Microstructure Evolution and Failure Behavior of Stellite 6 Coating on Steel after Long-Time Service

The microstructure evolution, elements diffusion and fracture behavior of the Stellite 6 weld overlay, deposited on 10Cr9Mo1VNbN (F91) steel by the tungsten inert gas (TIG) cladding process, were investigated after long-time service. Obvious diffusion of Fe occurred from the steel and fusion zone to the Stellite overlay, resulting in the microstructure evolution and hardness increase in the coating, where hard Co–Fe phases, σ phases (Fe–Cr metallic compounds) and Cr-rich carbides (Cr18.93Fe4.07C6) were formed. Besides, the width of the light zone, combined with the fusion zone and diffusion zone, increased significantly to a maximum value of 2.5 mm. The fracture of the Stellite coating samples mainly occurred in the light zone, which was caused by the formation and growth of circumferential crack and radial crack under high temperature and pressure conditions. Moreover, the micro-hardness values in the light zone increased to the maximum (470–680 HV) due to the formation and growth of brittle Co–Fe phases. The formation of these cracks might be caused by formed brittle phases and changes of micro-hardness during service.


Introduction
10Cr9Mo1VNbN (F91) steel, as a martensitic, heat-resistant steel, has outstanding high-temperature performance and corrosion resistance, and is massively applied in manufacture of steam boiler, valve body, tube and turbine components [1]. The valve discs and seats mainly made of F91 steels are easily subjected to the severe erosion and wear of solid particles under high temperature and pressure in the service process. To enhance the service life of the valve discs and seats, protective coatings with adequate mechanical properties, wear and spalling resistance, are usually required. Stellite alloy (Co-based alloy), especially Stellite 6 alloy, is often used as a coating material in valve disc manufacturing, due to its excellent wear resistance, corrosion resistance and high temperature properties [2][3][4].
In previous work, different overlaying methods have been performed to investigate the properties of the wearing coating, including laser cladding, arc welding, plasma transferred arc (PTA) welding and gas dynamic cold spray [5][6][7][8][9][10][11][12][13][14]. Kusmokoet et al. [15] deposited the Stellite 6 alloy on P22 steel and P91steel plates by laser cladding and investigated the sliding wear characteristics of the coating. The worn surface of the coating on P91 steel was much rougher compared to that on P22 steel plate, and fewer strong carbide-forming elements resulted in the reduction of the amount of carbon loss in the coating on P22 steel. Cincaet et al. [16] investigated the properties of Stellite 6 deposition on low alloy carbon steel made by cold gas spraying and found that the increase of gas pressure and distance could led to the reduction of deposition efficiencies.
Ferozhkhanet et al. [17] investigated the microstructure, hardness and wear mechanism of Stellite 6alloy coating deposited on 9Cr-1Mo steel by plasma transferred arc welding process, and 309-16L was used as the interlayer between the coating and 9Cr-1Mo steel. It was found that the amount of alloying elements (Cr, W and Co) in the Stellite 6 coating was higher than that in the nominal composition, and the dilution of Fe in Stellite coating was below 2%. Besides, the wear mechanism of the coating was the combination of delamination and abrasive wear.
Tungsten inert gas (TIG) welding has the advantages of convenience operations, excellent arc stability and welding quality. Wear-resistant layer surfacing in the valve body is often performed by TIG cladding in the production of power generation assembly. Mirshekariet et al. [18] and Molleda et al. [19] both characterized the microstructure and phase features of Stellite 6 surfacing layers deposited on steel substrates. However, the Stellite coating deposited on the steel components by TIG cladding could fracture after a long period of service process in power plants, due to the effect of the high temperature and pressure environment.
The fracture and spalling of the Stellite layer is one of the key problems for the manufacturing of the power generation assembly, especially a valve body; however, there are no systematic researches to be reported about the fracture mechanism and alloying elements diffusion of Stellite alloy-deposited steel parts after long-time service in power plants. In this paper, the TIG cladding method is used to deposit the surfacing resistance coating (Stellite 6 alloy) onto the F91 steel, and the microstructure evolution, alloying the elements diffusion and failure behavior of Stellite coating samples after long-time service process, were investigated.

