Facet-dependent adsorbate-mediated strong metal-support interaction in Ni/TiO_2

Domain: chemistry
Published: 2025-12-18
Editorial outcome: minor_revision

Abstract
Despite the growing significance of adsorbate-mediated strong metal-support interaction (A-SMSI) in various catalytic processes, a comprehensive mechanistic understanding of its formation and effective strategies for its precise modulation remain elusive. Herein, by constructing three well-defined model Ni/TiO2 catalysts with distinct exposed facets, we directly visualize a facet-dependent A-SMSI behavior in CO2 hydrogenation via in situ environmental transmission electron microscopy (ETEM) at the atomic level. The in situ results reveal distinct formation behaviors of TiO2-x overlayers: complete, partial, and no encapsulation of Ni nanoparticles (NPs) on the {100}, {101}, and {001} TiO2 facets, respectively. Complementary in situ diffuse reflectance infrared Fourier transform spectroscopy (DRIFTS), X-ray absorption spectroscopy (XAS) experiments and density functional theory (DFT) calculations further reveal a dual induction mechanism, where adsorbates both participate in TiO2-x formation and stabilize the encapsulating overlayer. Moreover, TiO2-x transfers electrons to Ni and stabilizes COOH* intermediates that dehydroxylate to form CO. Guided by these mechanistic insights, facet-dependent A-SMSI enables remarkable selectivity modulation, yielding CH4 selectivity exceeding 88% on Ni/TiO2-{001} and CO selectivity over 83% on Ni/TiO2-{100}. These findings advance the fundamental understanding of A-SMSI and offer a rational framework for designing oxide-supported catalysts with tailored interfacial properties.

Introduction
The strong metal–support interaction (SMSI), characterized by the migration of the support over metal nanoparticles (NPs) on a supported-metal catalyst1, offers a powerful strategy in heterogeneous catalysis for engineering new catalytic sites through interfacial metal-support interactions2,3. Classical SMSI typically proceeds via the formation of mobile species on the support surface, interfacial mass transfer, and the development of encapsulating overlayers, which governed by the thermodynamic imperative to reduce the overall surface energy of the catalyst system. However, uncontrolled migration of the support may obstruct the metal active sites, and the high reduction temperature conditions required for thermally induced SMSI (H2, ≥500 °C) can also lead to sintering of the metal NPs4.

A distinct variant of SMSI—adsorbate-induced SMSI (A-SMSI) — has recently emerged, wherein reactive intermediates (e.g., alcohols5,6, levulinic acid7, or CO2-derived species8–10) drive the migration of the support. In contrast to SMSI, A-SMSI occurs at markedly lower temperatures (~250 °C), owing to the reduced energy input required for its formation reactions11. This strategy effectively circumvents catalyst sintering and deactivation commonly associated with SMSI that necessitates high temperatures. Moreover, A-SMSI facilitates the formation of permeable overlayers that retain accessibility to active sites, enabling favorable modulation of reaction pathways without significantly compromising overall catalytic activity8,11.

Inspired by these advances, considerable efforts have been devoted to manipulating A-SMSI overlayer structures and their associated catalytic performances across various systems and applications12,13. For instance, Matsubu et al. demonstrated that the nature of the support plays a critical role: Rh catalysts supported on reducible oxide supports (TiO2, Nb2O5) could facilitate the formation of HCOx-functionalized encapsulation that benefits the tuning of CO2 reduction selectivity8. Furthermore, Li et al. reported that pretreatment temperature affects the concentration of oxygen vacancy in oxide supports, thereby critically influencing the stability of A-SMSI overlayer and the performance of Ru/TiO2 catalysts14. Recently, by using in situ environmental transmission electron microscopy (ETEM), Deng et al. further demonstrated that pretreatment in CO and H2 could dynamically and reversibly alter the morphology of Cu@TiO2-x catalysts, emphasizing that metal–support interfaces are highly sensitive to the chemical environment. These findings highlight the inherently dynamic nature of A-SMSI and the importance of in situ characterization15. Despite these pioneering efforts, many fundamental questions regarding the creation and manipulation of A-SMSI remain unresolved16,17. In particular, it remains unclear whether and how the atomic structure of the exposed support surface affects the formation, structure, and catalytic performance of A-SMSI. This knowledge gap largely stems from the lack of atomic-level, in situ insights into the overlayer dynamics.

Herein, taking the Ni/TiO2 as a model catalyst18–20, we carry out a systematic study of the facet-dependent A-SMSI behavior in CO2 hydrogenation. To fairly compare the Ni catalyst supported on different TiO2 facets, three well-defined Ni/TiO2 model catalysts with distinct exposed TiO2 facets are constructed: TiO2 nanorods (TiO2-R, {100}), bipyramidal nanocrystals (TiO2-B, {101}), and nanosheets (TiO2-S, {001}). In situ ETEM directly visualizes the dynamic evolution of A-SMSI overlayers, revealing complete and partial encapsulation of Ni NPs on {100} and {101} facets, respectively, but none on {001}. The exposed facets regulate both CO2 activation into induction intermediates (HCOx*) and TiO2-x formation of at the Ni-TiO2 interface. Complementary in situ diffuse reflectance infrared Fourier transform spectroscopy (DRIFTS) and in situ X-ray absorption spectroscopy (XAS) further unveil an interfacial electron transfer from TiO2-x to Ni, stabilizing COOH* intermediates that dehydroxylate to form CO. This facet-dependent A-SMSI fundamentally reshapes product selectivity: from 88% CH4 on Ni/TiO2-{001} to over 83% CO on Ni/TiO2-{100}. These findings deepen the mechanistic understanding of A-SMSI and offer a rational strategy for tailoring interfacial properties in oxide-supported catalysts.

Results
Characterization of Ni/TiO2 catalysts with different support facets
To investigate the effect of TiO2 crystal facets on A-SMSI, three well-defined Ni/TiO2 model catalysts were prepared by loading Ni NPs onto TiO2 supports with distinct morphologies. TiO2-R predominantly exposes the surface of {100} facet, with diameters of 40-60 nm and lengths of 200–300 nm; TiO2-B primarily exposes the {101} facet, with lengths of 100-120 nm; TiO2-S mainly exposes the {001} facet, with lengths of 60-110 nm and thicknesses of 5-10 nm, as shown in the transmission electron microscopy (TEM) images in Fig. 1a–c. Atomic-resolution high-angle angular dark field scanning TEM (HAADF STEM) images further identify the interplanar spacings of the three TiO2 supports as 3.78 Å (Fig. 1d), 3.51 Å (Fig. 1e), 2.38 Å (Fig. 1f), corresponding to the {100}, {101}, {001} facets, respectively. These results are consistent with the high-resolution TEM (HRTEM) images shown in Supplementary Fig. 1. The Ti columns could also be clearly identified due to the Z contrast21, which well matches the bulk-truncated surfaces of each facet. The thermodynamically stable {100} facet is exposed by 5-fold coordinated Ti atoms (Ti5c), with 2- and 3-fold oxygen atoms (O2c and O3c); the {101} facet is additionally exposed by Ti6c; while the {001} facet is only exposed by Ti5c and O2c22,23. Given the sensitivity of TiO2 to electron beam irradiation, all subsequent experiments were conducted under relatively low-dose conditions, referring to previous studies24–27. Moreover, control experiments confirmed that the electron beam did not obviously alter the relative stability and structure of different TiO2 facets (Supplementary Figs. 2, 3).

Ni/TiO2-S, Ni/TiO2-B, and Ni/TiO2-R catalysts were prepared by loading Ni onto the corresponding supports, with Ni loadings of 9.2 wt.%, 9.6 wt.%, and 9.4 wt.%, respectively, as determined by inductively coupled plasma-optical emission spectrometer (ICP-OES). HAADF STEM imaging and energy dispersive X-Ray (EDX) spectroscopy elemental mappings confirm the uniform distribution of Ni NPs on TiO2 (Fig. 1g–i). The size distribution violin plots of Ni NPs, as determined by analyzing 100 NPs per sample, are shown in Fig. 1j–l: 14.66 nm on Ni/TiO2-R, 16.33 nm on Ni/TiO2-B, and 12.55 nm on Ni/TiO2-S. The comparable sizes in the three samples provide an ideal platform to study the facet-dependent behavior28.

Evolution of encapsulation overlayer in various Ni/TiO2 catalysts
To explore the facet-dependent behavior, the ETEM experiments were employed to observe the structural evolution of various Ni/TiO2 catalysts under reactive gas environments. Before the performance test, reduction at elevated temperature (typically 400–500 °C) is a general routine for Ni/TiO2 catalyst activation to obtain metallic Ni. While in the ex-situ characterization, this active metallic Ni catalysts are prone to surface oxidation upon ambient air exposure29,30, forming superficial NiO layers, as revealed in Supplementary Fig. 4. Therefore, the ex-situ observation of Ni/TiO2 catalysts could not represent the realistic state of the catalysts in the reaction. To precisely probe the intrinsic physicochemical properties of Ni-based catalysts, the in situ study of the Ni/TiO2 catalysts under the pretreatment and reaction environments are indispensable for constructing a reliable structure-performance relationship. The corresponding in situ ETEM, HAADF STEM, and electron energy loss spectroscopy (EELS) results for Ni/TiO2-R, Ni/TiO2-B, and Ni/TiO2-S are presented in Fig. 2 and Supplementary Fig. 5.

ETEM imaging after the reduction pretreatment confirmed the complete removal of the NiO overlayer for all the catalysts (Fig. 2a and Supplementary Fig. 5a, i), and fast fourier transform (FFT) analysis further validated the metallic state of Ni, as exemplified by Ni/TiO2-R (Supplementary Fig. 6a). The TiO2-R support was imaged along the [010] direction with (100) facet exposure, while the located Ni was oriented along the [011] direction, exposing the (−200), (1–11), and (−1–11) facets (Supplementary Fig. 6a).

Under CO2 hydrogenation conditions, the three catalysts exhibited distinct structural evolution behaviors, as observed from the in situ ETEM images in Fig. 2b–d, and Supplementary Fig. 5. In particular, after 60 min of reactions, the Ni NPs in Ni/TiO2-R were completely covered by encapsulation layers (Fig. 2d), while partial encapsulation was observed in Ni/TiO2-B (Supplementary Fig. 5d) and no encapsulation layer formed in Ni/TiO2-S (Supplementary Fig. 5l).

