Multi-element collaboration in Cr₂TiAl_1−xSi_xC₂ MAX for the oxide barrier formation in a 550 °C LBE environment

Abstract

MAX phase ceramics have good oxidation and corrosion resistance, while the controllable formation of the protective oxide layer is still a challenge for their application in lead-cooled fast reactors. Herein, the oxide barrier formation mechanism of three Cr₂TiAl_1−xSi_xC₂ MAXs (x = 0, 0.2, 0.4) with different Si contents exposed to lead-bismuth eutectic (LBE) at 550 °C for 500–3000 hours was systematically studied. Based on the evolution of the structure and composition of the oxide layer, a diffusion rate-controlled oxide layer formation mechanism is proposed. The collaboration of Cr, Al, and Si regulates the structure and composition of the oxide layer. Particularly, the existence of Si facilitates the transport of Al across the oxide layer, which builds a continuous oxide double-layer with mainly Al₂O₃ as the outer layer in LBE. Eventually, the multi-component oxide layer regulates the corrosion resistance of the Cr₂TiAl_1−xSi_xC₂ MAXs, which paves the way for the performance-oriented design of the MAX phase.

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Cheng-Feng Du

Cheng-Feng Du is currently a distinguished research fellow in the School of Materials Science and Engineering at Northwestern Polytechnical University (NPU). He received his PhD degree in 2016 from the Fujian Institute of Research on the Structure of Matter, Chinese Academy of Sciences (CAS), and then worked at Nanyang Technological University (NTU), as a research fellow from 2016 to 2018. In 2018, he joined NPU, affiliated with the Center of Advanced Lubrication and Seal Materials. His current research interests focus on the synthesis, structure, and failure behaviors of MAX and MXene-based materials.

1. Introduction

The lead-cooled fast reactor (LFR) is considered to be one of the most promising reactors due to its advantages of safety and efficiency,^1,2 and uses lead-bismuth eutectic (LBE) as a coolant. The LBE has good neutron properties, excellent radiation damage resistance, excellent heat transfer properties, a low melting point, and a high boiling point.^3,4 However, LBE usually gives rise to complex corrosion effects on cladding structural materials at operating temperatures (550 °C or higher), such as dissolution and erosion of the component elements.^5–8 For instance, alloys including stainless steel,^9,10 FeCrAl,^11–13 and HEAs,^14,15 have excellent machinability and damage tolerance, and mainly suffered from dissolution and oxidative corrosion in LBE. In contrast, although ceramics like Al₂O₃,^16,17 SiC,¹⁸ and AlTiN,¹⁹ are less impacted by the LBE, their intrinsic brittleness makes them hard to solely act as the structural material. Therefore, it is particularly important to develop novel structural materials with integrated advantages from alloys and ceramics.

Recently, MAX phase materials have attracted a lot of attention for industrial applications due to their good mechanical properties, high-temperature oxidation resistance, and corrosion resistance.^20–22 The formula of a ternary MAX phase material is M_n+1AX_n (n = 1–3), in which M is an early transition metal (Ti, Cr, Zr, etc.), A is mainly IIIA or IVA elements (Al, Si, Sn, etc.), and X can be C or N.^23–25 The crystal structure of the MAX phase can be described as the interweaving of different MX layers with A atomic layers.^26,27 The unique layered structure and bonding characters thus endow the MAX phase with properties from both alloys and ceramics.²⁸ Meanwhile, the weakly bonded A atoms and MX layers make the A atoms easy to extract and react, which plays a role in oxidation and corrosion resistance of the material.^29,30 At present, the research of MAX in anti-LBE corrosion mainly focuses on ternary MAXs, such as Cr₂AlC,³¹ Zr₂AlC,³² Ti₂AlC,³³ Ti₃AlC₂,³⁴ Ti₃SiC₂,³⁵etc.³⁶ In these pioneering studies, the capability to form protective oxides (e.g., Al₂O₃, Cr₂O₃, and SiO₂) during corrosion is revealed to be critical for the corrosion resistance of MAXs,³¹ which is mainly beneficial due to the dense and continuous structure of the above oxides.^37,38 Therefore, the incorporation of these elements in the MAX phase lattice is believed a promising way to achieve good LBE corrosion resistance.

On the other hand, among the currently studied MAXs, the weaker Cr–Al bonding than Cr–C and the “third element effect” of Cr endows Cr₂AlC with a preferential role in anti-LBE corrosion compared to the other MAXs.³¹ However, given the sluggish growth rate of Al₂O₃, the protective layer is the thinnest on Cr₂AlC as well, which is vulnerable to the threat of accidental scratches. Interestingly, although the Ti-containing MAXs deliver a thick and quick-build TiO₂ layer at the early stage of LBE corrosion, the higher diffusion rate of Al and Cr atoms in the TiO₂ lattice endows the MAXs with potential capability of Al₂O₃ formation.^39,40 Thus, the introduction of Ti to collaborate with Cr₂AlC is promising for improving the Al₂O₃ formability of the MAX. Fortunately, a recently developed quaternary MAX, Cr₂TiAlC₂, is demonstrated to consist of Cr₂AlC and TiC components,⁴¹ which is an ideal precursor for the synergism of Ti with Cr₂AlC. Moreover, the cooperation of Al and Si on the LBE corrosion resistance of the Ti₃Al(Si)C₂ system was demonstrated very recently as well.³⁰ The existence of Si can facilitate the formation of a protective Al₂O₃ layer at 550 °C in oxygen-poor LBE, which proved the feasibility of the multi-component design of MAX again for an improved anti-LBE corrosion performance.

In this study, by integrating Ti, Cr, Al, and Si in a single MAX phase, three Cr₂TiAl_1−xSi_xC₂ (x = 0, 0.2, 0.4) MAXs with different Si contents were prepared by hot pressing sintering. The static corrosion behaviour of the MAXs in LBE at 550 °C for 500–3000 h was systematically studied. The surface morphology, microstructure, and chemical composition of the oxide layer after corrosion were characterized. The effects of synergistic regulation of the composition and structure of the oxide layer on improving the anti-LBE corrosion performance of MAX were explored, which provided more possibilities for the development of multi-component anti-LBE materials.

