Introduction

High-entropy alloys (HEAs) contain five or more principle elements in equiatomic or near equiatomic percentage (at.%). The atomic percentage of each principle element may vary between 5 and 35 (Ref 1, 2). Recently, a lot of interest in HEA is seen mainly due to its high entropy of mixing \((\Delta S_{{{\text{mix}}}} \ge 1.5 R\); where R is the universal gas constant), which is attributed to the formation of a single solid solution having either face-centered cubic (FCC) or body-centered cubic (BCC) or hexagonal closed pack (HCP) phase without any intermetallic phase formation. HEA is also known for their excellent mechanical properties such as wear resistance and microhardness (Ref 3,4,5,6), making it a promising candidate for various applications such as aerospace, defense, cryogenics, and surface engineering (Ref 7,8,9). The use of transition metal in the alloy system prevents the formation of the intermetallic phase and complex phases in the alloy (Ref 10). It has been observed that the addition of W and C in HEAs enhances the wear resistance and hardness (Ref 11, 12). Moreover, the segregation of the refractory element (W) and interstitial element (C) may occur at the grain boundary using the casting route (Ref 11).

The CoCrFeNi quaternary alloy consists of a single-phase FCC structure with low hardness, which limits its wide application (Ref 13). In order to overcome this limitation, the composition and crystal structure of the system is changed by addition of various elements such as Al, Cu, Mn, V, Pd, B, Mo and Ti (Ref 14,15,16,17,18,19,20,21). It has been observed that incorporating Al in CoCrFeNi alloy transforms the FCC to BCC structure. The movement of dislocation lattice in the BCC structure is less as compared to the FCC structure; therefore, BCC structure has lower plasticity and an improved hardness. However, addition of a high amount of Al leads to the formation of a brittle BCC phase making it undesirable for wear resistance (Ref 22). Furthermore, a substantial increase in hardness can be attained due to the formation of the intermetallic phase by addition of Ti in CoCrFeNi (Ref 23). The enhanced wear resistance and improved hardness have also been observed for addition of Ti (Ref 24, 25).

Currently, HEAs are being produced by various thermal spray techniques, such as high-velocity oxy-fuel (HVOF), high velocity air fuel (HVAF), Arc spray, Flame spray (Ref 26, 27), plasma spray (Ref 28,29,30), detonation gun spray and cold spray (Ref 29, 31). Recently, there have been multiple reports on comparing the mentioned methods and understating the advantages as well as the limitations of each of these methods. It has been observed that the various thermal spray techniques was used to synthesized variety of alloys such as CrMnFeCoNi (Ref 27), FeCoCrAlNi (Ref 28), CoCrFeMnNi (Ref 29), NiCoCrAlSi (Ref 30), AlCrFeCoNi (Ref 31) and AlCoCrFeNiTi (Ref 24) to name a few. Among all, the APS coated high-entropy alloy exhibits the superior hardness and wear resistance properties as compared to substrate material. Additionally, APS is a superior method to produce HEAs coatings for high-temperature applications (Ref 32, 33). APS facilitates a higher working temperature (~ 16,000 °C), excellent wear resistance and higher tensile strength (8000–10,000 psi) (Ref 32, 34). APS sprayed AlCoCrFeNiTi exhibits a higher yield strength, ultimate tensile strength, and wear resistance (Ref 35, 36). Moreover, As casted FeCrCoNiW0.3 + 5 at.% C alloy displays good oxidation resistance (Ref 11).

As a surface modification technique, laser cladding is used to improve surface defects and wear resistance (Ref 17). Furthermore, laser surface processing (LSP) has recently been applied to enhance the surface properties such as decreasing surface roughness and porosity, increasing wear resistance, corrosion resistance, and adhesion between surfaces of metallic components (Ref 37,38,39,40,41,42). LSP leads to an enhanced homogeneity of phase distribution, a reduced porosity and formation of new phases. The process being surface-sensitive does not affect the properties of the bulk material (Ref 38). By changing the operation mode of the laser, between pulsed or continuous, the penetration depth and dilution level can be accurately controlled. Additionally, by choosing appropriate laser processing conditions, the material properties on the surface can be modified for any desired applications (Ref 39). Moreover, the grain boundary strengthening mechanism enhances the wear characteristics of the coatings. It has been observed that laser remelting of HEAs leads to a significant improvement in hardness and wear resistance. These coating have lower porosity with the formation of a single solid solution phase (Ref 40,41,42).

The effect of laser surface processing on AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEAs is yet to be explored to the best of our knowledge. In the present study, LSP is used to improve the surface property of the coating produced by atmospheric plasma spraying (APS). Initially both the HEA, i.e., AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C was coated on the stellite substrate using the APS technique. Subsequently, APS coated HEA sample is subjected to laser surface processing. The effects of laser surface processing on surface oxidation, bonding, surface roughness, microstructure, microhardness, physical properties and wear properties were investigated in detail and compared with as-sprayed alloys.

Experimental Procedure

Preparation of High Entropy Alloy Powder

High purity powder 99.5% for Al, 99.8% for Co, 99% for Cr, 98% for Fe, 99.8% for Ni, 99.5% for Ti, 98% for W and 99% for C was procured from Orion Metal Powder Pvt. Ltd., India, having a particle size of 10–50 µm. Powders with a spherical form were favored over other shape for better alloy deposition efficiency. In order to keep powder particles from aggregating together, stearic acid was used. A table-top mixer (ALPHIE, ALPHIE-3) is used to pre-mix the powders at 40 rpm for 12 h for both alloy compositions. The detailed chemical composition of both alloy coatings is given in our previous work (Ref 34). Moreover, to reduce the moisture content and impurities, the powder particle is heated up before it is deposited. A process chart of the preparation and test carried out in the current manuscript is given in Fig. 1.