In situ ETEM imaging of the Ni/TiO2-R interface upon 15 min exposure to reaction conditions (Fig. 2b) revealed the formation of a 1.69 nm-thick amorphous overlayer on the Ni surface, consistent with the dynamic interfacial restructuring previously reported by Matsubu et al. for Rh/TiO28. The Ni NP showed slight morphology changes, with the length parallel to the support surface increasing from 13.98 nm to 14.46 nm (Fig. 2a, b), confirming slight flattening of the Ni NP due to overlayer growth. The FFT patterns of the dashed region in Fig. 2b showed no additional diffraction spots compared to the initial state (Supplementary Fig. 6b), indicating the absence of Ni oxidation. After 30 min of exposure (Fig. 2c), the overlayer grew further and induced a slight counterclockwise rotation of the Ni particle, exposing the {111} facet. Complete encapsulation of the particle was achieved after 60 min, as observed from Fig. 2d. The formation of the encapsulating overlayer requires the synergy of CO2 and H2 rather than either gas alone. In contrast to H2, which caused no significant morphological change (Fig. 2a), CO2 alone induced surface amorphization of the Ni nanoparticles but did not lead to encapsulation. This amorphization, evidenced by the appearance of blurred lattice fringes (Supplementary Fig. 7), aligns with established understanding, which attributes it primarily to the dissociative adsorption of CO2 that reconstructs the Ni surface atoms31. As TiO2 is known to be sensitive to electron beam irradiation in TEM32, the total exposure dose on the sample was carefully limited to below 6.14 × 103 e Å−2 to minimize beam damage.

To further identify the composition of the encapsulating species, HAADF STEM imaging of Ni/TiO2-R was conducted after reaction under atmospheric conditions (CO2/H2/N2 = 15/15/70, 350 °C) (Fig. 2e). The enlarged image (Fig. 2f) revealed a ~1.7 nm-thick overlayer with brighter contrast relative to the Ni core, consistent with the ~1.69 nm amorphous layer observed via in situ ETEM (Fig. 2b). EELS mappings confirmed that this overlayer was composed of titanium species (Fig. 2g). To probe its electronic properties, EELS line scan of the Ti L2,3-edge was collected from bulk TiO2 to vacuum (indicated in Fig. 2e). As shown in Fig. 2h, the L2,3-edge of Ti species at the edges showed a slight shift (~0.6 eV) toward lower energy relative to TiO2, indicating the stabilization of Ti3+ species33,34 and providing direct spectroscopic evidence for the formation of SMSI-associated TiO2-x under reaction conditions.

A TiO2-x overlayer was observed over Ni/TiO2-B as well (Supplementary Fig. 5a–d), whereas no analogous interfacial structure emerged on Ni/TiO2-S under identical reaction conditions (Supplementary Fig. 5i–l). Notably, EELS analysis of the Ti L-edge further confirmed partial encapsulation of Ni particles in Ni/TiO2-B (Fig. 2i, j), which were absent in Ni/TiO2-S (Fig. 2k, l).

Collectively, these results suggest that TiO2-x overlayer formation is more favorable on Ni/TiO2-R and Ni/TiO2-B than on Ni/TiO2-S. Consequently, facet-dependent differences in TiO2 lead to distinct encapsulation structural behaviors: full encapsulation on {100}-faceted Ni/TiO2-R, partial encapsulation on {101}-faceted Ni/TiO2-B, and an unencapsulated metallic surface on {001}-faceted Ni/TiO2-S.

Revelation of catalyst adsorbates and the mechanism of encapsulation induction
It has been reported that a high surface coverage of HCOO* adsorbates under CO2/H2 atmospheres could promote H2O desorption8,14, which in turn creates oxygen vacancies and forms non-stoichiometric TiO2-x species that encapsulate the metal35. The facet-dependent nature of TiO2 leads to variations in the adsorption of carbonaceous species, due to differences in atomic arrangements and coordination environments across exposed facets36. To acquire solid evidence for surface species and to elucidate the facet-dependent encapsulation mechanism, in situ diffuse reflectance infrared Fourier transform spectroscopy (DRIFTS) was employed on three Ni/TiO2 catalysts under reaction conditions.

As shown in Fig. 3a, Ni/TiO2-R exhibited absorption bands at ~1595 and ~1380 cm−1, corresponding to a high coverage of adsorbed HCOO*37,38, along with bands at ~1428 and ~1248 cm−1 assigned to HCOx*8,39. In contrast, Ni/TiO2-B initially formed monodentate carbonate (~1530 cm−1 and ~1325 cm−1) and HCOx* (~1440 cm−1) after ~5 min of CO2 pre-activation, which gradually transformed into HCOO* (~1585 cm−1 and ~1380 cm−1) and other carbonate species (Fig. 3b)40. These spectral evolutions highlight the pivotal role of HCOO* among the HCOx* species. Compared to the other two catalysts, Ni/TiO2-S showed significantly weaker bands for surface carbonaceous species (Fig. 3c), and HCOO* species disappeared rapidly upon heating (Supplementary Fig. 8). These results indicate that HCOO* accumulates and stabilizes only on Ni/TiO2-R and Ni/TiO2-B. Further comparison under H2 purge (Fig. 3d) revealed that after 16 min, the HCOO* signal on Ni/TiO2-B (~1585 cm−1) decreased by ~74%, whereas only a ~42% reduction was observed for Ni/TiO2-R (~1595 cm−1), indicating relatively weaker adsorption on the former. This trend is consistent with CO2 temperature-programmed desorption (CO2-TPD) results (Fig. 3e), which show that CO2 adsorption at moderate basic sites (150–300 °C) — precursors to HCOO* formation41—follows the order Ni/TiO2-R (55.8 μmol g−1) > Ni/TiO2-B (48.3 μmol g−1) > Ni/TiO2-S (36.7 μmol g−1). The greater CO2 uptake capacity of Ni/TiO2-R thus explains its enhanced ability to accumulate HCOO* and thereby contribute to stronger A-SMSI behavior. Furthermore, H2 temperature-programmed reduction (H2-TPR) profiles showed higher temperature shifts in Ni/TiO2-R and Ni/TiO2-B compared to Ni/TiO2-S, indicating stronger metal-support interactions (Fig. 3f)37,42.

The density functional theory (DFT) calculations further rationalize these observations using a TiO2-x/Ni(111) heterostructure model (Supplementary Fig. 9)3. The calculations showed that HCOO* can be stably adsorbed on the TiO2 surface (Fig. 3g), with adsorption energies (Eads) of −4.60 eV, −4.70 eV and −3.72 eV at three representative sites. In all cases, the negative Eads indicate thermodynamically favorable binding, suggesting that HCOO* adsorption on the TiO2-x overlayer of Ni/TiO2-R provides additional stabilization to the reduced surface, which in turn favor the formation of TiO2-x overlayers and promotes encapsulation.

Combined DRIFTS and DFT results indicate that the adsorption strength and coverage of HCOO* species critically affect the extent of A-SMSI. On Ni/TiO2-R, strongly bound HCOO* stabilizes the TiO2-x overlayer, while on Ni/TiO2-B, weaker adsorption only enables partial encapsulation, as evidenced in Fig. 2i, j and Supplementary Fig. 5. In contrast, the absence of stable HCOO* intermediates on Ni/TiO2-S eliminates this stabilizing effect, preventing overlayer formation.

As noted, once TiO2 is reduced to non-stoichiometric TiO2-x species, migration of the reduced oxide can lead to encapsulation of the active metal35. Previous studies have shown that formic acid (HCOOH*) can create oxygen vacancies on TiO2 via H2O desorption, leaving a HCOO-covered reduced surface (A-SMSI)8,43. To probe this mechanism, DFT calculations were performed for the HCOOH* intermediate on Ni/TiO2-x-R. Although HCOOH* can decompose into HCO* and OH* during CO2 hydrogenation over Ni-based catalysts (with HCO* subsequently hydrogenating to CH4)44, the calculations indicate a more favorable dissociation into HCOO* and H* on Ni/TiO2-R.

Two competing pathways were evaluated: (path a) the conventional hydrogenation sequence and (path b) the vacancy formation route. Dissociation along path a requires ΔE = 0.57 eV, whereas path b proceeds with a much lower barrier of ΔE = 0.19 eV (Fig. 3h). These results clearly demonstrate that the oxygen-vacancy formation pathway is kinetically highly favorable, thereby providing direct support for the proposed mechanism. Moreover, DFT calculations revealed that interfacial oxygen vacancy formation is most favorable on Ni/TiO2-R (2.67 eV) among three catalysts, providing a fundamental basis for TiO2-x overlayer formation (Supplementary Figs. 10–15).

Notably, both gaseous CO (~2162 cm−1) and COOH* (1544 cm–1) bands45,46 were observed over Ni/TiO2-R (Fig. 3a), aligning with its CO-rich selectivity (discussed later). Moreover, COOH* (61%) decayed faster than HCOO* (42%) after 16 min in H2 (Fig. 3d), suggesting COOH* is the reactive intermediate for CO formation47,48, while HCOO* mainly stabilizes the TiO2-x overlayer. In contrast, no CO-related bands were observed over Ni/TiO2-B and Ni/TiO2-S (Fig. 3b, c). Instead, new peaks at 3064 and 3014 cm−1 corresponding to C-H stretching vibrations of CH4 appeared on Ni/TiO2-B (Fig. 3b)49, while gaseous CH4 over Ni/TiO2-S was confirmed by bands at ~ 3014 and 1304 cm-1 (Fig. 3c). These observations are consistent with CO2 hydrogenation results, where CH4 is the dominant product over Ni/TiO2-B and Ni/TiO2-S (details provided later).

Collectively, these results demonstrate that facet-dependent structural configurations dictate the formation and stabilization of HCOO* species, which play a dual role in the induction A-SMSI. Specifically, HCOO* not only facilitates vacancy formation—contributing to the TiO2-x generation and subsequent encapsulation—while also stabilizes the TiO2-x overlayer through adsorption once encapsulation has occurred.

Electronic structure evolution of Ni induced by A-SMSI
The presence of A-SMSI on Ni/TiO2-R has been conclusively demonstrated by ETEM and DRIFTS, with a key feature being electron transfer between the metal and the support8,11. EELS measurements confirm the presence of Ti3+ species within the TiO2-x overlayer (Fig. 2h). To further elucidate the evolution of the electronic structure of Ni NPs, in situ X-ray absorption spectroscopy (XAS) was performed (Fig. 4). The normalized Ni K-edge X-ray absorption near edge structure (XANES) spectra of Ni/TiO2-S and Ni/TiO2-R after in situ hydrogen reduction exhibits light edge shifts toward higher photon energy relative to Ni foil (Fig. 4a), indicating a comparable Ni oxidation state, with the coexistence of metallic Ni and NiO. These findings imply that the valence state of Ni alone does not govern the onset of A-SMSI.