2. Experimental

2.1. Materials

The purity of Cr, Ti, Al, and Si elemental powder is 99.9%, all purchased from Qinghe Benyu Metal Materials Co., Ltd, graphite (C) powder (99.9%) was purchased from Alfa Aesar, and LBE ingots with Pb : Bi in a weight ratio of 45 : 55 were purchased from Dongguan Wochang Metal Products Co., Ltd. All powders are used as received.

2.2. Preparation of Cr₂TiAl_1−xSi_xC₂ (x = 0–0.4) MAXs

To synthesize Cr₂TiAl_1−xSi_xC₂ MAXs, Cr, Ti, Al, Si and C powders were first mixed by ball milling according to the molar ratio of 2 : 1 : 1.2 − x : x : 2. Particularly, 20% excess of Al was added as an auxiliary reagent to compensate for the high-temperature volatilization of Al. The mixed powder was hot pressed at 1400 °C and 10 MPa for 3 h. Ar protection is used during ball milling and hot pressing. The sintered sample was machined into a square sample with a size of about 15 mm × 4 mm × 3 mm and polished to 2000^# with SiC paper. The resulting specimens were ultrasonically treated with ethanol and then dried in air.

2.3. Characterization

All X-ray diffraction (XRD, Bruker D8 ADVANCE) patterns were recorded with Cu Kα radiation (λ = 1.5404 Å) in the 2θ range of 5°–120°, accompanied by a step size of 0.02°. GSAS II software was adopted for further Rietveld refinement of the XRD patterns.^42,43 The standard crystallographic information files (.cif) of different phases for XRD refinement were downloaded from Springer Materials, which were used to calculate the reference XRD patterns. The morphology of pristine and corroded surfaces was investigated via a scanning electron microscope (SEM, FEI Helios G4 CX), and the corresponding elemental distribution was probed by energy dispersive X-ray energy spectrometry (EDS). The microhardness of all samples was measured with a Vickers diamond indenter at a constant load of 500 gf and residence time of 5 s, respectively. The test was repeated at least 10 times for each sample. The molar ratios of dissolved Cr, Al, and Si elements in post-reacted LBE were determined by inductively coupled plasma emission spectrometry (ICP-OES, Agilent 5110).

2.4. Static liquid LBE corrosion tests

LBE corrosion was carried out in a sealed quartz tube. The experimental method was based on the previous report.¹⁴ Briefly, the polished MAX specimen is individually packaged with approximately 11 g LBE in a quartz tube with an inner diameter of 11 mm. The LBE ingot was cut into strips with a size of about 10 mm × 5 mm × 50 mm and polished to remove the surface oxides. The LBE strips were stored in an Ar-filled glove box with an oxygen and water content of less than 0.1 ppm, and further sub-packaged in a quartz tube before the test. The quartz tube is vacuumed to an internal pressure of about −1 bar and sealed with an oxyhydrogen flame. In order to simulate the corrosion of the steam generator heat exchange tube in the LFR system, static corrosion tests were carried out at 550 °C for 500, 1000, 2000, and 3000 h in a KSL-1200X box furnace. Based on the capacity of the quartz tube, the vacuum measurement during the experiment, and the Orlov relation,^44–46 the oxygen content dissolved in LBE at 550 °C is calculated to be in the range between 8.6 × 10⁻⁹ wt% (based on the vacuum volume) and 1.6 × 10⁻³ wt% (maximum solubility).

3. Results and discussion

3.1. Phase composition and lattice structure of the MAXs before and after LBE corrosion

To determine the phase composition and lattice structure of Cr₂TiAl_1−xSi_xC₂ MAXs before and after corrosion, XRD refinement was performed using GSASII software. The results are shown in Table S1 and Fig. S1.† As can be seen, the Cr₂TiAlC₂ phase (312 phase) is the main component of the specimens, among which the purity of Cr₂TiAlC₂ is 91.3 wt% (Fig. S1a†), Cr₂TiAl_0.8Si_0.2C₂ is 92.6 wt% (Fig. S1c†), and Cr₂TiAl_0.6Si_0.4C₂ is 80.0 wt% (Fig. S1e†). When the additional amount of Si increases to 0.4, the purity of the MAX phase decreases significantly, with 15.5 wt% Cr₅Si₃C and 4.5 wt% TiC as impurities.

After the corrosion of Cr₂TiAlC₂ in LBE for 3000 h, in addition to the LBE attached to the surface and the impurity phase contained in the material, the oxides on the corrosion surface are mainly alumina and a small amount of titania (Fig. S1b†). Compared with the phase composition before corrosion, it can be found that the contents of the MAX phase and intermetallic compound Cr₅Al₈ are significantly reduced, and finally converted to oxides during corrosion. For Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.4C₂, the oxides on the corroded surface are mainly alumina, silica, and titania (Fig. S1d and f†), whereas the content of the MAX phase and impurity Cr₅Si₃C decreased. Moreover, based on the phase composition of the three samples, the generation of TiC after LBE corrosion is detected.

With the increase in Si content, the a- and c-axes of the MAXs show a gradually decreasing trend (Table S2†). Correspondingly, the changes in bond lengths obtained by XRD refinement were summarized in Table S3.† Further calculation of bond lengths reveals that the shortening of the lattice parameter c is mainly due to the shorter length of the Cr–Si bond than that of the Cr–Al bond. In Cr₂TiAlC₂, the lattice parameter a is mainly determined by the length of metal–C bonds, while the c parameter is mainly controlled by the Cr–Al bond. Upon Si replacement, the Cr–Al bond is impacted, thus affecting mainly the lattice constant c. With the increased content of Si, the length of the Cr–A bond is gradually shortened, resulting in the shortening of the c-axis.

The changes in lattice parameters of Cr₂TiAl_1−xSi_xC₂ after LBE corrosion were determined as well, and are shown in Table S4† and illustrated in Fig. 1a and b. After 3000 h corrosion in LBE, the a- and c-axes of Cr₂TiAlC₂ expand by 0.033% and 0.030%, respectively. When the Si addition amount is 0.2, the LBE corrosion-induced expansion of the c-axis is significantly greater than that of the a-axis, which is 0.469% for the c-axis and 0.106% for the a-axis. However, after 3000 h corrosion of Cr₂TiAl_0.6Si_0.4C₂, although the difference between the expansion of the a- and c-axes is about 4 times (0.012% vs. 0.002%), the overall increment is significantly reduced compared with that of Cr₂TiAl_0.8Si_0.2C₂. Therefore, the level of Si doping might change the behaviour of the A-layer against LBE corrosion.

