Microstructural and Wear Properties of Annealed Medium Carbon Steel Plate (EN8) Cladded with Martensitic Stainless Steel (AISI410)

Limited work on the wear properties of martensitic stainless-steel weld clads initiated this work which included investigations on microstructural and wear properties of cladded AISI 410 (filler wire)/EN 8 plates (substrate). Three layers of martensitic stainless steel (AISI 410) were deposited using metal inert gas (MIG) welding on medium carbon steel (EN 8) achieving a 51.5 ± 2.35 HRC of top layer. The elemental and phase fractions of the cladded layers indicated 98% martensite phase and retained austenite (2%). About 40% dilution was observed between EN 8 and the first weld layer. The results of tests carried out on pin on disc tribometer revealed an enhancement of anti-wear life of the martensitic weld cladded EN 8 by three times that of uncladded EN 8. The uncladded EN 8 plate suffered severe damage and high wear, leading to its failure at 478 s. The failure of the uncladded EN 8 sample was identified by the occurrence of high vibration of the pin on disc tribometer which ultimately stopped the tribometer. On the other hand, the cladded EN 8 sample continued running for 3600 s, exhibiting normal wear. After the tribo test, the surfaces of the pins of both cladded and uncladded EN 8 were analyzed using scanning electron microscope (SEM) and 3D profilometer. The surface characterization of tribo pairs indicated ploughing and galling to be the primary wear mechanisms. The average grain size of top and middle layer was in the range of 2–3.5 μm, while the base metal showed 5.02 μm mean grain size, resulting in higher hardness of clad layers than base metal, also favoring better wear resistance of the cladded EN 8 samples as compared to uncladded EN 8 samples.


Introduction
Cladding is a common surface coating technique in which a thick layer of material covers a relatively low-cost structural material to enhance its mechanical properties, corrosion resistance properties, and to improve the service life [1]. The deposition is typically of superior quality, enhancing the mechanical and microstructural properties of base metal. With a refined microstructural arrangement in the top surface, the cladding material usually has the properties to withstand high temperature, impact loads, and high wear resistance.
with martensitic stainless steel, conducting to an enhancement of mechanical properties [25]. In order to prevent the impact of dilution, a thickness of 15 mm with five layers of deposition is generally practiced [19]. It is to be noted that wear of the top cladding surface not only depends on hardness of the top layer but also on its microstructure [26]. For applications such anti-ballistics, the strength of the weld deposits or weld joints is an important parameter [27] and thus, microstructure plays an important role. Rao et al [28] reported that the long-term exposure of the interface to high temperature will lead to grain coarsening and a brittle interface was found due to cast dendritic structure in the clad [28]. Thus, the interlayer claddings are used to reduce the dilution effect and to enhance wear resistance, hardness etc. However, chances of metallic carbides formation also increases in interlayer claddings [29]. Modern techniques such as explosive cladding and microwave cladding are being used in several manufacturing units. In explosive weld zones, fatigue life varies with temperature of exposure and time of exposure, resulting in grain growth and recrystallization in the sample as well as decarburization occurring at the interface zone [30]. Microwave cladding is one of the most economical and proficient ways of cladding where uniform surface texture and properties can be obtained with uniform heat distribution throughout the clad, with good tribological properties [31].
From the above discussion, it can be observed that many filler materials have been used for cladding process, from austenite stainless steel to martensitic stainless steel. Nonetheless, very limited work on the minimal layer thickness, wear characteristics and microstructural behavior of martensitic welded clads have been reported until now. In this paper, the microstructure, mechanical and tribological properties of AISI 410 (martensite) cladded on EN 8 surface were analyzed. As less brittleness compared to solid-state welding procedures can be achieved by MIG [4], this technique was employed for cladding. Results of pin on disk tribological tests, 3D profilometer, scanning electron microscope (SEM), X-ray diffraction (XRD), and energy dispersive spectroscopy (EDS) analysis are presented, underlining the efficiency of the proposed cladding process.

