Phase interface engineering enables state-of-the-art half-Heusler thermoelectrics

Domain: earth_science
Published: 2024-07-16
Editorial outcome: minor_revision

Abstract
In thermoelectric, phase interface engineering proves effective in reducing the lattice thermal conductivity via interface scattering and amplifying the density-of-states effective mass by energy filtering. However, the indiscriminate introduction of phase interfaces inevitably leads to diminished carrier mobility. Moreover, relying on a singular energy barrier is insufficient for comprehensive filtration of low-energy carriers throughout the entire temperature range. Addressing these challenges, we advocate the establishment of a composite phase interface using atomic layer deposition (ALD) technology. This design aims to effectively decouple the interrelated thermoelectric parameters in ZrNiSn. The engineered coherent dual-interface energy barriers substantially enhance the density-of-states effective mass across the entire temperature spectrum while preser carrier mobility. Simultaneously, the strong interface scattering on phonons is crucial for curtailing lattice thermal conductivity. Consequently, a 40-cycles TiO2 coating on ZrNi1.03Sn0.99Sb0.01 achieves an unprecedented zT value of 1.3 at 873 K. These findings deepen the understanding of coherent composite-phase interface engineering.

Introduction
The usage of thermoelectric materials holds great promise for the direct conversion of waste heat into useful electricity, making them valuable candidates for energy conversion applications and sustainable energy development. The efficiency of thermoelectric device depends significantly on the material’s dimensionless Fig of merit, denoted as zT = \sigma\alpha^2T/\kappa, where \alpha, \sigma, \kappa, and T are the Seebeck coefficient, electrical conductivity, total thermal conductivity (including lattice contribution \kappa_L and carrier contribution \kappa_e), and the absolute temperature, respectively1. A more detailed evaluation of zT involves the quality factor \beta/\mu_H m*^3/2 = \kappa_L, which reveals the utmost importance of maintaining high carrier mobility (\mu_H) in materials with low \kappa_L. Recent efforts toward high zT are in line with the quality factor, which boost the density-of-states effective mass (m*) by band structure manipulation3–7, as well as reducing \kappa_L via multiscale microstructures8–16.

As an emerging paradigm-shifting strategy, phase interface engineering plays a pivotal role in optimizing thermoelectric transport parameters. This strategy addresses two key aspects: enhancing m* via energy filtering effect and shrinking \kappa_L through dense phonon scattering at high-density interfaces17–22. However, leveraging phase interface engineering to achieve superior thermoelectric performance faces two significant challenges. Firstly, interface potential barriers between nanoinclusions and the matrix effectively filter out low-energy charge carriers, thereby augmenting the density of states (DOS) proximal to the Fermi level. Yet, this enhancement of m* tends to diminish at elevated temperatures due to the limitations imposed by a single energy potential barrier (\Delta E)23–25. Secondly, arbitrarily introduced phase interfaces can scatter both phonons and carriers, leading to a detrimental effect on carrier mobility (\mu_H) and hence the zT26,27. For example, the introduction of Al2O3 into ZrNiSn or the creation of amorphous ZrNiSn have both successfully led to a significant reduction in \kappa_L. However, these approaches did not effectively enhance PF due to the decrease in \mu_H. Consequently, the coupling between electronic and phonon transport properties has limited the potential for improving the figure of merit zT28,29. To date, achieving a significant enhancement in \mu_H m*^3/2 = \kappa_L over the entire temperature range of thermoelectric materials remains a challenging and elusive goal30. Hence, a judicious and rational phase interface design is crucial for implementing phase interface engineering in thermoelectrics.