Fig. 1 (a) Parameters of the a-axis for the three MAXs before and after 3000 h LBE corrosion, (b) parameters of the c-axis for the three MAXs before and after 3000 h LBE corrosion, and (c) bond length in Cr₂TiAl_1−xSi_xC₂ based on the refined lattice structure before and after 3000 h corrosion; the solid line is the change in bond length before corrosion. The dashed line shows the change in bond length after 3000 h of corrosion. (d) Schematic illustration of the lattice variation after LBE corrosion according to (c).

The changed bond lengths after corrosion are shown in Table S5† and Fig. 1c. After corrosion, the Ti–C bond in Cr₂TiAlC₂, Cr₂TiAl_0.8Si_0.2C₂, and Cr₂TiAl_0.6Si_0.4C₂ elongates by 1.59%, 2.24% and 1.87%, respectively. The elongation of the Ti–C bond in Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.4C₂ is higher than that in Cr₂TiAlC₂, whereas for the Cr–A and Cr–C bonds, the bond lengths are all contracted. Fig. 1d schematically illustrates the lattice variation after LBE corrosion according to the XRD refinements. Thermodynamically, the corrosion and decomposition of Cr₂TiAlC₂ in LBE can be explained by the competitive formation of the competing phases.⁴⁷ Generally, the stable competing phases for traditional ternary MAX are dissolved A elements and the corresponding MC carbides. However, due to the heterogeneous M–C bonds in Cr₂TiAlC₂ (Ti–C and Cr–C bonds), the thermodynamically preferential competing phases for its decomposition are Cr₂AlC and TiC before the extraction of A elements, which has been demonstrated in a recent study under high-temperature oxidation conditions.⁴¹ Therefore, by combining the XRD bond length results and the existing theories, the elongation of the Ti–C bond in the Cr₂TiAl_1−xSi_xC₂ MAXs may be the result of the decomposition of Cr₂TiAl_1−xSi_xC₂ into Cr₂Al_1−xSi_xC and TiC during the corrosion in LBE.

3.2. Surface and cross-section morphology after corrosion

The morphology and element distribution of the corroded surfaces from the three MAXs under different corrosion times were analysed by SEM. Fig. 2 and S2–S7† show the morphology and corresponding EDS analyses of the corroded surfaces from Cr₂TiAl_1−xSi_xC₂ MAXs at 550 °C in LBE for 500, 1000, 2000, and 3000 h.

Fig. 2 Enlarged surface morphology of Cr₂TiAl_1−xSi_xC₂ MAXs after different corrosion times in LBE at 550 °C: (a)–(d) Cr₂TiAlC₂, (e)–(h) Cr₂TiAl_0.8Si_0.2C₂, and (i)–(l) Cr₂TiAl_0.6Si_0.4C₂.

As can be seen from Fig. 2a and S2a,† after Cr₂TiAlC₂ was corroded in LBE for 500 h, bright and dark contrast on the corroded surface were observed, in which the bright area is attributed to residual LBE, and the enlarged SEM image shows that square particles with a regular shape and length of about 50 nm are distributed on the surface. According to the corresponding EDS distribution diagram (Fig. S3a†), the distribution of Cr, Ti, Al, C, O, Pb, and Bi elements is relatively uniform, which may be related to the fine oxide particles. After 1000–2000 h of corrosion, the surface morphology of corroded Cr₂TiAlC₂ is still covered by a granular oxide layer with even distribution of all elements (Fig. S2 and S3†). After 3000 h corrosion, the square particles on the corroded surface evolve into spherical nanoparticles. Meanwhile, a small number of flakes appear (Fig. 2d). The EDS diagram in Fig. S3d† confirms the enrichment of Pb and Bi in the bright spheres with a diameter of about 200 nm, which can be ascribed to residual LBE.

In order to further clarify the interaction of MAX in LBE, the cross-section of corroded Cr₂TiAlC₂ was analysed by SEM and EDS. Fig. 3 shows the cross-section of Cr₂TiAlC₂ after different exposure times and the corresponding EDS diagram. As shown in Fig. 3a, after 500 h of corrosion, an oxide layer of about 2.0 μm is formed on the surface of Cr₂TiAlC₂ (within the yellow dashed line range). Determined from the EDS element distribution and line scan diagram, the oxide layer mainly consisted of Cr and Ti, mixed with a small amount of Al-oxides. After 1000 h exposure, as shown in Fig. 3b, based on the elemental distribution, a double-layer structure with an outer oxide layer of about 5.1 μm and an inner layer of about 2.6 μm is detected in the scanned region. EDS analysis showed that the outer oxide layer is rich in Cr and Al elements, while the inner layer is mainly Al-rich species. The oxide double-layer evolves into a Cr and Al mixed oxide layer again when the exposure time extends to 2000 h, with a thickness of around 4.5 μm (Fig. 3c). Meanwhile, a Ti-enriched sublayer with a thickness of about 2.4 μm appears beneath the oxide layer. When the corrosion time reaches 3000 h, as shown in Fig. 3d, a stable mixed oxide layer containing Cr and Al forms on the surface, while the Ti-enriched sublayer disappears. When the Si content increases to 0.2, the surfaces of Cr₂TiAl_0.8Si_0.2C₂ corroded at 550 °C for 500–3000 h are shown in Fig. 2e–h and S4.† The corresponding EDS characterization is shown in Fig. S5.† As can be seen from the surface morphology (Fig. 2e), exposure to LBE first results in a closely cross-linked architecture consisting of spherical particles and a filamentous structure (500 h), and the spheres gradually grow into irregular polygonal grains accompanied by the disappearance of the filamentous parts (3000 h, Fig. 2f–h). The corresponding EDS analyses suggest the inhomogeneous distribution of Ti, Al, and Si on the corroded surface after 500 h. During the exposure time of 1000–2000 h, the elemental distribution on the corroded surface is relatively even. When the corrosion time reaches 3000 h, the elemental distribution of Ti, Al, Si, and O becomes inhomogeneous again, which might be due to the uneven surface consisting of abnormally grown large irregular polygonal grains.

Fig. 3 Cross-section and corresponding EDS diagram of Cr₂TiAlC₂ corroded in LBE for (a) 500, (b) 1000, (c) 2000, and (d) 3000 h at 550 °C. Al, Ti, and Cr elements in the EDS mapping are coloured in red, green, and blue, respectively.