Sample Preparation and Specification
Annealed EN 8 (rectangular flat bars) and AISI 410 filler material (spool form) were procured from local market in Chennai, India. The cladding was carried out on the EN 8 base metal parallelepipedal sample of 100 mm × 70 mm × 20 mm using MIG welding. Argon gas was used during cladding. The cladding was done on surface area of 61 mm × 32 mm. Before the cladding process, the impurities and unevenness on top surface the plate was thoroughly removed by sand papers and pencil grinder. The composition (weight percentage, wt.%) of the base metal material (uncladded EN 8) and the AISI 410 filler material were analyzed using an optical emission spectroscope (chemical) (OES Foundry Master, WAS, Germany) as per ASTM E 415, being presented in Tables 1 and 2, respectively. The filler material was chosen in order to achieve high wear resistance and hardness of the cladded surface.

Material Deposition
The cladding parameters were selected such that a standard medium heat input was achieved. AISI 410 wire (Φ1.2 mm), was used as the filler material. In order to obtain the desired medium heat input (0.55-0.60 kJ/mm), the cladding was done with an arc current of 110-115 A, arc voltage of 25 V, and welding speed of 4 mm/s. The heat input (H.I) of 0.55-0.60 kJ/mm was chosen in order to get bead height of above 3 mm with minimum penetration. During the process, it was observed that if the bead or deposition height was below 3 mm then it was difficult for grinding and making surface flat [32]. The Ador CHAMP MIG400 (Ador: Chennai, India) was used to conduct the cladding process. After the first weld deposition, the surface of the weld deposit was hand grinded to get a flat upper surface for the next deposition. The first deposition thickness was maintained at 3 mm, after hand grinding. Repeated welding and grinding were carried out for two more times and the thickness achieved for those two layers were 3 and 4 mm, respectively, making the total weld deposition thickness of 10 mm (Figure 1a) on the substrate of 20 mm thickness. A typical weld cladding done in this work have been shown in Figure 1b. In order to minimize the dilution from the substrate to the top layer, the heat input and the deposition thickness were kept at the medium level and constant, respectively. and welding speed of 4 mm/s. The heat input (H.I) of 0.55-0.60 kJ/mm was chosen in order to get bead height of above 3 mm with minimum penetration. During the process, it was observed that if the bead or deposition height was below 3 mm then it was difficult for grinding and making surface flat [33]. The Ador CHAMP MIG400 (Ador: Chennai, India) was used to conduct the cladding process. After the first weld deposition, the surface of the weld deposit was hand grinded to get a flat upper surface for the next deposition. The first deposition thickness was maintained at 3 mm, after hand grinding. Repeated welding and grinding were carried out for two more times and the thickness achieved for those two layers were 3 and 4 mm, respectively, making the total weld deposition thickness of 10 mm (Figure 1a) on the substrate of 20 mm thickness. A typical weld cladding done in this work have been shown in Figure 1b. In order to minimize the dilution from the substrate to the top layer, the heat input and the deposition thickness were kept at the medium level and constant, respectively.

Metallographic Analysis
A small specimen from the cladded plate was extracted for the metallographic analysis of three weld layers and dilution. The extracted specimen were grinded with different grades of emery paper (such as grit 120, 180, 240, 300, 400, 600, 800, 1000, and 1200) followed by cloth polishing with

Metallographic Analysis
A small specimen from the cladded plate was extracted for the metallographic analysis of three weld layers and dilution. The extracted specimen were grinded with different grades of emery paper (such as grit 120, 180, 240, 300, 400, 600, 800, 1000, and 1200) followed by cloth polishing with diamond paste of 0.15 µm. Then the specimen was etched with Kalling's reagent for nearly 10 s and immediately rinsed with acetone to identify the weld layers. The light microscope (Olympus: Tokyo, Japan) and the scanning electron microscope (SEM) (FEI Quanta FEG 200: Eindhoven, Netherlands) were used to identify the layers and dilution zones. The energy dispersive spectroscopy (EDS) was carried out on the specimen to understand the depth of carbon diffusion from the substrate. The grain size was measured across the top surface of the cladded plate to the substrate using line intercept method following a scale 1:1 [33]. In to use this method the micrograph was imported in AutoCAD 2017 software (N 52.0.0, Chennai, India). A random line was drawn using the drafting tools in AutoCAD software after scaling to 1:1. After the one was drawn, the number of grain boundaries intersecting the lines were counted. The average grain size was found by dividing the number of intersections with the actual length of the line.
A separate specimen of 15 mm × 4 mm × 2 mm was extracted for the X-ray diffraction (XRD) (Bruker, D8 Advance, Billerica, Massachusetts, USA) to identify the major phase fractions at weld deposits. The phase fraction was determined by considering d values from the peak intensity and corresponding 2θ values. The data base of the JCPDS (Joint Committee on Powder Diffraction Standards) was used to confirm the phases. The working principle of the XRD measurement is mainly based on the famous Brag's law, i.e., nλ = 2dsinθ (n = order of fringe, λ = wavelength of ray, d = lattice distance, θ = glancing angle.