In this context, the construction of a coherently composite-phase interface is proposed as a promising strategy for the synergistic optimization of \mu_H m*^3/2 = \kappa_L. The connotation of a composite phase interface contains two fundamental principles. Initially, theoretical considerations suggest that the optimal energy barrier should lie between 1 to 10 k_B T (k_B and T stand for Boltzmann’s constant and absolute temperature)31, a range that varies with temperature. To ensure a consistent enhancement of m* across varying temperatures, the establishment of multiple energy barriers is crucial, ensuring the filtration of low-energy electrons throughout the entire temperature spectrum23. Additionally, an ideal phase interface should scatter phonons more effectively than electrons. Analysis of the phase interface structure reveals that both coherent and incoherent interfacial structures lead to an obvious decrease in \kappa_L. However, materials with coherent interfaces exhibit markedly higher \mu_H than those with the incoherent interfaces32,33. Despite these insights, comprehensive studies that concurrently address these aspects remain scarce. The construction of coherent phase interfaces harboring multiple \Delta E is pivotal in disentangling the intricately linked thermoelectric parameters18,23,34. However, the primary challenge lies in the systematic engineering of these coherent composite phase interfaces featuring multiple \Delta E in experimental setups to optimize the zT. Historically, methods such as ball milling35, hydrothermal synthesis36, self-precipitation37,38, and melt spinning39 have been prevalent for introducing secondary phases into the matrix. Regrettably, these methods fall short in precisely constructing coherent phase interfaces with multiple \Delta E. To surmount this challenge, employing atomic layer deposition (ALD) for the manipulation of coherent composite-phase interfaces emerges as a strategically effective pathway. ALD, renowned for its self-limiting and self-saturating characteristics, is ideal for the layer-by-layer growth of 2D films. Crucially, the unique capability of ALD technology to uniformly coat diverse 3D bulk materials enables meticulous control over the thickness of the coating layer40. At elevated sintering temperatures, the amorphous nature of these coating enhances chemical reactivity, potentially initiating in-situ chemical reactions between the substrate and the ALD-coated layer. In this study, we present a unique methodology for constructing coherent composite-phase interfaces with multiple \Delta E through the in-situ reflection between the matrix and the ALD coating. For this purpose, the half-Heusler (HH) ZrNiSn alloy was chosen as an exemplary template for implementing phase interface engineering due to its high sintering temperature. We employed an ALD coating of amorphous TiO2, chosen for its high chemical reactivity, as a precursor layer. The coherent composite-phase interfaces within the n-type ZrNi1.03Sn0.99Sb0.01 matrix (referred to ZNSS) is thereafter formed after high-temperature sintering. The initially amorphous TiO2 coating undergoes a comprehensive transformation, being replaced by a Ti-ZNSS layer (aka Zr1-xTixNi1.03Sn0.99Sb0.01) along with uniformly dispersed ZrO2 nanoparticles, as depicted in the schematic diagram in Fig. 1a.

The in-situ formation of the composite phase interfaces, consisting of the Ti-ZNSS layer and ZrO2 nanoparticles, impeccably orchestrates the simultaneous optimization of electrical and thermal transport through two synergistic mechanisms. Firstly, the augmentation of the energy filtering effect on m* facilitated by various \Delta E strategically created between the matrix and composite phase, as illustrated in the schematic diagram in Fig. 1c. Secondly, the introduction of coherent phase interfaces substantially reduces the \kappa_L while maintaining \mu_H unchanged, as depicted in the schematic diagram in Fig. 1d. Furthermore, we are pleasantly surprised to discover that there are a significant number of twins within the ZrO2 nanoparticles. This twin structure provides an additional scattering center for phonon transport at the nanoscale. Consequently, the state-of-the-art zT of 1.3 is attained in the ZNSS sample coated with 40 cycles of TiO2 (i.e., TiO2 = 3.2 nm) (Fig. 1b). These findings underscore the pivotal role of coherent composite-phase interface engineering in achieving cutting-edge zT in ZrNiSn-based and other thermoelectric materials.

Results
Design principle
Designing composite phase structures with coherent phase interfaces and multiple \Delta E is aimed at enhancing m* and phonon scattering while minimally impacting the \mu_H. Given the necessity for high-temperature sintering to facilitate the interface chemical reaction between matrix and amorphous ALD coating layer, ZrNiSn-based materials, known for their high melting point, are selected as an ideal platform for constructing the composite phase interfaces. The initial step involves the synthesis of a high purity ZNSS matrix. To this end, the ZNSS matrix is fabricated using levitation melting, a technique chosen for its efficiency in producing high-purity materials. The room temperature powder X-ray diffraction (PXRD) patterns, as illustrated in Fig S1, reveal the diffraction peaks of the ZNSS matrix that are exclusively indexed to the cubic MgAgAs-type crystal structure without any detectable second phases. A critical prerequisite for the construction of coherent composite-phase interfaces is the preparation of ZNSS powders uniformly coated with amorphous TiO2 layers. The dynamic formation process of depositing ultrathin TiO2 layers onto the ZNSS powders through ALD is meticulously captured in Fig. 2a. During the ALD process, conducted at 423 K, tetrakis(dimethylamido)titanium (TDMAT) and H2O serve as precursors for TiO2 deposition41,42. Notably, the PXRD pattern of the ALD-coated samples (Fig. S2) reveals the absence of crystalline TiO2 diffraction peaks, as the TiO2 layer obtained is actually amorphous and less than 10 nm, which is beyond the detection limit of XRD. Further validation of the successful coating of the ultrathin TiO2 layer on the surface of ZNSS powder is obtained through X-ray photoelectron spectroscopy (XPS). The XPS analysis, particularly for the coated 40 cycles TiO2 sample, demonstrates a distinct Ti 2p peak, unequivocally indicating the presence of the amorphous TiO2 layers on the ZNSS surfaces (Fig S3)43.