Fig. 1
figure 1

Process chart for preparation and test carried out in current work

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APS Coating Operation

A 100 mm × 100 mm × 3 mm stellite substrate was ultrasonically cleaned and then immersed in an ethanol bath before the deposition. APS method was used to deposit AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEAs on the substrate. Over 20 µm of sample roughness was used to ensure good coating-to-substrate adhesion (Ref 31). The deposition is accomplished with the help of an APS radial spray gun (METCO 9 MB, India). The robotic arm moves the plasma torch assembly transversely over the substrate holder. The plasma torch moved at a speed of 100 m min−1 in the transverse direction. Spray velocity and temperature have an impact on the splat behavior, which results in a change in microstructure. As the main plasma gas, Argon (Ar) flowed at 42.1 slpm (standard litre per minute), while H2 gas flowed at 2.35 slpm to obtain the targeted plasma jet temperature. In this experiment, the voltage and current used in the spraying were 60 V and 530 A, respectively, with a transverse speed of 100 m/min. A 6-slpm high-pressure carrier gas (Argon) transports powder from the hopper to the convergent portion of the GH-type thermal plasma nozzle. The detailed process parameters used for deposition are given in our previous work (Ref 34). The deposited layer’s thickness was determined to be 200 µm using a ten-point calibration system.

Laser Surface Processing Operation

The sample was cleaned with ethanol before laser surface processing after AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA alloys were APS sprayed on a stellite substrate. The Nd: YAG laser (Quanta-Ray, India) was utilized in Q-switched mode for single-track LSP, with the process parameter used for LSP shown in Table. 1. An example of a pulsed laser is a Q-switched laser. Pulses of energy released at a specific pulse repetition frequency define its output. The Q-switch, a nonlinear crystal, is located within the laser’s cavity. After a certain number of repetitions, it will finally open. When the Q-switch is opened, a powerful laser pulse is generated, increasing the depth of penetration. The Q-switched pulse laser was shown to be preferable to the continuous laser in the present research because it provides double the cooling time, thus reducing the risk of laser ablation. A 210 W power laser used for processing both the APS coating with a scanning speed of 6 mm/s. In order to keep the surface from oxidizing, argon gas is used as a shielding gas. In Fig. 2, a schematic representation of the LSP process is presented.

Table 1 Optimized laser processing parameters for AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA alloy
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Fig. 2
figure 2

Schematic representation of laser surface processing

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Coating Characterization

After the laser surface processing sample was polished and etched, with prepared according to ASTM: E1920 standard. The sample mold is made using commercially available cold mounting and then grinded using a belt grinder BGGP (METCO Ltd., India). The microstructure and microhardness indentation were revealed after mechanical polishing on the laser-processed materials using a variable speed polisher BAINPOL-VTD individual drive (METCO Ltd., India).

The surface morphology and microstructure of the HEAs coating were examined using a field emission scanning electron microscope (FESEM) (JEOL, JSM 7900F) before and after laser processing at a working voltage of 15 kV. Energy dispersion x-ray spectroscopy (Oxford Instruments) was used to examine the chemical composition of the coated AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEAs and the XPS data were obtained utilizing (Nexsa™, Thermo Scientific,) instrument. The collected data were calibrated by assigning the most intensive peak to adventitious carbon at 284.5 eV. The XPS data were analyzed using WinSpec software. The background used for the fitting was Shirley and subtracted for ease of fitting. All peaks were fit using mixed singlet. The carbon content in the coating was determined using the CHNS analyser (Elementor, UNICUBE). The functional groups of both LSP HEAs coatings were confirmed by FTIR spectroscopy (Nicolet iS50, Thermo Scientific). The optical surface profilometer (OLYMPUS, DSX510) was used to evaluate the surface roughness of both HEAs before and after laser processing with a cutoff length of 300 µm and a surface area of 399 µm2. X-ray diffraction spectroscopy (Empyrean, Malvern Panalytical,) Bragg–Brentano method, Cu- source was used to determine the coating’s crystal structure and phase orientation with 10 s dwell time and 0.5-degree scanning step. The corresponding peak has been identified and compared to the standard from the powder diffraction dataset from the ICDD (International Centre for Diffraction Data). Microhardness of both HAEs was determined by applying a 200 gf load for 15 s of dwell time using the Vickers microhardness tester (Falcon 400, Innova Test). Nanoindentation was performed using (FoundatiONE-Multiscale testing system, India) under a load of 5000 µN.

Samples were polished with sandpaper of various grit sizes and cleaned in ethanol before the wear test analysis. An Al2O3 ball with a diameter of 10 mm was used as a counter body against a smooth coated surface in a ball-on-disc tribometer (POD 4.0, Ducom Instruments). The experiment is carried out under dry sliding conditions at 25 °C with a load of 5 N. The speed of the disc was 600 rpm with a wear track diameter of 10 mm. The Raman analysis was performed on the worn surface post-wear test using Raman Spectrometer (R532, Enspectr) with a working wavelength of 532 nm. The background for the Raman data was subtracted using a liner background and the data were normalized for proper identification of peaks.

Results and discussion:

Powder Characterization

To produce a pore-free, higher-density coating, spherical-shaped powder particles were chosen over flaked-shaped particles. Figure 3(a) and (b) shows the powder morphology AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C alloys, respectively. Figure 3 reveals that the majority of powder particles are spherical.

Fig. 3
figure 3

FESEM powder particle morphology of (a) AlCrCoFeNiTi and (b) FeCrCoNiW0.3 + 5 at.% C

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The size of the powder particles ranged from 10 to 50 µm. Figure 4 shows the particle distribution of different AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA constituents, which shows that the majority of the powder particle size ranges from 30 to 40 µm.

