Root-like crack propagation: Synergistic of gradient structure and diffusion barrier in TBCs on TiAl Alloy

Root-like crack propagation: Synergistic of gradient structure and diffusion barrier in TBCs on TiAl Alloy | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Root-like crack propagation: Synergistic of gradient structure and diffusion barrier in TBCs on TiAl Alloy Wufan Li, Chen Hua, Yanhong Zhuang, Tao Wang, Taihong Huang, peng song This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-9191760/v1 This work is licensed under a CC BY 4.0 License Status: Under Revision Version 1 posted 11 You are reading this latest preprint version Abstract In this study, MYC (CoNiCrAlY-Y 2 O 3 -Cr 3 C 2 ) coatings and MYC/8YSZ gradient coatings were fabricated by atmospheric plasma spraying (APS) technology. During the oxidation cycle of the MYC coating, an in-situ diffusion barrier was formed at the interface between the coating and the substrate. The Vickers hardness test and nanoindentation experiments verified that this low-modulus, high-hardness diffusion barrier structure played the role of an "energy buffer layer". In the 900°C water quenching thermal shock test of the MYC/8YSZ gradient coating, after 960 cycles, it was found that the gradient structure extended the crack propagation path, improved the energy dissipation capacity of the coating, and formed root-like cracks inside the coating, which increased the thermal shock cycle life of the coating. The coating did not experience failure. This study demonstrated the advantages of gradient structures in improving coating performance, especially in terms of crack resistance under thermal shock conditions. energy buffer layer gradient structure thermal shock root-like crack Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Figure 6 Figure 7 Figure 8 Figure 9 Figure 10 Figure 11 1 Introduction TiAl alloys are lightweight high-temperature structural materials with high specific strength, excellent creep resistance, and oxidation resistance, and they have demonstrated significant application potential in aerospace engines[ 1 – 4 ]. This intermetallic alloy was first introduced in aeroengines in 2011 (alloy Ti-48A1-2Nb-2Cr, or ‘48-2-2’) for the low-pressure turbine (LPT) blades[ 5 , 6 ]. As aircraft engines continue to evolve towards higher thrust-to-weight ratios and lower emissions, higher requirements are placed on high-temperature components[ 7 ]. Unfortunately, the operating temperatures of most commercially available TiAl alloys are below 750°C[ 8 , 9 ]. Numerous studies have shown that the high-temperature performance of TiAl alloys can be improved by applying protective coatings[ 10 , 11 ]. Thermal barrier coatings (TBCs) play a critical role in safeguarding the integrity of superalloys, significantly affecting the reliability and operational lifespan of essential components[ 12 – 15 ].8YSZ, as a typical thermal-protective coating material, significantly influences the thermal protection effect of the surface coating on the TiAl substrate, and is a key part of the TBC structure[ 16 , 17 ].Currently, due to its high deposition efficiency and cost-effectiveness, Atmospheric Plasma Spraying (APS) has become one of the main technologies for preparing TBCs[ 18 ]. Since the early 1990s, APS has been extensively utilized for the fabrication of bilayer and gradient coatings. The traditional double-layer 8YSZ TBCs consist of the ceramic coating and MCrAlY bond coating[ 19 ]. However, under high-temperature conditions, severe interdiffusion occurs between the MCrAlY coating and the TiAl substrate, leading to the formation of Kirkendall voids[ 20 ]. Furthermore, the difference in composition between the coating and the TiAl alloy causes elemental diffusion near the interface[ 11 ]. Consequently, interdiffusion zone (IDZ) is brittle, further deteriorating the adhesion at the coating/substrate interface[ 11 , 21 , 22 ]. To mitigate this issue, the introduction of an intermediate layer has been reported as an effective strategy to suppress elemental diffusion at the interface between the coating and the substrate[ 23 ]. For example, Al 2 O 3 films and Cr 2 O 3 have been used as intermediate layers between TiAl and NiCrAlY to inhibit the diffusion of Ni from the NiCrAlY bond coating into the TiAl substrate[ 23 , 24 ]. Although the direct preparation of a diffusion barrier layer on TiAl can effectively slow down interdiffusion, it introduces a hetero-phase, which may have potential impacts on the overall long-term stability[ 25 , 26 ]. Therefore, the development of in-situ diffusion barriers has attracted increasing attention as a promising strategy to improve interfacial performance[ 21 ]. Hua[ 27 ] developed a CoNiCrAlY-Y 2 O 3 -Cr 3 C 2 (MYC) gradient coating on the surface of TiAl alloy via plasma spraying technology. During the subsequent high-temperature oxidation process, a Ti 2 AlN phase was formed at the interface between the substrate and the bond coating. This in-situ formed phase further enhanced the overall bonding performance. In addition to interface weakening caused by diffusion, TBCs on TiAl alloys are also susceptible to thermal stress-induced failure. For the double-layer coating, there are differences in the elastic modulus and thermal expansion coefficient between the metal and ceramic coatings, which makes the thermal barrier coating susceptible to the thermal stress caused by temperature fluctuations during operation. These stresses can initiate and propagate cracks at the coating interface or within the layers, leading to delamination and functional failure[ 28 , 29 ]. Previous studies have shown that the bonding strength of the gradient coating is 2.5 times that of the double-layer coating[ 30 ]. This is because the gradient coating can achieve a continuous transition of composition and properties, consequently reducing the risk of coating peeling caused by the coefficient of thermal expansion (CTE) mismatch. During the plasma spraying process, by adjusting the spraying power and the particle size of the powder, the layer-by-layer transition of the coating can be achieved, resulting in a structure with different hardness and toughness[ 30 ]. Hua reported that continuous compositional transitions in gradient coatings result in hardness gradients and enhanced energy dissipation capability[ 31 ]. In the gradient region, the expansion of cracks will occur along the interface direction of the bond coating, while those cracks extending to the inner layer are more likely to expand laterally towards the interface between the bond coating and the substrate, rather than continuing to extend along the thickness direction to the substrate[ 32 ]. Chen further demonstrated that under the thermal shock condition of 900°C, the initiation and propagation characteristics of cracks in the gradient coating are closely related to the thickness of the gradient region and sample with the thickness of the gradient region was 150µm exhibited greater thermal shock tolerance[ 33 ].This paper reports the evolution of cracks in the MYC/YSZ gradient coating prepared by atmospheric plasma spraying. The crack propagation and failure mechanism of this gradient coating under mechanical and thermal stress were studied. Vickers hardness tests and thermal shock tests were conducted on the coating/ substrate specimens to simulate different service times and study the crack propagation process. 2 Experimental Details 2.1 Base materials The TiAl alloy substrate (nominal composition: Ti 51.4 at. %, Al 43.5 at. %, Nb 4 at. %, Mo 1 at. %, B 0.1 at. %) in this study was produced by Beijing Yanbang New Materials Technology Co., Ltd. with Vacuum Suction Melting Technology (unit type YBXF-3). The substrate was processed into cylindrical block with dimensions of Φ25*5mm using Computerized Numerical Control machine. The specimens were cleaned with anhydrous ethanol in ultrasonic cleaner (DR-MS20, Shenzhen Derui Ultrasonic Equipment Co., Ltd.) to thoroughly remove surface impurities and grease before sandblasting. Then, the surface of the base material was subjected to sandblasting treatment until the surface of the substrate is completely roughened and the roughness is uniform by CS-600 sandblasting machine (Shanghai Jichuan Machinery Technology Co., Ltd.). 2.2 Coatings Two coatings were applied to the TiAl alloy: CoNiCrAlY-30Y 2 O 3 -20Cr 3 C 2 Ceramic-metal coating (MYC), CoNiCrAlY-Y 2 O 3 -Cr 3 C 2 /8YSZ gradient coating (MYC/YSZ). The MYC powder is composed of CoNiCrAlY (nominal composition: Co balance, Ni 32 wt.%, Cr 21 wt.%, Al 8wt.%, Y 0.5 wt.%) Y 2 O 3 , and Cr 3 C 2 in a mass ratio of 5:3:2. The powder particle size is controlled to be between 15–45 µm using a sieve. The specific experimental procedures have been described in the previous report[ 27 ]. The TiAl-MYC and TiAl-MYC/YSZ were prepared by APS technology, the spraying parameters for MYC and MYC/YSZ were described in our previous study[ 27 ]. TiAl-MYC was obtained by APS, on the grit blasted surface of an TiAl sample. The crucial step in the process involves using a two-way powder feeding method, an alternating and superimposed MYC/YSZ layered structure was formed at the microscopic level, and a gradient transition in composition and structure was achieved at the macroscopic level. The TiAl-MYC/YSZ was fabricated following three steps. First, the MYC coating with a deposition thickness of approximately 50 µm. This not only provides a good interface between the coating and the substrate, but also serves as a carbon source for subsequent interface reactions, helping to form an in-situ diffusion barrier. Then, by dynamically adjusting the powder feeding rates of the two powders (from a single MYC gradually increasing to a single 8YSZ), a gradient region with a thickness of approximately 150 µm is deposited, allowing the composition to gradually transition from the metallic phase to the ceramic phase, avoiding interface stress concentration. After the feeding of MYC powder, 8YSZ surface layer with a thickness of approximately 200 µm is deposited to provide good thermal insulation and thermal shock resistance. During the spraying process, to ensure uniform coating thickness and consistency of microstructure in each area, the spraying path design uses a 3 mm step distance and multiple passes of spraying to avoid structural defects caused by uneven powder deposition density. At the same time, parameters such as the spray gun movement speed, base preheating temperature, plasma gas composition, and total power have also been systematically optimized. 