The ALD boasts significant advantages, particularly in its capability to homogeneously coat complex 3D structures and precisely control phase thickness. The quality of the ALD coating layer is validated through high-resolution transmission electron microscopy (HRTEM) images of a representative coated TiO2 powder coated for 40 cycles, as illustrated in Fig. 2b. It is evident that TiO2 with 40 cycles is uniformly enveloped by an ultrathin amorphous TiO2 layer, approximately 3 nm thick. The calculated growth rate of TiO2, is around 0.8 Å per cycle, which aligns with previous literatures40,44. The resulting samples, differentiated by the number of ALD-coated TiO2 layers, are designated as uncoated, TiO2 = 0.8 nm (10 cycles), TiO2 = 1.6 nm (20 cycles), TiO2 = 3.2 nm (40 cycles), TiO2 = 4.8 nm (60 cycles) and TiO2 = 6.4 nm (80 cycles). Such meticulous characterizations underscore the successful deposition of high-quality, ultrathin amorphous TiO2 layers on the surface of ZNSS powder, laying a robust foundation for the construction of composite phase interfaces.

The choice of amorphous TiO2 for the ALD coating is strategic, considering high chemical activity and susceptibility to oxygen deficiency at elevated temperatures, which are conducive to fostering the desired interfacial reactions. To expedite the chemical reaction at the interface between the TiO2 coating layer and ZNSS powder, a densification process is undertaken using spark plasma sintering (SPS) at 1373 K for 10 min. It is worth noting that no significant second phase was observed in the XRD patterns after sintering (Figs. S4 and S5). A comprehensive schematic depicting this interface chemical reaction is presented in Fig. 2c. In the initial stage, the O atoms within the amorphous TiO2 layer bond with the Zr atoms from the matrix, resulting in the formation of ZrO2 nanoparticles. Simultaneously, the residual Ti atoms occupy vacant Zr sites, creating a Ti-ZNSS layer. This sophisticated interfacial reaction mechanism is further substantiated by scanning transmission electron microscopy (STEM) images, including both annular bright field (ABF and corresponding EDS mapping) and atomically-resolved high angle annular dark field (HAADF) images, as shown in Fig. 2d. There images validate the uniform distribution of nanoparticles at the grain boundaries of the substrate, consistent perfectly with the hypothesized interface design. This intricate interplay at the nanoscale level highlights the ingenuity and precision of the phase interface engineering approach employed in this study.

Microscopic characterization of composite phase interfaces
In order to achieve a more precise elucidation of the reaction product of the SPSed TiO2 = 3.2 nm sample, we employed scanning electron microscope (SEM) and electron probe microanalysis (EPMA) to unveil the intricate composite phase compositions. Firstly, the X-ray energy dispersive spectrum (EDS) mapping and spot scanning results (Tables S1–S4) disclosed a noteworthy observation: Ti atoms did not enter into the ZNSS matrix. Instead, they replaced the pristine amorphous TiO2 layer, culminating in a uniform Ti-ZNSS layer, formally identified as Zr1-xTixNi1.03Sn0.99Sb0.01, as depicted in Fig. 3a. This uniform Ti-ZNSS layer encapsulates the ZNSS matrix. Further scrutiny, as shown in Figs. S6–S10, Supporting Information, revealed a profusion of nanoparticles uniformly dispersed around the ZNSS matrix. The excess of Ni in both the matrix and the Ti-ZNSS layer is a result of two factors: firstly, it is a consequence of our experimental design, and secondly, it reflects an inherent characteristic of ZrNiSn itself45,46. The EDS mapping images corroborate that these nanoparticles predominantly consist of ZrO2. To cross-check these findings, we also conducted a thorough TEM-EDS analysis, focusing on the size and composition of these nanoparticles (Fig. 3b), which affirmed their ZrO2 constitution. From these observations, we infer the formation of multiscale composite phase structures, which encompasses ZrO2 nanoparticles, ranging from 50–100 nm in size, and the Ti-ZNSS layer, with a thickness of 50 to 500 nm (Fig. S8). These features are predominantly derived from the chemical interaction between the highly reactive amorphous TiO2 layer and the ZNSS matrix during the SPS sintering process. Analogous phenomena have been observed in Al2O3-coated ZrNiSnSb0.99Sb0.01 and TiO2-coated Bi2Te2.7Se0.3 systems25,44. The underlying mechanism of this interface chemical reaction can be understood in terms of electronegativity and thermodynamic considerations. The electronegativity difference between Zr and O (2.11) is greater than that between Ti and O (1.9), hence favoring the bond formation between Zr and O over Ti and O. The formation energies of various oxides during high-temperature processes are critical determinants of their stability and presence in the final product. ZrO2 exhibits a formation energy of −3.801 eV, which is lower than that of TiO2 (−3.3 eV), SnO2 (−2.123 eV), NiO (−1.761 eV), and Sb2O3 (−1.728 eV). This indicates that ZrO2 is thermodynamically the most favorable oxide under the experimental conditions. The detailed chemical reaction formula is delineated as follows:

ZrNi1.03Sn0.99Sb0.01 + xTiO2 1373 K Zr1-xTixNi1.03Sn0.99Sb0.01 + xZrO2 (1)

Atomic scale characterization of the composite phase interface (ZNSS/ZrO2)
The microstructural intricacies at the phase interfaces play a critical role in disentangling the electrical and thermal transport properties. To this end, a comprehensive microstructural examination of two key interfaces of the SPSed TiO2 = 3.2 nm sample was conducted. Figure 4a–e illustrate the interface between the ZNSS matrix and ZrO2 nanoparticles. Employing aberration-corrected STEM with atomic resolution, we discerned a pristine coherent interface between the ZrO2 nanoparticles [100] and the ZNSS matrix [100]. The clarity of this interface is further validated by the inverse fast Fourier transform (IFFT) images, which exhibit a defect-free phase boundary. This immaculately coherent interface is crucial for sustaining a high carrier mobility \mu_H, thereby leading to excellent electrical properties. Furthermore, detailed examination of high-resolution ZrO2 atomic images, alongside Fast Fourier Transform (FFT) data (Fig. S13), confirms the crystalline structure of ZrO2 as monoclinic. This finding lays a solid theoretical foundation for investigating the interface potential barriers. A deeper analysis, integrating data from Fig. 4c and Fig. 4d, suggests that the impeccable coherence of the interface can be attributed to the closely matched lattice parameters of ZNSS (a=3.105 Å) and the monoclinic ZrO2 (a=3.005 Å).

Combining geometric phase analysis (GPA) with ZNSS as the reference matrix, we were able to spatially capture the high \epsilon_xy strain within the ZrO2 nanoparticles (Fig. 4e). Notably, the strain distribution in this context deviates from the norm observed in traditional nanocomposites, where strain typically accumulates around the phase boundary. Instead, in our sample, the strain is distributed throughout the entire particle. Further exploration through STEM HADDF (Fig. 4g) and ABF (Fig. 4h) phase images revealed abundant twin boundaries within the second-phase ZrO2. Twinning within the ZrO2 structure was confirmed by lattice splitting observed in the FFT analysis. Concurrently, GPA analysis demonstrated substantial strain in the \epsilon_xx direction at these twin boundary positions (Fig. 4i). This strain effectively impedes phonon propagation, resulting in a significant reduction in \kappa_L47. Remarkably, the presence of two distinct coherent twin boundaries within the ZrO2, each characterized by different rotation angles, as shown in Fig. 4j and k, underscores the sophisticated microstructural engineering. This intricate design is instrumental in achieving the decoupling of carrier and phonon transport.

Atomic scale characterization of the composite phase interface (ZNSS/Ti-ZNSS)
Figure 5a–d presents a detailed examination of the interface between the ZNSS matrix and the Ti-ZNSS layer. Elemental line scanning (Fig. 5a) indicates a heightened concentration of Ti around the grain boundaries, consistent with previous EDS results from EPMA. In regions with diminishing Ti concentration along the line scan (i.e., the interface between the Ti-ZNSS layer and the ZNSS matrix), a meticulous Z-contrast (Fig. 5b) discerns differences in mass between atoms, showcasing a flawlessly coherent interface established between the Ti-ZNSS layer and the ZNSS matrix, as confirmed by IFFT information48. As shown in Fig. 5c, d, advanced Z-contrast analysis of the HADDF images highlights pronounced darker features on the left side, conspicuously absent on the right. This phenomenon is attributed to the smaller atomic contrast exhibited by lighter Ti atoms compared to Zr. Consequently, through meticulous Z-contrast examination, a well-defined boundary delineates the Ti-ZNSS layer on the left from the ZNSS matrix on the right. To rigorously corroborate the accuracy of the data in Fig. 5d, a meticulous Z-contrast line scan is performed on the corresponding region in Fig. 5c. The observed trend reveals a sequential increase in atomic contrast from left to right, aligning precisely with the contrast patterns depicted in Fig. 5d. This meticulous analysis serves as a stringent cross-verification, ensuring the reliability and consistency of the presented results.