Fig. 4
figure 4

Particle distribution of different AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA constituents

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Surface Characterizations

Samples of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings were subjected to optical profilometer studies to determine the difference in line surface roughness (Ra) and areal surface roughness (Sa) after and before laser processing. Table 2 shows the pre- and post-laser sample areal surface roughness values. Both alloys line roughness profiles are shown in Fig.5. Variations in the line roughness profile (e.g., peaks and valleys) were identified before and after laser processing. Figure 5(a) shows the 2D line roughness profile of the AlCrCoFeNiTi alloy and the LP_ AlCrCoFeNiTi sample. The AlCrCoFeNiTi and LP_ AlCrCoFeNiTi samples exhibit line roughness of 6.05 µm, and 2.52 µm, respectively. In addition, FeCrCoNiW0.3 + 5 at.% C and LP_ FeCrCoNiW0.3 + 5 at.% C sample displays a substantial variation in the 2D line profile, which is equivalent to 12.07 µm and 3.47 µm, respectively, (Fig. 5(b)). The results show that laser surface processing improves homogeneity and reduces alloy surface undulation in both alloys by localized heating, melting, and solidification of the surface.

Table 2 Surface roughness of the deposited AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA APS coatings before and after laser processing
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Fig. 5
figure 5

2D line roughness profile of (a) AlCrCoFeNiTi and, (b) FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser processing

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The areal roughness value (Sa) after laser processing is significantly reduced in both the APS coatings, with the decrease in roughness attributed to localized heating, melting and solidification of the surface. Figure 6 shows the 3D surface roughness profiles of both samples. The AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coating exhibits needle-type spikes before laser processing, as illustrated in Fig.  6(a) and  (c). The 3D roughness profile homogeneity is smoothened and homogenized after laser processing, as illustrated in Fig. 6(b) and   (d). Whereas, the surface roughness of LP_AlCrCoFeNiTi sample decreases by 29.03%, that of LP_FeCrCoNiW0.3 + 5 at.% C alloy decreases by 29.59% after laser processing. Laser surface processing facilitates localized heating and rapid solidification attributed to lower surface defects, porosity and lower surface roughness. Both alloys reduced surface roughness improves wear resistance by reducing friction and asperity contact in sliding environments. The surface roughness analysis results are listed in Table 2.

Fig. 6
figure 6

3D surface roughness profile of the (a) as sprayed AlCrCoFeNiTi, (b) laser processed AlCrCoFeNiTi, (c) as sprayed FeCrCoNiW0.3 + 5 at.% C, and (d) laser processed FeCrCoNiW0.3 + 5 at.% C APS coatings

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FESEM cross-sectional image of laser processed AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coating is shown in Fig. 7(a) and (b), respectively. Both APS coating exhibit superior metallurgical bonding between the coating and substrate. There are no cracks or voids observed at the interface in both the APS coating. The dilution percentage is computed using the following equation to investigate the coating and substrate dilution levels.

$${\text{Dilution}} \left( D \right)\% = \frac{{A_{{{\text{ms}}}} }}{{A_{{\text{o}}} + A_{{{\text{ms}}}} }} \times 100$$
(1)

where, \(A_{{{\text{ms}}}}\) represents the area of the melted substrate during laser processing and \(A_{O}\) is area of overlay coating. The melted substrate thickness is measured as 20.8 μm, and 18.5 μm in the case of laser processed AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coating, respectively. The overlay coating thickness is measured as 199.54 μm and 201.32 μm. The dilution percentage is evaluated over a width of 1 mm. The dilution percentage is calculated as 9.43% and 8.41% in case of laser processed AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coating, respectively. The coating dilution level is controlled by the laser power, scanning speed, stand-off distance, and mode of operation of the laser. In general, high-temperature spraying procedures increase porosity in coatings, resulting in a large number of voids, which act as fracture initiation sites. However, the formation of pores in the coating can be reduced by utilizing laser processing and laser-optimized process parameters. In the interface region, the coating’s decreased porosity and fewer interlaminar cracks near the interface region. However, a dilution level study revealed a strong metallurgical bonding between coating and substrate.

Fig. 7
figure 7

FESEM image of coating cross section of (a) as spreyed AlCrCoFeNiTi, (b) as sprayed FeCrCoNiW0.3 + 5 at.% C (c) laser processed AlCrCoFeNiTi, and (d) laser processed FeCrCoNiW0.3 + 5 at.% C APS coatings

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Phase Analysis

Figure 8 shows the XRD results of as-sprayed and laser processed AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C HEA APS coatings. The diffraction results of the LP_ AlCrCoFeNiTi coating show two BCC solid solution phases of FeCr (BCC1) and AlNi (BCC2); both peaks were separated by the shift in the lattice parameter from 2.93 to 2.94 Å attributed to the addition of Ti. The BCC (BCC1) phase denotes diffraction peaks with high intensities, whereas the BCC (BCC2) phase denotes peaks with lesser intensities. In XRD, the Al-rich phase may be associated with the BCC1 phase (ICDD card no – 00–006-0694). The high-intensity BCC2 phase is overlapped and widens, making it difficult to see the low-intensity BCC2 phase (ICDD card no: 03–065-4198). As compared to the as-sprayed coating, more BCC phases developed in the LP_AlCrCoFeNiTi alloy, whereas minimal FCC phases were eliminated.

Fig. 8
figure 8

XRD patterns of (a) LP_ AlCrCoFeNiTi, (b) AlCrCoFeNiTi, (c) LP_ FeCrCoNiW0.3 + 5 at.% C and (d) FeCrCoNiW0.3 + 5 at.% C APS coatings

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X-ray diffraction pattern for LP_ FeCrCoNiW0.3 + 5 at.% C HEA coating exhibit the one FCC (ICDD card no – 00–047-1417, 00–033-0397 and 44–0962) solid solution phase containing two carbide phases, W-rich carbide represented as (WC) (ICDD card no – 003–1096 and 020–1315) and Cr rich carbide (ICDD card no – 035–0783, 35–0804 and 036–1482) indicated as (CrC). It is clear from the XRD results that there is no change in the crystals structure of the APS coated and LP samples. Because of the low carbon content, the peaks of W and Cr-rich carbides closely resemble those of W2C type carbide and Cr23C6 type carbide. On comparing the APS-sprayed coating with the laser-treated coating it was observed that there is an increase in the carbides phase in the laser-treated coatings. This may be due to the diffusion of carbon in Cr and W (Ref 11).