SEM and EDS line scanning of the cross-section from corroded Cr₂TiAl_0.8Si_0.2C₂ further demonstrated the changes in elements near the surface. At the initial corrosion stage (500 h, Fig. S8a†), the oxide layer on the surface of Cr₂TiAl_0.8Si_0.2C₂ is dominated by Cr and Ti, with a small amount of Al and Si intermixed in the oxide layer as well. According to the EDS line scanning, the scanned oxide region has a thickness of about 2.7 μm. After 1000 h exposure, as shown in Fig. S8b,† mixed oxides rich in Cr, Al, and Si are formed on the surface of Cr₂TiAl_0.8Si_0.2C₂. Meanwhile, as shown by the EDS line scanning, the Ti element is mainly beneath the oxide layer. At a corrosion time of 2000 h, the mixed oxide layer rich in Cr, Al, and Si thickens to about 2.9 μm (Fig. S8c†), and the enrichment of Ti beneath the oxide layer can still be detected. When the corrosion time extends to 3000 h, the oxide layer with a double-layer structure is formed (Fig. S8d†). According to the EDS line scanning, a strong signal of Al overlapping with a weak signal of Si is detected on the top region, and the enrichment signal of Cr mainly exists beneath the Al-rich layer. The structure of the oxide layer on Cr₂TiAl_0.8Si_0.2C₂ is different from that of Cr₂TiAlC₂, in which the Al-rich and Cr-rich oxide layers are gradually stratified after long-term LBE corrosion.

For Cr₂TiAl_0.6Si_0.4C₂, the exposure to LBE results in the formation of particle-like products with sizes of about 100 nm in the first stage (less than 1000 h, Fig. 2i, j and S6†). Prolonging the exposure time increases the coverage of the particles while keeping the size nearly constant. The corresponding EDS diagrams exhibit a homogeneous element distribution on the corroded surface, which might be due to the small size and homogeneous coverage of the corroded products on the surface (Fig. S7†). When the corrosion time increases to 2000 h, the spherical particles gradually disappear, and whisker-like structures are formed on the surface (Fig. 2k). On further extending the corrosion time to 3000 h, the whiskers grow into nanorods with a length of about 500 nm (Fig. 2l). According to the EDS mapping, enrichment of Si can be found on the corroded surface.

Fig. 4 shows the cross-section SEM and EDS images of the Cr₂TiAl_0.6Si_0.4C₂ after soaking in LBE for 500–3000 h. As can be seen from Fig. 4a, after 500 h interaction with LBE, Cr and Al are well-overlapped in the surface mixed oxide layer; meanwhile Ti and Si also can be detected in the oxide layer. When the exposure time extends to 1000 h, the thickness of the oxide layer reduces from about 5.5 μm to about 2.0 μm and mainly consists of Cr, Al, and Si elements (Fig. 4b). From the EDS line scanning shown in Fig. 4b, a small amount of Ti oxide exists in the mixed oxide layer. As the corrosion progressed to 2000 h, the surface oxide layer thickens to about 4.0 μm, with an Al-rich layer on the top (2.0 μm) and a Ti-rich sublayer (2.0 μm) beneath it (Fig. 4c). However, although the oxide layer retains a mixed composition containing Cr, Al, Si, and Ti, the signal intensity of Cr is further reduced. When the corrosion time reaches 3000 h, a continuous oxide layer with a thickness of about 1.5 μm dominated by Al oxides is built on the top surface (Fig. 4d), beneath which a Ti-rich sublayer (1.0 μm) can still be observed.

Fig. 4 Cross-section and corresponding EDS diagram of Cr₂TiAl_0.6Si_0.4C₂ corroded in LBE for (a) 500, (b) 1000, (c) 2000, and (d) 3000 h at 550 °C. Si, Al, Ti, and Cr elements in the EDS mapping are coloured in yellow, red, green, and blue, respectively.

3.3. The elemental content in the oxide layer and LBE residue

To elucidate the synergistic effect of multiple elements on LBE corrosion, the contents of Cr, Al, and Si in the surface oxide layer and LBE residue after corrosion were studied. Fig. 5a–c show the ratio of EDS element contents on the corroded surface of Cr₂TiAl_1−xSi_xC₂ over Al or Si on the pristine surface of the sample (Fig. S9†), which is used to evaluate the relative loss or accumulation of elements on the surface during corrosion. As can be seen from Fig. 5a the relative content of Cr in Cr₂TiAlC₂ first decreased and then remained stable within the 3000 h of reaction in LBE. Meanwhile, the relative content of Al is also at a stable level, indicating that the composition of the oxide layer is relatively stable during the corrosion process. Fig. 5b shows the relative contents of Cr, Al, and Si on the corrosion surface of Cr₂TiAl_0.8Si_0.2C₂ over time. When a small amount of Si is introduced, the content of Cr on the surface keeps decreasing over time. Al gradually decreased in the period of 500–2000 h, and increased in 2000–3000 h. Meanwhile, the content of Si is relatively stable over the 3000 h of corrosion. The relative content of Cr on Cr₂TiAl_0.6Si_0.4C₂ dramatically drops in the first 500 h of corrosion (Fig. 5c), while gradually increasing with time during the following 2500 h. In contrast with Cr, the relative content of Al and Si is at a stable level during the 3000 h of corrosion.

Fig. 5 The EDS surface atomic ratio of (a) Cr/Al₀ and Al/Al₀ on Cr₂TiAlC₂ over the corrosion time, (b) Cr/Si₀, Al/Si₀, and Si/Si₀ on Cr₂TiAl_0.8Si_0.2C₂ over the corrosion time, (c) Cr/Si₀, Al/Si₀, and Si/Si₀ on Cr₂TiAl_0.6Si_0.4C₂ over the corrosion time. (d) The atomic ratio of Cr/Ti and Cr/Al on the corroded surface of Cr₂TiAlC₂. (e) The atomic ratio of Cr/Ti, Cr/Al, and Al/Si on the corroded surface of Cr₂TiAl_0.8Si_0.2C₂. (f) The atomic ratio of Cr/Ti, Cr/Al, and Al/Si on the corroded surface of Cr₂TiAl_0.6Si_0.4C₂.