Hardness Test
The hardness test of all the cladded layers was performed using Rockwell hardness testing machine (Fuel Instrument & Engineers Pvt. Ltd. YADRA, India). A conical base diamond indenter with tip including angle of 120 • was used and 100 Kgf static load was applied for 10 s on each layer during the test. The test was repeated four times on each layer and the average of all the readings have been reported in this work.

Wear Test
The anti-wear properties of the cladded pins were investigated using a pin on disc tribometer (DUCOM, TR100, Bengaluru, India). Cylindrical pins of 10 mm diameter pins and height 30 mm were made from the cladded plate as shown in Figure 2. The counter surface to these cladded pins was an EN 31 disc with hardness 58-60 HRC which was rotated at 500 rpm. A normal load of 100 N, (contact pressure 1.273 N/mm 2 ) was applied on the pin. The tests were carried for 3600 s. The coefficient of friction and wear rate reported in the work are the average of three readings and were automatically calculated by the software of the tribometer (V 1.0.0, WINDUCOM, Bengaluru, India). In order to understand the severity of the wear, the surface of the pins were analyzed using a scanning electron microscope (SEM) and a 3D surface profilometer (Taylor Hobson, Leicester, UK).
Metals 2020, 10, x FOR PEER REVIEW 5 of 16 diamond paste of 0.15 µm. Then the specimen was etched with Kalling's reagent for nearly 10 s and immediately rinsed with acetone to identify the weld layers. The light microscope (Olympus: Tokyo, Japan) and the scanning electron microscope (SEM) (FEI Quanta FEG 200: Eindhoven, Netherlands) were used to identify the layers and dilution zones. The energy dispersive spectroscopy (EDS) was carried out on the specimen to understand the depth of carbon diffusion from the substrate. The grain size was measured across the top surface of the cladded plate to the substrate using line intercept method following a scale 1:1 [34]. In to use this method the micrograph was imported in AutoCAD 2017 software (N 52.0.0, Chennai, India). A random line was drawn using the drafting tools in AutoCAD software after scaling to 1:1. After the one was drawn, the number of grain boundaries intersecting the lines were counted. The average grain size was found by dividing the number of intersections with the actual length of the line. A separate specimen of 15 mm x 4 mm x 2 mm was extracted for the X-ray diffraction (XRD) (Bruker, D8 Advance, Billerica, Massachusetts, USA) to identify the major phase fractions at weld deposits. The phase fraction was determined by considering d values from the peak intensity and corresponding 2θ values. The data base of the JCPDS (Joint Committee on Powder Diffraction Standards) was used to confirm the phases. The working principle of the XRD measurement is mainly based on the famous Brag's law, i.e., nλ = 2dsin (n = order of fringe, λ = wavelength of ray, d = lattice distance, = glancing angle.

Hardness Test
The hardness test of all the cladded layers was performed using Rockwell hardness testing machine (Fuel Instrument & Engineers Pvt. Ltd. YADRA, India). A conical base diamond indenter with tip including angle of 120° was used and 100 Kgf static load was applied for 10 s on each layer during the test. The test was repeated four times on each layer and the average of all the readings have been reported in this work.