Electrical and thermal transport properties
In the context of the formation of coherent composite phase interfaces in TiO2-coated ZNSS samples, a pertinent question emerges: Can these specialized phase interfaces facilitate the decoupling of carrier and phonon transport? Fig. 6e elucidates the electrical and thermal transport properties of these as-fabricated samples. A striking observation from Fig. 6a is the substantial enhancement of the Seebeck coefficient (\alpha) in ALD-coated ZNSS samples, compared to the initial ZNSS sample (Fig. 6a). In typical thermoelectric materials, an increase in the \alpha is commonly accompanied by a decrease in the carrier concentration (n_H)49. However, as indicated by the Pisarenko curve of \alpha-n_H at room temperature (Fig. S16), the nearly constant n_H before and after ALD coating indicates that the increase in \alpha can only be attributed to the improvement in m* via the energy filtering effect50,51. Moreover, under the influence of coherent composite phase interfaces, the electrical conductivity (\sigma) of ALD-coated ZNSS samples does not show significant variation from the pristine uncoated sample (Fig. 6b). This stability in \sigma can be predominantly attributed to the coherent phase interfaces, which are integral in maintaining \mu_H. In the case of TiO2 = 4.8 nm and TiO2 = 6.4 nm samples, the slight increase in the \sigma is primarily due to the elevated n_H. Benefiting from the enhanced \alpha and retained \sigma via the coherent composite phase interfaces with multiple energy barrier, the power factor (PF) has been significantly enhanced after ALD coating (Fig. 6c). Taking the TiO2 = 3.2 nm sample as an example, the simultaneously achieved high \alpha of 204.1 \mu V K^-1 and \sigma of 1448 S cm^-1 at 873 K give rise to a large PF of 60.24 \mu W cm^-1 K^-2. This attests to the efficacy of our phase interface engineering to improve the electrical properties of ZrNiSn-based materials.

Figure 6d depicts the temperature dependence of the thermal conductivity (\kappa) and lattice thermal conductivity (\kappa_L) for all the samples. A marked reduction in \kappa is observed with increasing ALD coating layers. For instance, the room temperature \kappa quickly drops from 5.41 W m^-1 K^-1 for C=0 sample to 4.78 W m^-1 K^-1 for TiO2 = 6.4 nm sample. Since the rational coherent phase interface design results in negligible changes in \sigma across all the samples, this decline in \kappa as increasing ALD coating layers can be largely ascribed to the diminished \kappa_L. Herein, \kappa_e is evaluated by the Wiedemann-Franz law \kappa_e = L\sigma T and \kappa_L is obtained by subtracting the \kappa_e from \kappa, where the Lorenz number L is estimated via L=1.5+exp(-|\alpha|/116)52. As expected, the \kappa_L substantially declines with increasing ALD coating layers over the entire temperature range (Fig. 6e). In particular, the lowest \kappa_L of 1.72 W m^-1 K^-1 is attained in the TiO2 = 6.4 nm sample, which is 20% lower than that of the pure ZNSS matrix. Ultimately, all ALD-treated samples exhibit superior zT values compared to the pristine ZNSS sample. Notably, the TiO2 = 3.2 nm sample achieves an advanced peak zT value of approximately 1.3 (Fig. 6f).