Materials and Methods
Stellite6 alloy coating was deposited onto a F91martensitic heat-resistant steel valve disc (240 mm × 107 mm) by a multi-pass tungsten inert gas (TIG) cladding process with AWS ERCoCr-A (Shanghai, China) consumable (4.0 mm). The TIG welding machine is theTSP-300 (Panasonic, Kadoma, Japan) equipped a positioner and an automatic wire feeder. The nominal composition of F91 steel and ERCoCr-A wire were shown in Table 1. In addition, the micro-hardness of the Stellite 6 alloy and F91 steel are about 380-450 HV and 200 HV, respectively. Before cladding welding, the F91 steel should undergo a preheating treatment (200 • C for 1 h) in the electric furnace. During the cladding process, the welding current (200-290 A), arc voltage (22-30 V), wire feed speed (50-70 mm/min) and travel speed (80-180 mm/min) were kept constant. The shielding gas is argon, with a flow rate of 15-20 L/min. There were three layers in the multi-pass surfacing, and the thickness of the surfacing is about 4 mm. The extension length and diameter of the tungsten electrode was around 8 and 4.0 mm, respectively. After cladding welding, the weldments were subject to post-welding heat treatment, which was performed at 400 • C for 1 h. Two components with Stellite 6/F91 TIG cladding structure were obtained. The component#1 is on the pre-service condition, while the component #2 has been used in the main stop valve of the steam turbine for about 26,280 h, operating at 566 • C and 9.94 MPa. The weldments before and after long-time service in a power plant were marked as #1 and #2, respectively. The main microstructures of F91 steel are tempered martensite (Figure 1a), and the Stellite 6 alloy consists of Co-rich dendrites matrix and eutectic carbides within grain boundaries (GBs) (Figure 1b). The samples were cut from weldment #1 and #2 for microstructure analysis, and then they were etched for 20 s with FeCl 3 solution after sectioning, grinding and polishing. The microstructure features were studied using an optical microscope (OM, Leica, DM2700M, Wetzlar, Germany). In order to observe the microstructure evolution, the samples were examined by scanning electron microscopy (SEM, VEGA3 SBH, TESCAN, Brno, Czech Republic), and the micro-hardness profile across the sample was measured. Moreover, the composition in different regions of samples was measured with energy dispersive spectrum (EDS, VEGA3 SBH, TESCAN), and the phase identification of the coating layer was performed by X-ray diffraction (XRD, D/MAX-1200, Rigaku Industrial Corporation, Osaka, Japan). In addition, the microstructure of Stellite-deposited steel samples was investigated by transmission electron microscopy (TEM, Tecnai G2 20 S-TWIN, FEI, Hillsboro, OR, USA), and the TEM samples were prepared by focused ion beam (FIB, AURIGA, Zeiss, Oberkochen, Germany). Figure 1c,d illustrates the schematic of the Stellite 6 alloy weld overlay on F91 steel by the TIG cladding process and the samples for microstructural and micro-hardness analyses. The main microstructures of F91 steel are tempered martensite (Figure 1a), and the Stellite 6 alloy consists of Co-rich dendrites matrix and eutectic carbides within grain boundaries (GBs) (Figure 1b). The samples were cut from weldment #1 and #2 for microstructure analysis, and then they were etched for 20 s with FeCl3 solution after sectioning, grinding and polishing. The microstructure features were studied using an optical microscope (OM, Leica, DM2700M, Wetzlar, Germany). In order to observe the microstructure evolution, the samples were examined by scanning electron microscopy (SEM, VEGA3 SBH, TESCAN, Brno, Czech Republic), and the micro-hardness profile across the sample was measured. Moreover, the composition in different regions of samples was measured with energy dispersive spectrum (EDS, VEGA3 SBH, TESCAN), and the phase identification of the coating layer was performed by X-ray diffraction (XRD, D/MAX-1200, Rigaku Industrial Corporation, Osaka, Japan). In addition, the microstructure of Stellite-deposited steel samples was investigated by transmission electron microscopy (TEM, Tecnai G2 20 S-TWIN, FEI, Hillsboro, OR, USA), and the TEM samples were prepared by focused ion beam (FIB, AURIGA, Zeiss, Oberkochen, Germany). Figure 1c,d illustrates the schematic of the Stellite 6 alloy weld overlay on F91 steel by the TIG cladding process and the samples for microstructural and micro-hardness analyses.