Wear Test
The anti-wear properties of the cladded pins were investigated using a pin on disc tribometer (DUCOM, TR100, Bengaluru, India). Cylindrical pins of 10 mm diameter pins and height 30 mm were made from the cladded plate as shown in Figure 2. The counter surface to these cladded pins was an EN 31 disc with hardness 58-60 HRC which was rotated at 500 rpm. A normal load of 100 N, (contact pressure 1.273 N/mm 2 ) was applied on the pin. The tests were carried for 3600 s. The coefficient of friction and wear rate reported in the work are the average of three readings and were automatically calculated by the software of the tribometer (V 1.0.0, WINDUCOM, Bengaluru, India). In order to understand the severity of the wear, the surface of the pins were analyzed using a scanning electron microscope (SEM) and a 3D surface profilometer (Taylor Hobson, Leicester, UK).   Figure 3a shows the microstructure of the base metal, medium carbon steel (EN 8), where discrete pearlitic structures (black spots) are visible on the ferritic matrix. The base metal microstructure is primarily (~90%) dominated by α phase at annealed condition. Figure 3b shows the top layer which was partially reheated during the deposition of adjacent passes in the same layer. The grain size was measured using line intercept method after proper scaling (1:1) of the microstructural images [33]. The carbon percentage of AISI 410 (top layer) is suitable for the formation of lath martensite [34], hence the microstructure showed the presence of the lath martensite transformed from austenite during cooling. As welding passes were drawn consecutively at the top layer, this resulted in slow cooling, which allowed more time for grain growth and resulted in~4 µm of coarser grain ( Table 3). The middle layer ( Figure 3c) was a reheated zone having semi-equiaxed grains of~2 µm. The middle layer grain sizes were approximately two times smaller than the top layer, as shown in Table 3. The middle layer was subjected to continuous thermal stress from both the top and bottom layers resulting in heterogeneous stress distribution. Additionally, the cooling in the middle layer was comparatively faster compared to the top layer because the heat coming from the top layer was easily escaped to the colder bottom layer and prevented coarsening of the grain. Therefore, the prior austenitic morphology was changed from columnar to equiaxed due to reheating of the deposited layer.  Figure 3a shows the microstructure of the base metal, medium carbon steel (EN 8), where discrete pearlitic structures (black spots) are visible on the ferritic matrix. The base metal microstructure is primarily (~90%) dominated by α phase at annealed condition. Figure 3b shows the top layer which was partially reheated during the deposition of adjacent passes in the same layer. The grain size was measured using line intercept method after proper scaling (1:1) of the microstructural images [34]. The carbon percentage of AISI 410 (top layer) is suitable for the formation of lath martensite [35], hence the microstructure showed the presence of the lath martensite transformed from austenite during cooling. As welding passes were drawn consecutively at the top layer, this resulted in slow cooling, which allowed more time for grain growth and resulted in ~4 µm of coarser grain ( Table 3). The middle layer (Figure 3c) was a reheated zone having semi-equiaxed grains of ~2µm. The middle layer grain sizes were approximately two times smaller than the top layer, as shown in Table 3. The middle layer was subjected to continuous thermal stress from both the top and bottom layers resulting in heterogeneous stress distribution. Additionally, the cooling in the middle layer was comparatively faster compared to the top layer because the heat coming from the top layer was easily escaped to the colder bottom layer and prevented coarsening of the grain. Therefore, the prior austenitic morphology was changed from columnar to equiaxed due to reheating of the deposited layer.   The bottom layer microstructure, presented in Figure 3d and e, showed normally larger grain structure, martensite, and traces of austenite. The presence of larger grains was due to the diffusion of carbon from base plate, which enhanced the stability of austenite at higher temperature and restricted the grain refinement on cooling. Furthermore, the presence of high carbon with 12-13% of chromium was likely to cause precipitation of metal (M) carbides. The M23C6 type of carbide is the most common at the temperature between 950-1050 °C [36][37][38], as also observed from the micrograph in Figure 3d, and it was seen that the bottom layer had higher fraction of carbon that led to the formation of carbides. This formation of carbides, in spite of having larger grains than top layer and the bottom layer, resulted in higher average hardness. The rapid dissipation of heat through the base metal did not allow sufficient time of alloy segregation and solid-solution, which might be a probable cause of carbide precipitation. XRD analysis of the top layer, shown in Figure 3f, revealed the presence of martensite as the major phase fraction. Martensite has body centric tetragonal structure and is a super saturated form of carbon with high bond energy [39].