Analysis of electrical and thermal transport mechanisms
To elucidate the synergistic mechanism underlying the observed enhancement of thermoelectric performance, we employed the single parabolic band model (SPB) to generate Pisarenko curves for samples with different numbers of ALD cycles (Fig. 7a), with experimental data at 300 K and 800 K distinctly annotated. Initially, an analysis of the variation in m*, as depicted in Fig. 7b, reveals a crucial insight. The intrinsic band structure of the ZNSS matrix remains highly stable across temperatures, with m* being 2.7 m_e at both 300 K and 800 K. This finding aligns with the prior research conducted by Zhu et al.53. Furthermore, an intriguing trend is observed: regardless of at the temperature (300 K or 800 K), the m* demonstrates a gradual increase with an escalating number of ALD cycles. Most intriguingly, at 800 K, the m* for all ALD-treated samples is significantly elevated compared to their 300 K counterparts. According to the research conducted by Snyder et al.50, the interface potential barriers between nanoinclusions and the matrix can selectively filter out low-energy charge carriers, thereby enhancing the DOS near the Fermi level (i.e., m*) and consequently improving the \alpha. Moreover, Zhang et al. proposed that the enhancement in m* resulting from a single interface potential barrier tends to diminish at elevated temperatures23. The theoretically optimal estimation of \Delta E~1-10 k_B T supports notion that constructing hierarchical potential barriers could effectively mitigate this issue31. Therefore, the construction and optimization of potential energy barriers at the interfaces between the composite phases and the ZNSS matrix are critical to the enhancement of the electrical transport properties.

Given the composite phase structures prepared by ALD, it is imperative to further investigate the details of the two distinct types of \Delta E at the Ti-ZNSS/ZNSS and ZrO2/ZNSS interfaces. After calibrating the work function for each band structure, an interface \Delta E of approximately 0.2 eV was observed between the conduction band minimum (CBM) at the X-point between ZrNiSn, Zr0.67Ti0.33NiSn and TiNiSn (Fig. S13). This observation aligns with the the chemical reaction dynamics outlined in Equation 1, revealing that with an increasing number of ALD coating layers, there is a corresponding rise in Ti content in the Ti-ZNSS layer. Consequently, this leads to a tunable \Delta E between the Ti-ZNSS layer and the ZNSS matrix, ranging approximately from 0 to 0.20 eV. In contrast, the band structure analysis of the monoclinic ZrO2 indicated a higher \Delta E, approximately 0.6 eV (Fig. S14). This differential in \Delta E values for the two interfaces, at approximately 0.2 eV and 0.6 eV respectively (Fig. 7c), fits well within the theoretically optimal range of \Delta E~1-10 k_B T (0.25 eV/300 K – 0.75 eV/873 K)31.

Furthermore, Fig. 7d distinctly demonstrates the remarkable elevation in the m* for the ALD-treated samples, observable at both ambient and elevated temperatures. The enhancement is more pronounced at higher temperatures. The underlying rationale for this phenomenon becomes apparent when juxtaposed with the semiconducting behavior of single \Delta E-induced Half-Heusler (HH)/Full Heusler (FH) samples54. Particularly at 800 K, the ratio of m* (with \Delta E) to m* (without \Delta E) displays a diminishing trend relative to 300 K, highlighting a limitation inherent in a singular energy barrier system. Hence, the PF (Fig. 7e) reveals a clear and consistent enhancement across all ALD-treated samples than that in pristine ZNSS sample, evident at both 300 K and 873 K. Such a phenomenon is a testament to the strategic construction of interfaces through ALD, offering a compelling insight into the intricate dynamics governing the thermoelectric enhancement observed in our experimental samples.

From the aspect of defect physics, the introduction of a second phase with a distinct interface from the matrix is conventionally perceived as a two-dimensional defect, imparting scattering interactions on both phonons and carriers55,56. However, coherent phase interfaces are recognized as a distinctive subclass of two-dimensional defects57,58. Notably, these interfaces are remarkable for their ability to scatter phonons while simultaneously preserving high \mu_H, rendering them a specialized and nuanced facet in the landscape of material defects. Therefore, the coherent composite phase interfaces constructed via ALD exhibit a comprehensive integration of this particular attribute. As depicted in Fig. 7f, a discernible pattern is observed with the continual increase in ALD cycles. At 300 K, \mu_H remarkably consistent, while there is a notable successive reduction in \kappa_L. This observation underscores the meticulous orchestration of these interfaces in simultaneously upholding \mu_H and modulating \kappa_L, highlighting the efficacy of ALD-driven fabrication of coherent composite phases.