Microstructure
The typical microstructures at cross sections of sample #1 are shown in Figure 2. The boundary between the base material and coating was obvious, and a wide fusion zone was observed in Figure  2a. Moreover, the microstructure of the Stellite weld overlay zone (WOZ) was a typical hypoeutectic dendritic in Figure 2b, which consisted of Co solid solution and a network of small carbides particles. The Co solid solution consisted of fcc γ and hcp ε phases, and the carbides particles were mainly M23C6, Cr7C3 and Co3W intermetallic phases, where M was (Cr, Co, W, Ni, Fe) [2,20].During the TIG cladding process, thermal cycle could result in the formation of microstructures with different features in the heat-affected zone (HAZ), which was divided into a coarse grain heat-affected zone (CG-HAZ), fine grain heat-affected zone (FG-HAZ) and partial normalized zone (PNZ), as shown in Figure 2c-e. Obviously, grain coarsening occurred in the CG-HAZ, where the austenitizing induced by the high peak temperature happened during TIG cladding, following a significant growth of austenite grains. Finally, coarse original austenite grains remained after cooling. The fine austenite grains formed in FG-HAZ due to relatively low heat input compared with CG-HAZ, in which the pinning effect of undissolved precipitated phases would restrain the grain boundary migration. Furthermore, a part of the microstructure in PNZ

Microstructure
The typical microstructures at cross sections of sample #1 are shown in Figure 2. The boundary between the base material and coating was obvious, and a wide fusion zone was observed in Figure 2a. Moreover, the microstructure of the Stellite weld overlay zone (WOZ) was a typical hypoeutectic dendritic in Figure 2b, which consisted of Co solid solution and a network of small carbides particles. The Co solid solution consisted of fcc γ and hcp ε phases, and the carbides particles were mainly M 23 C 6 , Cr 7 C 3 and Co 3 W intermetallic phases, where M was (Cr, Co, W, Ni, Fe) [2,20].During the TIG cladding process, thermal cycle could result in the formation of microstructures with different features in the heat-affected zone (HAZ), which was divided into a coarse grain heat-affected zone (CG-HAZ), fine grain heat-affected zone (FG-HAZ) and partial normalized zone (PNZ), as shown in Figure 2c-e. Obviously, grain coarsening occurred in the CG-HAZ, where the austenitizing induced by the high peak temperature happened during TIG cladding, following a significant growth of austenite grains. Finally, coarse original austenite grains remained after cooling. The fine austenite grains formed in FG-HAZ due to relatively low heat input compared with CG-HAZ, in which the pinning effect of undissolved precipitated phases would restrain the grain boundary migration. Furthermore, a part of the microstructure in PNZ underwent an austenitizing transformation, while the other part remained of martensite structure, which eventually generated a mixture structure which consisted of un-tempered martensite and over-tempered martensite.
Coatings 2019, 9, x FOR PEER REVIEW 4 of 11 underwent an austenitizing transformation, while the other part remained of martensite structure, which eventually generated a mixture structure which consisted of un-tempered martensite and over-tempered martensite. The fusion zone had broadened in the Stellite 6 coating layer near F91 steel due to the diffusion of elements from the steel to the Stellite 6 coating after the service process. The elements diffusion eventually made a wider light zone form, consisting of the original fusion zone and new mutual diffusion zone, as shown in Figure 3. The light zone still showed a dendritic microstructure as same as that in our Stellite weld layer of sample #2. The width of the fusion zone was measured as 31, 28 and 49 μm at different locations of sample #1 in Figure 3a, indicating the width of the fusion zone was non-uniform. After long-time service, the width of the fusion zone increased significantly to form a large diffusion zone as marked in Figure 3b (sample #2), which could even reach to around 2.5 mm, indicating that a high temperature environment could promote the extent of the light zone through influencing the diffusion of alloying elements and the formation and growth of metallic compound phases and carbides. Figure 3c,d shows the XRD results of phases in the light zone of samples #1 and #2. It can be seen from the XRD results in Figure 3c that the phases in the light zone mainly consist of Co-Fe substitution solid solutions and σ phases (Fe-Cr metallic compounds), and the two elements, Co and Fe, have good solid solubility, thus they could dissolve with each other by almost any proportion. The carbides were not found from the XRD results due to the smaller quantity and small size. The type of phases in the light zone of sample #2 was similar to that of sample #1; however, there existed a more evident diffraction peak of carbides in Figure 3d, which were Cr18.93Fe4.07C6 carbides. It meant that the formation and growth of the precipitated phases happened, and the Co matrix had evolved to form Co-Fe phases during the service, which could result in the increasing of the micro-hardness and brittleness of this light zone. The fusion zone had broadened in the Stellite 6 coating layer near F91 steel due to the diffusion of elements from the steel to the Stellite 6 coating after the service process. The elements diffusion eventually made a wider light zone form, consisting of the original fusion zone and new mutual diffusion zone, as shown in Figure 3. The light zone still showed a dendritic microstructure as same as that in our Stellite weld layer of sample #2. The width of the fusion zone was measured as 31, 28 and 49 µm at different locations of sample #1 in Figure 3a, indicating the width of the fusion zone was non-uniform. After long-time service, the width of the fusion zone increased significantly to form a large diffusion zone as marked in Figure 3b (sample #2), which could even reach to around 2.5 mm, indicating that a high temperature environment could promote the extent of the light zone through influencing the diffusion of alloying elements and the formation and growth of metallic compound phases and carbides. Figure 3c,d shows the XRD results of phases in the light zone of samples #1 and #2. It can be seen from the XRD results in Figure 3c that the phases in the light zone mainly consist of Co-Fe substitution solid solutions and σ phases (Fe-Cr metallic compounds), and the two elements, Co and Fe, have good solid solubility, thus they could dissolve with each other by almost any proportion. The carbides were not found from the XRD results due to the smaller quantity and small size. The type of phases in the light zone of sample #2 was similar to that of sample #1; however, there existed a more evident diffraction peak of carbides in Figure 3d, which were Cr 18.93 Fe 4.07 C 6 carbides. It meant that the formation and growth of the precipitated phases happened, and the Co matrix had evolved to form Co-Fe phases during the service, which could result in the increasing of the micro-hardness and brittleness of this light zone.