2.3 High-temperature Test Static cyclic oxidation tests were conducted using a box-type electric furnace (SX-B01123, Tianjin Zhonghuan Electric Furnace Co., Ltd.). The samples were held at 900°C in static air for 24 hours, then cooled in air for 30 minutes, which was defined as one cycle. Samples were taken out at 96 hours (4 cycles), 240 hours (10 cycles), and 504 hours (21 cycles). The main focus was on the changes in the coating structure at high temperatures. The thermal shock test was also carried out using the same type of box-type electric furnace. The samples were held at 900 ℃ in static air for 15 minutes and then rapidly quenched in water to room temperature, which was defined as one thermal shock cycle. Thermal shock cycles of 384 and 960 times were conducted on the MYC/YSZ gradient coating (lasting for 96 hours and 240 hours respectively). Before the high-temperature test, the temperature of the box-type furnace was calibrated using a multi-channel temperature tester (HPS3032, Changzhou Haierpa Electronic Technology Co., Ltd.). The specific operation was as follows: the furnace was set at 900°C, and the temperature at the center of the furnace cavity was measured using the temperature tester. The furnace setting temperature was adjusted until the measured temperature stabilized at 900 ± 1 ℃. 2.4 Mechanics Performance Test and Characterization The Vickers Hardness tests were conducted in a Microhardness tester (DHV-1000Z, Jiangsu Nan Guang Electronic Technology Co., Ltd.) to test the hardness of the coating cross-section. the cold-embedded samples were ground and polished with SiC sandpaper on their cross-sections to ensure the flatness met the testing requirements. The indenter was loaded along the direction from the bond coating to the ceramic coating at different positions of the coating. The loading force of the indenter was set at 0.5N, and the holding time was 10 seconds. For the Nanoindentation tests, the maximum applied load in the experiment was 30 mN, the displacement noise level was 0.2 nm, and the thermal drift parameter was less than 0.05 nm/s. In this paper, the nanoindentation tests were all conducted using a Berkovich Triangular conical indenter. The microstructure and composition distribution of the coating cross-section and fracture surface were analyzed using a field emission scanning electron microscope (SEM, Tescan-MIRA LMS + Bruker Quantax 200 XFlash 6|6) and the crack propagation was observed. 3 Results and discussion 3.1 microstructure of coating cross section Figure 1 shows the SEM image of the cross-sectional morphology of MYC after 96h, 240h and 504h of oxidation at 900 ℃. From Fig. 1 a to Fig. 1 c, no obvious diffusion layer is seen. Furthermore, after prolonged high-temperature exposure, a layer gradually forms and its thickness gradually increases, showing a stable growth trend (Fig. 1 d-f). This phenomenon reveals that the MYC coating prepared by APS can effectively inhibit the diffusion of metal elements such as Co and Ni to the substrate under high temperature conditions, inhibit the formation of the CoNiTi 2 brittle phase, and thus maintain excellent high-temperature phase stability. During the initial stages of oxidation, high temperature is likely to prompt chemical reactions at the material interface, leading to diffusion barriers formation in situ. Its existence can not only effectively hinder the diffusion of Co and Ni to the substrate, inhibit the formation of intermetallic compounds, but also reduce thermal stress mismatch. Additionally, in an air environment of 900°C, the diffusion barrier continues to generate and grow continuously during the thermal exposure process, showing good high temperature stability. This feature not only helps to improve the coating's ability to hinder element diffusion during long-term service, but also provides an important basis for the subsequent analysis of the evolution and fracture behavior of the interface microstructure, especially the interface mechanical properties. It should be noted that some black particles can be observed at the interface between the coating and the substrate (Fig. 1 b). These particles are corundum (Al 2 O 3 ) particles embedded in the surface of the substrate during sandblasting pretreatment. Regarding element distribution, Ti, C and N elements in the Layer 1 is highly enriched (Fig. 2 ). This compositional feature suggests that Layer 1 is primarily composed of the Ti 2 CN phase[ 27 ]. This phase not only forms rapidly during the early heat treatment process but also maintains structural stability during subsequent high-temperature exposure, demonstrating excellent anti-diffusion performance and thermal stability. In contrast, the Layer 2 shows signals of elements such as Ti, Al, and N, suggesting that this region may have generated another phase structure mainly composed of Ti, Al and N, which is most likely Ti 2 AlN[ 27 ]. This discovery indicates that during the evolution of the in-situ diffusion barrier, a double-layer diffusion barrier composed of Ti 2 CN and Ti 2 AlN was formed[ 27 ]. 3.2 Interface Vickers hardness analysis Researchers have found that the length of crack propagation is closely related to fracture toughness and interfacial bonding strength. The interfaces of the diffusion barrier were observed by SEM, Fig. 3 a and d shows the interface after 96 hours of oxidation, the thickness of the diffusion barrier is approximately 2 µm. Under cyclic oxidation conditions of 96h and 240h (Fig. 3 b and e), this diffusion barrier gradually thickens with time. However, when the high-temperature oxidation time is extended to 504h (Fig. 3 c and f), the growth trend of the diffusion barrier thickness slows down significantly, which means that the diffusion barrier is close to growth saturation after entering the 504h stage. There was no significant interfacial lateral crack propagation in the interface area, which indicated that a firm bonding interface was formed between the diffusion barrier and the substrate and coating, which could withstand higher local stress concentration. Only when the interface is affected by the hardness indentation head, short longitudinal cracks appear in the vertical direction, and these cracks have limited lengths of expansion (about 3 ~ 4µm). In addition, the diffusion barrier is caused by elements in the coating to diffusion into the inside of the substrate. The growth direction is ingrown and will not bring too much growth stress to the interface area. This characteristic means that the interface diffusion barrier has a high fracture toughness and interface combination strength, which can effectively relieve stress concentration under local loading conditions and avoid large-scale expansion of cracks. 3.3 Interface nanoindentation test As shown in Fig. 4 , the nanoindentation test results and morphological characteristics of the interface area after 96 hours, 240 hours, and 504 hours of oxidation are presented. Figures a to c and b to f respectively illustrate the three-dimensional distribution maps of the elastic modulus and hardness in this area. After 96h oxidation, the modulus and hardness of the interface area have increased compared to the substrate, but the growth amplitudes of the two are significantly different. The increase in hardness is much greater than that of the modulus. By further observing the nanoindentation morphology shown in Fig. 4 g, it can be clearly seen that the indentation size at the interface is significantly smaller than that in the substrate region, indicating a substantial increase in hardness. However, the modulus is relatively low, effectively alleviate stress concentration and inhibit crack initiation and propagation under local loading, thereby improving the interface toughness and stability. As the oxidation time increases, the thickness of the diffusion barrier gradually increases, and the mechanical properties also undergo evolution. As shown in Fig. 4 b and e, after 240 hours of oxidation, the difference in modulus between the interface and the susbstrate did not increase significantly compared to that at 96 hours, but the increase in hardness was more obvious. The modulus mismatch remained significantly lower than the hardness. This suggests that the diffusion barrier maintains a relatively high hardness while still having moderate modulus matching and good crack resistance. When the heat exposure time was extended to 504 hours, as shown in Fig. 4 c and f, the diffusion barrier further thickened, and the modulus difference between the interface and the substrate increased. However, the modulus growth was not significant, and the mismatch amplitude was smaller than the hardness, maintaining a good elastic adjustment function. This helps to alleviate the residual stress at the interface under high-temperature cyclic loads and avoids interface peeling and cracking[ 34 ]. The hardness of the interface area has significantly increased (Fig. 4 d-f), which ensures that the interface layer can effectively resist mechanical loads in the high-temperature oxidation and thermal cycling environment, maintaining the overall integrity of the coating/substrate system. The significant increase in hardness also inhibited the initiation of microcracks, fundamentally extending the service life. On the other hand, from the modulus perspective, although the elastic modulus of the diffusion barrier also increased with the extension of the oxidation time, its relative modulus mismatch with the substrate was still significantly lower than the hardness mismatch, indicating that the low modulus characteristic represents that the diffusion barrier has a certain deformation ability, which can alleviate the thermal physical property differences between MYC and the substrate, improve the damage tolerance, and this high hardness and low modulus combination is very crucial for interface performance. The low modulus mismatch between the Ti 2 CN and Ti 2 AlN double-layer diffusion barrier and the substrate is crucial as it acts as an "energy transition buffer layer" at the interface. The existence of the Ti 2 CN and Ti 2 AlN double-layer diffusion barrier can form a gradually decaying and transitional stress buffer zone in this region, relieve the thermal physical properties differences between the coating and the substrate, thereby effectively reducing residual stress and potential cracking risks caused by thermal mismatch. Although the MYC coating has solved the problem of interfacial mutual diffusion, the double-layer coating in practical applications still has the issue of poor adhesion. Inspired by this gradient structure, by combining MYC with YSZ, a gradient structure was designed, the element distribution in the initial gradient region is shown in Fig. 5 . The distribution of Ni and Cr is completely opposite to that of Zr. The outer layer is observed to the YSZ thermal protection coating, while within the gradient region, both CoNiCrAlY and YSZ are present. 