Fig. 5d–f show the atomic ratio on the surface of Cr₂TiAl_1−xSi_xC₂ after 500–3000 h of LBE corrosion, which illustrates the competition of metal composition in the oxide layer over time. As can be seen from Fig. 5d, the variation trend of Cr/Al in Cr₂TiAlC₂ generally shows a downward trend, while the Cr/Ti ratio mainly increases. Therefore, during the corrosion process of 500–3000 h, Al-rich oxides might accumulate on the corroded surface, while Ti atoms are continuously lost from the oxide layer. For Cr₂TiAl_0.8Si_0.2C₂, as shown in Fig. 5e, the change in Cr/Al and Cr/Ti atomic ratios on the corroded surface during 3000 h of corrosion is similar to that of Cr₂TiAlC₂. Meanwhile, the Al/Si ratio shows an upward trend against the corrosion time, which demonstrates the continuous accumulation of Al-rich oxides on the corroded surface as well. Unlike Cr₂TiAl_0.8Si_0.2C₂, the Cr/Al ratio on Cr₂TiAl_0.6Si_0.4C₂ drops in the first 500 h of corrosion (Fig. 5f), while gradually increasing over time within the following corrosion up to 3000 h, indicating the accumulation of Cr oxides. Correspondingly, the Cr/Ti ratio also remains stable, which indicates the potentially improved protecting performance of the oxide layer. Moreover, the Al/Si ratio maintains a relatively stable level during the first 2000 h of corrosion and reduces at 3000 h, which is in accordance with the observation of a continuous Al oxide layer from cross-section EDS mapping in Fig. 4d.

Fig. 6a shows the contents of Cr, Al, and Si elements in the LBE residue after 3000 h interaction with Cr₂TiAl_1−xSi_xC₂ measured by ICP-OES. The dissolution of the three elements in LBE is of the same order of magnitude. In the sample Cr₂TiAlC₂, the dissolved Cr and Al in LBE are 1.75 and 3.00 mg kg⁻¹, respectively. In the LBE residue of Cr₂TiAl_0.8Si_0.2C₂, the contents of Cr and Al increase slightly to 3.85 and 7.40 mg kg⁻¹, respectively, and the dissolved Si is 9.3 mg kg⁻¹. When the added amount of Si reaches 0.4, the dissolved Cr and Al in the LBE residue reduces to 1.95 and 5.60 mg kg⁻¹, respectively, while the dissolved Si dramatically increases to 276 mg kg⁻¹. According to the XRD refinement, the Cr₂TiAl_0.6Si_0.4C₂ specimen contains a high content of Cr₅Si₃C impurity. Given the Cr–Si intermetallic nature of Cr₅Si₃C and its much higher stoichiometric Si/Cr ratio (0.6) than that of Cr₂TiAl_0.6Si_0.4C₂ (0.2), the corrosion of Cr₅Si₃C might offer more dissolved Si in LBE. Meanwhile, the facilitated formation of SiO₂ species can also reduce the dissolution of Cr and Al. Therefore, the higher dissolution amount of Si can be ascribed to the high content of Cr₅Si₃C impurity. Fig. S10† compares the atomic ratios of Al/Si and Cr/Si in the raw materials, on the pristine surface, and in the LBE residue after 3000 h corrosion of Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.4C₂, respectively. As shown, the ratio of Al/Si and Cr/Si in the LBE residue is much smaller than that on the corroded surface, which demonstrates that the Cr and Al in Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.4C₂ are mainly deposited on the corroded surface in forms of oxides than dissolved in the LBE when compared to Si.

Fig. 6 (a) The contents of Cr, Al, and Si elements in the LBE residue after 3000 h corrosion of Cr₂TiAl_1−xSi_xC₂ measured by ICP-OES. (b) The average thickness of the oxide layer on Cr₂TiAl_1−xSi_xC₂ during the corrosion at 550 °C in LBE for 3000 h.

3.4. Thermodynamics and kinetics of oxidations in LBE

During vacuum heat treatment, Cr₂TiAlC₂ began to decompose at 500 °C, and the products were Cr₂AlC and TiC.⁴¹ By comparing the phase composition in Fig. S1,† the TiC content in the three samples increases after 3000 h corrosion, and it can be found from Tables S3† that the Ti–C in Cr₂TiAl_1−xSi_xC₂ MAX is elongated after 3000 h of LBE corrosion. Combining these two results, Cr₂TiAl_1−xSi_xC₂ is possibly decomposed in liquid LBE (Cr₂TiAl_1−xSi_xC₂ → Cr₂Al_1−xSi_xC + TiC) before oxidation. Therefore, the Ellingham diagram of the oxides was plotted based on the oxidation of Cr₂Al_1−xSi_xC at 550 °C in liquid LBE (Fig. S11†). For the calculation in Fig. S11,† the Gibbs free energy (ΔG_fθ) of Cr₂Al_1−xSi_xC is estimated by using the formation energy of metal carbide with the corresponding number of atoms: ΔG^θ_f(M_n+1AX_n) = (n + 1)ΔG^θ_f(MX),⁴⁸ ΔG^θ_f for other compounds were obtained according to the table (NIST-JANAF Thermochemical Tables). At 550 °C, the Gibbs free energies of Al₂O₃, SiO₂, Cr₂O₃, TiO₂, and PbO are calculated successively as follows: Δ_rG[Al₂O₃] < Δ_rG[SiO₂] < Δ_rG[Cr₂O₃] < Δ_rG[TiO₂] < Δ_rG[PbO], which indicate a thermodynamic formation order of oxides with Al₂O₃ > SiO₂ > Cr₂O₃ > TiO₂ > PbO. The oxygen concentration required for the reaction to occur at 550 °C is calculated according to eqn (1)–(3).^49,50

Δ_rG = Δ_rG^θ + 2.303RT lg k (1)

Lg(P[O₂]/P^o) = 10.96 − 2.259 × 10⁴/T (720 ≤ T ≤ 1098) (2)

Δ_rG = RT ln[C_o² exp(13.558 − 32005/T)] (3)

where Δ_rG^θ is the change in standard Gibbs free energy, R is the gas constant, T is the thermodynamic temperature (K), k is the reaction equilibrium constant, P[O₂] is the oxygen partial pressure, and C_o is the oxygen concentration (wt%) formed by oxides in liquid LBE.