Distribution of Alloying Elementsand Phases
The microstructure of the boundary between the Stellite alloy and F91 steel of sample #1 and #2 were shown in Figure 4, and the major composition of different positions in Figure 4 was measured by EDS, as shown in Table 2. It can be seen that the microstructure evolution was obvious on the F91 steel side, where a mass of carbides within grains dissolved, precipitated and then grew along GBs, forming coarse prior austenite grain boundaries (PAGBs). Comparing the composition in position A with that in position C, it was found that the content of Fe elements increased obviously, while the content of Co, Cr and W decreased. It indicated that the diffusion process of Fe from the fusion zone to the Stellite layer predominated. The decrease of Cr content might be attributed to the diffusion of Fe, the formation of a Co-Fe matrix and a composition fluctuation of alloying elements. Thus, the content of Cr decreased in the matrix and formed Cr-rich second phases in some positions. Moreover, the content of Co in positions B and D were similar, demonstrating that the diffusion rate of Co elements from the Stellite layer to the steel was relatively low under the high temperature and pressure condition.

Distribution of Alloying Elementsand Phases
The microstructure of the boundary between the Stellite alloy and F91 steel of sample #1 and #2 were shown in Figure 4, and the major composition of different positions in Figure 4 was measured by EDS, as shown in Table 2. It can be seen that the microstructure evolution was obvious on the F91 steel side, where a mass of carbides within grains dissolved, precipitated and then grew along GBs, forming coarse prior austenite grain boundaries (PAGBs). Comparing the composition in position A with that in position C, it was found that the content of Fe elements increased obviously, while the content of Co, Cr and W decreased. It indicated that the diffusion process of Fe from the fusion zone to the Stellite layer predominated. The decrease of Cr content might be attributed to the diffusion of Fe, the formation of a Co-Fe matrix and a composition fluctuation of alloying elements. Thus, the content of Cr decreased in the matrix and formed Cr-rich second phases in some positions. Moreover, the content of Co in positions B and D were similar, demonstrating that the diffusion rate of Co elements from the Stellite layer to the steel was relatively low under the high temperature and pressure condition.  The microstructure around the boundary between the Stellite and the steel of samples #1 and #2 are shown in Figure 5a,e, and the distributions of Co, Cr and Fe in the rectangle area are displayed in Figure 5b-d,f-h. It can be seen that Co mainly formed a Co matrix in the Stellite layer, and Fe mainly distributed in the steel region to form Fe-base phases and Fe-rich carbides in Figure  The microstructure around the boundary between the Stellite and the steel of samples #1 and #2 are shown in Figure 5a,e, and the distributions of Co, Cr and Fe in the rectangle area are displayed in Figure 5b-d,f-h. It can be seen that Co mainly formed a Co matrix in the Stellite layer, and Fe mainly distributed in the steel region to form Fe-base phases and Fe-rich carbides in Figure 5b,d. Besides, the distribution of Cr was uniform from the Stellite layer to the steel in Figure 5c. It was found that changing of the Co and Cr content around the Stellite-steel interface was not significant after the service process in Figure 5f,h. However, as shown in Figure 5h, an apparent increase of the Fe content was happening in the Stellite layer near the interface, indicating that a mass of Fe had diffused from the steel into the Stellite because of the large diffusion rate of Fe. Obvious mutual diffusion of major elements happened near the boundary during the service process due to the high temperature condition, thereby leading to the formation of a mutual diffusion zone in the Stellite layer near the fusion zone. 5b,d. Besides, the distribution of Cr was uniform from the Stellite layer to the steel in Figure 5c. It was found that changing of the Co and Cr content around the Stellite-steel interface was not significant after the service process in Figure 5f,h. However, as shown in Figure 5h, an apparent increase of the Fe content was happening in the Stellite layer near the interface, indicating that a mass of Fe had diffused from the steel into the Stellite because of the large diffusion rate of Fe. Obvious mutual diffusion of major elements happened near the boundary during the service process due to the high temperature condition, thereby leading to the formation of a mutual diffusion zone in the Stellite layer near the fusion zone. The rapid diffusion of Fe could change the structure and induce the formation of new phases, resulting in the property variation of the Stellite coating components after the service process. The distribution curve of Fe, Co and Cr content perpendicular to the fusion line of