To capture the dynamic evolution of the Ni electronic structure, Ni/TiO2-R was sequentially exposed to different gas exposures, each held for ~40 min to allow stabilization (Fig. 4b). Under Ar at room temperature, the extended X-ray absorption fine structure (EXAFS) spectra of unreduced NiO/TiO2-R as well as the Fourier transform and wavelet transform (WT) results (Fig. 4c, d and Supplementary Fig. 16) confirm that NiO dominates on the catalyst. Note that the distances observed in the Fourier transformed spectra (Fig. 4c) are systematically shorter than the actual absorber-backscatter distances (Supplementary Table 1), owing to phase shift effects. After H2 reduction at 400 °C (green curve in Fig. 4b), a mixed Ni/NiO phase was obtained, as discussed above. Subsequent introduction of CO2 (CO2/Ar = 10/90) at 350 °C resulted in a notable edge shift to higher photon energy (blue curve in Fig. 4b), indicative of partial reoxidation of Ni.

Upon exposure to H2/CO2 (CO2/H2/Ar = 10/10/80), the edge clearly shifted back toward lower energy, approaching that of Ni-foil (yellow curve) shown in Fig. 4b, suggesting that Ni gains electrons under reaction conditions. X-ray photoelectron spectroscopy (XPS) analysis further corroborated this conclusion, as evidenced by the shift of the Ni⁰ peak in Ni/TiO2-R after reaction to a lower binding energy (851.9 → 851.5 eV) compared to the reduced sample (Supplementary Fig. 17). This is likely driven by the accumulation of adsorbed HCOO* intermediates, which induce the formation of a TiO2-x encapsulating layer. However, compared to the Ni foil (Fig. 4c, e), the Ni particle size over Ni/TiO2-R is too large to notably contribute to metal-support bonding (Ni-Ti) (Fig. 4c, f)50. Therefore, changes in disorder (σ²) were analyzed using EXAFS, as tabulated in Supplementary Table S1. The highest static disorder was observed under reaction conditions (11.6 × 10⁻2 Å2), implying strong interaction between Ni particles and decorating TiO2-x species present in Ni/TiO2-R50,51. The observed reduction in Ni valence state is consistent with electron transfer from Ti to Ni, facilitated by the modification of A-SMSI.

To quantitatively elucidate this electron transfer mechanism, the reactive intermediate COOH* was introduced onto a pre-adsorbed TiO2-x/Ni(111)–HCOO* structure (representative of A-SMSI), as shown in Fig. 3g. Differential charge density and Bader charge analyses of the resulting TiO2-x/Ni(111)–HCOO*–COOH* configuration revealed a net electron transfer of 0.17 e⁻ from the TiO2-x overlayer to Ni (Supplementary Figs. 18, 19), consistent with the XAS measurements. Furthermore, the calculated Eads of COOH* on TiO2-x/Ni (111)–HCOO* confirmed that the overlayer provides a favorable environment for intermediate stabilization, consistent with the active role of A-SMSI systems proposed by Matsubu et al.8.

Collectively, these XAS and DFT findings elucidate the electron transfer mechanism governing A-SMSI evolution under CO2 hydrogenation conditions. On Ni/TiO2-R, HCOO* induces the formation of a TiO2-x encapsulation layer (EELS Ti L2,3-edge shift: ΔE = −0.6 eV, Fig. 2h), enabling directional electron transfer through TiO2-x to Ni. The non-stoichiometric TiO2-x thus acts as the functional component for COOH* adsorption, as supported by both DRIFTS and DFT. This process simultaneously decreases the surface coverage of atomic hydrogen and thereby suppressing the further hydrogenation of CO generated via the dehydroxylation step of COOH*8,14. These findings provide mechanistic insight into the evolution of A-SMSI and the induced selectivity shift in CO2 hydrogenation, as elaborated in the following section.

Tuning catalytic performance by A-SMSI
The influence of A-SMSI on the catalytic performance of CO2 hydrogenation (CO2/H2/N2 = 15/15/70) was further evaluated over different Ni/TiO2 catalysts across 200–500 °C. As shown in Fig. 5a–c, the product selectivity is strongly facet-dependent.

Generally, increasing temperature thermodynamically favors CO formation via the RWGS reaction. Nevertheless, Ni/TiO2-R consistently exhibited high CO selectivity (83% at 350 °C, even 100% at 230 °C (Fig. 5a), along with the highest CO production rate across the temperature range among three catalysts (Supplementary Fig. 20). In contrast, Ni/TiO2-S strongly favored CH4 formation, reaching selectivity of up to 88% at 350 °C and 98% at 230 °C under identical conditions (Fig. 5c), with the highest CH4 production rate (Supplementary Fig. 20). Ni/TiO2-B exhibited an intermediate behavior with a relatively stable CO and CH4 selectivity profile, with CH4 selectivity remaining above ~60% throughout the reaction (Fig. 5b). Importantly, the carbon balance was consistently >98% across all catalytic experiments (Supplementary Table 2), indicating negligible carbon deposition, in agreement with both TEM (Fig. 2) and DRIFTS observations (Fig. 3).

Kinetic experiments were carried out to quantify intrinsic activity of different Ni-TiO2 catalysts (Fig. 5d, e). Reaction temperature of below 400 °C was selected to have the reaction rates determined at relatively low CO2 conversions (i.e., <20%), avoiding the effect of diffusion control51. Arrhenius plots of Ln(rCH4) and Ln(rCO) versus T-1 plotted by production rates presented good linear correlation for all three catalysts. The apparent activation energies (Ea) of CH4 formation followed the order of Ni/TiO2-R (90.0 kJ mol−1) > Ni/TiO2-B (75.4 kJ mol−1) > Ni/TiO2-S (67.4 kJ mol−1), while those for CO formation were: Ni/TiO2-S (100.3 kJ mol−1) > Ni/TiO2-B (80.7 kJ mol−1) > Ni/TiO2-R (68.1 kJ mol−1). These results indicate that the presence of A-SMSI in Ni/TiO2-R effectively reduces the Ea for CO formation, thereby favoring the RWGS pathway.

Interestingly, for Ni/TiO2-S, CH4 yield remained nearly constant between 350 and 400 °C (17.9 mmol·gcat−1·h−1 → 18.1 mmol·gcat−1·h−1), while dropped to 14.3 mmol·gcat−1·h−1 at 450 °C, indicating a thermal optimum (Supplementary Fig. 20). A 100 h stability test confirmed that Ni/TiO2-S maintained >88% CH4 selectivity and 20% CO2 conversion, approaching the 23.6% thermodynamic limit of the Sabatier reaction under equimolar CO2/H2 (Fig. 5f). The experimental evidence confirms the absence of A-SMSI in Ni/TiO2-S over long-term reaction.

Collectively, these results unequivocally demonstrate that the extent of A-SMSI—governed by the crystallographic facet of the TiO2 support—plays a pivotal role in steering catalytic pathways and product distribution. Ni/TiO2-S, lacking A-SMSI, catalyzes nearly exclusive CH4 formation. In contrast, the fully developed A-SMSI in Ni/TiO2-R, characterized by complete TiO2-x encapsulation, strongly promotes CO formation. Ni/TiO2-B, with partial encapsulation, exhibits a balanced CO/CH4 output, bridging the two extremes.

Methods
Catalyst synthesis
Three Ni/TiO2 model catalysts were prepared using supports with different well‑defined morphologies.
TiO2-R (nanorod): A two-step hydrothermal approach was employed. First, P25 TiO2 (1000 mg) was converted into Na-titanate by stirring in NaOH (10 M 40 mL) for 10 min, then autoclaving at 120 °C for 24 h. The resulting white precipitate was washed (pH 10.5) and centrifuged (10000 r min−1, 10 min). Next, the wet Na-titanates (4000 mg) were dispersed in deionized water (152 mL) with H2O2 (8 mL), stirred for 15 min, and hydrothermally treated at 180 °C for 20 h, yielding anatase TiO2 nanorods. The precipitate was washed with deionized water/ethanol, then vacuum-dried (60 °C, 12 h).
TiO2-B (bipyramidal nanocrystal): Wet Na-titanate (3000 mg) was mixed with LiCl (393.2 mg) and deionized water (120 mL), stirred for 15 min, and hydrothermally treated at 200 °C for 4 h. The precipitate was washed, centrifuged, and vacuum-dried (60 °C, 12 h).
TiO2-S (nanosheet): Tetrabutyl titanate (10 mL) and HF (1.8 mL, 40 wt%) were mixed, stirred for 5 min, and hydrothermally treated at 200 °C for 20 h. The precipitate was washed with deionized water/ethanol, and NaOH (1 M), then vacuum-dried (60 °C, 12 h).
Ni/TiO2-R, Ni/TiO2-B, and Ni/TiO2-S catalysts were prepared by the wet impregnation method. Each TiO2 support was dispersed in Ni(NO3)2·6H2O solution (30 mL), stirred at 50 °C until solvent evaporation, then dried at 60 °C for 12 h. The solids were calcined in air at 500 °C for 2 h to afford fresh NiO/TiO2, and subsequently reduced in 10% H2/N2 (400 °C, 100 min) to yield the Ni/TiO2-R, Ni/TiO2-B, and Ni/TiO2-S catalysts with a nominal Ni loading of 10 wt%.

Catalytic evaluation
CO2 hydrogenation was carried out in a continuous-flow fixed-bed reactor with a quartz tube (inner diameter, 8 mm; length, 650 mm). Prior to reaction, fresh catalysts (NiO/TiO2-S, NiO/TiO2-B, and NiO/TiO2-R) were in situ reduced under 10% H2/N2 (flow rate: 100 mL min−1) at 400 °C for 100 min. The reaction was conducted in 15% CO2, 15% H2, and 70% N2 at a total flow rate of 200 mL min−1 at ambient pressure, corresponding to a GHSV of 12000 mL gcat−1 h−1. Gas flows were regulated by mass flow controllers, and the outlet flow rate was measured by a membrane flow meter. The outlet gas composition (CO2, H2, CO, CH4) was analyzed online using a NDIR gas analyzer (GASBOARD-3100, Wuhan cubic optoelectronic Co., Ltd.).
The CO2 conversion (X(CO2)), CH4 selectivity (S(CH4)), CO selectivity (S(CO)), CH4 production rate (r(CH4)), CO production rate (r(CO)), and carbon balance were calculated using the following equations:
X(CO2)% = (Fin*Cin(CO2) - Fout*Cout(CO2)) / (Fin*Cin(CO2)) * 100% (1)
S(CH4)% = n(CH4) / Σn(products) * 100% (2)
S(CO)% = n(CO) / Σn(products) * 100% (3)
r(CH4) (mmol gcat−1 h−1) = (Fin*Cin(CO2) * X(CO2) * S(CH4) * 60) / (22.4 * mcat) (4)
r(CO) (mmol gcat−1 h−1) = (Fin*Cin(CO2) * X(CO2) * S(CO) * 60) / (22.4 * mcat) (5)
Carbon balance% = (Fout*Cout(CO) + Fout*Cout(CH4) + Fout*Cout(CO2)) / (Fin*Cin(CO2)) * 100% (6)
Here, Fin and Fout (mL min−1) denote the inlet and outlet gas flow rates, respectively. C denotes the concentration (%), mcat denotes the weight of catalyst (g); and n denotes the molar quantity of the products.

