Effect of autogenous GTAW on the reciprocating sliding wear behavior of a carbon martensitic steel

Martensitic steels have been successfully employed in resource-based industries where components must endure aggressive conditions. In industrial practice, many parts of these components are joined by welding techniques. The aim of this work was to understand the influence of welding on the wear resistance of quenched and tempered carbon martensitic steel subjected to dry linear reciprocating sliding micro-wear tests. Weld-joints were produced using autogenous Gas Tungsten Arc Welding process (GTAW). Micro-wear tests were performed at base metal (BM), weld metal (WM), coarse grained heat affected zone (CG-HAZ) and lowest hardness region of heat affected zone (LHR-HAZ). LHR-HAZ was softened during welding process so plastic deformation was facilitated, and consequently adhesion, material displacement and micro-ploughing. WM and CG-HAZ presented a similar martensitic structure, which explain the similarities found on wear behavior. These regions presented the lowest worn volume average values (w). It was interesting to note that despite its highest microhardness value, the highest w was observed for BM. For some BM samples, debris had a key role promoting material loss by micro-cutting which causes great extent of material removal compared to other micro-wear mechanisms as micro-ploughing and adhesion. Due to debris action BM also presented a great dispersion in w results. The results suggest that material loss of welded joint and BM was strongly controlled by micro-wear mechanisms.


Introduction
The usage of martensitic steels in resourced-based industries can be effective for minimizing costs due to wear, thus extending the service life (Mindivan, 2013). The fine grains o f their martensitic structure create characteristics features of high strength, toughness, and resistance to abrasive wear (Białobrzeska, Dziurka, Żak, & Bała, 2018;Białobrzeska & Kostencki, 2015;Frydman, Konat, & Pękalski 2008;Mendez et al., 2014). These properties are obtained through a specific combination of thermomechanical processes and a strictly selected chemical composition that includes low quantities of phosphorus and sulphur (Białobrzeska & Kostencki, 2015). Mishra et al. (2019), highlighted that the addition of low quantities of boron in martensitic steels promotes grain refinement, thus improving strength levels and toughness. In order to improve the wear resistance, these martensitic low -alloyed steels are generally tempered to modify the balance between mechanical properties, and in particular to improve toughness and ductility while keeping a relatively high hardness (Xu, Zwaag, & Xu, 2016).
A significant amount of studies on wear resistance of martensitic steels were also developed (Białobrzeska & Kostencki, 2015;Hacısalihoğlu, Yildiz, & Çelik, 2018;Hernandez, Hardell, Winkelmann, Ripoll, & Prakash, 2015;Mindivan, 2013;Pawlak, Białobrzeska, & Konat, 2016;). Białobrzeska and Kostencki (2015) investigated the abrasive wear characteristics of selected low alloy boron steels in both field experiments and laboratory tests. The obtained results showed that hardness had a secondary role in controlling the wear rate. Results also indicated that abrasive wear rate depends strongly on the predominant mechanism of wear, which in turn is dependent on the microstructure. The authors highlighted that in their study, micro-cutting was the predominant micro-wear mechanism in material removal.
Welding techniques are widely used in industrial practice. The quality of welded joints mainly depends on mechanical properties of the weld zone and HAZ, which in turn is influenced by chemical composition and metallurgical variations of the weld zone. In other words, the quality of welded joints depends on welding process (Pujari, Patil, & Mewundi, 2018). Gas Tungsten Arc Welding (GTAW) and Gas Metal Arc Welding (GMAW) are the most used welding processes. However, GMAW offers problems such as melt and distortion, especially for thin workpieces (Pujari, Patil, & Mewundi, 2018). The welded joint between thin plates are typically made by GTAW. One of its main advantages is its relatively shallow penetration achievable during autogenous welding (Vasantharaja, Vasudevan, & Parameswaran, 2019).
Despite its significance in wear behavior, the effect of weld resulting microstructure and HAZ softening has not received adequate attention. In this study, worn volume and micro-wear mechanisms were evaluated in the weld metal (WM), coarse grained heat affected zone (CG -HAZ), lowest hardness region of the heat affected zone (LHR-HAZ) and base metal (BM) of a wear-resistant steel welded joint. This work aims to analyze the influence of weld resulting microstructure on the wear resistance of a selected martensitic steel subjected to a reciprocating sliding micro-wear test.