sample #1 and #2 is presented using EDS, as shown in Figure 6. It can be seen that the Co content in the steel was very low. The amount of Fein the Stellite coating near the fusion line was around 30 mol %-60 mol %, indicating the diffusion process had been occurred. The amount of Fe reached 30 mol %-35 mol % at the distance of 2.2 mm from the Stellite-steel interface in the Stellite coating. Besides, the variation of the amount of Cr was not significant. Therefore, the diffusion of Fe was the main alloy elements migration in the weld overlay, resulting in microstructure evolution in the light zone. Comparing the distribution of Fe, Co and Cr of sample #2 with that of sample #1, the Fe increased obviously in the light zone near the fusion line due to diffusion of Fe during the service. The difference of Fe content between sample #1 and #2 decreased gradually when away from the fusion line in the Stellite coating. In addition, the variation tendency of Cr and Co was not very significant after the service. The rapid diffusion of Fe could change the structure and induce the formation of new phases, resulting in the property variation of the Stellite coating components after the service process. The distribution curve of Fe, Co and Cr content perpendicular to the fusion line of sample #1 and #2 is presented using EDS, as shown in Figure 6. It can be seen that the Co content in the steel was very low. The amount of Fein the Stellite coating near the fusion line was around 30 mol %-60 mol %, indicating the diffusion process had been occurred. The amount of Fe reached 30 mol %-35 mol % at the distance of 2.2 mm from the Stellite-steel interface in the Stellite coating. Besides, the variation of the amount of Cr was not significant. Therefore, the diffusion of Fe was the main alloy elements migration in the weld overlay, resulting in microstructure evolution in the light zone. Comparing the distribution of Fe, Co and Cr of sample #2 with that of sample #1, the Fe increased obviously in the light zone near the fusion line due to diffusion of Fe during the service. The difference of Fe content between sample #1 and #2 decreased gradually when away from the fusion line in the Stellite coating. In addition, the variation tendency of Cr and Co was not very significant after the service. 5b,d. Besides, the distribution of Cr was uniform from the Stellite layer to the steel in Figure 5c. It was found that changing of the Co and Cr content around the Stellite-steel interface was not significant after the service process in Figure 5f,h. However, as shown in Figure 5h, an apparent increase of the Fe content was happening in the Stellite layer near the interface, indicating that a mass of Fe had diffused from the steel into the Stellite because of the large diffusion rate of Fe. Obvious mutual diffusion of major elements happened near the boundary during the service process due to the high temperature condition, thereby leading to the formation of a mutual diffusion zone in the Stellite layer near the fusion zone. The rapid diffusion of Fe could change the structure and induce the formation of new phases, resulting in the property variation of the Stellite coating components after the service process. The distribution curve of Fe, Co and Cr content perpendicular to the fusion line of sample #1 and #2 is presented using EDS, as shown in Figure 6. It can be seen that the Co content in the steel was very low. The amount of Fein the Stellite coating near the fusion line was around 30 mol %-60 mol %, indicating the diffusion process had been occurred. The amount of Fe reached 30 mol %-35 mol % at the distance of 2.2 mm from the Stellite-steel interface in the Stellite coating. Besides, the variation of the amount of Cr was not significant. Therefore, the diffusion of Fe was the main alloy elements migration in the weld overlay, resulting in microstructure evolution in the light zone. Comparing the distribution of Fe, Co and Cr of sample #2 with that of sample #1, the Fe increased obviously in the light zone near the fusion line due to diffusion of Fe during the service. The difference of Fe content between sample #1 and #2 decreased gradually when away from the fusion line in the Stellite coating. In addition, the variation tendency of Cr and Co was not very significant after the service. TEM investigations were performed to analyze the phases in the diffusion zone. Figure 7 shows the TEM images and diffraction patterns of the microstructure in the light zone between the Stellite and the steel of sample #2. The TEM image of the matrix in the diffusion zone was displayed in Figure 7a, which was cubic structure Co-Fe phase, according to the diffraction pattern in Figure 7b, as well as XRD results and EDS analysis. It meant that Fe gradually diffused from the fusion zone and the steel to the Stellite layer in the high temperature environment, and then the Fe and Co were mutually soluble with each other to form the Co-Fe matrix in the coating layer.