The minimum oxygen concentration required by the reaction to generate Al₂O₃, SiO₂, Cr₂O₃, and TiO₂ is 1.32 × 10⁻²⁴, 7.69 × 10⁻²⁰, 1.27 × 10⁻¹⁸ to 7.35 × 10⁻¹⁶, and 2.08 × 10⁻¹⁰ wt%, respectively. Therefore, within the liquid LBE environment in this study, C_o meets the minimum oxygen concentration required for all reactions. Meanwhile, the oxidation of the MAX phase was preferentially started at the edge sites. Thus, although the oxidation of Al is thermodynamically more favourable, kinetically a mixed oxide might be eventually obtained.

The thickness of the oxide layer on Cr₂TiAl_1−xSi_xC₂ was calculated by measuring five different positions in the corroded cross-section and is plotted in Fig. 6b. As shown, during the first 500 h of corrosion in LBE at 550 °C, Cr₂TiAlC₂ delivers the thinnest oxide layer. After 1000 h, both the oxide layers on Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.6C₂ decrease while the oxide layer on Cr₂TiAlC₂ thickens. With a further increase in the exposure time, the thickness of the oxide layer on Cr₂TiAlC₂ is nearly stable. Whereas for the two Si incorporated MAXs, the thickness of the oxide layer on Cr₂TiAl_0.8Si_0.2C₂ increases, while it slightly increases and then reduces on Cr₂TiAl_0.6Si_0.4C₂.

As shown in Fig. 3, the thickening of the oxide layer on Cr₂TiAlC₂ after interacting with LBE for 1000 h is closely related to the evolution in the composition of the oxide layer. The initial outer layer of the oxide mainly consists of Cr and Ti. On extending the corrosion time, the initially formed Ti oxide may peel off or dissolve in LBE,⁵¹ and thus Cr and Al dominated in the outer layer during the following interaction. Eventually, a stable (Cr, Al)₂O₃ mixed oxide layer was formed after 3000 h, which is in accordance with the results of XRD refinement as well. Meanwhile, the inner layer transfers from Al-rich to a Ti-rich composition, indicating the continuous supply of Al from the inner layer to the outer layer. When turning to Cr₂TiAl_0.8Si_0.2C₂ and Cr₂TiAl_0.6Si_0.6C₂, the initially built oxide layer consists of Cr, Ti, Al, and Si (Fig. S8† and 4). Along with the extending of exposure time to 1000 h, the outwards growth of a composite oxide layer containing Cr, Al, and Si is detected. The further variation of the oxide layer thickness can be ascribed to the dissolution of Si into the LBE, which converts the composition of the oxide layer to an Al-dominated outer layer and Cr/Ti-rich inner layer.

3.5. Synergism of Ti, Cr, Al, and Si during LBE corrosion

Fig. 7 schematically shows the corrosion process of Cr₂TiAl_1−xSi_xC₂ in LBE. For Cr₂TiAlC₂, at the initial stage of corrosion, a mixed oxide layer composed of Cr and Ti builds on the surface. The lower content of Al in the oxide layer might be ascribed to the leaching by LBE. Due to a lower diffusion coefficient of oxygen in metal oxides than that of metal ions in metal oxides, Cr₂O₃ with higher thermodynamic favourability than TiO₂ continuously forms at the top surface of the oxide layer. This oxidation process is mainly controlled by the outward diffusion of Cr. Meanwhile, the generated TiO₂ species might offer a barrier towards LBE corrosion as well, like what was observed from Cr₂AlC and Ti₃AlC₂ under liquid Pb corrosion conditions.³¹ On the other hand, the faster diffusion rate of Al and Cr atoms in rutile TiO₂ can facilitate the outward transportation of Al/Cr atoms across the TiO₂ grains,^39,40 which facilitates the continuous growth of Al₂O₃/Cr₂O₃ on the top surface on further extending the exposure time and thus improves the anti-corrosion performance. With the prolongation of corrosion time, Al atoms not only diffuse across TiO₂ grains, but also accumulate beneath the Cr-rich oxide layer and mix in it. The reaction progress is similar to that predicted in the FeCrAl alloy by the first-principles study, in which the oxygen adsorption triggered formation of the Cr₂O₃ layer is gradually transformed into an Al₂O₃ film.⁵² The transformation of oxide species is mainly because the Cr-rich oxide layer formed at the early stage can barrier the inward diffusion of oxygen, and the oxidation process has gradually changed from Cr outward diffusion control to Cr and Al outward diffusion control. Given the higher diffusion rate of Al in Cr₂O₃ than that of Cr in Cr₂O₃ (Al diffusion rate: 10⁻¹⁰ m² s⁻¹; Cr < 10⁻¹² m² s⁻¹),⁵³ the content of Al in the oxide layer increases (Fig. 5), and eventually forms the mixed oxide layer.

Fig. 7 Schematic illustration of the corrosion process of Cr₂TiAl_1−xSi_xC₂ at 550 °C in LBE.

In the process of oxidative corrosion of Cr₂TiAl_0.8Si_0.2C₂, the initial oxidation layer formation process is similar to that of Cr₂TiAlC₂. However, due to a faster diffusion rate of Al in SiO₂ (about 10⁻²² m² s⁻¹) than that of Al in Al₂O₃ (<10⁻²⁴ m² s⁻¹),^54,55 SiO₂ formed in the earlier stage becomes a channel for Al diffusion and promotes its diffusion. Thus, the existence of Si can facilitate the formation of an Al-rich oxide layer, as probed in the Ti₃Al_1−xSi_xC₂ system.³⁰ Moreover, due to the higher diffusion rate of Al in Cr₂O₃ than that of Cr, the compositional transformation of the outer layer from Cr, Al, and Si into Al and Si was observed as well (Fig. S8†).

When the Si content further increases to 0.4, the initial oxidation process of Cr₂TiAl_0.6Si_0.4C₂ is slightly different from that of Cr₂TiAlC₂ and Cr₂TiAl_0.8Si_0.2C₂. On the one hand, the formation of mixed oxides containing Cr and Al occurs at the early stage. On the other hand, the thickness of the oxide layer is higher than that of the other two MAXs. Meanwhile, as revealed in the SIMP steel,⁵⁶ the formed SiO₂ nanoparticles in the oxide layer can act as the heterogeneous site for Cr₂O₃ nucleation, which promotes the formation of a Cr-rich oxide layer with Al and Si. Accompanied by the much slower oxygen diffusion rate in SiO₂ than that in Cr₂O₃, the inward diffusion of oxygen can be hindered by the built oxide layer.⁵⁷ In the later stage of corrosion, the outwards diffusion of Al through the Cr₂O₃ grain becomes the primary mechanism, which results in the transformation of a Cr and Al-rich oxide layer into an Al-rich oxide layer (Fig. 4).