Metallographic Analysis and Hardness
The SEM analysis of the cladded and uncladded EN 8 (base metal) samples also showed the presence of martensite lath structures in cladded layers and ferritic-pearlitic structure in base plate, as shown in Figure 4a-d. The EDS element mapping of bottom layer around fusion boundary depicted strong evidence of carbon diffusion across the boundary in the fusion zone ( Figure 5). The carbon diffusion at the bottom layer was reached up to ~30-40%, however, this diffusion of carbon was exceptionally reduced at the middle and top layers due to the pre-deposition grinding of several millimeters. It was observed that the top layer had minimum carbon diffusion of ~3-5% only. The bottom layer microstructure, presented in Figure 3d and e, showed normally larger grain structure, martensite, and traces of austenite. The presence of larger grains was due to the diffusion of carbon from base plate, which enhanced the stability of austenite at higher temperature and restricted the grain refinement on cooling. Furthermore, the presence of high carbon with 12-13% of chromium was likely to cause precipitation of metal (M) carbides. The M 23 C 6 type of carbide is the most common at the temperature between 950-1050 • C [35][36][37], as also observed from the micrograph in Figure 3d, and it was seen that the bottom layer had higher fraction of carbon that led to the formation of carbides. This formation of carbides, in spite of having larger grains than top layer and the bottom layer, resulted in higher average hardness. The rapid dissipation of heat through the base metal did not allow sufficient time of alloy segregation and solid-solution, which might be a probable cause of carbide precipitation. XRD analysis of the top layer, shown in Figure 3f, revealed the presence of martensite as the major phase fraction. Martensite has body centric tetragonal structure and is a super saturated form of carbon with high bond energy [38].
The SEM analysis of the cladded and uncladded EN 8 (base metal) samples also showed the presence of martensite lath structures in cladded layers and ferritic-pearlitic structure in base plate, as shown in Figure 4a-d. The EDS element mapping of bottom layer around fusion boundary depicted strong evidence of carbon diffusion across the boundary in the fusion zone ( Figure 5). The carbon diffusion at the bottom layer was reached up to~30-40%, however, this diffusion of carbon was exceptionally reduced at the middle and top layers due to the pre-deposition grinding of several millimeters. It was observed that the top layer had minimum carbon diffusion of~3-5% only.  The bottom layer microstructure, presented in Figure 3d and e, showed normally larger grain structure, martensite, and traces of austenite. The presence of larger grains was due to the diffusion of carbon from base plate, which enhanced the stability of austenite at higher temperature and restricted the grain refinement on cooling. Furthermore, the presence of high carbon with 12-13% of chromium was likely to cause precipitation of metal (M) carbides. The M23C6 type of carbide is the most common at the temperature between 950-1050 °C [36][37][38], as also observed from the micrograph in Figure 3d, and it was seen that the bottom layer had higher fraction of carbon that led to the formation of carbides. This formation of carbides, in spite of having larger grains than top layer and the bottom layer, resulted in higher average hardness. The rapid dissipation of heat through the base metal did not allow sufficient time of alloy segregation and solid-solution, which might be a probable cause of carbide precipitation. XRD analysis of the top layer, shown in Figure 3f, revealed the presence of martensite as the major phase fraction. Martensite has body centric tetragonal structure and is a super saturated form of carbon with high bond energy [39].
The SEM analysis of the cladded and uncladded EN 8 (base metal) samples also showed the presence of martensite lath structures in cladded layers and ferritic-pearlitic structure in base plate, as shown in Figure 4a-d. The EDS element mapping of bottom layer around fusion boundary depicted strong evidence of carbon diffusion across the boundary in the fusion zone ( Figure 5). The carbon diffusion at the bottom layer was reached up to ~30-40%, however, this diffusion of carbon was exceptionally reduced at the middle and top layers due to the pre-deposition grinding of several millimeters. It was observed that the top layer had minimum carbon diffusion of ~3-5% only.  