Furthermore, there existed some Cr-rich phases in the diffusion zone, as shown in Figure 7c, which were black knife-like particles with a sharp edge. It can be determined that these Cr-rich phases were (Cr, Fe) 23 C 6 carbides (Cr 18.93 Fe 4.07 C 6 ) through indexing the corresponding diffraction pattern. The black particles in the light zone were σ phases (Fe-Cr phases) with a size of about hundreds of nanometers. According to the Fe-Cr phase diagram, these σ phases formed intemperatures ranging from 450-830 • C, and were also detected in alloys with 23.4 wt %-33.7 wt % Cr during plastic deformation followed by aging due to a composition fluctuation of alloying elements [21,22]. Thus, under the condition of in-service temperature (566 • C) and a large stress load, the formation of multiple brittle σ phases could be predicted, which could lead to deterioration of both corrosion resistance and ductility of the coatings. Figure 7g,h shows the TEM images of the microstructure in the light zone between the Stellite and the steel. It can be seen that there were many small, black precipitate particles in the matrix. These nano-scaled carbides were M 6 C carbides according to their small size and dispersive distribution. M 6 C carbides could significantly inhibit the dislocation motion within the matrix, thereby inducing a stronger strengthening effect than that of M 23 C 6 carbides. Therefore, the strength and hardness of the matrix was improved. TEM investigations were performed to analyze the phases in the diffusion zone. Figure 7 shows the TEM images and diffraction patterns of the microstructure in the light zone between the Stellite and the steel of sample #2. The TEM image of the matrix in the diffusion zone was displayed in Figure 7a, which was cubic structure Co-Fe phase, according to the diffraction pattern in Figure  7b, as well as XRD results and EDS analysis. It meant that Fe gradually diffused from the fusion zone and the steel to the Stellite layer in the high temperature environment, and then the Fe and Co were mutually soluble with each other to form the Co-Fe matrix in the coating layer.
Furthermore, there existed some Cr-rich phases in the diffusion zone, as shown in Figure 7c, which were black knife-like particles with a sharp edge. It can be determined that these Cr-rich phases were (Cr, Fe)23C6 carbides (Cr18.93Fe4.07C6) through indexing the corresponding diffraction pattern. The black particles in the light zone were σ phases (Fe-Cr phases) with a size of about hundreds of nanometers. According to the Fe-Cr phase diagram, these σ phases formed intemperatures ranging from 450-830 °C, and were also detected in alloys with 23.4 wt %-33.7 wt % Cr during plastic deformation followed by aging due to a composition fluctuation of alloying elements [21,22]. Thus, under the condition of in-service temperature (566 °C) and a large stress load, the formation of multiple brittle σ phases could be predicted, which could lead to deterioration of both corrosion resistance and ductility of the coatings. Figure 7g,h shows the TEM images of the microstructure in the light zone between the Stellite and the steel. It can be seen that there were many small, black precipitate particles in the matrix. These nano-scaled carbides were M6C carbides according to their small size and dispersive distribution. M6C carbides could significantly inhibit the dislocation motion within the matrix, thereby inducing a stronger strengthening effect than that of M23C6 carbides. Therefore, the strength and hardness of the matrix was improved.  Figure 8a shows the micro-hardness profiles perpendicular to the fusion line along the cross-section of the samples. It can be seen that the micro-hardness in the weld overlay was higher than that in the steel. Meanwhile, the micro-hardness in the HAZ was higher compared with F91 steel for the sample #1. After a long period of service, the micro-hardness values of sample #2 increased to the maximum (470-680 HV) in the weld overlay near fusion zone due to the formation and growth of hard matrix phases. Moreover, as for sample #2, the micro-hardness values in the HAZ dropped to be the same as that in the base metal, and the micro-hardness values in the Stellite  Figure 8a shows the micro-hardness profiles perpendicular to the fusion line along the cross-section of the samples. It can be seen that the micro-hardness in the weld overlay was higher than that in the steel. Meanwhile, the micro-hardness in the HAZ was higher compared with F91 steel for the sample #1. After a long period of service, the micro-hardness values of sample #2 increased to the maximum (470-680 HV) in the weld overlay near fusion zone due to the formation and growth of hard matrix phases. Moreover, as for sample #2, the micro-hardness values in the HAZ dropped to be the same as that in the base metal, and the micro-hardness values in the Stellite overlay had increased (450-500 HV). In the HAZ of F91 steel, the micro-hardness decreased because of the coarsening of carbides.