Microstructural analysis

Laser-treated AlCrCoFeNiTi HEA coating deposited by the APS technique, (Fig. 9a, b) has a lamellar microstructure. FESEM images showed fewer pores on the laser-processed surface of the sample than on the as-sprayed one (Fig. 9a, b). Generally, laser processing is known to cause (CrMnFeCoNi) coating to hot crack (Ref 42); however, in the present case, FESEM examination shows minimum interlaminar cracks in the laser-treated sample.

Fig. 9
figure 9

FESEM cross-sectional image of (a) as sprayed AlCrCoFeNiTi, (b) LP_ AlCrCoFeNiTi, (c) as sprayed FeCrCoNiW0.3 + 5 at.% C, and (d) LP_FeCrCoNiW0.3 + 5 at.% C APS coatings

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There are three distinct regions present in the laser-processed AlCoCrFeNiTi APS coating (Fig. 9b). The point EDS confirmed that the a2 region is associated with Al-rich phase, whereas the a3 region indicates Cr-rich phase with Al depletion, and the a1 region denotes rich in Ti-rich oxides.

As-sprayed FESEM image of the FeCrCoNiW0.3 + 5 at.% C APS coating is shown in Fig. 9(c). It comprises three distinct regions. Point EDS confirm that the a2 region is rich in CrC and Fe phase, whereas the a3 region is rich in WC phase and the a1 region is rich in Fe, Ni and Co phase. Due to the dispersion of carbon in Cr and W, WC and CrC phases are found. Figure 9(d) shows a FESEM image of a laser treated FeCrCoNiW0.3 + 5 at.% C HEA alloy coating denoted as (LP_FeCrCoNiW0.3 + 5 at.% C) with same distinct phases as-sprayed one. LP_ FeCrCoNiW0.3 + 5 at.% C APS coating microstructure revealed the more homogeneous distribution of phase, as shown in Fig. 9(d).

Prior to and after laser processing, the chemical constituent elements of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coating was identified using the area EDS and point EDS. Table 3 shows the EDS findings at several locations marked in Fig. 9. Moreover, the carbon content in the FeCrCoNiW0.3 + 5 at.% C and LP_FeCrCoNiW0.3 + 5 at.% C is found to be 7 and 6.2 at.%, respectively. The EDS results demonstrate a perfect correlation between the nominal compositions of each element used in coating synthesis. The homogenous dispersion of elements attributed to a little change in the atomic percentage of Al, Cr, Fe, Ni, Ti, W, and C after laser processing. Because of the high melting temperature and oxygen formation on the grain boundary, the atomic percentage of oxides was increased in laser-treated samples for both alloys. In addition, XPS tests were carried out to discover more about oxygen’s significance.

Table 3 Chemical composition of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser processing
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Figure 10 shows the results of an elemental mapping study performed using EDS on the alloy coatings before and after laser processing. The distribution of all of the elements in the area may be seen via the use of elemental mapping. Figure 10(b) and (d) demonstrates that LP_AlCrCoFeNiTi and LP_FeCrCoNiW0.3 + 5 at.% C coatings constituents have a more homogenous distribution and are mixed well. However, LSP does not affect W due to its higher melting temperature, and W can be seen as a particle phase. Due to the fast heating and solidification effects during LSP, a more uniform distribution of constituents is observed in contrast to the as-sprayed alloys.

Fig. 10
figure 10

EDS elemental mapping of (a) AlCrCoFeNiTi, (b) LP_AlCrCoFeNiTi, (c) FeCrCoNiW0.3 + 5 at.% C and (d) LP_FeCrCoNiW0.3 + 5 at.% C APS coatings

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Chemical Analysis of the Surface

AlCrCoFeNiTi alloy’s carbon 1 s spectrum can be split into three distinct peaks at 284.5, 285.9, and 288.2 eV, correspondingly as shown in Fig. 11(a). C 1 s peaks at 285.9 eV and 288.2 eV indicate the alloy contains hydrocarbon compounds. The Al 2p peak at 73.0 eV may be attributed to metallic Al, whereas the peak at 73.5 eV indicates the existence of an AlNiO (Ref 43) alloy, and the peak at 74.5 eV can be attributed to Al2O3 (Ref 44). Ni 2p has three peaks at 854.0, 855.0, and 861.23 eV, respectively. NiO is responsible for the peak at 854.0 eV (Ref 46), whereas Ni(OH)x is responsible for the peak at 855.0 eV (Ref 49). There is a Ni 2psat peak at 861.2 eV (Ref 46). Both the peaks at 575.6 and 576.8 eV of Cr 2p have been attributed to Cr oxides (Ref 45). The 779.8 eV Co peak may be attributed to Co2O3, whereas the 781.8 and 785.5 eV Co peaks can be attributed to Co3O4 and CoO, respectively (Ref 45). Both the Fe 2p peaks, 711.7 and 709.7 eV may be attributed to Fe2O3 and FeO, respectively (Ref 46). TiO2 has two oxidation states, with peak energies of 457.9 and 456.2 eV, respectively (Ref 47). O 1 s peak at 529.5 eV indicates the presence of several metal oxides, and the peak at 530.8 eV indicates the presence of both AlNiO and NiOH.

Fig. 11
figure 11

XPS spectra of AlCrCoFeNiTi APS coating (a) before laser processing and, (b) after laser processing

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Post laser processing a minor changes are observed in LP_AlCrCoFeNiTi coating in contrast to as sprayed coating, as shown in Fig. 11(b). C 1 s shows an additional peak at 283.5 eV suggests the presence of carbides. Al 2p peak at 71.7 eV can be assigned to metallic Al, 73.9 eV suggests the presence of AlNiO alloy (Ref 43), whereas the peak at 77.4 eV can be assigned to Al2O3 (Ref 44). The Ni 2p peak at 856.2 eV can be assigned to AlNiO (Ref 43), whereas the peak at 852.0 eV is assigned to NiTi (Ref 45), and 852.7 eV can be assigned to Ni0 (Ref 46). The Ti 2p peaks 453.6 eV and 454.1 eV can be assigned to Ti0 (Ref 47) and NiTi, respectively (Ref 45). The peaks of Cr 2p at 573.3 eV, 574.0 eV and 576.26 eV can be assigned to CrC, FeCr (Ref 48) and Cr2O3 respectively (Ref 46). The peaks of Fe 2p of 706.0 eV, 707.46 eV and 710.9 eV can be assigned to Fe0, FeCr (Ref 48) and Fe2O3 respectively (Ref 46). Co peaks at 77.6 eV, 782.2 eV and 793.5 eV are assigned to Co0, CoO, and Cosat, respectively (Ref 46).