Experimental details
The chemical composition of the base material used in this study and its carbon equivalent (CE) according to International Institute of Welding (IIW) is presented in Table 1 and Equation 1, respectively. Its microstructure is fully composed by small grains of tempered martensite as showed in Figure 1. CE formula is used to estimate the weldability of steels; it determines if preheat or post -welding heat treatments are necessary. For the selected martensitic steel used in this work preheat and post -welding treatments was not necessary (SSAB, 2019). (1) The weld-joint was produced by single pass butt-welding of 1 mm thick samples of approximately 300 x 100 mm, using a MA1400 Yaskawa welding robot and a Trans Plus Synergic 4000R power source. The joints were produced using autogenous GTAW, which is a high-quality joint process employed in thin workpieces where the usage of filler metal is not required. The welding parameters used in this work and the corresponding welding energy (E) are listed in Table 2. E can be determined by Equation 2 (SSAB, 2019), where arc efficiency ( ) was considered to be 0.7 (Hackenhaar, Mazzaferro, Gonzalez, & Machado, 2016). The welding was performed using direct current electrode negative (DCEN) and the shielding gas adopted was pure argon. The welding parameters were selected through several exper iments until the weld joint was acceptable (low distortions, full penetration, and absence of discontinuities). To minimize distortion the plates were held in position by clamping devices during the welding.
A schematic drawing of the welded joint is presented in Figure 2. The welding was performed in the same direction of rolling. The normal, transversal and longitudinal directions of the rolling process are referred as ND, TD and LD, respectively. Macro and micro observations, sliding wear tests and microhardness measures were executed in samples prepared with standard metallographic practice, followed by etching in 10% Nital. The microhardness measurements were carried out on ND plane (normal direction of rolling process) using a semi-automatic microhardness tester. These tests were performed at a load of 300g during 10s using a Vickers pyramid micro indenter. Microhardness profiles were determined by taking microhardness readings along lines perpendicular to the welded joint as shown in Figure 2. These profiles were used to identify the gradient of microstructures (HAZ sub-zones) and the locations where reciprocating sliding micro-wear tests were later performed. Microhardness measures were carried out as per ASTM E384-17. The microhardness value obtained for as-received condition was 499 ± 6 HV.
In order to evaluate the influence of the weld resulting microstructure and HAZ softening on wear resistance, reciprocating sliding micro-wear tests were performed at BM, WM, CG-HAZ and LHR-HAZ sub-zone using a tribometer developed at Welding & Related Techniques Laboratory of Federal University of Rio Grande do Sul. The reciprocating micro-wear tests were performed on ND plane using Al 2 O 3 balls with 4.762 mm diameter in ball-on-flat configuration. The tribometer device was able to record the coefficient of friction (CoF) continuously during micro-wear tests. CoF medium values reported in this work were calculated excluding first 15 minutes data, to provide that Al 2 O 3 balls have settled on the worn surface. All tests were carried out in dry condition under 10 N load. The linear sliding speed was 4 mm/s and the reciprocating stroke was set as 2 mm, with oscillation frequency of 1 Hz. The reciprocating micro-wear tests were performed during 2 h totaling a sliding distance of 28.8 m. The reciprocating micro-wear tests were repeated at least three times for each area.
After the wear tests, the worn tracks developed on the surface of the samples were analyzed by a 3Dprofilometry, SEM and OM. In this study, the wear test results were quantified in terms of worn volume, which was calculated by integrating the area across the wear scar profile, and then multiplying the length of the tracks. Each worn trail had 10 cross-sectional profiles evaluated. ImageJ was used for computing the worn volumes of Al 2 O 3 measuring the area of worn surfaces on OM images as per ASTM G133-05.