Failure Analysis
There were two kinds of cracks existing in the weld overlay of sample #2 as shown in Figure 8b. One kind of them is the circumferential crack paralleling to the fusion line, and the other one is a radial crack perpendicular to the fusion line. When the two kinds of cracks merged gradually, the coating could fracture and fall off. Figure 8c shows the failure locations of sample #2, and it can be seen that the fracture mainly occurred in the light zone. overlay had increased (450-500 HV). In the HAZ of F91 steel, the micro-hardness decreased because of the coarsening of carbides. There were two kinds of cracks existing in the weld overlay of sample #2 as shown in Figure 8b. One kind of them is the circumferential crack paralleling to the fusion line, and the other one is a radial crack perpendicular to the fusion line. When the two kinds of cracks merged gradually, the coating could fracture and fall off. Figure 8c shows the failure locations of sample #2, and it can be seen that the fracture mainly occurred in the light zone.

Discussion
After the service process, hard Co-Fe matrix and σ phases formed, caused by the diffusion of Fe from steel into the Stellite layer, resulting in the formation of a wide light zone, which consisted of the original fusion zone and new mutual diffusion zone. The growth of the Co-Fe matrix phases with high hardness in the light zone led to formation of an interface between the light zone and Stellite layer. The micro-hardness in the light zone was higher than that in the Stellite layer significantly, and the maximum difference in micro-hardness between the two zones was even up to 230 HV. Moreover, from the Stellite layer to the light zone near the interface, the micro-hardness suddenly changed from 450 to 680 HV as displayed in Figure 8a. It indicated that the changes of micro-hardness between the light zone and Stellite layer happened after a long period of service process. Large stress concentrations could easily occur on this interface because of changes in micro-hardness, resulting in the formation of a vast nucleation site of micro-cracks in the interface as shown in Figure 9a. In addition, the σ phases, which are hard and brittle, harmfully affect the mechanical properties of the coatings by creating local embrittlement and by forming micro-cracks at the γ/σ phase interfaces during loading. After nucleating at these interfaces, the cracks aggregated and grew gradually in Figure 9b. Finally the cracks propagated into the light zone and became large cracks. The large cracks gradually increased, and thus led to the fracture and fall off of the coating.

Discussion
After the service process, hard Co-Fe matrix and σ phases formed, caused by the diffusion of Fe from steel into the Stellite layer, resulting in the formation of a wide light zone, which consisted of the original fusion zone and new mutual diffusion zone. The growth of the Co-Fe matrix phases with high hardness in the light zone led to formation of an interface between the light zone and Stellite layer. The micro-hardness in the light zone was higher than that in the Stellite layer significantly, and the maximum difference in micro-hardness between the two zones was even up to 230 HV. Moreover, from the Stellite layer to the light zone near the interface, the micro-hardness suddenly changed from 450 to 680 HV as displayed in Figure 8a. It indicated that the changes of micro-hardness between the light zone and Stellite layer happened after a long period of service process. Large stress concentrations could easily occur on this interface because of changes in micro-hardness, resulting in the formation of a vast nucleation site of micro-cracks in the interface as shown in Figure 9a. In addition, the σ phases, which are hard and brittle, harmfully affect the mechanical properties of the coatings by creating local embrittlement and by forming micro-cracks at the γ/σ phase interfaces during loading. After nucleating at these interfaces, the cracks aggregated and grew gradually in Figure 9b. Finally the cracks propagated into the light zone and became large cracks. The large cracks gradually increased, and thus led to the fracture and fall off of the coating. Coatings 2019, 9, x FOR PEER REVIEW 9 of 11 As shown in Figure 8b,c, the formation of circumferential cracks on the interface was related to the hardness changes between the light zone and Stellite layer, which could be separated from the Stellite weld layer and the light zone, and fracture under the combined effect of high temperature and impact load under service conditions.
Besides, inter-dendritic regions were weak zones, where impurities, coarse carbides and micro-cracks caused by thermal stress and residual stress existed, which could induce the formation of radial cracks. The high hardness made it easier for the micro cracks to form and propagate in the light zone. The cracks which initiated on the interface between the light zone and Stellite layer could propagate into the light zone gradually and form a large macro-crack. Thus, the light zone becomes the weakest region and fracture failure could happen in this zone.