3.4 Thermal shock test After 384 thermal shock cycles (Fig. 6 a and c), no transverse cracks were observed on the cross-section of the coating, and the ceramic coating and bond coating maintained a good adhesion. This demonstrates that the gradient structure can maintain high structural stability and crack resistance under multiple alternating cold and hot impacts. However, when the thermal shock cycle number increased to 960 (Fig. 6 b and d), a large number of transverse micro-cracks appeared in the gradient region, and some vertical cracks formed step-like cracks in different directions with the transverse cracks. Most of the cracks stopped expanding in the ceramic layer. The deflection and bridging of the cracks increased the energy required for their expansion, improve the overall toughness of the coating significantly. From the mechanism perspective, introducing an interface between the CoNiCrAlY, Cr 3 C 2 and YSZ phase leads to a certain degree of thermal expansion coefficient mismatch. During the thermal shock test, the vertical cracks first extend from the surface of the ceramic coating and extend towards the interface between the ceramic coating and the gradient region. When the vertical crack is far away from the crack band structure, the expansion of the horizontal cracks within the crack band will change the distribution of internal stress, thereby exerting an inhibitory effect on the downward extension of the vertical crack. When the vertical cracks extend towards the horizontal cracks until they merge, the energy in the coating is released, and finally, transverse cracks parallel to the interface are formed, which can prevent the surface vertical cracks from extending to the interface between the coating and the substrate. As shown in Fig. 7 a and b, the morphology of the interface between the coating and the substrate and EDS mapping after 384 thermal shock cycles are presented. It can be observed that in the thermal shock environment, the formation of the in-situ diffusion barrier is not affected. It forms a good bond between the coating and the substrate, and no signs of cracking or peeling such as cracks are observed at the interface. Through EDS analysis of Fig. 7 b, a very thin Al 2 O 3 layer can be seen above the diffusion barrier, while the Ti 2 CN diffusion barrier has been formed. Additionally, a small amount of Ni and Co diffusion products appear below the diffusion barrier, but Ti has not diffused into the coating. After the thermal shock test was conducted for 960 times, the interface structure remained intact, as shown in Fig. 7 c and d. Through EDS analysis of Fig. 7 d, the composition of the diffusion barrier has shown a layer structure namely Ti 2 CN on the upper layer and Ti 2 AlN on the lower layer. Particularly, Ni and Co diffusion products were detected in the interface area. The analysis suggests that part of the reason lies in the extreme temperature cycling and rapid cooling and heating alternation in the water quench thermal shock test, which causes the formation of the diffusion barrier to be constantly starting and stopping. Furthermore, the thermal cycling process will result in a significant temperature gradient and thermal expansion differences. These differences introduce a pronounced multi-directional stress field within the coating. This stress state is extremely complex, including both thermal stress and the high strain rate cyclic loading caused by thermal shock, causing the diffusion barrier to be repeatedly subjected to tension and compression, shear, and thermal fatigue. As a result, vertical cracks form at the interface. These minor cracks observably enhance the permeability and the number of diffusion channels in the interface area, promoting the penetration of Ni and Co through the in-situ diffusion barrier and entering the TiAl substrate to form a diffusion layer. Figure 8 shows the cross-sectional morphologies of the MYC/8YSZ coating after 384 and 960 thermal shock cycles. The results indicate that vertical cracks appeared at the top of the ceramic coating and the gradient region. After 384 thermal shock cycles, vertical cracks appeared at the top of the ceramic layer, but the cracks were blocked when they extended to the top of the gradient region. Crack is always under compressive stress, preventing it from extending downward. However, when the thermal shock cycles increased to 960 (Fig. 9 ), vertical cracks in the ceramic layer extended downward from the top. When they reached the gradient region, they were hindered by the metallic phase, causing the cracks to deflect. Some of them formed horizontal cracks, while others continued to extend downward. There are intrinsic thermal physical property parameters (such as thermal expansion coefficient, elastic modulus, fracture toughness, etc.) differences among different phases within the gradient region. This results in the cracks not extending in the same direction but instead preferentially expanding along the weak interface with the lowest local fracture energy, demonstrating the characteristic of selective fracture. These weak interfaces are not uniformly distributed in the gradient region, they exhibit a certain degree of randomness and gradient characteristics. This distribution pattern leads to multiple deflections, splits, and branching of the crack propagation path, thereby delaying the formation and propagation of the through-crack. The further propagation of the crack requires a greater amount of energy to drive, thus preventing the coating from peeling off. Figure 10 shows the cross-sectional morphology and EDS mapping after 960 thermal shock tests of the coating. It is unable to provide sufficient energy to sustain crack propagation. The horizontal cracks extend to the CoNiCrAlY metallic phase and no longer grow. This indicates that the YSZ ceramic phase has a synergistic blocking effect on crack propagation with CoNiCrAlY, and the local stress generated by the ceramic phase can be effectively alleviated through the stress transfer between adjacent metal phases, thereby achieving effective fracture energy transfer between the phases[ 35 , 36 ]. The compressive stress of the CoNiCrAlY metallic phase has an effect of "bridging" and "passivation" on the cracks tip[ 35 ], inhibiting the driving force for the cracks to penetrate the metallic phase. Only when the energy accumulates to a sufficient level will the cracks continue to deflect and extend through the metal substrate. Only when the energy accumulates to a sufficiently high level or encounters other cracks will the cracks continue to deform and extend through the metal substrate. Such cracks do not exist within the gradient region of 384 thermal shocks, and the cracks in the 960 thermal shock case are also few. The interface bonding of the coating was relatively strong, the crack propagation mechanism is matched, and there is no obvious stress concentration[ 35 ]. 3.5 Crack Propagation Mechanism and Failure Model under Thermal Shock Under thermal shock conditions, the crack propagation process is root-like, that is, the crack starts to grow from the external ceramic coating and then extends along the interior of the coating. It delays failure through multiple bends and branching. As shown in Fig. 11 a to d, the initial initiation of the crack usually occurs in the ceramic layer, which has a low fracture toughness and is prone to crack initiation at weak interfaces. At this time, the cracks mainly extend along the surface of the ceramic coating, usually manifesting as short vertical micro-cracks. The initial propagation of the crack is usually confined to the upper ceramic region and is not likely to rapidly reach the ceramic bottom layer. As the number of thermal shock impact cycles increases, the surface-initiated cracks gradually extend to the gradient coating (Fig. 11 e and g). For the traditional double-layer coating, when cracks propagate at the interface between the coating and the bonding layer, they will deflect along the interface and form transverse cracks, resulting in the detachment of the coating. In the gradient region, however, due to the presence of the metal phase, the macroscopic mechanical properties of this area can be adjusted to exhibit a gradient change, thereby prolonging the crack propagation path and increasing energy dissipation. In detail, the adjacent phases have different Young’s modulus and thermal expansion coefficients, which lead to a complex stress field at the crack tip, thus causing the crack to exhibit a root-like branching and deflection (Fig. 11 g). When the crack contacts the metal phase, due to the higher hardness of the metal phase, the crack will deviate along the metal-ceramic interface or temporarily stop. With the increase in the number of thermal shock impacts, the crack then extends along the metal phase. When the driving force is strong enough, the crack can break through the obstruction of the metal phase and continue to extend (Fig. 11 f and h). This repeated deflection and stopping increase the possible extension path length of the crack and also increase the energy required for crack propagation. The discontinuous metal layer can convert the main crack into multiple dispersed micro-cracks, effectively preventing the propagation of through-cracks, significantly enhancing the coating's toughness and thermal shock resistance. In addition, cracks can form bridging bands in the gradient region. These bridging bands reduce the stress at the crack tip. During the repeated thermal shock cycles, the plastic deformation of the bridging bands leads to continuous energy dissipation, significantly extending the crack propagation path and dispersed the driving force for crack propagation. 4 Conclusion In this study, the mechanical properties of the MYC coating after 96 hours, 240 hours, and 504 hours of high-temperature oxidation were tested and analyzed through scanning electron microscopy observations and EDS spectra. Through 384 and 960 thermal shock tests, the crack propagation mechanism of the MYC/YSZ gradient coating under high-temperature oxidation conditions was revealed. The main conclusions are as follows: The MYC coating and the interface underwent in-situ reactions, forming a dense double-layer diffusion barrier composed of Ti 2 CN and Ti 2 AlN. The Vickers hardness test and nanoindentation experiments indicated that this low modulus and high hardness diffusion barrier played the role of an "energy buffer layer" at the interface, effectively reducing residual stress and the potential crack risk caused by thermal mismatch. After 384 thermal shock cycles, no horizontal cracks appeared in the MYC/YSZ gradient coating, while when the number of thermal shock cycles increased to 960, layer-like crack bands appeared in the gradient zone. The deflection and bridging of cracks increased the energy required for their expansion, significantly improving their toughness. The crack propagation process within the MYC/YSZ coating is a root-like mechanism. This root-like mechanism effectively prolongs the possible crack propagation path and enhances the thermal shock resistance of the coating. The component gradient effectively alleviated the stress concentration at the ceramic/metal interface, causing the cracks to exhibit branching and bending characteristics, continuously consuming the fracture energy and effectively preventing the expansion of through-cracks, significantly improving the toughness of the coating. Declarations Author Contribution Wufan Li: Conceptualization, Investigation, Methodology, Validation, Writing-original draft.Chen Hua: Data curation, Project administration, Supervision, Writing-review and editing.Yanhong Zhuang: Data curation, Software, Validation.Tao Wang: Methodology, Visualization.Taihong Huang: Data curation, Project administration, Funding acquisition, Supervision, Writing-review and editing. Peng Song: Funding acquisition, Project administration, Supervision. Acknowledgement I would like to express my heartfelt gratitude to supervisor Huang Taohong and senior brother Hua Chen. From the selection of the topic, the framework construction, to the revision and improvement, they all provided me with patient guidance and professional suggestions, enabling me to successfully complete the research. Meanwhile, I would like to thank team member of research program for helping and supporting each other during the learning and writing process. References Y. Lu, X. Wang, Z. Yuan, P. He, P. Wang, Y. He, G. Zhang, Y. Peng, Advancement of the behavior and regulation at heterogeneous brazing interfaces of TiAl alloys, J. Mater. Sci. 60 (2025) 24103–24128. https://doi.org/10.1007/s10853-025-11795-5 . H. Xue, Y. Liang, H. Peng, Y. Wang, J. 