Welded joint: microscopy and hardness
The macrostructures of welded joint on TD and ND (polished and etched) planes are presented in Figure 3. The samples were successfully autogenous welded in full penetration. Some distortions along the weld bead were observed. No weld defects such as pores and cracks were found. The microhardness profile across the weld-joint is shown in Figure 4. The highest microhardness average value was obtained on BM (499 ± 6 HV) followed by WM (452 ± 10 HV) and CG-HAZ (419 ± 9 HV). At LHR-HAZ microhardness value was 253 ± 6 HV, where hardness reduction reached a range of 50% compared to BM. HAZ extension was approximately 11 mm. BM structure is fully martensitic which is a thermally unstable phase, because of that it will tend to transform during welding, resulting in microstructural changes and softening of the local microstructure along the welded joint.
The microstructure of WM is composed by martensite -in this region the initial structure (BM) was melted and rapidly re-solidified resulting in a homogenous distribution of martensite presented in Figure  5. The microstructure of CG-HAZ was also martensite -this region is located immediately adjacent to the fusion boundary and was heated to temperatures high enough to enable the austenitic transformation, the generated austenite transforms to microstructures with less imperfect density and coarser grain size (Hochhauser, Ernst, Rauch, Vallant, & Enzinger, 2012). During subsequent cooling the coarse austenite is transformed into the martensitic structure presented in Figure 6. The p resence of martensite at WM and CG-HAZ is controlled by cooling rate, which decreases while moving from the weld bead center. The martensite structure in these regions are coarser than BM, the hardness reduction on WM and CG-HAZ were respectively 10 and 16%.   At LHR-HAZ the microstructure was prevailingly over tempered martensite (M) with a mixture of ferrite (F) and carbides (C), as presented in Figure 7. These microstructure features indicate that heating during welding provoked the intensification of carbides precipitation from the initial tempered martensite (BM). Still during heating these carbides coalesced and some of them were dissolved to originate small austenite grains, which during cooling decompose to ferrite (Unfried, Garzon, & Giraldo, 2009). Softening phenomena was observed in this region, which is located at IC-HAZ/SC-HAZ interface. This result is in agreement with other authors who showed in their works that HAZ softening in martensitic steel weld-joints was predominantly controlled by over tempering of martensite and formation of fresh ferrite in the SC-HAZ and in IC-HAZ (Biro et al., 2014;Frydman, Konat, & Pękalski, 2008;Guzman-Aguilera et al., 2018;Krishnan et al., 2006;Lu et al., 2018;Pandey et al., 2019;Pouranvari et al., 2018;Wang et al., 2016).  Table 3 presents the worn volumes of trails (w t ) and Al 2 O 3 balls (w b ) used in the respective reciprocating sliding micro-wear tests. For the welded joint, the lowest worn volume average value was found in WM, followed by BM and CG-HAZ. The highest worn volume among the welded regions was obtained in LHR-HAZ. At this region the wear mechanisms related to plastic deformation were facilitated, which increased material removal by adhesion. The WM worn trail surface is presented in Figure 8 and 9a. In this region the material loss is mostly due to abrasion. Some detachment sites were also identified in Figure 8, however, adhesion occurred in much less pronounced form. The worn surface of the Al 2 O 3 balls used in the reciprocating micro-wear tests in WM is presented in Figure 9b, the balls presented a circular smooth worn surface delimited by debris. The cross-sectional profile of the worn trail indicates that most of the abrasion grooves was shallow and their density per unit area was low, as showed in Figure 9c. According to Bhushan (2000) and Stachowiak and Batchelor (1993), these features suggest that most abrasive wear occurs on microploughing form. Compared to micro-cutting, micro-ploughing leads to discreet wear damage, which resulted in the lowest worn volume. In this case, most of material removed from the grooves was only shifted to its borders, and then plastically deformed and dislocated.