Catalyst characterizations
ICP-OES. Ni loadings of the catalysts were quantified via ICP-OES (Agilent 720ES). Approximately 45–50 mg of each sample was digested in acid, diluted to 10 mL (V0), and measured. The elemental concentration (C0, mg L−1) obtained from the instrument was used to calculate the actual Ni mass percentage using the following formula:
W(Ni)% = (C0 * V0 * f) / (m * 10^9) * 100% (7)
Here, f represents the dilution ratio of C0 (set to 50), and m denotes the samples mass (mg).
TEM characterizations. The morphology, structure and composition of the catalysts were examined by TEM using an FEI Tecnai G2 F20 and an FEI Titan G2 60–300 KV Cs-corrected TEM with HAADF and EDX detectors to determine the distribution of Ni.
STEM characterizations. The HAADF STEM and EELS characterizations were performed on an FEI Titan G2 80–200 TEM/STEM equipped with a spherical aberration corrector (operated at 200 kV).
In situ ETEM experiments. In situ ETEM observations were performed using a Hitachi H-9500 microscope operated at 300 KV, equipped with a heating holder (up to 1000 °C) and enabled pressure around the sample reaching 5 × 10−2 Pa. All data were collected with a Oneview (Gatan) CMOS camera, which was processed using Digital Micrograph software. To minimize electron beam–induced damage, especially considering the high sensitivity of TiO2 supports to beam irradiation, strict low-dose conditions were employed, with the electron dose limited to approximately 6.14 × 103 e Å−2 throughout the ETEM experiments. Furthermore, to avoid cumulative exposure, the beam was activated only during image acquisition and the beam valve remained closed at all other times. For Ni/TiO2-R, -B, -S, individual Ni NPs with representative sizes were selected for tracking the structural evolution. In the in situ experiments, the catalysts were firstly reduced under H2 (5 × 10−2 Pa) at 400 °C to eliminate surface oxides. Subsequently, a CO2/H2/N2 gas mixture (15/15/70) was introduced to simulate CO2 hydrogenation conditions (5 × 10−2 Pa, 350 °C).
In situ DRIFTS experiments. The experiments were performed on a Nicolet iS50 spectrometer (Thermo Scientific) equipped with an in situ IR reaction cell, a gas flow system, and a liquid nitrogen-cooled mercury cadmium telluride (MCT) detector. Spectra were collected with 64 scans at a resolution of 4 cm–1. Prior to measurements, the Ni/TiO2 catalysts were pre-reduced under H2 (H2/N2 = 10/90) at 400 °C, to remove the NiO overlayer formed upon air exposure. The catalysts were then stabilized in N2 at 350 °C to ensure complete removal of adsorbed and absorbed hydrogen species. After acquiring a background spectrum under N2, a reactive gas mixture (CO2/H2/N2 = 15/15/75, 100 mL/min) was introduced at 350 °C to initiate catalytic reactions for the measurements. After a 30 min purge in the reactant mixture, the gas composition was switched to H2 (H2/N2 = 15/85), while maintaining the same flow rate to assess the stability of the adsorbed species.
CO2-TPD. CO2-TPD experiments were conducted on a Microtrac BELCat II. Each sample (100 mg) was loaded into a quartz reaction tube and pretreated by heating from room temperature to 400 °C at a rate of 1 °C min−1 under a H2 flow (H2/Ar = 10/90, 30 mL min−1) for 1 h to remove the surface NiO layer. After cooling to 50 °C, the sample was saturated with a CO2/He mixture (CO2/He = 10/90, 30 mL min–1) for 1 h. Subsequently, the system was purged with pure He (30 mL min−1) for another hour to remove weakly physically adsorbed CO2. Finally, the desorption profile was recorded by heating the sample to 500 °C at 10 °C min−1 under He flow, with CO2 desorption monitored using a thermal conductivity detector (TCD).
H2-TPR. H2-TPR experiments were carried out on a Microtrac BELCAT II instrument. Fresh NiO/TiO2-R, NiO/TiO2-B, and NiO/TiO2-S samples (each 50 mg) were loaded into a U-shaped quartz reactor. The samples were first pretreated by heating from room temperature to 300 °C at a ramp rate of 10 °C min−1 under a H2 flow (30 mL min−1) for 1 h. After cooling to 50 °C, a H2 mixture (H2/Ar = 10/90, 30 mL min−1) was introduced for 30 min to stabilize the baseline. The reduction profiles were then recorded by heating the sample to 600 °C at 10 °C min−1 under the same H2/Ar flow, while monitoring H2 consumption using a thermal conductivity detector (TCD).
In situ XAS experiments. The XAS of Ni K-edge experiments were performed at the BL20U1 Hard branch of energy material beamline (E-line) of the Shanghai Synchrotron Radiation Facility (SSRF) operating at the 3.5 GeV with an injection current of 220 mA. A Si (111) double-crystal monochromator was used for energy selection. Ni foil was used as a reference, and the measurements were conducted in transmission mode. Data processing was performed using the Demeter software package.
XPS. The XPS analysis was implemented on Thermo Scientific K-Alpha using the Al-Kα X-rays source. The treated catalysts were made into small tablets and placed on the sample holder in a glove box without exposure to air. The sample was then introduced into the ultra-high vacuum chamber for XPS measurements at room temperature.
DFT calculations
All DFT calculations were conducted using the projector augmented wave (PAW) method and the Perdew-Burke-Ernzerhof exchange-correlation functional, as implemented in the Vienna ab initio simulation package (VASP) 5.4.452,53. The long-range van der Waals (vdW) interaction was described by the DFT-D3 method54. The Brillouin-zone integration was performed using Gama k-mesh to ensure computational accuracy55. No Hubbard correction (LDAU) was applied to the Ni/TiO2 system. Full computational parameters are provided in Section Calculation method and Supplementary Figs. 10–15 of SI.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "This is a very interesting study that can have a strong impact in the way in which we think about metal-support interactions in catalysis. There are several key points that need to be taken into consideration or polished.\n1) Lines 49-68: From previous E-TEM studies, it is already known the CO2 and H2 can induce big changes in the morphology of metal-titania catalysts (Deng et al, ACS Catal. 2024, 14, 11822-11844). This is important and needs to be mentioned.\n2) Lines 70-90: Did the authors check that the relative concentration of the different titania facets was not affected by the electron beam?\n3) Figure 1, in the Ni panels, it is very hard to separate some of the dark red and black features.\n4) Lines 135-140, did the authors examine the effects of CO2 and H2 alone on the morphology of the Ni particles in Figure 2?\n5) Lines 170-190, in ref 2 (Monai et al, Science 380, 644–651 (2023)), the deposition of a lot of carbon was reported in similar experiments. Did the authors detect such a behavior?", "reviewer_2": "Manuscript titled ‘Facet-Dependent Adsorbate-Mediated Strong Metal-Support Interaction in Ni/TiO2’ presents in-situ investigations on the behavior of TiO2 supported Ni model catalysts. Authors report CO2 mediated strong metal-support interaction on three TiO2 facets, namely {100}, {101}, and {001}, which lead to complete, partial, and no encapsulation of Ni nanoparticles (NPs), respectively. The extent of encapsulation was found to modulate selectivity of CO2 hydrogenation dramatically from CH4 to CO, going from no encapsulation to complete encapsulation. The facet-dependency of the A-SMSI offers an explanation for the strategy to tune the selectivity towards either CH4 or CO by changing the TiO2 support facet.\nDue to major shortcomings in calculations discussed in detail below, experimental evidence on the role of adsorbates on encapsulation is not strong enough, and the discussion of results is mainly focused on technical detail, rather than giving chemical insight on the A-SMSI phenomenon, therefore the publication of the manuscript in Nature Communications is not recommended.\nPlease find a point-by-point review below:\n1. The authors refer to Density Functional Theory calculations on line 240 of the manuscript, and report some results in the Supplementary Information. However, the computational methods section seems to be missing from the manuscript. Therefore, there is not enough detail provided for the computational work to be reproduced. To properly evaluate the soundness of the DFT data, a method section including software, exchange correlation functional, Hubbard correction scheme, reciprocal space sampling etc. must be included. Further details are needed especially to address the following questions:\na. How were the Ni/TiO2 models selected and constructed?\ni. The Ni/TiO2 model systems seem to comprise small Ni clusters on TiO2 slabs. The manuscript lacks the justification of computational catalysts models both for the oxide support and metal component. How well do small clusters like these represent the real (pre-encapsulation) catalysts that have large ~13-15 nm Ni particles? How was the TiO2 models selected?\nii. How were the cluster geometries optimized?\niii. The supported Ni clusters seem to have different number of Ni atoms. If so, what is the justification for changing the size of the cluster? These systems seem to be in the ‘every atom counts’ size regime, which means changing the number of atoms from one model to the next makes comparisons between the models less straightforward.\nb. How were the possible oxygen vacancy locations sampled and the vacancy formation energies computed?\nc. How was the electronic structure of the oxygen vacancy containing structures analyzed? Are the vacancy electrons localized (and if so, where) or delocalized?\n2. The authors show in-situ DRIFT data which they interpret as evidence for HCOx surface species being responsible for the formation of the TiO2-x overlayers on the Ni/TiO2-R and Ni/TiO2-B samples, and state that the TiO2-S sample is unable to stabilize these species\na. How does the inability to stabilize those species explain why there is no encapsulation on Ni/TiO2-S?\nFurthermore, the authors draw a direct connection between the strength of the HCOO adsorption and the extent of encapsulation, i.e. stronger HCOO adsorption is related to greater encapsulation.\nb. What is the mechanistic explanation for this effect?\nc. It would be useful to calculate the HCOO* binding energy on each TiO2 facet (or compare to values in literature) to support this claim\nHowever, the mechanism for how exactly the HCOO* induces the migration of TiO2 onto the Ni particles remains unclear. The experimental evidence seems to be only circumstantial.