The bottom layer of the cladding has large carbon diffusion of up to 30% in places, as calculated from the SEM-EDS analysis is shown in Figure 5. The micrographs of bottom layer (Figure 3d) also depict the formation of micro-cracks oriented towards the middle layer due to the high thermal stress generated from the consecutive passes and partial entrapment of hydrogen [36]. The fast heat dissipation from the bottom layer is probably the cause of hydrogen entrapment within the molten pool, enhancing the cracking susceptibility of the cladded layer [36]. The bottom layer of the cladding has large carbon diffusion of up to 30% in places, as calculated from the SEM-EDS analysis is shown in Figure 5. The micrographs of bottom layer (Figure 3d) also depict the formation of micro-cracks oriented towards the middle layer due to the high thermal stress generated from the consecutive passes and partial entrapment of hydrogen [35]. The fast heat dissipation from the bottom layer is probably the cause of hydrogen entrapment within the molten pool, enhancing the cracking susceptibility of the cladded layer [35]. The bottom layer of the cladding has large carbon diffusion of up to 30% in places, as calculated from the SEM-EDS analysis is shown in Figure 5. The micrographs of bottom layer (Figure 3d) also depict the formation of micro-cracks oriented towards the middle layer due to the high thermal stress generated from the consecutive passes and partial entrapment of hydrogen [36]. The fast heat dissipation from the bottom layer is probably the cause of hydrogen entrapment within the molten pool, enhancing the cracking susceptibility of the cladded layer [36].  Table 4 exhibits the hardness of each layer of the clad. Compared with base metal (uncladded EN 8) sample (22)(23), the clads exhibited higher hardness. Among all the layers, the middle layer displayed the highest hardness. Compared to the base metal, the higher hardness of top and bottom layer is probably due to the presence of martensitic phase with small amount of retained austenite. In addition to the appearance of these phases, the hardness of the bottom layer often depends on the presence of precipitated carbides resulted from base metal dilution. The hardness of the cladded layers is also dependent on the average grain size and can be estimated by the Hall-Petch equation: where, Hv = average hardness value, H0 = material constant, Ky = coefficient, and d =average grain size (μm). From Table 3, it was seen that the middle layer had the least average grain size and thus resulted in the highest hardness (Table 4). Similar trend of hardness results were also reported earlier [40].    (22)(23), the clads exhibited higher hardness. Among all the layers, the middle layer displayed the highest hardness. Compared to the base metal, the higher hardness of top and bottom layer is probably due to the presence of martensitic phase with small amount of retained austenite. In addition to the appearance of these phases, the hardness of the bottom layer often depends on the presence of precipitated carbides resulted from base metal dilution. The hardness of the cladded layers is also dependent on the average grain size and can be estimated by the Hall-Petch equation: where, H v = average hardness value, H o = material constant, K y = coefficient, and d = average grain size (µm). From Table 3, it was seen that the middle layer had the least average grain size and thus resulted in the highest hardness (Table 4). Similar trend of hardness results were also reported earlier [39]. Figure 6 shows that the grain size is inversely proportional to hardness and thus, the layer hardness increased with decreasing grain size. Relatively larger grain size and traces of retained austenite reduced the extreme brittleness and induced slight ductility in the top and bottom layers [40,41]. Nevertheless, the top layer maintained good hardness due to the presence of martensite. The middle layer (Figure 3c) showed highest hardness due to its refined martensite grain structure. In spite of having larger grains in bottom layer due to carbon diffusion from base metal, which in turn resulted in carbide formation, the bottom layer exhibited higher average hardness value than the top layer.
Base metal 23 ± 1.2 Figure 6 shows that the grain size is inversely proportional to hardness and thus, the layer hardness increased with decreasing grain size. Relatively larger grain size and traces of retained austenite reduced the extreme brittleness and induced slight ductility in the top and bottom layers [41,42]. Nevertheless, the top layer maintained good hardness due to the presence of martensite. The middle layer (Figure 3c) showed highest hardness due to its refined martensite grain structure. In spite of having larger grains in bottom layer due to carbon diffusion from base metal, which in turn resulted in carbide formation, the bottom layer exhibited higher average hardness value than the top layer.