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Cai, S.-D. Zhao, Influences of the near-spherical 3D pore on failure mechanism of atmospheric plasma spraying TBCs using a macro-micro integrated model, Surf. Coat. Technol. 437 (2022) 128375. https://doi.org/10.1016/j.surfcoat.2022.128375 . Z. Yan, Q. Huang, Z. Guo, Y. Song, C. Li, S. Liu, Q. Han, C. Deng, Vacuum plasma sprayed FeAl/Al2O3 functionally graded coatings for fusion reactor applications, Fusion Eng. Des. 85 (2010) 1542–1545. https://doi.org/10.1016/j.fusengdes.2010.04.039 . C. Hua, T. Huang, T. Ma, G. Yin, R. Zhou, C. Li, X. Sun, B. Wang, R. Zhai, P. Song, Based on energy dissipation investigation of preparation and toughening mechanism of continuous transition coating, Surf. Coat. Technol. 465 (2023) 129606. https://doi.org/10.1016/j.surfcoat.2023.129606 . X. He, P. Song, X. Yu, C. Li, T. Huang, Y. Zhou, Q. Li, K. Lü, J. Lü, J. Lu, Evolution of cracks within an Al2O3–40 wt%TiO2/NiCoCrAl gradient coating, Ceram. Int. 44 (2018) 20798–20807. https://doi.org/10.1016/j.ceramint.2018.08.081 . R. Chen, F. Wan, X. He, T. Huang, P. Song, Structural characteristics and evolution in the crack propagation process of NiCrAlY/YSZ gradient coatings through thermal shock at 900°C, Mater. Sci. Technol. 41 (2024) 810–823. https://doi.org/10.1177/02670836241272085 . Y. Yue, G. Song, L. Li, J. Zhao, X. Li, G. Cao, X. Luo, B. Yu, M. Fang, Y. Li, G. Wu, L. Ma, Collaborative improvement of interfacial properties of carbon fiber/epoxy resin composites through modulus/toughness matching and gradient interface, Composites, Part B 298 (2025) 112398. https://doi.org/10.1016/j.compositesb.2025.112398 . J. Zhao, L. Meng, B. Ya, H. Ren, S. Li, B. Zhou, Effect of layered structural parameter on microstructure and mechanical properties of ti/Al3Ti laminated composites, Mater. Charact. 230 (2025) 115818. https://doi.org/10.1016/j.matchar.2025.115818 . M. Li, K. Wang, Q. Guo, X. Tian, Y. Liu, K. Wang, Y. Wang, H. Hou, Z. Xiong, Y. Zhao, Synergistic crack inhibition by lamellar structure and graphene nanoplatelets in 2024 al-GNPs/TC4 laminated metal composites, Mater. Sci. Eng., A 901 (2024) 146347. https://doi.org/10.1016/j.msea.2024.146347 . Additional Declarations No competing interests reported. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-9191760","acceptedTermsAndConditions":true,"allowDirectSubmit":false,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":618516802,"identity":"83c93ed1-ef47-48d5-aeef-55c7f00d3d60","order_by":0,"name":"Wufan Li","email":"","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Wufan","middleName":"","lastName":"Li","suffix":""},{"id":618516805,"identity":"991493cb-faec-4cdb-822e-838ed0cd8d75","order_by":1,"name":"Chen Hua","email":"","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Chen","middleName":"","lastName":"Hua","suffix":""},{"id":618516806,"identity":"8559059e-80f0-4d45-9868-f1307b6ee21d","order_by":2,"name":"Yanhong Zhuang","email":"","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Yanhong","middleName":"","lastName":"Zhuang","suffix":""},{"id":618516808,"identity":"d70eec18-49ce-44b1-8bd0-82a55867285e","order_by":3,"name":"Tao Wang","email":"","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"Tao","middleName":"","lastName":"Wang","suffix":""},{"id":618516810,"identity":"f1a99eab-e7b7-4b8b-9bde-0be3d9d60114","order_by":4,"name":"Taihong Huang","email":"data:image/png;base64,iVBORw0KGgoAAAANSUhEUgAAAZAAAAAyAQMAAABI0h/eAAAABlBMVEX///8AAABVwtN+AAAACXBIWXMAAA7EAAAOxAGVKw4bAAAA5klEQVRIiWNgGAWjYFACxjYQmQBmfzCwsSOogQdZC+OMgrRkIrQwsMG1MPN8OMTYQEiLPfvhtgcfd9TmGRw/e/i1jcEBZgb2w0c34LWFJ7HdcOaZ48UGZ/LSrHMM7vAx8KSl3cDvsMQ2ad62Y4kbDuSYGecYPGNmkOAxw6+F/yFUy/k3ZsYWBocZGwhqkQDbUpO44UaO8WMGorTceNgmObPtQOLMG2/MGHsM0pLZCPmFvT/9mcTHtrrEvvM5xh9+/LGx42c/fAyvFig4DCLYJMAkEcpBoA5EMH8gUvUoGAWjYBSMMAAAg/9O7whcgSIAAAAASUVORK5CYII=","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":true,"prefix":"","firstName":"Taihong","middleName":"","lastName":"Huang","suffix":""},{"id":618516812,"identity":"07e8db98-6870-45c6-8eb2-e373b1f7c8f7","order_by":5,"name":"peng song","email":"","orcid":"","institution":"Kunming University of Science and Technology","correspondingAuthor":false,"prefix":"","firstName":"peng","middleName":"","lastName":"song","suffix":""}],"badges":[],"createdAt":"2026-03-22 14:08:37","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-9191760/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-9191760/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":106323701,"identity":"507bcfbf-a430-468f-904a-1a3a5166a5ce","added_by":"auto","created_at":"2026-04-07 12:43:27","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":397224,"visible":true,"origin":"","legend":"\u003cp\u003e(a) and (d) cross-sectional morphology of MYC-96h, (b) and (e) cross-sectional morphology of MYC-240h, (c) and (f) cross-sectional morphology of MYC-504h.\u003c/p\u003e","description":"","filename":"image1.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/627635a4967cca6d3f49c862.png"},{"id":106323671,"identity":"4e6865f4-0a66-418b-8238-98856ca9c2fb","added_by":"auto","created_at":"2026-04-07 12:43:16","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":366224,"visible":true,"origin":"","legend":"\u003cp\u003eIn-situ diffusion barrier morphology and EDS line scan results of the interface of MYC coating oxidized for 504h.\u003c/p\u003e","description":"","filename":"image2.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/1078be9b031aa522b5f9de35.png"},{"id":106323650,"identity":"20de6837-cbd5-472c-979d-3316691e25f2","added_by":"auto","created_at":"2026-04-07 12:43:13","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":301196,"visible":true,"origin":"","legend":"\u003cp\u003eVickers hardness indentation and crack growth morphology of the single-layer MYC coating after oxidation for different time periods: (a) and (d) 96 h; (b) and (e) 240 h; (c) and (f) 504 h.\u003c/p\u003e","description":"","filename":"image3.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/46232a5735b353fd8b8180be.png"},{"id":106323649,"identity":"25c8cf79-f130-4b08-9569-61f53c8fbcc2","added_by":"auto","created_at":"2026-04-07 12:43:13","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":460028,"visible":true,"origin":"","legend":"\u003cp\u003eSpatial distribution of the elastic modulus (a-c) and hardness (d-f) of MYC coatings after oxidation at different time periods: (a), (d) and (g) nano-indentation test of MYC-96h;(b), (e) and (h) nano-indentation test nano-indentation test of MYC-240h; (c), (f) and (i) nano-indentation test of MYC-504h.\u003c/p\u003e","description":"","filename":"image4.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/f009913b17cc4f17faf57308.png"},{"id":106323697,"identity":"a16fa68f-0821-403d-b3c1-5ec786c8387c","added_by":"auto","created_at":"2026-04-07 12:43:22","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":617586,"visible":true,"origin":"","legend":"\u003cp\u003eThe interface morphology of the gradient region of the MYC/YSZ coating and EDS line scanning.\u003c/p\u003e","description":"","filename":"image5.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/763fc64c18fa572c49a8c7e7.png"},{"id":106323681,"identity":"2de92c77-9bce-40e1-aed4-9e07f79974fb","added_by":"auto","created_at":"2026-04-07 12:43:18","extension":"png","order_by":6,"title":"Figure 6","display":"","copyAsset":false,"role":"figure","size":326484,"visible":true,"origin":"","legend":"\u003cp\u003eCross-sectional morphology of MYC/8YSZ gradient structure coating after thermal shock test: (a) and (c) thermal shock 384 times, (b) and (d) thermal shock 960 times.\u003c/p\u003e","description":"","filename":"image6.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/c751bb3613cc549271a2bc74.png"},{"id":106323700,"identity":"ca14e974-7d0c-4be1-bab6-bf6b66ab455c","added_by":"auto","created_at":"2026-04-07 12:43:27","extension":"png","order_by":7,"title":"Figure 7","display":"","copyAsset":false,"role":"figure","size":647101,"visible":true,"origin":"","legend":"\u003cp\u003e(a) Interface morphology and (b) EDS mapping of MYC/8YSZ gradient structure coating after thermal shock 384 times between coating and substrate. (c) Interface morphology and (d) EDS mapping of MYC/8YSZ gradient structure coating after thermal shock 960 times between coating and substrate.\u003c/p\u003e","description":"","filename":"image7.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/2655c40828b81dccdb135b2d.png"},{"id":106323654,"identity":"70f13f13-3859-4790-b40e-a39f863760ac","added_by":"auto","created_at":"2026-04-07 12:43:13","extension":"png","order_by":8,"title":"Figure 8","display":"","copyAsset":false,"role":"figure","size":217146,"visible":true,"origin":"","legend":"\u003cp\u003eThe propagation morphology of vertical cracks in the cross section of MYC/8YSZ gradient structure coating after 384 thermal shock cycles.\u003c/p\u003e","description":"","filename":"image8.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/4348d4f9cfc510e9ecd5cf5c.png"},{"id":106323672,"identity":"c0fe9f92-fc9c-4f6a-aed1-8587501b95e4","added_by":"auto","created_at":"2026-04-07 12:43:16","extension":"png","order_by":9,"title":"Figure 9","display":"","copyAsset":false,"role":"figure","size":489690,"visible":true,"origin":"","legend":"\u003cp\u003eThe propagation morphology of vertical cracks in the cross section of MYC/8YSZ gradient structure coating after 960 thermal shock cycles.\u003c/p\u003e","description":"","filename":"image9.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/28c14769d9fc2d464e88e264.png"},{"id":106323665,"identity":"77685db9-8cbb-49b5-b529-9d9f8479b0c2","added_by":"auto","created_at":"2026-04-07 12:43:14","extension":"png","order_by":10,"title":"Figure 10","display":"","copyAsset":false,"role":"figure","size":480515,"visible":true,"origin":"","legend":"\u003cp\u003eCross-sectional morphology of MYC/8YSZ gradient structure coating after thermal shock test: thermal shock 960 times.\u003c/p\u003e","description":"","filename":"image10.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/f0072079e1e91dd208e5887b.png"},{"id":106403967,"identity":"5efd53a4-c000-492e-8096-5fdfa9c5dc48","added_by":"auto","created_at":"2026-04-08 09:15:17","extension":"png","order_by":11,"title":"Figure 11","display":"","copyAsset":false,"role":"figure","size":245950,"visible":true,"origin":"","legend":"\u003cp\u003eThe mechanism of root-like crack propagation within the MYC/8YSZ coating.