\n3. The DFT calculations showed the trend Ni/TiO2-S (3.865 eV) > Ni/TiO2-B (3.555 eV) > Ni/TiO2-R (2.575 eV) in oxygen vacancy formation energies, and the authors state that the easy formation of oxygen vacancies in the Ni/TiO2-R sample “facilitates the transformation of the TiO2 interface into TiO2-x, providing a foundational driving force for A-SMSI formation”. They also speculate that oxygen vacancies can more readily be induced by HCOx* on Ni/TiO2-R, as the formation energy is the lowest. While it is known that non-stoichiometric TiO2 substrates more readily encapsulate metals. (Fu & Wagner, 2007, https://doi.org/10.1016/j.surfrep.2007.07.001), this is a rather strong claim drawn only from formation energies.\na. How exactly do the HCOx adsorbates induce oxygen vacancy formation? Can this claim be supported by DFT calculations?\nb. Why aren’t the oxygen vacancies formed already during the H2 pre-treatment?\nc. Is the role of the adsorbates merely to induce the creation of oxygen vacancy, which drives the encapsulation, or are they active participants in the migration of substrate onto the metal?\n4. How can the modification of Ni electronic structure be responsible for adsorbate stability if the particles are encapsulated by a ~1.7 nm thick amorphous TiO2 layer? What are the proposed active sites of the reaction?\na. Is it possible to construct models of these active sites and perform DFT calculations to confirm the effect on intermediate (such as COOH) stabilities?\n5. Overall, the results section lacks coherence, and it is not written in a clear way to make the connection between the experimental data and the conclusions drawn from them convincing enough.", "reviewer_3": "I co-reviewed this manuscript with one of the reviewers who provided the listed reports. This is part of the Nature Communications initiative to facilitate training in peer review and to provide appropriate recognition for Early Career Researchers who co-review manuscripts."}, "author_response": "Point-by-point response to Reviewers\nManuscript Number: NCOMMS-25-40560\nTitle: Facet-Dependent Adsorbate-Mediated Strong Metal-Support Interaction in Ni/TiO2\nReviewer comments:\nReviewer #1:\nGeneral Comments: This is a very interesting study that can have a strong impact in the way in which we think about metal-support interactions in catalysis. There are several key points that need to be taken into consideration or polished.\nResponse: We sincerely thank you for your positive and insightful comments on our work, describing it as \"very interesting\" and acknowledging its potential for \"strong impact\" in the field of metal-support interactions. According to your advice, we have carefully revised the manuscript. We believe the point-by-point responses below and the updated version of the manuscript addressed all the specific points you raised, and your valuable suggestions have significantly improved the quality and clarity of our manuscript.\nComment #1: Lines 49-68: From previous E-TEM studies, it is already known the CO and H can 2 2 induce big changes in the morphology of metal-titania catalysts (Deng et al, ACS Catal. 2024, 14, 11822-11844). This is important and needs to be mentioned.\nResponse: We thank you for highlighting this important reference. We have now cited this work and incorporated a brief discussion in the revised manuscript:\nPage 2, Lines 57–61: “Recently, by using in-situ environmental transmission electron microscopy (ETEM), Deng et al. further demonstrated that pretreatment in CO and H could dynamically and 2 2 reversibly alter the morphology of Cu@TiO catalysts, emphasizing that metal–support interfaces are x highly sensitive to the chemical environment. These findings highlight the inherently dynamic nature of A-SMSI and the importance of in-situ characterization15.”\nComment #2: Lines 70-90: Did the authors check that the relative concentration of the different titania facets was not affected by the electron beam?\nResponse: Thank you for this critical point regarding the potential influence of the electron beam on the TiO facets. We agree that the beam effect is a fundamental issue that should be addressed to ensure 2 1 reliable observations. Actually, our group has extensive prior experience and well-established methodologies for minimizing electron beam irradiation effects during in-situ ETEM studies of TiO 2 and TiO -supported metal catalysts1–6, which were systematically applied in the present work. 2 Specifically:\na) All ETEM experiments were conducted under a relatively low electron-dose conditions (~ 6.14× 103 e Å-2) using a Oneview (Gatan) CMOS camera to minimize electron beam damage to the catalysts. According to previous studies3,5–8, a dose rate below the threshold value of ~ 1×104 e Å- 2 is within the safe range to avoid electron-beam-induced damage of TiO . Moreover, the electron 2 beam was applied only during image acquisition, to avoid prolonged exposure or cumulative exposure. Under these conditions, the titania facets did not show obvious structural change.\nb) Control experiments were conducted in which the {100}, {101}, and {001} facets were continuously exposed to the same electron dose rate (~ 4.91×103 e Å-2 s-1) in vacuum (10-5 Pa) for 30 minutes. The surfaces remained flat and well-defined, and no obvious structural change has been observed, as shown in Fig. R1.\nFigure R1 (Supplementary Fig. 2 in Supplementary Information (SI)). The in-situ TEM images show the structural stability of different TiO facets under identical electron beam dose (~ 4.91×103 e 2 Å-2 s-1 for 30 min). (a, d) TiO -R; (b, e) TiO -B; (c, f) TiO -S. 2 2 2\nc) In the in-situ experiments, we also examined regions that were not exposed to the electron beam at low magnification. The morphologies of the three typical samples kept unchanged (Fig. R2), which further confirmed that, under our experimental conditions, the relative concentration of the different exposed titania facets remained unaffected. 2 Figure R2 (Supplementary Fig. 3 in SI). Low-magnification in-situ ETEM images of the three Ni/TiO samples recorded from regions not exposed to the electron beam before imaging. (a) Ni/TiO - 2 2 R; (b) Ni/TiO -B; (c) Ni/TiO -S. 2 2 2\nA corresponding clarification has now been added to Sections Results and Methods, and the supporting comparison images (Figs. R1, R2 above) are provided in the revised SI (Supplementary Figs. 2, 3).\nPage 3, Lines 96-99: “Given the sensitivity of TiO to electron beam irradiation, all subsequent 2 experiments were conducted under a relatively low-dose conditions referring to previous studies24-27. Moreover, control experiments confirmed that the electron beam did not obviously alter the relative stability and structure of different TiO facets (Supplementary Figs. 2, 3).” 2\nPage 15, Lines 429-433: “To minimize electron beam–induced damage, especially considering the high sensitivity of TiO supports to beam irradiation, strict low-dose conditions were employed, with 2 the electron dose limited to approximately 6.14×103 e Å-2 throughout the ETEM experiments. Furthermore, to avoid cumulative exposure, the beam was activated only during image acquisition and the beam valve remained closed at all other times.”\nComment #3: Figure 1, in the Ni panels, it is very hard to separate some of the dark red and black features.\nResponse: Thank you for the valuable comment regarding the clarity of the Ni elemental maps in Figure 1. To improve clarity, we have reprocessed the EDX elemental maps in Fig. 1(g-i) using an enhanced colormap and optimized brightness/contrast (Fig. R3(g-i) below).\nComment #4: Lines 135-140, did the authors examine the effects of CO and H alone on the 2 2 morphology of the Ni particles in Figure 2?\nResponse: As shown in Fig. 2a of the manuscript, after a H atmosphere (H /N = 10/90) pretreatment 2 2 2 at 400 °C, Ni/TiO -R displayed no noticeable changes in Ni morphology and lattice features. To further 2 clarify the effect of CO alone, we conducted additional in-situ experiments on pre-reduced Ni/TiO -R 2 2 under a CO atmosphere (CO /N = 10/90, 5×10-2 Pa) at 350 °C. Under these conditions, amorphous 2 2 2 surface features appeared, consistent with preview work9, but no significant overlayer growth or TiO 2 encapsulation was observed (see Fig. R4 below). It was explained that the dissociative adsorption of CO on Ni nanoparticles can produce amorphous surface layers with blurred or indiscernible lattice 2 fringes, attributed to a rearrangement of surface Ni atoms.\nFigure R4 (Supplementary Fig. 7 in SI). The morphology of Ni/TiO -R catalysts under CO 2 2 atmospheres. (a) Ni/TiO under CO atmosphere (CO /N = 10/90, 5×10-2 Pa, 350 °C) and (b, c) FFT 2 2 2 2 patterns of the boxed region in (a).\nPage 6, Lines 150-155: “The formation of the encapsulating overlayer requires the synergy of CO and 2 H rather than either gas alone. In contrast to H , which caused no significant morphological change 2 2 (Fig. 2a), CO alone induced surface amorphization of the Ni nanoparticles but did not lead to 2 encapsulation. This amorphization, evidenced by the appearance of blurred lattice fringes (Supplementary Fig. 7), aligns with established understanding which attributes it primarily to the dissociative adsorption of CO that reconstructs the Ni surface atoms31.” 2\nComment #5: Lines 170-190, in ref 2 (Monai et al, Science 380, 644–651 (2023)), the deposition of a lot of carbon was reported in similar experiments. Did the authors detect such a behavior?\nResponse: We thank you for the insightful comment. In our experimental conditions, no notable carbon deposition was observed. It could be attributed to the difference in reaction atmospheres (CO /H in our 2 2 work vs. CO/CO /H in Monai et al., Science 2023, 380, 644-651). The presence of CO favored carbon 2 2 deposition via the Boudouard reaction (2CO ⇌ CO + C), as suggested by Monai et al. in their work. 2\nIn-situ ETEM imaging (Fig. 2(a-d) in the manuscript) showed clean surfaces without amorphous carbon layers or carbon filaments and ex-situ EELS mapping (Fig. 2(e-l) in the manuscript) further confirmed the absence of carbon accumulation. Moreover, the in-situ DRIFTS spectra (Figs. 3 (a-c) in the manuscript) detected carbonate and formate intermediates (e.g., CO * and HCO *), but no features 3 2 of polymeric carbon species.\nFurthermore, carbon balance calculations (Eq. (1)) showed >98% (mostly >99%) across all catalysts 5 (Table R1), confirming negligible carbon deposition in our experiments.\n𝐹 ∗𝐶 (𝐶𝑂)+𝐹 ∗𝐶 (𝐶𝐻 )+𝐹 ∗𝐶 (𝐶𝑂 ) 𝐶𝑎𝑟𝑏𝑜𝑛 𝑏𝑎𝑙𝑎𝑛𝑐𝑒 %= 𝑜𝑢𝑡 𝑜𝑢𝑡 𝑜𝑢𝑡 𝑜𝑢𝑡 4 𝑜𝑢𝑡 𝑜𝑢𝑡 2 (1) 𝐹 ∗𝐶 (𝐶𝑂 ) 𝑖𝑛 𝑖𝑛 2\nTable R1 (Table 2 in SI). Carbon balance under CO hydrogenation conditions over different catalysts. 2\nTemperature Carbon balance over Carbon balance over Carbon balance over (°C) Ni/TiO -R (%) Ni/TiO -B (%) Ni/TiO -S (%) 2 2 2\n230 98.9 99.9 99.8\n290 99.1 100.0 99.3\n350 100.0 99.8 98.0\nA clarification has been added to the revised manuscript to address this point explicitly.\nPage 13, Lines 338-340: “Importantly, the carbon balance was consistently >98% across all catalytic experiments (Supplementary Table 2), indicating negligible carbon deposition, in agreement with both TEM (Fig. 2) and DRIFTS observations (Fig. 3).” 