Thermoelectric performance comparison
The remarkable advancements in elevating m*, preserving \mu_H, and attenuating \kappa_L through coherent composite phase interface engineering have unequivocally established this methodology as a cornerstone in advancing thermoelectric performance. Figure 8a, b vividly illustrate that the PF achieved via the coherent composite interfaces significantly outperforms those observed in n-type ZrNiSn systems. Concurrently, the multidimensional defects meticulously constructed via composite phase interface engineering exert a profound constraint on phonon propagation, resulting in a marked reduction in \kappa_L. Consequently, this sample manifests a remarkable average zT value of 0.73 within the temperature range of 300 to 873 K (Fig. 8c). The relevance of this finding is amplified when considering the critical benchmarks for assessing thermoelectric materials: the average zT and average power factor. The average zT dictates the device’s conversion efficiency, whereas the average power factor governs the device’s output power. In this context, as illustrated in Fig. 8d, the TiO2 = 3.2 nm sample not only demonstrates a significant improvement, almost twofold, compared to the original ZrNiSn sample, but also stands out remarkably against other high-performance n-type half-Heusler thermoelectric materials reported in contemporary research. These results demonstrate the effectiveness of coherent composite phase engineering in decoupling carrier and phonon transport in thermoelectric materials.

Methods
Polycrystalline samples ZrNi1.03Sn0.99Sb0.01 matrix was prepared by the levitation melting-ball milling-spark plasma sintering (SPS). Alloys were first prepared by levitation melting the stoichiometric amounts of Zr (piece, 99.99%), Ni (rods, 99.99%), Sn (block, 99,99%), and Sb (block, 99.99%) under an argon atmosphere for 2 min, and the melt was quenched in a water-cooled copper crucible. The ingots were remelted six times to ensure homogeneity. The ball milling process for 1 h at 800 rpm under an argon atmosphere. The fine powders were to use as the matrix of ALD.

A TiO2 layer was coated on the surface of the ZrNi1.03Sn0.99Sb0.01 matrix in a homemade continuous-flow ALD reactor operated under a base pressure of ~1 Torr. Typically, the 8 g as-prepared ZrNi1.03Sn0.99Sb0.01 matrix was transferred into an ALD chamber equipped with a vertical stainless rotating sample chamber. The TiO2 layer was deposited using tetrakis(dimethylamido)titanium (TDMAT) and H2O as ALD precursors at 423 K. The bubbler containing TDMA was heated to 323 K and the delivery line was heated to 343 K. The precursor dose and purge time were 30 seconds and 60 seconds, respectively. Ar gas served as both a carrier and a purging gas. Then, the coated fine powders were sintered by Spark Plasma Sintering (SPS) at 1373 K under 60 MPa in vacuum for 10 min. The as-sintered samples, of which the relative density was >98%, were cut for thermoelectric property measurement and characterization.

The phase structure of coated powders and sintered samples were measured by X-ray diffraction (XRD) on a Rigaku Smartlab 9kw (tube voltage: 45 kV, tube current: 200 mA) diffractometer using Cu K a radiation (\lambda=1.5406 Å) and X-ray photoelectron spectra (XPS) were acquired using a Microlab 350 surface analysis system equipped with a monochromatized Al anode X-ray source, pass energy is 20.0 eV. The chemical composition was obtained during electron probe microanalysis (EPMA, JOEL, JXA-8100) using wavelength dispersive spectroscopy (WDS), acceleration voltage of 300 kV and test beam current: 50 nA for surface analysis, 20 nA for quantitative testing. The TEM specimens of the sintered bulk materials were meticulously prepared utilizing a state-of-the-art focused ion beam system (FIB; JIB4601F, JEOL) to ensure precise cross-sectioning. Subsequently, the surface morphology and detailed crystallographic orientation were elucidated through high-resolution field-emission transmission electron microscopy (HR-TEM; JEM-ARM300F), operated at an acceleration voltage of 300 kV to achieve optimal imaging conditions.

The sintered specimen was carefully sectioned into two distinct shapes: a rectangular prism measuring 3 mm by 3 mm by 11 mm, and a square prism with dimensions of 9.8 mm by 9.8 mm by 2 mm. The rectangular prisms underwent meticulous polishing to facilitate the measurement of the Seebeck coefficient (\alpha) and electrical conductivity (\sigma) using the ULVAC ZEM-3 apparatus, within a temperature span of 300 to 873 K. Subsequently, the overall thermal conductivity (\kappa) was deduced through the formula \kappa=DC_p \rho, where D denotes the thermal diffusivity, C_p is the specific heat at constant pressure, and \rho is the material density. The thermal diffusivity coefficient (D) was ascertained using a Netzsch LFA 467-HT laser flash analyzer, across the temperature gradient of 300 to 873 Kelvin. Concurrently, the specific heat capacity (C_p) was determined employing a differential scanning calorimeter (DSC404F3), with a controlled heating rate of 10 K per minute. Sapphire comparison method, calibration procedure: blank test, sapphire test and sample test, each DSC curve is subtracted from the blank curve. The carrier density of sintered samples were measured using a physical property measurements system (PPMS, Quantum Design).