Investigating the Wear Properties of the Clad
The wear test was conducted using a pin on disc tribometer (DUCOM, Bengaluru, India). During the tribo test, it has been observed that EN 8 pin (base metal) without any clad failed within 475-500 s whereas the cladded EN 8 pin ran up to 3600 s, as shown by the plot in Figure 7a. The failure of the pin was noticed with high wear rate resulting in high vibration of the pin holder and ultimately tipping the tribometer. The wear rate of uncladded EN 8 was found to be 0.0003 mm 3 /Nm within a short period of time (478 s) whereas, the cladded EN 8 was found to be 0.00011 mm 3 /Nm in 3600 s. The cladded pin sample ran smoothly for 3600 s without any kind of vibration and tipping of the tribometer, indicating the occurrence of normal wear. Figure 7b exhibits the better wear resistance of the cladded pins as compared to the base metal (EN 8 uncladded). Furthermore, the cladded EN 8 pin lasted about nine times longer than the uncladded EN 8 pin. Even after 3600 s of run time the cladded pin did not completely failed, whereas the scar depth of base metal was found to be much higher for 478 s.

Investigating the Wear Properties of the Clad
The wear test was conducted using a pin on disc tribometer (DUCOM, Bengaluru, India). During the tribo test, it has been observed that EN 8 pin (base metal) without any clad failed within 475-500 s whereas the cladded EN 8 pin ran up to 3600 s, as shown by the plot in Figure 7a. The failure of the pin was noticed with high wear rate resulting in high vibration of the pin holder and ultimately tipping the tribometer. The wear rate of uncladded EN 8 was found to be 0.0003 mm 3 /Nm within a short period of time (478 s) whereas, the cladded EN 8 was found to be 0.00011 mm 3 /Nm in 3600 s. The cladded pin sample ran smoothly for 3600 s without any kind of vibration and tipping of the tribometer, indicating the occurrence of normal wear. Figure 7b exhibits the better wear resistance of the cladded pins as compared to the base metal (EN 8 uncladded). Furthermore, the cladded EN 8 pin lasted about nine times longer than the uncladded EN 8 pin. Even after 3600 s of run time the cladded pin did not completely failed, whereas the scar depth of base metal was found to be much higher for 478 s. The Scanning electron microscope images presented in Figure 7c and d shows grooves on the surfaces of the uncladded and cladded pins. It is to be noted that the surface shown in Figures 7c is not the actual surface of the pin which was in contact with EN 31 disc. This surface is the left-over surface of the pin. The base metal (uncladded EN 8) pin broke after 478 s and the left-over surface seems to have undergone a polishing effect. The failure was so catastrophic that portion of the pin which was out of the pin holder wore rapidly and the left-over surface inside the holder was taken for analyzing the wear. On other hand, the cladded EN 8 ran for 3600 s indicating high durability of the clads. As the tests were conducted under dry condition, the resultant wear mechanism was a combination of both adhesion and ploughing [43]. Grooves and galling were observed in both cladded and uncladded surfaces, indicating adhesion and ploughing wear mechanisms involved during the tribotest. The 3D surface characteristics of both uncladded and cladded EN 8 also indicated high wear in the uncladded leftover pin surface (Figures 8 and 9). The Scanning electron microscope images presented in Figure 7c and d shows grooves on the surfaces of the uncladded and cladded pins. It is to be noted that the surface shown in Figure 7c is not the actual surface of the pin which was in contact with EN 31 disc. This surface is the left-over surface of the pin. The base metal (uncladded EN 8) pin broke after 478 s and the left-over surface seems to have undergone a polishing effect. The failure was so catastrophic that portion of the pin which was out of the pin holder wore rapidly and the left-over surface inside the holder was taken for analyzing the wear. On other hand, the cladded EN 8 ran for 3600 s indicating high durability of the clads. As the tests were conducted under dry condition, the resultant wear mechanism was a combination of both adhesion and ploughing [42]. Grooves and galling were observed in both cladded and uncladded surfaces, indicating adhesion and ploughing wear mechanisms involved during the tribotest. The 3D surface characteristics of both uncladded and cladded EN 8 also indicated high wear in the uncladded leftover pin surface (Figures 8 and 9). Metals 2020, 10, x FOR PEER REVIEW 12 of 16  The average scar width and depth of uncladded and cladded surfaces were calculated from the results of texture analysis and represented in Table 5. The scar width of the cladded specimen is  The average scar width and depth of uncladded and cladded surfaces were calculated from the results of texture analysis and represented in Table 5. The scar width of the cladded specimen is The average scar width and depth of uncladded and cladded surfaces were calculated from the results of texture analysis and represented in Table 5. The scar width of the cladded specimen is significantly lower than the uncladded EN 8 even after 3600 s of sliding time. However, due to higher sliding time of the cladded specimen, the scar depth is~1 µm higher than the uncladded specimen. The volumetric loss of the cladded and uncladded specimens were calculated using the Archard equation as given below.
where, V = volumetric loss, K = constant, S = sliding distance, N = normal load, and H = hardness of pin.
where, k = specific wear rate (K/H). The volumetric loss for uncladded EN 8 (base metal) was found to be 48.80 mm 3 for S = 1.626 km and N = 100 N. Similarly, the volumetric loss for cladded EN 8 was found to be 134.774 mm 3 for S = 12.252 km at same applied load. For same sliding time of 478 s, the volumetric loss of the cladded specimen was significantly lower than the base metal (uncladded EN 8) specimen as shown in Table 6. The graph has been plotted between volumetric loss of cladded EN 8 and Speed-normal load(SN) at different time intervals, shown in Figure 10a. The graph was found to be a straight line, which is in line with other authors [43]. significantly lower than the uncladded EN 8 even after 3600 s of sliding time. However, due to higher sliding time of the cladded specimen, the scar depth is ~1 µm higher than the uncladded specimen. The volumetric loss of the cladded and uncladded specimens were calculated using the Archard equation as given below.
where, V = volumetric loss, K = constant, S = sliding distance, N = normal load, and H = hardness of pin. = where, k = specific wear rate (K/H).  The volumetric loss for uncladded EN 8 (base metal) was found to be 48.80 mm 3 for S = 1.626 km and N = 100 N. Similarly, the volumetric loss for cladded EN 8 was found to be 134.774 mm 3 for S = 12.252 km at same applied load. For same sliding time of 478 s, the volumetric loss of the cladded specimen was significantly lower than the base metal (uncladded EN 8) specimen as shown in Table  6. The graph has been plotted between volumetric loss of cladded EN 8 and Speed-normal load(SN) at different time intervals, shown in Figure 10a. The graph was found to be a straight line, which is The ratio of volumetric loss of base metal to that of clad was calculated by Archard equation (Equation (2)). The energy dissipation was also calculated, using Equation (4).
where, ∆E = energy dissipation, F = force, ∆t = time interval, v = velocity, and F avg. = mean force. The energy dissipation value for the cladded pin was 30.497 × 10 4 J after 3600 s, while the total energy dissipation for base metal was found to be 4.803 × 10 4 J after 478 s. Comparing the energy dissipation of base metal EN 8 (uncladded) with the energy dissipation value for cladded material after the same period of 478 s, it was found to be lesser for cladded material (Figure 10b), indicating that AISI 410 cladded material possessed superior wear resistance properties than uncladded EN 8 sample.