Pre laser processing FeCrCoNiW0.3 + 5 at.% C alloy’s carbon 1 s peak is associated with carbide at 282.8 eV, whereas chromium hydrocarbon complexes may be found at 286.3 and 287.9 eV as shown in Fig. 12(a). Three significant peaks at 530.18, 531.28, and 533.4 eV can be seen in the O 1 s spectrum. A variety of oxides of Cr, Fe, Ni, Co, and W may be responsible for the 530.18 eV peak (Ref 46). On the other hand, the peak at 531.2 eV suggests the alloy contains hydroxyl species. In the peak of Cr 2p peak at 573.6 eV, is correspond to chromium carbide (Ref 48), while, peak at 578.5 eV is associated with chromium hydrocarbon complex (Ref 47). The presence of CrO is inferred from the peak at 576.3 eV. The 777.8, 778.58, and 781.9 eV Co 2p peaks may be attributed to metallic Co, CoNi alloy (Ref 52), and Co(OH)x, respectively (Ref 46). Three peaks of Ni 2p spectrum at 852.9 eV is allocated to CoNi alloy (Ref 51), whereas 853.8 and 856.1 eV are assigned to NiO and Ni(OH)x, respectively (Ref 52). Fe 2p peak at 706.4 eV may be attributed to metallic Fe, while the peak at 707.9 eV can be attributed to FeC, which corresponds to the 282.8 eV C 1 s peak. The other two peaks at 710.7 eV and 714.4 eV may be ascribed to Fe2O3 and FeOOH, respectively (Ref 46, 51). W 4f has two peaks, one at 30.6 eV and the other at 34.9 eV may be assigned to metallic W and WC, respectively (Ref 52).

Fig. 12
figure 12

XPS spectra of FeCrCoNiW0.3 + 5 at.% C APS coating (a) before laser processing and (b) after laser processing

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Post laser processing FeCrCoNiW0.3 + 5 at.% C alloy exhibits minor changes in comparison to as sprayed counterpart, as shown in Fig. 12(b). The additional peak of carbon 1 s peak at 283.1 eV suggests the presence of carbides. The Cr 2p peaks at 573.4 and 577.3 eV can be assigned to chromium carbide (Ref 50), CrFe alloy (Ref 48) and chromium oxide, respectively. The Fe 2p peak at 707.2 eV confirms the presence of CrFe alloy (Ref 48), whereas the peaks at 706.3 eV can be assigned to metallic Fe (Ref 46, 51). One peak is observed for metallic Co at 777.8 eV and metallic W at 30.4 eV (Ref 52). Ni 2p shows the presence of metallic Ni at 869.41 eV and NiO at 872.91 eV (Ref 52).

AlCrCoFeNiTi and FeCCoNiW0.3 + 5 at.% C HEA alloy’s surface before and after laser processing was observed using Fourier transform infrared spectroscopy to detect molecular vibrational states as shown in Fig. 13. AlCrCoFeNiTi HEA coating’s vibrational characteristics spectrum, with peaks ranging from 444 to 681 cm−1, indicates the existence of different metal oxides such as Fe–O, Cr–O, Co–O, Ni–O, Ti–O, Al–O or Ni–O stretching mode on the sample surface as shown in Fig. 13 (Ref 53, 54). Similar findings were found in the preceding section’s XPS investigation; both confirm the formation of various metal oxides at the coating surface. Vibration spectra show a Ti-O-Ti bond at 670–450 cm−1 wavenumber (Ref 55). Characteristics spectrum presence at 1706 cm−1 attributed to C = O stretching mode, although the spectra present at 1210 cm−1 ascribed to the C = O stretching mode. Furthermore, the sequence of distinctive peaks between 1400 and 1600 cm−1 shows Ti-O-C molecular vibration, which verifies the bonding between Ti and C (Ref 56). Vibrational spectra displayed at 1568, 1690 and 1727 cm−1 were ascribed to C = O asymmetric stretching, C = O stretching and -C = O stretching correspondingly, and spectra present at 2350 cm−1 related to -O-C-O- bond (Ref 57). However, FTIR analysis of LP_AlCrCoFeNiTi sample does not have a significant deviation from the value assigned to the as-sprayed sample. Similarities can be seen between the FTIR and XPS findings. The quality of the coating was confirmed by the FTIR findings, which showed no impurities on the sample surface.

Fig. 13
figure 13

High resolution FTIR molecular vibrational spectra of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser processing

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It can be seen in Fig. 13 that the FeCrCoNiW0.3 + 5 at.% C sample’s distinctive vibrational spectrum has many peaks between wavenumbers 444 and 681 cm−1. These peaks have been linked to metal oxides in stretching mode, such as Fe–O, Cr–O, Co–O, and Ni–O (Ref 53). Shi et al. have revealed that WC exhibits some distinct bands from 1336 −1 cm to 1590 cm−1  Ref 58). The characteristics peak at 1496 cm−1 wave number represents the C=O stretching mode. The vibrational spectra at 1568 and 1690 cm−1 corresponded to C=O asymmetric stretching, and C=O stretching mode, respectively. The peaks present at 2350 cm−1 signify the presence of the –O–C–O– group (Ref 57). The LP_FeCrCoNiW0.3 + 5% C sample does not deviate from the value assigned to the as-sprayed sample. It shows a peak at 1727 cm  −1 associated with C=O asymmetric stretching. The XPS and FTIR results confirm that there are no compositional changes in the LP_FeCrCoNiW0.3 + 5% C sample before and after laser processing.