6\nReviewer #2:\nGeneral Comments: Manuscript titled “Facet-Dependent Adsorbate-Mediated Strong Metal-Support Interaction in Ni/TiO ” presents in-situ investigations on the behavior of TiO supported Ni model 2 2 catalysts. Authors report CO mediated strong metal-support interaction on three TiO facets, namely 2 2 {100}, {101}, and {001}, which lead to complete, partial, and no encapsulation of Ni nanoparticles (NPs), respectively. The extent of encapsulation was found to modulate selectivity of CO 2 hydrogenation dramatically from CH to CO, going from no encapsulation to complete encapsulation. 4 The facet-dependency of the A-SMSI offers an explanation for the strategy to tune the selectivity towards either CH or CO by changing the TiO support facet. 4 2\nDue to major shortcomings in calculations discussed in detail below, experimental evidence on the role of adsorbates on encapsulation is not strong enough, and the discussion of results is mainly focused on technical detail, rather than giving chemical insight on the A-SMSI phenomenon, therefore the publication of the manuscript in Nature Communications is not recommended.\nResponse: We sincerely thank you for the thorough and constructive feedback. We agree that our original submission was primarily designed as an experiment-driven study focusing on the in-situ investigation of facet-dependent A-SMSI, while the theoretical analysis was only briefly included as complementary support. This limited theoretical depth indeed result in insufficient chemical insight into the A-SMSI phenomenon.\nIn the revised manuscript, we have therefore undertaken substantial improvements to the theoretical analysis in response to the reviewer’s suggestions. Specifically, we have (i) supplemented complete computational methods to ensure full reproducibility; (ii) constructed more representative Ni/TiO 2 models with different cluster sizes; (iii) systematically calculated oxygen vacancy formation energies, vacancy-induced electronic structures, and electron localization functions; (iv) performed additional adsorption and transition-state calculations for key intermediates HCOO* to clarify their roles in stabilizing and inducing TiO overlayers; and (v) performed additional adsorption for COOH* to 2-x clarify the active role of TiO . These new calculations provide crucial mechanistic insights that were 2-x missing in the original version, including the dual role of adsorbates in both vacancy formation and overlayer stabilization, as well as the identification of TiO overlayers as the most probable active 2-x species.\nTogether with our extensive in-situ experimental evidence, the integration of these new theoretical results now establishes a coherent and chemically insightful picture of facet-dependent adsorbate- mediated SMSI. We believe that these major revisions have significantly strengthened the manuscript and think that the improved balance of experiment and theory will meet the high standards of Nature Communications.\nThe corresponding details and revisions are provided in the point-by-point responses below, as well as 7 in the revised manuscript and Supplementary Information (SI).\nComment #1: The authors refer to Density Functional Theory calculations on line 240 of the manuscript, and report some results in the Supplementary Information. However, the computational methods section seems to be missing from the manuscript. Therefore, there is not enough detail provided for the computational work to be reproduced. To properly evaluate the soundness of the DFT data, a method section including software, exchange correlation functional, Hubbard correction scheme, reciprocal space sampling etc. must be included. Further details are needed especially to address the following questions:\nResponse: We sincerely thank you for the thorough feedback and apologize for missing the detailed computational methods in the original submission. We have added sections in both the manuscript (Methods: DFT Calculations) and SI (Calculation Method) that provide full methodological details for both the calculations in the original manuscript (oxygen vacancy formation energy) and additional calculations in the revised manuscript (heterojunction model and electronic structure analysis of heterojunction model, adsorption energy of intermediates, and transition state calculations of HCOOH*).\nPage 17, Lines 472-479 in the manuscript:\nDFT calculations\nAll DFT calculations were conducted using the projector augmented wave (PAW) method and the Perdew-Burke-Ernzerhof exchange-correlation functional, as implemented in the Vienna ab initio simulation package (VASP) 5.4.454,55. The long range van der Waals (vdW) interaction was described by the DFT-D3 method56. The Brillouin-zone integration was performed using Gama k-mesh to ensure computational accuracy57. No Hubbard correction (LDAU) was applied to the Ni/TiO system. Full 2 computational parameters are provided in Section Calculation method and Supplementary Figs. 10-15 of SI.\nPage 9 in the Supplementary Information:\nCalculation Method:\nOxygen vacancy formation energy: The plane-wave cut-off energy was set to 400 eV and a vacuum layer with a thickness of 15 Å was introduced. The self-consistent-field iteration was well converged to 10-5 eV. A conjugate-gradient algorithm was used to relax the ions into their ground states, and the forces on each ion were converged within 0.01 eV Å-1. The oxygen vacancy formation E energy was 8 ov calculated as6-12.\n1 𝐸 = 𝐸 − 𝐸 + 𝐸 (1) 𝑜𝑣 𝑑𝑒𝑓𝑒𝑐𝑡 𝑝𝑒𝑟𝑓𝑒𝑐𝑡 2 𝑂2\nwhere 𝐸 and 𝐸 correspond to the total energies of the oxygen-vacancy model and the 𝑑𝑒𝑓𝑒𝑐𝑡 𝑝𝑒𝑟𝑓𝑒𝑐𝑡 initial perfect model, respectively. 𝑬 was the energy of an oxygen molecule. 𝑶𝟐\nHeterojunction model and transition state calculations: The cut-off energy for plane wave was set to 450 eV. A vacuum layer with a thickness of 12 Å was added perpendicular to the sheet and the lateral distances between adsorbed molecules were maintained at a minimum of 8 Å to avoid artificial interaction between periodic images. During geometry optimization, all atomic coordinates were relaxed with the convergence criteria of 0.03 eV·Å-1 for the forces on each atom, except for the atoms in the bottom layer of the support. For each self-consistent iteration, the total electronic energy was converged to 10−5 eV. Transition states were optimized using the climbing image nudged elastic band13 and dimer14 methods, and each saddle point was confirmed by the presence of a single imaginary vibrational frequency. The charge density difference was calculated as:\n𝑇𝑖𝑂 ∆𝜌(𝑟)= 𝜌𝑁𝑖 − 𝜌𝑇𝑖𝑂− 𝜌𝑁𝑖 (2) where 𝝆𝑇𝑖𝑂/𝑁𝑖 is the electron density of the interacting system, while 𝝆𝑇𝑖𝑂 and 𝝆𝑁𝑖 are the electron densities of the TiO film with adsorbates and Ni slab calculated for the isolated systems, respectively, 2 with the atomic positions fixed as in the composite cell. The electron density difference isosurfaces were visualized with VESTA15.\nComment #1-a: How were the Ni/TiO models selected and constructed? 2\nComment #1-a-i: The Ni/TiO model systems seem to comprise small Ni clusters on TiO slabs. The 2 2 manuscript lacks the justification of computational catalysts models both for the oxide support and metal component. How well do small clusters like these represent the real (pre-encapsulation) catalysts that have large ~13-15 nm Ni particles? How was the TiO models selected? 2\nResponse: Indeed, Ni nanoparticles of ~13-15 nm may contain hundreds of atoms. However, to balance computational cost with accuracy, we adopted spherical Ni nanoclusters with 13 atoms, which have been widely used in prior CO hydrogenation studies as representative models for nanoparticles of 10 2 to 15 nm in size10–12. To further validate this choice, we also performed calculations of oxygen vacancy formation on Ni and Ni clusters supported on TiO (100), which yielded consistent trends (see the 6 9 2 response to Comment #1-a-iii for details).\nFor the oxide support, anatase TiO unit cell model was selected to match the experimental material. 2 The optimized anatase TiO unit cell was cleaved and expanded in Material Studio 2020 to generate 2 9 slabs exposing the (100), (001) and (101) facets, corresponding to TiO -R, TiO -B, and TiO -S, 2 2 2 respectively. Consecutive slabs were separated by a vacuum region of at 15 Å. After placing the Ni 13 cluster, the bottom two layers of Ti and O atoms were fixed, followed by structural optimization. The plane-wave cut-off energy was set to 400 eV, which was found to be high enough for energy convergence, and the self-consistent-field iteration was well converged to 10-5 eV. A conjugate-gradient algorithm was used to relax the ions into their ground states, and the forces on each ion were converged within 0.01 eV Å-1. The above process follows procedures established in the literature10–16.\nFull parameter settings and model construction details, including facet generation and Ni cluster sizes, are now provided in the revised manuscript (Section Method: DFT calculations) and Supplementary Information (Supplementary Figs. 10,13).\nComment #1-a-ii: How were the cluster geometries optimized?\nResponse: In the computational model, Ni clusters were used as the metal component. For structural 13 optimization, each cluster was encapsulated in a 10×10×10 Å vacuum layer, with a 1×1×1 K-point mesh. The plane-wave cutoff energy, as well as the convergence criteria for force and energy, were kept consistent with those used in the optimization of TiO . These details are now included in the revised 2 manuscript (Section Method: DFT calculations) and Supplementary Information (Section Calculation Method).\nComment #1-a-iii: The supported Ni clusters seem to have different number of Ni atoms. If so, what is the justification for changing the size of the cluster? These systems seem to be in the ‘every atom counts’ size regime, which means changing the number of atoms from one model to the next makes comparisons between the models less straightforward.\nResponse: As mentioned above, all Ni clusters in the original manuscript consist of 13 atoms without changing in different models. To examine the possible size effect on oxygen vacancy formation, we additionally performed calculations on Ni and Ni clusters supported on TiO (100) (see Fig. R5 below). 6 9 2 The corresponding vacancy formation energies were 3.77 eV, 4.19 eV, and 2.67 eV for Ni , Ni , and 6 9 Ni /TiO -R, respectively (see Fig. R6 below), all lower than that of bare TiO (100) (4.48 eV) (see Fig. 13 2 2 R7 below). These results confirm a general trend that oxygen vacancies form more readily upon Ni loading on Ni/TiO -R, independent of the precise cluster size. 2\nThe rationale for selecting the Ni cluster has been clarified and is now included in Supplementary 13 Figs. 12 and 13 (and the corresponding notes).