4. Conclusion

In summary, the oxide barrier formation mechanism of Cr₂TiAl_1−xSi_xC₂ MAXs with different Si contents was studied at 550 °C in liquid LBE for 500–3000 h. According to the phase, chemical composition, and bond length analysis of Cr₂TiAl_1−xSi_xC₂ after corrosion, an elongation of the Ti–C bond in Cr₂TiAl_1−xSi_xC₂ MAXs during the corrosion in liquid LBE at 550 °C and an increase in TiC content were detected. This suggests that the interaction of Cr₂TiAl_1−xSi_xC₂ with LBE may involve the decomposition of Cr₂TiAl_1−xSi_xC₂ into Cr₂Al_1−xSi_xC and TiC. Moreover, the protective oxide layer on the MAXs is crucial for the corrosion resistance to LBE. The synergistic action of Cr, Al, and Si elements in Cr₂TiAl_1−xSi_xC₂ during LBE corrosion plays an important role in regulating the composition and structure of the oxide layer. The addition of Si leads to the formation of SiO₂ in the oxide layer. The formed SiO₂ can serve as the heterogeneous nucleation site of Cr₂O₃, and promotes the diffusion of Al to the top surface of the oxide layer with the channelling effect of SiO₂, increasing the content of Al oxides in the oxide layer. Eventually, the multi-element collaboration in Cr₂TiAl_1−xSi_xC₂ regulates the composition and structure of the oxide barrier and thus the corrosion resistance, which paves the way for the performance-oriented design of the MAX phase.

Data availability

The data supporting this article have been included as part of the ESI.†

Author contributions

Cheng-Feng Du: conceptualization, writing-review, funding acquisition. Qingyan Zeng: investigation, data curation, formal analysis, writing-original draft. Junjie Cha: resources. Hong Yu: writing-review, supervision. Hongwei Liang: investigation. Kun Liang: resources. Shiyao Lei: investigation. Lili Xue: investigation. Xian-Zong Wang: formal analysis, supervision.

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

The authors are grateful for the financial support from the National Natural Science Foundation of China (No. 52275212 and 52471094), the Aeronautical Science Foundation of China (2024Z046053001), the “Special Lubrication and Sealing for Aerospace” Shaanxi Provincial Science and Technology Innovation Team (No. 2024RSCXTD-63), and the Open Project of Zhejiang Key Laboratory of Data-Driven High-Safety Energy Materials and Applications.