\u003c/p\u003e","description":"","filename":"image11.png","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/6ea4e92fec249220269d37a8.png"},{"id":106405588,"identity":"3b048f15-0a7a-4168-a927-a88201a95913","added_by":"auto","created_at":"2026-04-08 09:27:40","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":5144062,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-9191760/v1/537f8002-d280-4b7e-a924-84501f3a727c.pdf"}],"financialInterests":"No competing interests reported.","formattedTitle":"Root-like crack propagation: Synergistic of gradient structure and diffusion barrier in TBCs on TiAl Alloy","fulltext":[{"header":"1 Introduction","content":"\u003cp\u003eTiAl alloys are lightweight high-temperature structural materials with high specific strength, excellent creep resistance, and oxidation resistance, and they have demonstrated significant application potential in aerospace engines[\u003cspan additionalcitationids=\"CR2 CR3\" citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e]. This intermetallic alloy was first introduced in aeroengines in 2011 (alloy Ti-48A1-2Nb-2Cr, or \u0026lsquo;48-2-2\u0026rsquo;) for the low-pressure turbine (LPT) blades[\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e]. As aircraft engines continue to evolve towards higher thrust-to-weight ratios and lower emissions, higher requirements are placed on high-temperature components[\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e]. Unfortunately, the operating temperatures of most commercially available TiAl alloys are below 750\u0026deg;C[\u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e, \u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e]. Numerous studies have shown that the high-temperature performance of TiAl alloys can be improved by applying protective coatings[\u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e, \u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e]. Thermal barrier coatings (TBCs) play a critical role in safeguarding the integrity of superalloys, significantly affecting the reliability and operational lifespan of essential components[\u003cspan additionalcitationids=\"CR13 CR14\" citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e].8YSZ, as a typical thermal-protective coating material, significantly influences the thermal protection effect of the surface coating on the TiAl substrate, and is a key part of the TBC structure[\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e, \u003cspan citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e].Currently, due to its high deposition efficiency and cost-effectiveness, Atmospheric Plasma Spraying (APS) has become one of the main technologies for preparing TBCs[\u003cspan citationid=\"CR18\" class=\"CitationRef\"\u003e18\u003c/span\u003e]. Since the early 1990s, APS has been extensively utilized for the fabrication of bilayer and gradient coatings. The traditional double-layer 8YSZ TBCs consist of the ceramic coating and MCrAlY bond coating[\u003cspan citationid=\"CR19\" class=\"CitationRef\"\u003e19\u003c/span\u003e]. However, under high-temperature conditions, severe interdiffusion occurs between the MCrAlY coating and the TiAl substrate, leading to the formation of Kirkendall voids[\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]. Furthermore, the difference in composition between the coating and the TiAl alloy causes elemental diffusion near the interface[\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e]. Consequently, interdiffusion zone (IDZ) is brittle, further deteriorating the adhesion at the coating/substrate interface[\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e, \u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e, \u003cspan citationid=\"CR22\" class=\"CitationRef\"\u003e22\u003c/span\u003e]. To mitigate this issue, the introduction of an intermediate layer has been reported as an effective strategy to suppress elemental diffusion at the interface between the coating and the substrate[\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e]. For example, Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e films and Cr\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e have been used as intermediate layers between TiAl and NiCrAlY to inhibit the diffusion of Ni from the NiCrAlY bond coating into the TiAl substrate[\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e, \u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e]. Although the direct preparation of a diffusion barrier layer on TiAl can effectively slow down interdiffusion, it introduces a hetero-phase, which may have potential impacts on the overall long-term stability[\u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e, \u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e]. Therefore, the development of in-situ diffusion barriers has attracted increasing attention as a promising strategy to improve interfacial performance[\u003cspan citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e]. Hua[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e] developed a CoNiCrAlY-Y\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e-Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e (MYC) gradient coating on the surface of TiAl alloy via plasma spraying technology. During the subsequent high-temperature oxidation process, a Ti\u003csub\u003e2\u003c/sub\u003eAlN phase was formed at the interface between the substrate and the bond coating. This in-situ formed phase further enhanced the overall bonding performance.\u003c/p\u003e \u003cp\u003eIn addition to interface weakening caused by diffusion, TBCs on TiAl alloys are also susceptible to thermal stress-induced failure. For the double-layer coating, there are differences in the elastic modulus and thermal expansion coefficient between the metal and ceramic coatings, which makes the thermal barrier coating susceptible to the thermal stress caused by temperature fluctuations during operation. These stresses can initiate and propagate cracks at the coating interface or within the layers, leading to delamination and functional failure[\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e, \u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e]. Previous studies have shown that the bonding strength of the gradient coating is 2.5 times that of the double-layer coating[\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. This is because the gradient coating can achieve a continuous transition of composition and properties, consequently reducing the risk of coating peeling caused by the coefficient of thermal expansion (CTE) mismatch. During the plasma spraying process, by adjusting the spraying power and the particle size of the powder, the layer-by-layer transition of the coating can be achieved, resulting in a structure with different hardness and toughness[\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e]. Hua reported that continuous compositional transitions in gradient coatings result in hardness gradients and enhanced energy dissipation capability[\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e]. In the gradient region, the expansion of cracks will occur along the interface direction of the bond coating, while those cracks extending to the inner layer are more likely to expand laterally towards the interface between the bond coating and the substrate, rather than continuing to extend along the thickness direction to the substrate[\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e]. Chen further demonstrated that under the thermal shock condition of 900\u0026deg;C, the initiation and propagation characteristics of cracks in the gradient coating are closely related to the thickness of the gradient region and sample with the thickness of the gradient region was 150\u0026micro;m exhibited greater thermal shock tolerance[\u003cspan citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e].This paper reports the evolution of cracks in the MYC/YSZ gradient coating prepared by atmospheric plasma spraying. The crack propagation and failure mechanism of this gradient coating under mechanical and thermal stress were studied. Vickers hardness tests and thermal shock tests were conducted on the coating/ substrate specimens to simulate different service times and study the crack propagation process.\u003c/p\u003e"},{"header":"2 Experimental Details","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003e2.1 Base materials\u003c/h2\u003e \u003cp\u003eThe TiAl alloy substrate (nominal composition: Ti 51.4 at. %, Al 43.5 at. %, Nb 4 at. %, Mo 1 at. %, B 0.1 at. %) in this study was produced by Beijing Yanbang New Materials Technology Co., Ltd. with Vacuum Suction Melting Technology (unit type YBXF-3). The substrate was processed into cylindrical block with dimensions of Φ25*5mm using Computerized Numerical Control machine. The specimens were cleaned with anhydrous ethanol in ultrasonic cleaner (DR-MS20, Shenzhen Derui Ultrasonic Equipment Co., Ltd.) to thoroughly remove surface impurities and grease before sandblasting. Then, the surface of the base material was subjected to sandblasting treatment until the surface of the substrate is completely roughened and the roughness is uniform by CS-600 sandblasting machine (Shanghai Jichuan Machinery Technology Co., Ltd.).\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec4\" class=\"Section2\"\u003e \u003ch2\u003e2.2 Coatings\u003c/h2\u003e \u003cp\u003eTwo coatings were applied to the TiAl alloy: CoNiCrAlY-30Y\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e-20Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e Ceramic-metal coating (MYC), CoNiCrAlY-Y\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e-Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e/8YSZ gradient coating (MYC/YSZ). The MYC powder is composed of CoNiCrAlY (nominal composition: Co balance, Ni 32 wt.%, Cr 21 wt.%, Al 8wt.%, Y 0.5 wt.%) Y\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e, and Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e in a mass ratio of 5:3:2. The powder particle size is controlled to be between 15\u0026ndash;45 \u0026micro;m using a sieve. The specific experimental procedures have been described in the previous report[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThe TiAl-MYC and TiAl-MYC/YSZ were prepared by APS technology, the spraying parameters for MYC and MYC/YSZ were described in our previous study[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. TiAl-MYC was obtained by APS, on the grit blasted surface of an TiAl sample. The crucial step in the process involves using a two-way powder feeding method, an alternating and superimposed MYC/YSZ layered structure was formed at the microscopic level, and a gradient transition in composition and structure was achieved at the macroscopic level. The TiAl-MYC/YSZ was fabricated following three steps. First, the MYC coating with a deposition thickness of approximately 50 \u0026micro;m. This not only provides a good interface between the coating and the substrate, but also serves as a carbon source for subsequent interface reactions, helping to form an in-situ diffusion barrier. Then, by dynamically adjusting the powder feeding rates of the two powders (from a single MYC gradually increasing to a single 8YSZ), a gradient region with a thickness of approximately 150 \u0026micro;m is deposited, allowing the composition to gradually transition from the metallic phase to the ceramic phase, avoiding interface stress concentration. After the feeding of MYC powder, 8YSZ surface layer with a thickness of approximately 200 \u0026micro;m is deposited to provide good thermal insulation and thermal shock resistance.