\n“Supplementary Note 8: To reduce the possible sensitivity of the results to a specific cluster geometry, we systematically varied the number of Ni atoms supported on TiO -R (Ni , Ni , and Ni ). Among them, 2 6 9 13 the Ni and Ni clusters were spherical, while the Ni cluster was disk-shaped. This size-dependent 6 13 9 comparison allows us to test the robustness of the calculated trends. The optimization parameters for Ni and Ni clusters were kept consistent with those used for Ni .” 6 9 13\nFigure R6 (Supplementary Fig. 13 in SI). Oxygen vacancy formation energies (E ) for Ni , Ni and 13 9 ov Ni /TiO -R. (a) Ni /TiO -R; (b) Ni /TiO -R; (c) Ni /TiO -R. 6 2 13 2 9 2 6 2\n11 “Supplementary Note 9: The calculated formation energies of surface three-coordinated O atoms are 3.77 eV, 4.19 eV, and 2.67 eV, for Ni , Ni , and Ni /TiO , respectively, all lower than that of bare TiO 6 9 13 2 2 (100) (4.48 eV). This confirms a general trend that oxygen vacancies form more readily after Ni loading. While different cluster sizes yield consistent results, the Ni cluster was selected as the representative 13 model due to its spherical morphology, computational tractability, and broad use in prior literature8-10.”\nComment #1-b: How were the possible oxygen vacancy locations sampled and the vacancy formation energies computed?\nResponse: In the original manuscript, we focused on the interfacial oxygen atoms for calculating oxygen vacancy formation energies on Ni/TiO , as they are directly involved in SMSI17,18. To provide 2 a more complete comparison, we further examined surface oxygen sites on the TiO (100) facet. The 2 surface oxygen vacancies on TiO (100) are classified into two-coordinate and three-coordinate 2 vacancies. For the unloaded TiO (100) slab, the oxygen vacancy formation energies are 4.48 eV (two- 2 coordinate O) and 4.86 eV (three-coordinate O), respectively.\nThe calculation results of oxygen vacancy formation energy (E ) on bare TiO (100) have been added ov 2 to Supplementary Fig. 14 (see Fig. R7 below). The calculation equation has been provided above in Eq. (1), and in SI (Calculation method).\nFigure R7 (Supplementary Fig.14 in SI). Oxygen vacancy formation energies (E ) for TiO (100). ov 2\nComment #1-c: How was the electronic structure of the oxygen vacancy containing structures analyzed? Are the vacancy electrons localized (and if so, where) or delocalized?\nThese ELF results have been added to Supplementary Fig. 15\nFigure R8 (Supplementary Fig. 15 in SI). ELF Calculations for TiO (100)-ov and Ni /TiO -R-ov. 2 13 2\nComment #2: The authors show in-situ DRIFT data which they interpret as evidence for HCO surface x species being responsible for the formation of the TiO overlayers on the Ni/TiO -R and Ni/TiO -B 2-x 2 2 samples, and state that the TiO -S sample is unable to stabilize these species. 2\nComment #2-a: How does the inability to stabilize those species explain why there is no encapsulation on Ni/TiO -S? 2\nResponse: We thank you for raising this important question. We have clarified this point in the revised manuscript. In short, both the formation of oxygen vacancies and the stabilization of TiO species by 2-x surface intermediates are less favorable on Ni/TiO -S than that on Ni/TiO -R, which disfavors 2 2 overlayers formation.\n1) High formation energy of oxygen vacancies on Ni/TiO -S: 2\nDFT calculations on three representative Ni/TiO interfacial models indicate that oxygen vacancy 2 formation is most favorable on Ni/TiO -R (2.67 eV), whereas the higher formation energy on Ni/TiO - 2 2 S (3.34 eV) suggests that fewer TiO species are generated initially (Supplementary Fig 11). Since 2-x non-stoichiometric TiO substrates are known to encapsulate metals more readily19, the scarcity of TiO 2 2- species rationalizes the absence of encapsulation on Ni/TiO -S. x 2\n13 2) Insufficient stabilization of TiO by adsorbates on Ni/TiO -S: 2-x 2\nOn Ni/TiO -R, adsorbates such as HCOO* stabilize TiO , via favorable adsorption (negative energies), 2 2-x lowering the total system energy and preserving TiO species. However, on Ni/TiO -S, these 2-x 2 intermediates are rapidly converted into subsequent species, as shown by in-situ DRIFTS, providing little stabilization for TiO and thereby preventing the formation of a stable TiO overlayer. 2-x 2-x\nTo elucidate the role and mechanism of these adsorbates, we performed DFT calculations on a TiO 2 (100)/Ni (111) heterostructure with minimized lattice mismatch (<1%), introducing oxygen vacancies to model TiO species20 (Fig. R9). This model represents the Ni/TiO -R catalyst, since TiO -R 2-x 2 2 primarily exposes the (100) facet experimentally. Considering that HCOO* has been previously demonstrated to form typically via transformation of HCO *21, we focused our calculations on HCOO* 3 adsorption. The calculations show that HCOO* can be stably adsorbed on the TiO surface of the TiO - 2-x 2 x/Ni (111) heterojunction (Fig. R10), with adsorption energies (E ) of -4.60 eV, -4.70 eV, and -3.72 eV ads at three representative sites.\nIn all cases, the negative E indicate thermodynamically favorable binding and consistent stabilization, ads suggesting that HCOO* adsorption on TiO sites of Ni/TiO -R likely further stabilize the reduced 2-x 2 surface, which could in turn favor the formation of TiO overlayers and encapsulation. In contrast, the 2-x lack of stable HCOO* intermediates on Ni/TiO -S removes this stabilizing effect. 2\nFigure R9 (Supplementary Fig. 9 in SI). Optimized structures of Ni (111) supported TiO (100) 2 overlayers: (a) stoichiometric surface TiO (100)/Ni (111) and (b) surface with oxygen vacancies TiO 2 2- /Ni (111). Red, gray, and blue atoms represent O, Ti, and Ni, respectively. x\nComment #2-b: What is the mechanistic explanation for this effect?\nResponse: As discussed above, the mechanistic explanation is that the stable adsorption of HCOO* intermediates on Ni/TiO -R provides essential thermodynamic stabilization to the reduced TiO 2 2-x species, which is a requisite for overlayer formation. Our in-situ DRIFTS results confirm that HCOO* persists on Ni/TiO -R but not on Ni/TiO -S. Complementary DFT calculations further demonstrate that 2 2 HCOO* adsorption on the TiO /Ni is strongly exothermic, lowering the system energy and stabilizing 2-x the otherwise metastable TiO overlayer. In contrast, the inability of Ni/TiO -S to stabilize these 2-x 2 intermediates leaves the reduced TiO species vulnerable, preventing the development of a stable 2-x encapsulating overlayer.\nComment #2-c: It would be useful to calculate the HCOO* binding energy on each TiO facet (or 2 compare to values in literature) to support this claim. However, the mechanism for how exactly the HCOO* induces the migration of TiO onto the Ni particles remains unclear. The experimental evidence 2 seems to be only circumstantial.\nResponse: Thank you for this valuable comment. Our in-situ DRIFTS experiments already indicated that HCOO* cannot be stabilized on Ni/TiO -S, and H -pulse experiments further revealed that the 2 2 adsorption stabilization follows the order Ni/TiO -R > Ni/TiO -B (Fig. 3d). Following your suggestion, 2 2 we performed additional calculations on the adsorption of HCOO* (adsorbed formate) on the (100), (001), and (101) facets of TiO , considering both monodentate and bidentate coordination adsorption 2 configurations (see Fig. R11 below). The results show that bidentate adsorption is more stable for HCOO*22,23, with TiO (001) exhibiting the strongest adsorption, followed by TiO (100), and TiO (101) 2 2 2 15 showing the weakest24.\nInterestingly, this trend on pristine TiO surface differs from the order observed in our in-situ DRIFTS 2 experiments for Ni/TiO (Ni/TiO -R (100) > Ni/TiO -B (101) > Ni/TiO -S (001)). To further probe this 2 2 2 2 discrepancy, we compared the adsorption on the isolated TiO (100) surface with that on a TiO (100) 2 2-x overlayer supported on Ni (111), as discussed above. In this case, the HCOO* adsorption energy was dramatically enhanced (-4.70 eV) compared to that on the isolated TiO (100) surface (-0.99 eV), 2 highlighting that a metal-support synergistic effect, rather than the intrinsic oxide properties, governs the adsorption behavior.\nMechanistically, we emphasize that HCOO* is unlikely to act as the direct driving force for TiO 2 migration. Rather, HCOO* facilitates the formation of oxygen vacancies, leading to the generation of TiO (as further elaborated in the response to Comment #3 below). In this process, the oxygen vacancy 2-x formation energy thus emerges as a contributing factor. Reduced TiO species are generally more 2-x prone to encapsulation, while HCOO* serves to stabilize the overlayer once it has formed. Therefore, what may appear as circumstantial experimental evidence actually provides a fundamental basis for our calculations and aligns with the proposed mechanism.\nFigure R11. Adsorption configurations and corresponding energies (E ) of HCOO* on different TiO ads 2 facets. (a-c) Monodentate adsorption configurations on the TiO (100), (101), (001). (d-f) Bidentate 2 adsorption configurations on the TiO (100), (101), (001). 2\nA corresponding clarification has now been added to the revised manuscript (Section: Revelation of catalyst adsorbates and the mechanism of encapsulation induction) to address this point explicitly.\nPage 9, Lines 223-234: “The density functional theory (DFT) calculations further rationalize these observations using a TiO /Ni (111) heterostructure model (Supplementary Fig. 9)3. The calculations 2-x 16 showed that HCOO* can be stably adsorbed on the TiO surface (Fig. 3g), with adsorption energies 2-x (E ) of -4.60 eV, -4.70 eV and -3.72 eV at three representative sites. In all cases, the negative E ads ads indicate thermodynamically favorable binding, suggesting that HCOO* adsorption on the TiO 2-x overlayer of Ni/TiO -R provides additional stabilization to the reduced surface, which in turn favor the 2 formation of TiO overlayers and promotes encapsulation. 2-x\nCombined DRIFTS and DFT results indicate that the adsorption strength and coverage of HCOO* species critically affect the extent of A-SMSI. On Ni/TiO -R, strongly bound HCOO* stabilizes the 2 TiO overlayer, while on Ni/TiO -B, weaker adsorption only enables partial encapsulation, as 2-x 2 evidenced in Figs. 2i, j and Supplementary Fig. 5. In contrast, the absence of stable HCOO* intermediates on Ni/TiO -S eliminates this stabilizing effect, preventing overlayer formation.” 2\nComment #3: The DFT calculations showed the trend Ni/TiO -S (3.865 eV) > Ni/TiO -B (3.555 eV) > 2 2 Ni/TiO -R (2.575 eV) in oxygen vacancy formation energies, and the authors state that the easy 2 formation of oxygen vacancies in the Ni/TiO -R sample “facilitates the transformation of the TiO 2 2 interface into TiO , providing a foundational driving force for A-SMSI formation”. They also speculate 2-x that oxygen vacancies can more readily be induced by HCO * on Ni/TiO -R, as the formation energy is x 2 the lowest. While it is known that non-stoichiometric TiO substrates more readily encapsulate metals. 