation of microcracks are influenced by grain size, type of crystalline structure, material texture, and obstacles to dislocation movement, such as carbides, precipitates, and inclusions [
70,
71,
72,
73]. Analyzing the fatigue resistance limit results from the axial fatigue tests, it is evident that complex phase steel has advantages over two-phase steel and TRIP. This superior performance is directly linked to the microstructural characteristics of complex phase steel, which has a much finer grain refinement compared to other advanced high-strength structural steels. This refinement is one of the hardening mechanisms for metals, specifically the grain size factor [
70,
71,
72,
73].
Fatigue resistance can be increased by reducing ductility, which is the material’s ability to deform. By avoiding the formation of persistent slip bands and increasing the material’s hardness, fatigue resistance can be significantly improved [
71]. The presence of an unwanted phase or a large difference in hardness values between two phases or particles (such as carbides, precipitates, or inclusions) can drastically reduce the fatigue resistance of a material [
73]. This explains the fatigue results observed for the material after the laser welding process, where there was an increase in ductility and the formation of a new microstructure in the molten and heat-affected zone.