References

1. X. Gong, M. P. Short, T. Auger, E. Charalampopoulou and K. Lambrinou, Prog. Mater. Sci., 2022, 126, 100920.
2. D. Zhang, X. Zhang, J. Zhang, H. Ren, Z. Liao, X. Zeng and Q. Yan, J. Mater. Sci. Technol., 2025, 222, 55–67.
3. J. Zhang, Adv. Eng. Mater., 2014, 16, 349–356.
4. D. Wang, S. Liu, C. Xiao, X. Ma, Y. Sun, G. Yuan, J. Zeng, Y. Liang, Y. Hu, F. Niu and X. Gong, Corros. Sci., 2023, 218, 111169.
5. J. Zhang and N. Li, J. Nucl. Mater., 2008, 373, 351–377.
6. E. Yamaki, K. Ginestar and L. Martinelli, Corros. Sci., 2011, 53, 3075–3085.
7. K. Lambrinou, E. Charalampopoulou, T. Van der Donck, R. Delville and D. Schryvers, J. Nucl. Mater., 2017, 490, 9–27.
8. C. Schroer, O. Wedemeyer, J. Novotny, A. Skrypnik and J. Konys, Corros. Sci., 2014, 84, 113–124.
9. G. Chen, J. Wang, H. Zhang, L. Li and H. Fan, Scr. Mater., 2021, 202, 114014.
10. C. Cionea, M. D. Abad, Y. Aussat, D. Frazer, A. J. Gubser and P. Hosemann, Sol. Energy Mater. Sol. Cells, 2016, 144, 235–246.
11. J.-S. Li, Y.-F. Wang, J. Chai, W. Gong and X.-Z. Wang, J. Nucl. Mater., 2025, 604, 155516.
12. W. Zhang, J. Deng, H. Yue, S. Hu, X. Qiu, H. Yin, Q. Li, H. Liu, M. Zhou and J. Yang, Corros. Sci., 2023, 225, 111590.
13. Z. Zhu, J. Tan, X. Wu, Z. Zhang, E.-H. Han and X. Wang, Corros. Sci., 2022, 209, 110767.
14. Y.-F. Wang, J.-S. Li, W. Xu and X.-Z. Wang, Corros. Sci., 2024, 227, 111693.
15. X. Gong, H. Chen, F. Zhang, W. Zhu, H. Ma, B. Pang and Y. Yin, J. Nucl. Mater., 2022, 558, 153364.
16. Y. Zhong, W. Zhang, Q. Chen, J. Yang, C. Zhu, Q. Li, J. Yang, N. Liu and J. Yang, Surf. Coat. Technol., 2022, 443, 128598.
17. X. Yin, Y. Wang, H. Wang, K. Zhao, Y. Sun, J. Xiao, Y. Zhao, F. Gong and Y. Chen, Vacuum, 2023, 215, 112251.
18. M. Takahashi and M. Kondo, Prog. Nucl. Energy, 2011, 53, 1061–1065.
19. Z. Y. Wu, X. Zhao, Y. Liu, Y. Cai, J. Y. Li, H. Chen, Q. Wan, D. Yang, J. Tan, H. D. Liu, Y. M. Chen, J. L. Guo, J. Zhang, G. D. Zhang, Z. G. Li and B. Yang, J. Nucl. Mater., 2020, 539, 152280.
20. C.-F. Du, L. Xue, C. Wang, Q. Zeng, Z. Yang, Y. Xue, X. Wang, Z. Wang and H. Yu, Wear, 2023, 520–521, 204671.
21. C.-F. Du, C. Wang, H. Liang, L. Xue, Y. Xue, Z. Wang, X. Wang and H. Yu, J. Eur. Ceram. Soc., 2023, 43, 7341–7353.
22. C.-F. Du, Y. Q. Xue, C. C. Wang, Q. Y. Zeng, J. J. Wang, X. M. Wang, L. L. Xue and H. Yu, J. Eur. Ceram. Soc., 2023, 43, 4684–4695.
23. C. F. Du, Z. Yang, Q. Zeng, L. Xue, C. Wang, J. Wang and H. Yu, Chem.–Eur. J., 2023, 29, e202203106.
24. C.-F. Du, C. Wang, M. Xu, Y. Xue, L. Xue, C. Meng, L. Wang, W. Qi, X. Liu and H. Yu, Nat. Commun., 2025, 16, 3011.
25. M. Dahlqvist, M. W. Barsoum and J. Rosen, Mater. Today, 2024, 72, 1–24.
26. H. Yu, L. Xue, Y. Xue, H. Lu, Y. Liu, L. Wang, C.-F. Du and W. Liu, Carbon Energy, 2024, 6, e597.
27. L. Xue, Q. Xu, C. Meng, S. Lei, G. Zhang, M. Tang, W. Zhai, H. Yu, X. Liu and C.-F. Du, Tribol. Int., 2024, 199, 110030.
28. Y. Xue, H. Yu, H. Liang, X. Wang, C. Meng, S. Lei, H. Wang, L. Wang and C.-F. Du, Tribol. Int., 2025, 204, 110525.
29. C.-F. Du, Y. Xue, H. Liang, C. Wang, Q. Zeng, J. Wang, L. Xue and H. Yu, J. Mater. Sci. Technol., 2024, 187, 49–62.
30. C.-F. Du, Q. Zeng, Y. Wang, H. Liang, C. Wang, Y. Xue, J. Wang, H. Yu and X.-Z. Wang, Corros. Sci., 2024, 233, 112068.
31. H. Shi, R. Azmi, L. Han, C. Tang, A. Weisenburger, A. Heinzel, J. Maibach, M. Stüber, K. Wang and G. Müller, Corros. Sci., 2022, 201, 110275.
32. B. Tunca, T. Lapauw, C. Callaert, J. Hadermann, R. Delville, E. a. N. Caspi, M. Dahlqvist, J. Rosén, A. Marshal, K. G. Pradeep, J. M. Schneider, J. Vleugels and K. Lambrinou, Corros. Sci., 2020, 171, 108704.
33. X. Ma, X. Shen, H. Tang, R. Qiu and K. Lyu, Mater. Lett., 2025, 392, 138534.
34. Z. Zhao, S. Shi, X. Gao, Q. Wang, Y. Zhang, Y. Wen, Y. Ling, X. Zeng and M. Ma, J. Nucl. Mater., 2025, 607, 155708.
35. A. Heinzel, A. Weisenburger and G. Müller, J. Nucl. Mater., 2016, 482, 114–123.
36. T. Lapauw, B. Tunca, J. Joris, A. Jianu, R. Fetzer, A. Weisenburger, J. Vleugels and K. Lambrinou, J. Nucl. Mater., 2019, 520, 258–272.
37. Z. Xu, L. Song, Y. Zhao and S. Liu, Corros. Sci., 2021, 190, 109634.
38. D. Zhang, J. Zhang, H. Ren, X. Zeng, Z. Liang, X. Zhang and Q. Yan, Corros. Sci., 2025, 244, 112643.
39. M. W. Barsoum, J. Electrochem. Soc., 2001, 148, C544–C550.
40. J. Sasaki, N. L. Peterson and K. Hoshino, J. Phys. Chem. Solids, 1985, 46, 1267–1283.
41. Z. Liu, J. Yang, H. Zhang, Z. Liu, H. Li, M. Li, Y. Qian and J. Xu, Ceram. Int., 2023, 49, 12034–12041.
42. B. H. Toby and R. B. Von Dreele, J. Appl. Crystallogr., 2013, 46, 544–549.
43. H. Yu, H. Jing, Y. Gao, X. Wang, Z.-Y. Gu, L. Li, J. Wang, S. Wang, X.-L. Wu, W. Qi, Q. Liang and C.-F. Du, Adv. Mater., 2025, 2400229,  DOI:10.1002/adma.202400229.
44. G. Müller, A. Heinzel, G. Schumacher and A. Weisenburger, J. Nucl. Mater., 2003, 321, 256–262.
45. H. Zhu, X. Liu, B. Chang, X. Li, M. Qi, Y. Wang, F. Niu, Y. Ma and L. Lyu, Mater. Corros., 2021, 73, 196–206.
46. R. Ganesan, T. Gnanasekaran and R. S. Srinivasa, J. Nucl. Mater., 2006, 349, 133–149.
47. M. Dahlqvist and J. Rosen, Nanoscale, 2020, 12, 785–794.
48. M. W. Barsoum, Prog. Solid State Chem., 2000, 28, 201–281.
49. C. Schroer and J. Konys, Physical chemistry of corrosion and oxygen control in liquid lead and lead-bismuth eutectic, Institut für Materialforschung - Werkstoffprozesstechnik (IMF3), 2007.
50. A. Kishimoto, A. Wada, T. Michimoto, T. Furukawa, K. Aoto and T. Oishi, Mater. Trans., 2006, 47, 122–128.
51. M. P. Popovic, D. L. Olmsted, A. M. Bolind, M. Asta, S. Sohn, J. Schroers, R. Shao and P. Hosemann, Mater. Des., 2018, 159, 240–251.
52. J. Wang, K. Yan, W. Huang and Z. Lu, Appl. Surf. Sci., 2024, 644, 158782.
53. I. Roy, P. K. Ray and G. Balasubramanian, Materialia, 2022, 24, 101497.
54. A. H. Heuer, J. Eur. Ceram. Soc., 2008, 28, 1495–1507.
55. R. Pankrath and O. W. Floerke, Eur. J. Mineral., 1994, 6, 435–457.
56. L. Zhang, W. Yan, Q. Shi, Y. Li, Y. Shan and K. Yang, Corros. Sci., 2020, 167, 108519.
57. J. Zhou, Z. Liu, C. Jiang, W. Peng, X. Wei, R. Luo, X. Deng, T. Tan, F. Chen, J. Liu, X. Luo and X. Xu, Surf. Coat. Technol., 2023, 470, 129864.

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Footnote

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d5ta02706c

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