\u003c/p\u003e \u003cp\u003eDuring the spraying process, to ensure uniform coating thickness and consistency of microstructure in each area, the spraying path design uses a 3 mm step distance and multiple passes of spraying to avoid structural defects caused by uneven powder deposition density. At the same time, parameters such as the spray gun movement speed, base preheating temperature, plasma gas composition, and total power have also been systematically optimized.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec5\" class=\"Section2\"\u003e \u003ch2\u003e2.3 High-temperature Test\u003c/h2\u003e \u003cp\u003eStatic cyclic oxidation tests were conducted using a box-type electric furnace (SX-B01123, Tianjin Zhonghuan Electric Furnace Co., Ltd.). The samples were held at 900\u0026deg;C in static air for 24 hours, then cooled in air for 30 minutes, which was defined as one cycle. Samples were taken out at 96 hours (4 cycles), 240 hours (10 cycles), and 504 hours (21 cycles). The main focus was on the changes in the coating structure at high temperatures.\u003c/p\u003e \u003cp\u003eThe thermal shock test was also carried out using the same type of box-type electric furnace. The samples were held at 900 ℃ in static air for 15 minutes and then rapidly quenched in water to room temperature, which was defined as one thermal shock cycle. Thermal shock cycles of 384 and 960 times were conducted on the MYC/YSZ gradient coating (lasting for 96 hours and 240 hours respectively). Before the high-temperature test, the temperature of the box-type furnace was calibrated using a multi-channel temperature tester (HPS3032, Changzhou Haierpa Electronic Technology Co., Ltd.). The specific operation was as follows: the furnace was set at 900\u0026deg;C, and the temperature at the center of the furnace cavity was measured using the temperature tester. The furnace setting temperature was adjusted until the measured temperature stabilized at 900\u0026thinsp;\u0026plusmn;\u0026thinsp;1 ℃.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec6\" class=\"Section2\"\u003e \u003ch2\u003e2.4 Mechanics Performance Test and Characterization\u003c/h2\u003e \u003cp\u003eThe Vickers Hardness tests were conducted in a Microhardness tester (DHV-1000Z, Jiangsu Nan Guang Electronic Technology Co., Ltd.) to test the hardness of the coating cross-section. the cold-embedded samples were ground and polished with SiC sandpaper on their cross-sections to ensure the flatness met the testing requirements. The indenter was loaded along the direction from the bond coating to the ceramic coating at different positions of the coating. The loading force of the indenter was set at 0.5N, and the holding time was 10 seconds.\u003c/p\u003e \u003cp\u003eFor the Nanoindentation tests, the maximum applied load in the experiment was 30 mN, the displacement noise level was 0.2 nm, and the thermal drift parameter was less than 0.05 nm/s. In this paper, the nanoindentation tests were all conducted using a Berkovich Triangular conical indenter.\u003c/p\u003e \u003cp\u003eThe microstructure and composition distribution of the coating cross-section and fracture surface were analyzed using a field emission scanning electron microscope (SEM, Tescan-MIRA LMS\u0026thinsp;+\u0026thinsp;Bruker Quantax 200 XFlash 6|6) and the crack propagation was observed.\u003c/p\u003e \u003c/div\u003e"},{"header":"3 Results and discussion","content":"\u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003e3.1 microstructure of coating cross section\u003c/h2\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e shows the SEM image of the cross-sectional morphology of MYC after 96h, 240h and 504h of oxidation at 900 ℃. From Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea to Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ec, no obvious diffusion layer is seen. Furthermore, after prolonged high-temperature exposure, a layer gradually forms and its thickness gradually increases, showing a stable growth trend (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ed-f). This phenomenon reveals that the MYC coating prepared by APS can effectively inhibit the diffusion of metal elements such as Co and Ni to the substrate under high temperature conditions, inhibit the formation of the CoNiTi\u003csub\u003e2\u003c/sub\u003e brittle phase, and thus maintain excellent high-temperature phase stability. During the initial stages of oxidation, high temperature is likely to prompt chemical reactions at the material interface, leading to diffusion barriers formation in situ. Its existence can not only effectively hinder the diffusion of Co and Ni to the substrate, inhibit the formation of intermetallic compounds, but also reduce thermal stress mismatch. Additionally, in an air environment of 900\u0026deg;C, the diffusion barrier continues to generate and grow continuously during the thermal exposure process, showing good high temperature stability. This feature not only helps to improve the coating's ability to hinder element diffusion during long-term service, but also provides an important basis for the subsequent analysis of the evolution and fracture behavior of the interface microstructure, especially the interface mechanical properties. It should be noted that some black particles can be observed at the interface between the coating and the substrate (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb). These particles are corundum (Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e) particles embedded in the surface of the substrate during sandblasting pretreatment.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eRegarding element distribution, Ti, C and N elements in the Layer 1 is highly enriched (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e). This compositional feature suggests that Layer 1 is primarily composed of the Ti\u003csub\u003e2\u003c/sub\u003eCN phase[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. This phase not only forms rapidly during the early heat treatment process but also maintains structural stability during subsequent high-temperature exposure, demonstrating excellent anti-diffusion performance and thermal stability. In contrast, the Layer 2 shows signals of elements such as Ti, Al, and N, suggesting that this region may have generated another phase structure mainly composed of Ti, Al and N, which is most likely Ti\u003csub\u003e2\u003c/sub\u003eAlN[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. This discovery indicates that during the evolution of the in-situ diffusion barrier, a double-layer diffusion barrier composed of Ti\u003csub\u003e2\u003c/sub\u003eCN and Ti\u003csub\u003e2\u003c/sub\u003eAlN was formed[\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec9\" class=\"Section2\"\u003e \u003ch2\u003e3.2 Interface Vickers hardness analysis\u003c/h2\u003e \u003cp\u003eResearchers have found that the length of crack propagation is closely related to fracture toughness and interfacial bonding strength. The interfaces of the diffusion barrier were observed by SEM, Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ea and d shows the interface after 96 hours of oxidation, the thickness of the diffusion barrier is approximately 2 \u0026micro;m. Under cyclic oxidation conditions of 96h and 240h (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb and e), this diffusion barrier gradually thickens with time. However, when the high-temperature oxidation time is extended to 504h (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec and f), the growth trend of the diffusion barrier thickness slows down significantly, which means that the diffusion barrier is close to growth saturation after entering the 504h stage. There was no significant interfacial lateral crack propagation in the interface area, which indicated that a firm bonding interface was formed between the diffusion barrier and the substrate and coating, which could withstand higher local stress concentration. Only when the interface is affected by the hardness indentation head, short longitudinal cracks appear in the vertical direction, and these cracks have limited lengths of expansion (about 3\u0026thinsp;~\u0026thinsp;4\u0026micro;m). In addition, the diffusion barrier is caused by elements in the coating to diffusion into the inside of the substrate. The growth direction is ingrown and will not bring too much growth stress to the interface area. This characteristic means that the interface diffusion barrier has a high fracture toughness and interface combination strength, which can effectively relieve stress concentration under local loading conditions and avoid large-scale expansion of cracks.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec10\" class=\"Section2\"\u003e \u003ch2\u003e3.3 Interface nanoindentation test\u003c/h2\u003e \u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003e, the nanoindentation test results and morphological characteristics of the interface area after 96 hours, 240 hours, and 504 hours of oxidation are presented. Figures a to c and b to f respectively illustrate the three-dimensional distribution maps of the elastic modulus and hardness in this area. After 96h oxidation, the modulus and hardness of the interface area have increased compared to the substrate, but the growth amplitudes of the two are significantly different. The increase in hardness is much greater than that of the modulus. By further observing the nanoindentation morphology shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eg, it can be clearly seen that the indentation size at the interface is significantly smaller than that in the substrate region, indicating a substantial increase in hardness. However, the modulus is relatively low, effectively alleviate stress concentration and inhibit crack initiation and propagation under local loading, thereby improving the interface toughness and stability. As the oxidation time increases, the thickness of the diffusion barrier gradually increases, and the mechanical properties also undergo evolution. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003eb and e, after 240 hours of oxidation, the difference in modulus between the interface and the susbstrate did not increase significantly compared to that at 96 hours, but the increase in hardness was more obvious. The modulus mismatch remained significantly lower than the hardness. This suggests that the diffusion barrier maintains a relatively high hardness while still having moderate modulus matching and good crack resistance. When the heat exposure time was extended to 504 hours, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec and f, the diffusion barrier further thickened, and the modulus difference between the interface and the substrate increased. However, the modulus growth was not significant, and the mismatch amplitude was smaller than the hardness, maintaining a good elastic adjustment function. This helps to alleviate the residual stress at the interface under high-temperature cyclic loads and avoids interface peeling and cracking[\u003cspan citationid=\"CR34\" class=\"CitationRef\"\u003e34\u003c/span\u003e]. The hardness of the interface area has significantly increased (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ed-f), which ensures that the interface layer can effectively resist mechanical loads in the high-temperature oxidation and thermal cycling environment, maintaining the overall integrity of the coating/substrate system. The significant increase in hardness also inhibited the initiation of microcracks, fundamentally extending the service life. On the other hand, from the modulus perspective, although the elastic modulus of the diffusion barrier also increased with the extension of the oxidation time, its relative modulus mismatch with the substrate was still significantly lower than the hardness mismatch, indicating that the low modulus characteristic represents that the diffusion barrier has a certain deformation ability, which can alleviate the thermal physical property differences between MYC and the substrate, improve the damage tolerance, and this high hardness and low modulus combination is very crucial for interface performance. The low modulus mismatch between the Ti\u003csub\u003e2\u003c/sub\u003eCN and Ti\u003csub\u003e2\u003c/sub\u003eAlN double-layer diffusion barrier and the substrate is crucial as it acts as an \"energy transition buffer layer\" at the interface. The existence of the Ti\u003csub\u003e2\u003c/sub\u003eCN and Ti\u003csub\u003e2\u003c/sub\u003eAlN double-layer diffusion barrier can form a gradually decaying and transitional stress buffer zone in this region, relieve the thermal physical properties differences between the coating and the substrate, thereby effectively reducing residual stress and potential cracking risks caused by thermal mismatch. Although the MYC coating has solved the problem of interfacial mutual diffusion, the double-layer coating in practical applications still has the issue of poor adhesion.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eInspired by this gradient structure, by combining MYC with YSZ, a gradient structure was designed, the element distribution in the initial gradient region is shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e. The distribution of Ni and Cr is completely opposite to that of Zr. The outer layer is observed to the YSZ thermal protection coating, while within the gradient region, both CoNiCrAlY and YSZ are present.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec11\" class=\"Section2\"\u003e \u003ch2\u003e3.4 Thermal shock test\u003c/h2\u003e \u003cp\u003eAfter 384 thermal shock cycles (Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003ea and c), no transverse cracks were observed on the cross-section of the coating, and the ceramic coating and bond coating maintained a good adhesion. This demonstrates that the gradient structure can maintain high structural stability and crack resistance under multiple alternating cold and hot impacts. However, when the thermal shock cycle number increased to 960 (Fig.\u0026nbsp;\u003cspan refid=\"Fig6\" class=\"InternalRef\"\u003e6\u003c/span\u003eb and d), a large number of transverse micro-cracks appeared in the gradient region, and some vertical cracks formed step-like cracks in different directions with the transverse cracks. Most of the cracks stopped expanding in the ceramic layer. The deflection and bridging of the cracks increased the energy required for their expansion, improve the overall toughness of the coating significantly.\u003c/p\u003e \u003cp\u003eFrom the mechanism perspective, introducing an interface between the CoNiCrAlY, Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e and YSZ phase leads to a certain degree of thermal expansion coefficient mismatch. During the thermal shock test, the vertical cracks first extend from the surface of the ceramic coating and extend towards the interface between the ceramic coating and the gradient region. When the vertical crack is far away from the crack band structure, the expansion of the horizontal cracks within the crack band will change the distribution of internal stress, thereby exerting an inhibitory effect on the downward extension of the vertical crack. When the vertical cracks extend towards the horizontal cracks until they merge, the energy in the coating is released, and finally, transverse cracks parallel to the interface are formed, which can prevent the surface vertical cracks from extending to the interface between the coating and the substrate.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eAs shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ea and b, the morphology of the interface between the coating and the substrate and EDS mapping after 384 thermal shock cycles are presented. It can be observed that in the thermal shock environment, the formation of the in-situ diffusion barrier is not affected. It forms a good bond between the coating and the substrate, and no signs of cracking or peeling such as cracks are observed at the interface. Through EDS analysis of Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003eb, a very thin Al\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e layer can be seen above the diffusion barrier, while the Ti\u003csub\u003e2\u003c/sub\u003eCN diffusion barrier has been formed. Additionally, a small amount of Ni and Co diffusion products appear below the diffusion barrier, but Ti has not diffused into the coating. After the thermal shock test was conducted for 960 times, the interface structure remained intact, as shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ec and d. Through EDS analysis of Fig.\u0026nbsp;\u003cspan refid=\"Fig7\" class=\"InternalRef\"\u003e7\u003c/span\u003ed, the composition of the diffusion barrier has shown a layer structure namely Ti\u003csub\u003e2\u003c/sub\u003eCN on the upper layer and Ti\u003csub\u003e2\u003c/sub\u003eAlN on the lower layer. Particularly, Ni and Co diffusion products were detected in the interface area. The analysis suggests that part of the reason lies in the extreme temperature cycling and rapid cooling and heating alternation in the water quench thermal shock test, which causes the formation of the diffusion barrier to be constantly starting and stopping. Furthermore, the thermal cycling process will result in a significant temperature gradient and thermal expansion differences. These differences introduce a pronounced multi-directional stress field within the coating. This stress state is extremely complex, including both thermal stress and the high strain rate cyclic loading caused by thermal shock, causing the diffusion barrier to be repeatedly subjected to tension and compression, shear, and thermal fatigue. As a result, vertical cracks form at the interface. These minor cracks observably enhance the permeability and the number of diffusion channels in the interface area, promoting the penetration of Ni and Co through the in-situ diffusion barrier and entering the TiAl substrate to form a diffusion layer.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig8\" class=\"InternalRef\"\u003e8\u003c/span\u003e shows the cross-sectional morphologies of the MYC/8YSZ coating after 384 and 960 thermal shock cycles. The results indicate that vertical cracks appeared at the top of the ceramic coating and the gradient region. After 384 thermal shock cycles, vertical cracks appeared at the top of the ceramic layer, but the cracks were blocked when they extended to the top of the gradient region. Crack is always under compressive stress, preventing it from extending downward.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eHowever, when the thermal shock cycles increased to 960 (Fig.\u0026nbsp;\u003cspan refid=\"Fig9\" class=\"InternalRef\"\u003e9\u003c/span\u003e), vertical cracks in the ceramic layer extended downward from the top. When they reached the gradient region, they were hindered by the metallic phase, causing the cracks to deflect. Some of them formed horizontal cracks, while others continued to extend downward. There are intrinsic thermal physical property parameters (such as thermal expansion coefficient, elastic modulus, fracture toughness, etc.) differences among different phases within the gradient region. This results in the cracks not extending in the same direction but instead preferentially expanding along the weak interface with the lowest local fracture energy, demonstrating the characteristic of selective fracture. These weak interfaces are not uniformly distributed in the gradient region, they exhibit a certain degree of randomness and gradient characteristics. This distribution pattern leads to multiple deflections, splits, and branching of the crack propagation path, thereby delaying the formation and propagation of the through-crack. The further propagation of the crack requires a greater amount of energy to drive, thus preventing the coating from peeling off.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig10\" class=\"InternalRef\"\u003e10\u003c/span\u003e shows the cross-sectional morphology and EDS mapping after 960 thermal shock tests of the coating. It is unable to provide sufficient energy to sustain crack propagation. The horizontal cracks extend to the CoNiCrAlY metallic phase and no longer grow. This indicates that the YSZ ceramic phase has a synergistic blocking effect on crack propagation with CoNiCrAlY, and the local stress generated by the ceramic phase can be effectively alleviated through the stress transfer between adjacent metal phases, thereby achieving effective fracture energy transfer between the phases[\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e, \u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e]. The compressive stress of the CoNiCrAlY metallic phase has an effect of \"bridging\" and \"passivation\" on the cracks tip[\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e], inhibiting the driving force for the cracks to penetrate the metallic phase. Only when the energy accumulates to a sufficient level will the cracks continue to deflect and extend through the metal substrate. Only when the energy accumulates to a sufficiently high level or encounters other cracks will the cracks continue to deform and extend through the metal substrate. Such cracks do not exist within the gradient region of 384 thermal shocks, and the cracks in the 960 thermal shock case are also few. The interface bonding of the coating was relatively strong, the crack propagation mechanism is matched, and there is no obvious stress concentration[\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec12\" class=\"Section2\"\u003e \u003ch2\u003e3.5 Crack Propagation Mechanism and Failure Model under Thermal Shock\u003c/h2\u003e \u003cp\u003eUnder thermal shock conditions, the crack propagation process is root-like, that is, the crack starts to grow from the external ceramic coating and then extends along the interior of the coating. It delays failure through multiple bends and branching. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ea to d, the initial initiation of the crack usually occurs in the ceramic layer, which has a low fracture toughness and is prone to crack initiation at weak interfaces. At this time, the cracks mainly extend along the surface of the ceramic coating, usually manifesting as short vertical micro-cracks. The initial propagation of the crack is usually confined to the upper ceramic region and is not likely to rapidly reach the ceramic bottom layer. As the number of thermal shock impact cycles increases, the surface-initiated cracks gradually extend to the gradient coating (Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ee and g). For the traditional double-layer coating, when cracks propagate at the interface between the coating and the bonding layer, they will deflect along the interface and form transverse cracks, resulting in the detachment of the coating. In the gradient region, however, due to the presence of the metal phase, the macroscopic mechanical properties of this area can be adjusted to exhibit a gradient change, thereby prolonging the crack propagation path and increasing energy dissipation. In detail, the adjacent phases have different Young\u0026rsquo;s modulus and thermal expansion coefficients, which lead to a complex stress field at the crack tip, thus causing the crack to exhibit a root-like branching and deflection (Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003eg). When the crack contacts the metal phase, due to the higher hardness of the metal phase, the crack will deviate along the metal-ceramic interface or temporarily stop. With the increase in the number of thermal shock impacts, the crack then extends along the metal phase. When the driving force is strong enough, the crack can break through the obstruction of the metal phase and continue to extend (Fig.\u0026nbsp;\u003cspan refid=\"Fig11\" class=\"InternalRef\"\u003e11\u003c/span\u003ef and h). This repeated deflection and stopping increase the possible extension path length of the crack and also increase the energy required for crack propagation. The discontinuous metal layer can convert the main crack into multiple dispersed micro-cracks, effectively preventing the propagation of through-cracks, significantly enhancing the coating's toughness and thermal shock resistance. In addition, cracks can form bridging bands in the gradient region. These bridging bands reduce the stress at the crack tip. During the repeated thermal shock cycles, the plastic deformation of the bridging bands leads to continuous energy dissipation, significantly extending the crack propagation path and dispersed the driving force for crack propagation.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e"},{"header":"4 Conclusion","content":"\u003cp\u003eIn this study, the mechanical properties of the MYC coating after 96 hours, 240 hours, and 504 hours of high-temperature oxidation were tested and analyzed through scanning electron microscopy observations and EDS spectra. Through 384 and 960 thermal shock tests, the crack propagation mechanism of the MYC/YSZ gradient coating under high-temperature oxidation conditions was revealed. The main conclusions are as follows:\u003c/p\u003e \u003cp\u003eThe MYC coating and the interface underwent in-situ reactions, forming a dense double-layer diffusion barrier composed of Ti\u003csub\u003e2\u003c/sub\u003eCN and Ti\u003csub\u003e2\u003c/sub\u003eAlN. The Vickers hardness test and nanoindentation experiments indicated that this low modulus and high hardness diffusion barrier played the role of an \"energy buffer layer\" at the interface, effectively reducing residual stress and the potential crack risk caused by thermal mismatch.\u003c/p\u003e \u003cp\u003eAfter 384 thermal shock cycles, no horizontal cracks appeared in the MYC/YSZ gradient coating, while when the number of thermal shock cycles increased to 960, layer-like crack bands appeared in the gradient zone. The deflection and bridging of cracks increased the energy required for their expansion, significantly improving their toughness.\u003c/p\u003e \u003cp\u003eThe crack propagation process within the MYC/YSZ coating is a root-like mechanism. This root-like mechanism effectively prolongs the possible crack propagation path and enhances the thermal shock resistance of the coating. The component gradient effectively alleviated the stress concentration at the ceramic/metal interface, causing the cracks to exhibit branching and bending characteristics, continuously consuming the fracture energy and effectively preventing the expansion of through-cracks, significantly improving the toughness of the coating.\u003c/p\u003e"},{"header":"Declarations","content":"\u003ch2\u003eAuthor Contribution\u003c/h2\u003e\u003cp\u003eWufan Li: Conceptualization, Investigation, Methodology, Validation, Writing-original draft.Chen Hua: Data curation, Project administration, Supervision, Writing-review and editing.Yanhong Zhuang: Data curation, Software, Validation.Tao Wang: Methodology, Visualization.Taihong Huang: Data curation, Project administration, Funding acquisition, Supervision, Writing-review and editing. Peng Song: Funding acquisition, Project administration, Supervision.\u003c/p\u003e\u003ch2\u003eAcknowledgement\u003c/h2\u003e\u003cp\u003eI would like to express my heartfelt gratitude to supervisor Huang Taohong and senior brother Hua Chen. From the selection of the topic, the framework construction, to the revision and improvement, they all provided me with patient guidance and professional suggestions, enabling me to successfully complete the research. Meanwhile, I would like to thank team member of research program for helping and supporting each other during the learning and writing process.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\u003cli\u003e\u003cspan\u003eY. Lu, X. Wang, Z. Yuan, P. He, P. Wang, Y. He, G. Zhang, Y. Peng, Advancement of the behavior and regulation at heterogeneous brazing interfaces of TiAl alloys, J. Mater. 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Eng., A 901 (2024) 146347. \u003cspan class=\"ExternalRef\"\u003e\u003cspan class=\"RefSource\"\u003ehttps://doi.org/10.1016/j.msea.2024.146347\u003c/span\u003e\u003cspan address=\"10.1016/j.msea.2024.146347\" targettype=\"DOI\" class=\"RefTarget\"\u003e\u003c/span\u003e\u003c/span\u003e.\u003c/span\u003e\u003c/li\u003e\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":false,"hideJournal":false,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"surface-science-and-technology","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"Learn more about [Surface Science and Technology](https://link.springer.com/journal/44251)","snPcode":"44251","submissionUrl":"https://submission.springernature.com/new-submission/44251/3","title":"Surface Science and Technology","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"stoa","reportingPortfolio":"Springer Open","inReviewEnabled":true,"inReviewRevisionsEnabled":true},"keywords":"energy buffer layer, gradient structure, thermal shock, root-like crack","lastPublishedDoi":"10.21203/rs.3.rs-9191760/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-9191760/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eIn this study, MYC (CoNiCrAlY-Y\u003csub\u003e2\u003c/sub\u003eO\u003csub\u003e3\u003c/sub\u003e-Cr\u003csub\u003e3\u003c/sub\u003eC\u003csub\u003e2\u003c/sub\u003e) coatings and MYC/8YSZ gradient coatings were fabricated by atmospheric plasma spraying (APS) technology. During the oxidation cycle of the MYC coating, an in-situ diffusion barrier was formed at the interface between the coating and the substrate. The Vickers hardness test and nanoindentation experiments verified that this low-modulus, high-hardness diffusion barrier structure played the role of an \"energy buffer layer\". In the 900\u0026deg;C water quenching thermal shock test of the MYC/8YSZ gradient coating, after 960 cycles, it was found that the gradient structure extended the crack propagation path, improved the energy dissipation capacity of the coating, and formed root-like cracks inside the coating, which increased the thermal shock cycle life of the coating. The coating did not experience failure. This study demonstrated the advantages of gradient structures in improving coating performance, especially in terms of crack resistance under thermal shock conditions.\u003c/p\u003e","manuscriptTitle":"Root-like crack propagation: Synergistic of gradient structure and diffusion barrier in TBCs on TiAl Alloy","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2026-04-07 12:43:01","doi":"10.21203/rs.3.rs-9191760/v1","editorialEvents":[{"type":"communityComments","content":0},{"type":"decision","content":"Revision requested","date":"2026-05-06T07:02:37+00:00","index":"","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2026-04-29T22:59:44+00:00","index":"hide","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2026-04-28T06:01:00+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"335521863427885123089380835182716745875","date":"2026-04-26T16:13:34+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"331496346054369164795041132536229910149","date":"2026-04-25T04:38:40+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"310171825993128695898605081548921111749","date":"2026-04-24T03:18:27+00:00","index":"hide","fulltext":""},{"type":"reviewerAgreed","content":"202678531139225334559146215004062659674","date":"2026-04-07T01:59:55+00:00","index":"hide","fulltext":""},{"type":"reviewersInvited","content":"","date":"2026-04-02T02:26:41+00:00","index":"","fulltext":""},{"type":"editorAssigned","content":"","date":"2026-03-31T05:50:16+00:00","index":"","fulltext":""},{"type":"checksComplete","content":"","date":"2026-03-31T05:49:30+00:00","index":"","fulltext":""},{"type":"submitted","content":"Surface Science and Technology","date":"2026-03-22T13:59:00+00:00","index":"","fulltext":""}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"surface-science-and-technology","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":false,"externalIdentity":"","sideBox":"Learn more about [Surface Science and Technology](https://link.springer.com/journal/44251)","snPcode":"44251","submissionUrl":"https://submission.springernature.com/new-submission/44251/3","title":"Surface Science and Technology","twitterHandle":"","acdcEnabled":true,"dfaEnabled":true,"editorialSystem":"stoa","reportingPortfolio":"Springer Open","inReviewEnabled":true,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"2b8f4432-23d3-45b5-b155-faba87281d56","owner":[],"postedDate":"April 7th, 2026","published":true,"recentEditorialEvents":[{"type":"decision","content":"Revision requested","date":"2026-05-06T07:02:37+00:00","index":"","fulltext":""},{"type":"editorInvitedReview","content":"","date":"2026-04-29T22:59:44+00:00","index":29,"fulltext":""}],"rejectedJournal":[],"revision":"","amendment":"","status":"in-revision","subjectAreas":[],"tags":[],"updatedAt":"2026-05-06T07:10:48+00:00","versionOfRecord":[],"versionCreatedAt":"2026-04-07 12:43:01","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-9191760","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-9191760","identity":"rs-9191760","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}