Conclusions
The microstructure, mechanical and tribological properties of AISI 410 (martensite) cladded on EN 8 surface were analyzed. Results of hardness, pin on disk tribological tests (friction coefficient and energy dissipation), 3D profilometry (wear rate and volumetric loss), scanning electron microscopy (SEM), X-ray diffraction (XRD), and EDS analysis underlined the efficiency of the proposed cladding process.
The results exhibited that cladded layers had maximum hardness of about 58 HRC in the middle layer, due to the presence of hard martensite phase, the middle clad layer being a reheated zone. Due to the presence of retained austenite, there was a slight decrease in the hardness of the top layer (49-53 HRC). The XRD results indicated the presence of martensitic phase too. In spite of having larger grain size, the average hardness of bottom layer (52-55 HRC) was slightly higher than the top layer due to the presence of metal carbides resulted from the carbon diffusion from base metal.
The average grain size of top and middle layer was between 2-3.5 µm, while the base metal showed 5.02 µm grain size, resulting in higher hardness of clad layers than base metal.
Some micro cracks were also observed in the bottom clad layers, generated by high thermal stresses, multiple passes, and partial entrapment of hydrogen. The bottom layer was the fusion zone of two different metals with different thermal coefficient of expansion, increasing the chances of crack generation. The EDS mapping also showed high diffusion of carbon (30-40%) in the bottom layer, which reduced to 3-4% in the top most layer as multiple passes and layers were deposited.
The volumetric wear loss from Archard's equation was found to be 48.80 mm 3 for uncladded EN 8 (base metal) and 17.90 mm 3 for AISI 410 cladded EN 8. The wear analysis through energy dissipation exhibited that the cladded EN 8 absorbed 30.497 × 10 4 J of energy and did not fail, whereas the uncladded EN 8 was able to absorb only 4.803 × 10 4 J of energy before failure.
The scanning electron microscope and 3D profilometer images exhibited ploughing and galling to be the main wear mechanisms, and the wear tests revealed better wear resistance of the cladded pins as compared to the uncladded pins. The wear rate after cladding reduced by almost three times as compared to the uncladded part.

Conflicts of Interest:
The authors declare no conflict of interest.