Conclusions
Overlaying of Stellite 6 alloy on F91 steel substrates was performed with multi-pass TIG cladding method. The effect of in-service environment on the element diffusion, microstructure evolution and fracture behavior of the resultant weld overlay were investigated. The conclusions are listed as follows: • The microstructure of the Stellite weld overlay near the fusion zone had changed to form a light zone, consisting of Co-Fe substitution solid solutions, σ phases (Fe-Cr metallic compounds) and Cr18.93Fe4.07C6 carbides. The high temperature environment could promote the extent of the light zone. After service, the width of the light zone, combined with fusion zone and diffusion zone, increased significantly to form a large diffusion zone, which could even reach to around 2.5 mm.

•
The obvious diffusion of Fe occurred from the steel and fusion zone to the Stellite overlay, resulting in the microstructure evolution and hardness increase in the weld overlay. The content of Fe increased intensively, but the content of Co decreased, which could eventually lead to the formation of hard and brittle Co-Fe phases.

•
The micro-hardness in the Stellite weld overlay was higher than that in the steel. After cladding, the micro-hardness in the HAZ increased. After the service process, the micro-hardness values in the Stellite overlay slightly increased to 450-500 HV, while those in the HAZ dropped where the precipitates had coarsened. Moreover, the micro-hardness values in the light zone increased to the maximum (470-680 HV), resulting in changes of micro-hardness between the base material and the Stellite weld overlay.

•
The fracture of the Stellite coating samples mainly occurred in the light zone (fusion zone + diffusion zone) after the service process. The formation of these cracks might be caused by formed brittle phases and changes of micro-hardness during service.  As shown in Figure 8b,c, the formation of circumferential cracks on the interface was related to the hardness changes between the light zone and Stellite layer, which could be separated from the Stellite weld layer and the light zone, and fracture under the combined effect of high temperature and impact load under service conditions.
Besides, inter-dendritic regions were weak zones, where impurities, coarse carbides and micro-cracks caused by thermal stress and residual stress existed, which could induce the formation of radial cracks. The high hardness made it easier for the micro cracks to form and propagate in the light zone. The cracks which initiated on the interface between the light zone and Stellite layer could propagate into the light zone gradually and form a large macro-crack. Thus, the light zone becomes the weakest region and fracture failure could happen in this zone.

Conclusions
Overlaying of Stellite 6 alloy on F91 steel substrates was performed with multi-pass TIG cladding method. The effect of in-service environment on the element diffusion, microstructure evolution and fracture behavior of the resultant weld overlay were investigated. The conclusions are listed as follows: • The microstructure of the Stellite weld overlay near the fusion zone had changed to form a light zone, consisting of Co-Fe substitution solid solutions, σ phases (Fe-Cr metallic compounds) and Cr 18.93 Fe 4.07 C 6 carbides. The high temperature environment could promote the extent of the light zone. After service, the width of the light zone, combined with fusion zone and diffusion zone, increased significantly to form a large diffusion zone, which could even reach to around 2.5 mm.

•
The obvious diffusion of Fe occurred from the steel and fusion zone to the Stellite overlay, resulting in the microstructure evolution and hardness increase in the weld overlay. The content of Fe increased intensively, but the content of Co decreased, which could eventually lead to the formation of hard and brittle Co-Fe phases.

•
The micro-hardness in the Stellite weld overlay was higher than that in the steel. After cladding, the micro-hardness in the HAZ increased. After the service process, the micro-hardness values in the Stellite overlay slightly increased to 450-500 HV, while those in the HAZ dropped where the precipitates had coarsened. Moreover, the micro-hardness values in the light zone increased to the maximum (470-680 HV), resulting in changes of micro-hardness between the base material and the Stellite weld overlay.

•
The fracture of the Stellite coating samples mainly occurred in the light zone (fusion zone + diffusion zone) after the service process. The formation of these cracks might be caused by formed brittle phases and changes of micro-hardness during service.