Physical Properties

As illustrated in Fig. 14, the microhardness is evaluated at the coating and substrate cross-section. The microhardness is measured at five different locations in the same region, and the average microhardness is reported. The average microhardness of LP_AlCrCoFeNiTi and LP_FeCoCrNiW0.3 + 5 at.% C is calculated as 792 ± 24 HV0.2, and 410 ± 16 HV0.2, respectively. An average microhardness is used since the phase in the alloys has a small size distribution. The LP_AlCrCoFeNiTi had substantially higher microhardness attributed to a sprayed sample due to many reasons. Laser processing reduces porosity and homogeneity in the microstructure, enhancing surface hardness. The increase in hardness is attributed to the development of the BCC phase and the dissolution of the FCC phase upon laser processing, as shown by the XRD findings. The BCC2 phase exhibits a greater weight ratio than the BCC1 phase, according to XRD findings. Greater BCC2 phase content in the alloy is associated with higher hardness (Ref 40). Al and Ti have a larger atomic radius than other elements, resulting in lattice distortion and an increase in hardness for these alloys. This is also true for the LP_FeCrCoNiW0.3 + 5 at.% C alloy, where the microhardness has increased because more CrC and WC carbides have formed compared with what was found in the original sprayed coating, as shown by the XRD findings. LP_ FeCrCoNiW0.3 + 5 at.% C alloy exhibit higher hardness due to diffusion of carbon in Cr and W (Ref 11) upon the laser heating and rapid cooling. It is observed that laser surface processing enhances the microhardness of the as-sprayed AlCrCoFeNiTi and as-sprayed FeCrCoNiW0.3 + 5 at.% C alloy From 754 ± 12 HV0.2 to 792 ± 24 HV0.2 (i.e., 5.03%) and 388 ± 16 HV0.2 to 410 ± 16 HV0.2 (i.e. 5.67%), respectively. Figure 14 shows the microhardness results of both the alloy before and after laser processing at the coating substrate cross-section region.

Fig. 14
figure 14

Microhardness results of the AlCrCoFeNiTi and FeCoCrNiW0.3 + 5 at.% C APS coatings before and after laser processing

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The nanoindentation technique is used to investigate the physical properties of AlCrCoFeNiTi and FeCoCrNiW0.3 + 5 at.% C HEA coating before and after laser processing. The physical properties of each phase are evaluated under a maximum load of 5000µN, 30 indents are made over each phase, and the average value is reported. Table 4 lists the results of physical properties such as contact depth (hc), contact stiffness (S), reduced elastic modulus (Er), and hardness (H).The reduced elastic modulus (\({E}_{r}\)) and hardness (H) is obtained using Oliver and Pharr method (Ref 59).

$$E_{r} = \frac{\sqrt \pi }{{\sqrt {A_{c} } }}\frac{S}{2}$$
(2)
$$H = \frac{{P_{{{\text{max}}}} }}{{A_{c} }}$$
(3)

where, S is the contact stiffness, Ac is the projected contact area, and Pmax is the maximum load applied. Load Vs depth curves for both the alloy before and after laser processing are shown in Fig. 15. The a2B region (Al-rich) exhibits the lower depth, followed by the a1 region (Ti rich) and a3 region (Cr-FeNi rich) in the case of AlCrCoFeNiTi and LP_AlCrCoFeNiTi alloy as shown in Fig. 15(a) and (b) respectively. It implies that the a2 region (rich in Al) is more resistant to plastic deformation in contrast to the a1 region (rich in Ti) and a3 region (rich in Cr-Fe–Ni). However, the a3 region shows lower depth, followed by a2 and a1 regions in the case of FeCoCrNiW0.3 + 5 at.% C and LP_FeCoCrNiW0.3 + 5 at.% C alloy, as shown in Fig. 15(c) and (d), respectively. It infers that the a3 region (rich in W) has more resistance to plastic deformation in comparison to the a2 region (rich in Cr) and a1 region (rich in Fe-Co–Ni). The laser processed both the APS coatings display lower contact depth in contrast to the as-sprayed alloy, which may be attributed due to the formation of various metal oxides at the surface as confirmed by XPS and FTIR results in the preceding section. The average reduced elastic modulus of 90 indents each sample is calculated as 131.42 ± 28 Gpa, 146.66 ± 23 Gpa, 97.51 ± 36 Gpa, and 114.33 ± 24 Gpa for AlCrCoFeNiTi, LP_AlCrCoFeNiTi, FeCoCrNiW0.3 + 5 at.% C and LP_ FeCoCrNiW0.3 + 5 at.% C HEAs, respectively. The average hardness is evaluated as 9.11 ± 5 Gpa, 11 ± 4 Gpa, 6.6 ± 3 and 8.23 ± 3 Gpa for AlCrCoFeNiTi, LP_AlCrCoFeNiTi, FeCoCrNiW0.3 + 5 at.% C and LP_FeCoCrNiW0.3 + 5 at.% C HEA, respectively.

Table 4 Nanoindentation results of individual phases observed in AlCrCoFeNiTi and FeCoCrNiW0.3 + 5 at.% C APS coatings coating before and after laser processing
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Fig. 15
figure 15

Load-depth curve of (a) AlCrCoFeNiTi, (b) LP_AlCrCoFeNiTi, (c) FeCoCrNiW0.3 + 5 at.% C and (d) LP_FeCoCrNiW0.3 + 5 at.% C APS coatings

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The LP_AlCrCoFeNiTi and LP_FeCoCrNiW0.3 + 5 at.% C HEA display an increase in the average reduced elastic modulus (Er) value with 10.39, 14.71% and average hardness with 17.18 and 19.80%, respectively. The reduced elastic modulus depends upon the crystal structure and the chemistry of the alloy (Ref 60). Adding elements like C changes the ideal shear strength resulting in a change in nanoindentation properties. The lower elastic modulus/hardness or increased elastic modulus/hardness is attributed due to shear instability (Ref 60). The increase in hardness of laser-processed alloy indicates a reduction in plasticity (Ref 32). The difference in hardness between the phases in both alloys is accredited due to composition, phase structure, solidification rate, and oxidation kinetics during the plasma spray and laser processing processes. (Ref 32, 60). The Vickers microhardness value and the nanoindentation hardness have a substantial difference. As compared to microindentation, nanoindentation shows higher hardness values owing to dislocations because of the small indentation area and a comparatively low load of (5000 µN) (Ref 32). Therefore, the nanoindentation findings revealed that laser-treated alloys display higher hardness values that are dependent on phase formation, oxidation state and solidification rate during laser processing (Ref 40).