2 (Fu & Wagner, 2007, https://doi.org/10.1016/j.surfrep.2007.07.001), this is a rather strong claim drawn only from formation energies.\nComment #3-a: How exactly do the HCO * adsorbates induce oxygen vacancy formation? Can this x claim be supported by DFT calculations?\nResponse: We appreciate your insightful comment. Previous studies have indeed shown that at high coverages of HCO * species, particularly formic acid (HCOOH*), and at temperatures above ~100 °C, x oxygen vacancies can form on TiO via H O desorption, leaving a HCOO-covered reduced TiO 2 2 2-x surface25,26. As noted, once TiO is transformed into non-stoichiometric TiO species, migration of the 2 2-x reduced oxide can lead to encapsulation of the active metal19. This prior knowledge underpins our mechanistic interpretation.\nTo further substantiate this mechanism, we performed DFT transition-state calculations on a Ni /TiO - 13 2 R model. HCOOH* is known as a key intermediate in CO hydrogenation over Ni-based catalysts, which 2 can further decompose into HCO* and OH*, with HCO* subsequently hydrogenating to CH 27. However, 4 our DFT calculations indicate that HCOOH* preferentially dissociates into HCOO* and H*, thereby following the oxygen-vacancy formation pathway described above.\nSpecifically, two competing pathways were computed: (path a) the conventional hydrogenation 17 sequence and (path b) the vacancy formation route. Dissociation along path a requires overcoming a barrier of ΔE = 0.573 eV, whereas path b proceeds with a much lower barrier of ΔE = 0.193 eV, producing HCOO* and H* (Fig. R12). These results clearly demonstrate that the oxygen-vacancy formation pathway is kinetically highly favorable over Ni/TiO -R, thereby providing direct support for 2 the proposed mechanism.\nFigure R12 (Fig. 3h in MS). DFT transition-state calculations for HCOOH* over Ni/TiO -R. Blue, gray, 2 red, black, and white spheres are Ni, Ti, O, C and H atoms, respectively.\nThis mechanism has been clarified in the revised manuscript:\nPage 9, Lines 235-248: “As noted, once TiO is reduced to non-stoichiometric TiO species, migration 2 2-x of the reduced oxide can lead to encapsulation of the active metal35. Previous studies have shown that formic acid (HCOOH*) can create oxygen vacancies on TiO via H O desorption, leaving a HCOO- 2 2 covered reduced surface (A-SMSI)8,44. To probe this mechanism, DFT calculations were performed for HCOOH* intermediate on Ni/TiO -R. Although HCOOH* can decompose into HCO* and OH* during 2 CO hydrogenation over Ni-based catalysts (with HCO* subsequently hydrogenating to CH )45, the 2 4 calculations indicate a more favorable dissociation into HCOO* and H* on Ni/TiO -R. 2\nTwo competing pathways were evaluated: (path a) the conventional hydrogenation sequence and (path b) the vacancy formation route. Dissociation along path a requires ΔE = 0.573 eV, whereas path b proceeds with a much lower barrier of ΔE = 0.193 eV (Fig. 3h). These results clearly demonstrate that the oxygen-vacancy formation pathway is kinetically highly favorable, thereby providing direct support for the proposed mechanism. Moreover, DFT calculations revealed that interfacial oxygen vacancy formation is most favorable on Ni/TiO -R (2.67 eV) among three catalysts, providing a fundamental 2 basis for TiO overlayer formation (Supplementary Figs. 10-15).” 2-x\nComment #3-b: Why aren’t the oxygen vacancies formed already during the H pre-treatment? 2\nFigure R13 (Fig. 3f in MS). H -TPR profiles of various Ni/TiO catalysts. 2 2\nComment #3-c: Is the role of the adsorbates merely to induce the creation of oxygen vacancy, which drives the encapsulation, or are they active participants in the migration of substrate onto the metal?\nResponse: As discussed above, the DFT results demonstrate a dual role of adsorbates such as HCOO*: (i) they facilitate the generation of TiO oxygen vacancies, in line with the well-established notion that 2 non-stoichiometric TiO substrates more readily encapsulate metals due to the associated lowering of 2 the system energy19,30; and (ii) they act as stabilizing factors for the encapsulating overlayer after its formation.\nWe hope this mechanistic clarification fully addresses the reviewer’s concern.\nResponse: We thank you for these insightful questions. The electronic modification of Ni primarily originates from the TiO overlayer, while the stabilization of adsorbates is also governed by this 2-x overlayer. Thus, the connection between the two is indirect, with the TiO overlayer itself acting as the 2-x most probable active component.\nTo clarify this point, additional DFT calculations were performed on a TiO /Ni (111)-HCOO* model 2-x (the most favorable configuration shown in Fig. R10). Based on this model, we examined adsorption of the key intermediate COOH* (confirmed by in-situ DRIFTS) at three representative sites on the TiO 2-x overlayer. As shown in Fig. R14, the calculated E are –2.49, –2.28, and –1.71 eV, suggesting that ads the overlayer provides a favorable environment for intermediate stabilization. This finding aligns with the proposed active sites in A-SMSI systems, as suggested by Matsubu et al.26.\nFigure R14 (Supplementary Fig. 18 in SI). Adsorption energies (E ) of COOH* adsorbed on three ads sites of the TiO /Ni (111)–HCOO* heterojunction. 2-x\nFurthermore, differential charge density and Bader charge analyses of the most stable configuration reveal that 0.17 e is transferred from TiO to Ni (Fig. R15), leading to an electron-enriched Ni state. 2-x This finding is consistent with our in-situ XANES measurements, where the Ni absorption edge shifts to slight lower energy under CO /H exposure compared with H alone (Fig. 4b in the manuscript). 2 2 2\nIn summary, our results suggest that (i) the stability of adsorbed intermediates primarily originates from the TiO overlayer, and (ii) Ni electronic modification arises from interfacial charge transfer (~0.17 e) 2-x from TiO to Ni, as corroborated by both DFT and XANES experiments. 2-x\nComment #4-a: Is it possible to construct models of these active sites and perform DFT calculations to confirm the effect on intermediate (such as COOH*) stabilities?\nResponse: As mentioned above, the TiO overlayer is proposed to stabilize key intermediates (COOH*) 2-x and likely play a direct role as an active site in A-SMSI — a statement consistent with the original proposal by Matsubu et al.26. A clarification addressing this point has been added to the revised manuscript (Section: Electronic structure evolution of Ni induced by A-SMSI).\nPage 11, Lines 303-310: “To quantitatively elucidate this electron transfer mechanism, the reactive intermediate COOH* was introduced onto a pre-adsorbed TiO /Ni (111)–HCOO* structure 2-x (representative of A-SMSI), as shown in Fig. 3g. Differential charge density and Bader charge analyses of the resulting TiO /Ni (111)–HCOO*–COOH* configuration revealed a net electron transfer of 2-x 0.17 e⁻ from the TiO overlayer to Ni (Supplementary Figs. 18, 19), consistent with the XAS 2-x measurements. Furthermore, the calculated E of COOH* on TiO /Ni (111)–HCOO* confirmed that ads 2-x the overlayer provides a favorable environment for intermediate stabilization, consistent with the active role of A-SMSI systems proposed by Matsubu et al.8”\nComment #5. Overall, the results section lacks coherence, and it is not written in a clear way to make the connection between the experimental data and the conclusions drawn from them convincing enough.\nResponse: We thank you for this constructive comment. As described above, in the revised manuscript, 21 we have thoroughly reorganized the Results section to improve both clarity and coherence. In brief:\ni. Integration of DFT with experiments:\nRather than presenting the computational and experimental results separately, we now integrate the DFT calculations into the relevant experimental discussions. For example:\n➢ Stabilization of the encapsulating overlayer: In-situ DRIFTS detection of strengthened HCOO* intermediates, together with new DFT calculations showing negative adsorption energies of HCOO* on TiO overlayers, confirm that HCOO* serves as a stabilizing factor for the 2-x encapsulating structure.\n➢ TiO formation: DFT calculations reveal that HCOOH* preferentially dissociates into HCOO* 2-x and H* on Ni/TiO -R. This low-barrier pathway makes oxygen-vacancy formation kinetically 2 highly favorable, thereby rationalizing the adsorbate-induced generation of TiO species. 2-x\n➢ Charge transfer and electronic modification: The in-situ XANES observation of a shift in the Ni absorption edge is corroborated by differential charge density and Bader charge analysis, which reveal ~0.17 e- transfer from TiO to Ni. 2-x\nii. Sequential narrative:\nThe revised Results section now follows a logical progression that links the different types of evidence into a coherent storyline:\n➢ Structural evolution (ETEM): In-situ ETEM directly visualizes the facet-dependent encapsulation of Ni by TiO overlayers. 2-x\n➢ Identification of intermediates (DRIFTS): In-situ DRIFTS identifies HCOO* intermediates, which are stabilized on Ni/TiO -R but rapidly converted on Ni/TiO -S due to facet effects. These 2 2 findings are further correlated with DFT calculations, which demonstrate that HCOO* induces the oxygen vacancy formation and stabilizes the TiO overlayer. 2-x\n➢ Electronic structure evolution: In-situ XAS, supported by differential charge density and Bader analyses, confirms interfacial charge transfer from TiO to Ni (~0.17 e), explaining the observed 2-x Ni edge shift and providing insight into how the support modifies the metal.\n➢ Catalytic performance: CO hydrogenation tests reveal facet-dependent selectivity shifts 2 between CO and CH , which correlate directly with the encapsulation extent and surface 4 intermediates observed above.\nWe believe these revisions substantially enhance the coherence of the Results section and make the interplay between experimental and theoretical evidence much more convincing. 22"}, {"round": 2, "reviewer_comments": {"reviewer_1": "The response of the authors to my previous comments is satisfactory. The authors have substantially improved the scientific level of the article and, therefore, I recommend acceptance for publication.", "reviewer_4": "Having considered the authors' rebuttal to the comments from Reviewer 2, I find their responses to be comprehensive and effective. The revisions have enhanced the quality of the paper, and I therefore recommend its publication."}, "author_response": ""}]

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