Sliding micro-wear behaviour
The CG-HAZ worn trail surface is presented in Figure 10 and 11a. On this region material removal was due mostly to abrasion and adhesion. Similarly, to WM, the Al 2 O 3 balls used in the reciprocating micro-wear tests, also presented a circular smooth worn surface delimited by debris, as showed in Figure 11b. The crosssectional profile of the worn surface presented in Figure 11c shows shallow abrasion grooves, suggesting that abrasive wear occurred on micro-ploughing form.
The worn surfaces of the trails at LHR-HAZ are presented in Figure 12 and 13a. On this region the main wear mechanisms were adhesion and abrasion on micro-ploughing form, as showed in the cross-sectional profile of the worn surface presented in Figure 13c. The Al 2 O 3 balls used as counter body presented a smooth circular surface delimited by debris, as showed in Figure 13b. LHR-HAZ was severely softened during welding, the ferrite Acta Scientiarum. Technology, v. 43, e50488, 2021 transformation and the over-tempering process of the initial martensitic structure increased ductility, which facilitated the wear mechanisms associated to plastic deformation. LHR-HAZ suffered massive adhesive damage that resulted on a worn volume over 70% greater than WM.     For the BM most of the specimens presented a slight amount of material loss ( = 3.36 10 −5 3 ) caused by micro-ploughing. The abrasion grooves were shallow and no evidence of wear damage by threebody abrasion was found, as showed in Figures 14 and 15a and c. Figure 15b shows the Al 2 O 3 balls surface. For these tests, the balls presented a smooth worn surface where no traces of debris were found.  Even though BM presents high hardness, some samples presented regions were massive wear damage by abrasion occurred. Compared to the BM samples where abrasion damage is uniform and slight along the full length of the trails, the material loss is up to 9 times greater. Figures 16 and 17a present the regions where severe wear damage occurred. According to other works (Guzman-Aguilera et al., 2018;Saroj, Sahoo, Tijo, Kumar, & Masanta, 2017), these deep grooves (Figure 17a) are usually caused by the presence of hard particles (debris) that slide over the surface causing the removal of material along its path.
Debris increase abrasive wear resulting in deeper grooves and massive wear damage. Besides that, according to Kirk, Shipway, Sun and Bennett (2019) debris can also adhere to the surfaces contributing to adhesive wear, which could explain the large detachments sites in Figure 16. The worn surface of the Al 2 O 3 balls presented in Figure 17b also shows evidences of debris action. In fact, these features indicate that abrasion occurred on three-body mode, which causes greater material loss (Bhushan, 2000;Stachowiak & Batchelor, 1993).  During reciprocating sliding micro-wear tests debris are formed with different sizes depending on the tribological conditions (Bhushan, 2000;Hokkirigawa, & Kato, 1988). It is likely that during reciprocating sliding some debris particles were trapped on the contact surface and hardened, increasing debris abrasiveness and promoting massive wear damage on micro-cutting form along contact surface. Al 2 O 3 balls and steel possess poor chemical affinity; because of that balls wear is due only to abrasion. However, debris accumulation on the worn surface can increase chemical affinity with worn trails surface, thus increasing adhesive wear on trails. Note that debris may increase materials loss due to adhesion and abrasion simultaneously.
Hardness is often used to predict the wear resistance of materials exposed to abrasion (Frydman et al., 2008;Mindivan, 2013;Sundström, Rendón, & Olsson, 2001) and according to Bhushan (2000) abrasive wear is proportional to the sliding distance and inversely proportional to the hardness of the wearing material. However, many authors (Dalbert et al., 2019;Mindivan, 2013;Sundström et al., 2001;Trevisiol, Jourani, & Bouvier, 2017) state that hardness must be used with caution when it comes to predicting the Acta Scientiarum. Technology, v. 43, e50488, 2021 wear resistance of materials exposed to abrasion. Wear resistance of a material in a given tribosystem is strongly influenced by the system parameters as well as toughness, work-hardening ability, ductility, microstructure, wear mechanism, etc. (Atkins, 2009;Shah & Bakshi, 2018;Sundström et al., 2001;Trevisiol et al., 2017).
In this study, wear resistance seems to be controlled by hardness. BM, WM and CG -HAZ presented a similar martensitic structure that was responsible for the discreet wear damage of trails. At LHR-HAZ the microstructure was severely softened leading to a greater material loss by plastic deformation related mechanisms.

Coefficient of friction
Coefficients of friction (CoF) values of the BM, WM, CG-HAZ and LHR-HAZ are presented in Table 4. Figure 18 shows the relationship between CoF values of the regions of interest on the welded joint and BM. CG-HAZ had the lowest CoF average value, followed by BM. WM and LHR-HAZ presented quite similar CoF, in the range of 0.43 -0.44. CoF is minimized by region with the highest volume fraction of martensite (CG-HAZ) where less plastic deformation occurred. The highest CoF was reached for the softest sample.
It is shown in Figure 18 that the initial stage of the friction is unstable, and the CoF value continuously increases during this stage with fluctuations. According to Li, Deng, Huang, Ji and Wang (2019), these fluctuations on CoF values at the beginning of sliding wear tests is due to the fact that the upper and lower friction pairs are uneven surfaces in the running-in process. The interaction between the friction pairs causes the settlement of the Al 2 O 3 balls on metal surface, which increases CoF. As the number of reciprocating slides increases, the CoF values begin to stabilize.  18. Examples of evolution of the coefficient of friction during reciprocating micro-wear tests.

Conclusion
In this study, the influence of welding on the reciprocating sliding wear under dry condition of quenched and tempered carbon martensitic steel was investigated. The main results are summarized as follow: • Microstructure features, debris action, hardness and toughness seem to have an important role on microwear mechanisms which in turn are key factors controlling material removal; • For the welded joint, the lowest worn volume average value was found in WM, followed by CG-HAZ, and both presented a similar martensitic structure after welding. For these regions, most material removal was due mostly by micro-ploughing; • In a very narrow area at IC-HAZ/SC-HAZ interface, HAZ softening was observed, caused by a degradation of the initial martensitic structure after welding process. The microhardness reduction reached 50% in comparison with BM. This region in the welded joint showed the highest volume loss during sliding wear tests; • Related to BM, debris had a key role on wear resistance. In some BM samples, debris acted as a third-body and promoted material loss by micro-cutting mechanism, which increased material loss in some regions. For those tests where only micro-ploughing was observed, worn volume value was smaller because most of the deformed material was smashed back into worn trails surface, decreasing material removal.