Tribological Properties

In the dry sliding environment with a load of 5 N, the wear behavior of both alloy samples i.e. before and after laser processing is studied at room temperature (25 °C). The wear volume loss is calculated from the wear track profile using optical microscope as 14 × 10–4 mm3, 8.9 × 10–4 mm3, 38.9 × 10–4 mm3 and 23.45× 10–4 mm3 in case of AlCrCoFeNiTi, LP_AlCrCoFeNiTi, FeCrCoNiW0.3 + 5 at.%. C and LP_FeCrCoNiW0.3 + 5 at.%. C APS coating, respectively. The volume wear rate of all the samples is calculated and shown in Fig. 16. The volume wear rate calculated is 2.8 × 10–7, 1.78 × 10–7, 7.78 × 10–7 and 4.69 × 10–7 mm3/N.m for AlCrCoFeNiTi, LP_AlCrCoFeNiTi, FeCrCoNiW0.3 + 5 at.%. C and LP_FeCrCoNiW0.3 + 5at.%. C APS coating, respectively. The LP_AlCrCoFeNiTi alloy has a 36.42% reduced wear volume loss in comparison to the as-sprayed sample. In addition, LP_FeCrCoNiW0.3 + 5 at.% C sample revealed a substantial reduction of 39.71% in wear volume loss compared to the APS sample, which can be attributed to the formation of hard carbides during the laser processing as suggested by the XPS and XRD results. Lower volume loss in the wear resistance test makes evident the superior wear resistance property of the LP_AlCrCoFeNiTi APS coating sample. The rapid phase change from solid to liquid to solid is facilitated by laser processing increasing the sample’s hardness and wear resistance. Laser processing is also attributed to the lower wear volume loss under dry sliding conditions because of the formation of new BCC phases and CrC and WC phases revealed by XRD. The average increase in the hardness value of 5.03 and 5.67% was attributed to the reduction in volume loss of 36.42 and 39.71% in both the laser-processed samples and follows Archard’s law (Ref 34). Laser surface processing increases the hardness of both alloys, and their wear resistance indicates that laser surface processing is an efficient way to improve the wear performance of coatings.

Fig. 16
figure 16

Volume wear rate of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser processing

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Figure 17 shows the coefficient of friction (COF) curve for as-sprayed and laser-treated AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings. The average friction coefficient of the AlCrCoFeNiTi alloy before and after laser processing is 0.46 and 0.45, respectively. Similarly, the average coefficient value for the FeCrCoNiW0.3 + 5 at.% C APS coating before and after laser processing was calculated as 0.53 and 0.52, respectively. Laser processed alloys sample exhibit a significant decrease in the coefficient of friction value attributed due to the higher hardness value (Ref 40). The initial rapid increase in the coefficient of friction observed in Fig. 17 can be due to the change in contact area between the tribo pair. After 230–330 s of testing the COF curve reaches equilibrium owing to the increased contact area and decreased contact pressure (Ref 61) between the tribo pair. The development and breakdown of the oxide layer on the surface of the samples leads to the observed fluctuations in the COF (Fig. 17). Over time, the tribo pair heats up, forming an oxide layer that reduces COF and the XPS and FTIR suggests the development of the oxide layer (Fig. 11, 12 and 13). As the test continues, the oxide layer thins and eventually breaks, increasing the COF value. Although oxide production and braking cause COF curve variation, sliding velocity and applied load also play a role. A low coefficient of friction was found in the laser processing of both alloys.

Fig. 17
figure 17

Coefficient of Friction curve of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser processing under a load of 5 N

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The results of FESEM analysis is presented in Fig. 18. It is used to investigate the wear mechanism in both alloy samples before and after laser treatment. AlCrCoFeNiTi alloy wear morphology showed the presence of an adhesive wear mechanism Fig. 18(a) with exfoliation of the hard particle. Micro scratches are a result of the coating’s abrasiveness. On the worn surface, there are a few grooves that may be ascribed to particle pull-out. Less surface damage is seen in LP_AlCrCoFeNiTi APS coated sample, as shown in Fig. 18(b), compared to the as-sprayed sample, as shown in Fig. 18(a). The decreased friction coefficient may be caused by the rolling of abrasive wear debris (Ref 40). An increase in temperature causes the oxide layer to develop and break, changing the friction coefficient abruptly. The wear debris moving between the ball and the sample reduces the coefficient of friction in LP_AlCrCoFeNiTi APS coating. Besides, wear debris accelerates the three-body abrasive wear. The micro scratch confirms the abrasive wear mechanism. Abrasion and adhesion are the primary wear mechanisms observed in LP_AlCrCoFeNiTi HEA. However, FeCrCoNiW0.3 + 5 at.% C APS coated samples exhibit adhesive wear with more surface damage attributed to the lower hardness value and confirm the higher coefficient of friction, as shown in Fig. 18(c). The presence of groves was attributed to the removal of the WC particle that was embedded in it. Abrasive wear is demonstrated by the micro scratches. In LP_FeCrCoNiW0.3 + 5 at.% C APS coated samples, adhesive, abrasive, and oxide wear mechanisms predominate, as shown in Fig. 18(d). Only a few grooves formed in the LP_FeCrCoNiW0.3 + 5 at.% C alloy after laser surface processing, which may have been produced by a hard particle (Carbides) pulled out during the wear test. The exfoliation of hard particles was also observed in the wear track, which can be attributed to abrasive wear. As a result of its lower surface roughness, higher hardness, superior wear resistance, and lower coefficient of friction, laser processing can be used to improve the surface and mechanical properties of coatings in aerospace and naval applications.

Fig. 18
figure 18

Worn surface morphology of (a) as sprayed AlCrCoFeNiTi, (b) LP_ AlCrCoFeNiTi, (c) as sprayed FeCrCoNiW0.3 + 5 at.% C alloy and (d) LP_ FeCrCoNiW0.3 + 5 at. % C APS coatings

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EDS analysis was conducted on worn surfaces at different locations highlighted in Fig. 18(a), (b), (c) and (d). The findings of the EDS analysis demonstrate the presence of metal oxides on the worn surfaces. The oxygen content in the adhesive layer of the as-sprayed alloy is 29.54%, with a high concentration of Al, Cr, and Ti and a low concentration of Co and Ni. Laser-processed alloys reveal the presence of metal oxide spots, as illustrated in Fig. 18(b) and (c), which indicate the development of oxide layers (d). The oxides spots of LP_ AlCrCoFeNiTi and LP_ FeCrCoNiW0.3 + 5 at.% C APS coatings, as shown in Fig. 18(b) and (d), respectively, have a higher oxygen concentration of 46.21 and 47.52%. The EDS analysis of the wear debris (Table 5) suggests the presence of all the metals as well as oxygen.

Table 5 EDS analysis result of the worn surface at various locations highlighted in Fig. 18
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Additionally, Raman spectroscopy was used to examine the worn surfaces of both APS coatings before and after laser processing to identify the oxidation states and the type of oxides developed. Figure 19 shows the Raman spectra of both APS coatings before and after laser processing. The Raman spectra of AlCrCoFeNiTi coating as shown in Fig. 19(a) display the four distinct peaks at 480.9, 553.04, 614.80 and 684.8 cm−1 can be allocated to the Cr2O3, Cr2O3, Fe2O4 and Fe2O4, respectively (Ref 27, Ref 62,63,64). The characteristics spectra of LP_ AlCrCoFeNiTi APS coating display five different peaks at 535.2, 565.4, 615.73, 651.43 and 674.43 can be associated with Cr2O3, Fe2O4, Fe2O4, Co3O4, and Fe2O3 correspondingly (Ref 62,63,64,65, Ref 27). The characteristics spectrum of LP_AlCrCoFeNiTi APS coating exhibit similar peaks as AlCrCoFeNiTi APS coating. However, AlCrCoFeNiTi alloy show the higher relative intensity of Fe2O4, and LP_AlCrCoFeNiTi APS coating shows the higher relative intensity of Cr2O3 in comparison to each other. Raman spectra of FeCrCoNiW0.3 + 5 at.% C APS coating exhibit the six distinct peaks at 303.8, 351.9, 528.8, 555.8, and 614.8 cm−1 can be assigned to Co3O4, Fe2O4, Cr2O3, Fe2O4, and Fe2O4, respectively (Ref 62,63,64) as shown in Fig. 19(b). The characteristics spectrum of LP_FeCrCoNiW0.3 + 5 at.% C APS coating exhibit similar peaks as FeCrCoNiW0.3 + 5 at.% C APS coating with higher relative intensity of Fe2O4 in compare to Cr2O3 and Co3O4. Therefore, Raman spectroscopy revealed the presence of Cr2O3, Fe2O4, TiO2, and Co3O4 kinds of oxides, which is consistent with EDS findings. Higher oxygen concentration in the laser-processed alloy improves the wear resistance and reduces the tribo pair adhesion due to self-lubricating properties (Ref 32).

Fig. 19
figure 19

Raman spectra of AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings before and after laser treatments on the worn surfaces

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Summary and Conclusions

AlCrCoFeNiTi and FeCrCoNiW0.3 + 5 at.% C APS coatings were treated utilizing the laser surface processing technique. The effect of laser surface processing on both materials was thoroughly investigated. The research came to the following outcome.

  • Laser processing attributed to decrease in the surface roughness (Sa) of LP_AlCrCoFeNiTi and LP_FeCrCoNiW0.3 + 5 at.% C alloy coatings by 29.03% and 29.59%, respectively.

  • After laser processing, more BCC phases evolved with the depletion of minor FCC phase in LP_AlCrCoFeNiTi HEA and more CrC and WC hard phases evolved in LP_FeCrCoNiW0.3 + 5 at.% C HEA coating.

  • The dilution percentage were calculated as 9.43% and 8.41% for LP_ AlCrCoFeNiTi and LP_FeCrCoNiW0.3 + 5 at.% C HEAs. They exhibit good metallurgical bonding between coating and substrate.

  • The average microhardness was improved after laser processing by 5.03% and 5.67% in LP_AlCrCoFeNiTi and LP_ FeCrCoNiW0.3 + 5 at.% C HEAs, respectively.

  • Nanoindentation results suggest that both the HEAs i.e. AlCrCoFeNiTi and FeCoCrNiW0.3 + 5 at.% C, have an increase in the average reduced elastic modulus (Er) of 10.39 and 14.71%, respectively. In addition, the average hardness also increased by 17.18 and 19.80%, respectively post laser processing. The Al-rich phase of AlCrCoFeNiTi exhibit a higher hardness value, and the W-rich phase of FeCoCrNiW0.3 + 5 at. % C exhibits a higher hardness

  • The wear volume is significantly reduced by 37.32 and 39.71% in LP_ AlCrCoFeNiTi and LP_FeCrCoNiW0.3 + 5 at.% C alloy post laser processing, the wear mechanism involved is a combination of adhesive and abrasive in nature with oxide formation. The results suggest that the alloy can be potentially used for wear resistance applications.

However, the high temperature phase stability of laser treated alloy is still not well understood. Studies in this area would help in further adapting the aforementioned alloys for the multitude of high temperature applications i.e. aerospace, metallurgy, automobile, etc.