Electronic structure calculations of ZrNiSn, TiNiSn and ZrO2 were carried out using the density functional theory (DFT) with the projector-augmented-wave (PAW) method as implemented in the Vienna Ab initio simulation package (VASP). The computational framework for our study is based on the Generalized Gradient Approximation (GGA), specifically incorporating the exchange-correlation functional formulated by Perdew, Burke, and Ernzerhof (PBE). The structural optimization of the models was performed using the conjugate gradient (CG) method, which was iterated until the maximum forces acting on the Hellmann-Feynman theorem and the total energy reached a threshold of 10^-4 eV, indicating a stable configuration. The simulation parameters were carefully chosen, with a plane-wave cutoff energy set at 450 electron volts to ensure the accuracy of the calculations. For the integration over the Brillouin zone (BZ) of the primitive unit cell, we employed a Monkhorst–Pack grid with a dense k-point sampling of 11×11×11, which provides a comprehensive representation of the reciprocal space.

Peer review rounds
[{"round": 1, "reviewer_comments": {"reviewer_1": "The manuscript submitted by Zhang et al. demonstrates a phase interface engineering\napproach to enhance the thermoelectric properties of ZrNiSn-based half-Heusler\ncompounds. The experimental design is quite adjective, and ALD-coated TiO2 as a source\nreacted with ZrNiSnSb and formed ZrNiSnSb-ZrTiNiSnSb-ZrO2 nanocomposite. In this case,\na unique interface ZrNiSnSb-ZrTiNiSnSb-ZrO2 with a suitable interface potential barrier is\nbuilt up to scatter low-energy electrons, enhancing the Seebeck coefficient. A huge effort\nhas been devoted to characterizing the interface microstructures, which helps understand\nthe structure-properties relation and supports the argument of the energy filtering effect. It\noffers a new technical route to design the interface mircostructure. However, I have several\nconcerns regarding technical aspects that need addressing before considering this\nmanuscript for publication in Nature Communications. My comments are as follows:\n1) Many references are cited incorrectly. The authors need to double-check all references\nand cite them correctly.\n2) In the introduction part, page 4, lines 89-90, “these methods fall short in precisely\nconstructing coherent phase interface with multiple delta E”. Well, it is hard to judge this\nargument. At least, all those cited methods (ball milling, hydrothermal, self-precipitation, and\nmelt-spinning) are not aimed at creating the coherent phase interface with multiple delta E.\nTo prove the authors’ statement, they can try to ball mill ZrNiSnSb-(Zr0.67Ti0.33)SnSb-ZrO2\n(same chemical composition as C=40, which I calculated according to the EMPA results in\nthe supporting information).\n3) To me, it is more a modulation doping. If TiO2 can coat all matrix particles, then what you\nget after the SPS is a core-shell structure. How can you make sure TiO2 can coat all matrix\nparticles? How long will it take to coat the 8 g matrix?\n4) It is nice to see the authors show that the matrix is a Ni-rich phase. The Ni-rich nature has\nbeen confirmed by neutron scattering experiments.\n5) If the authors measured DSC to get Cp, do they observe the Sn melting peak in the DSC\ncurve? Although the XRD does not show visible impurity, the SEM&EDS show the Sn/rich\narea (Figure 3a). Also, Ni is not uniformly distributed in Figure 3a.\n6) Ti EDS mapping is missing in Figure 3b.\n7) Temperature-dependent kL is missing in Figure 6. Figures 6 d and 6 e are the same.\n8) It is assumed that the density functional theory (DFT) calculations were conducted for\nZrNiSn, TiNiSn, and ZrO2 at 0 K. However, these data can not accurately represent the\nband structure at 300 K and 800 K.\n9) Besides, if we take the calculated band structures of ZrNiSn and TiNiSn, we can only\nconclude that the delta E between them is around 0.2. How can you conclude that the delta\nE between Zr0.67Ti0.33NiSn0.99Sb0.01 and ZrNiSn0.99Sb0.01 is in the range of 0-0.2?\n10) The conductivity data has a certain trend with increased c numbers. I guess the\nconductivity is a tradeoff of the Ti doping (increase the n) and interface scattering (decrease\nthe mobility)\n11) I think modulation doping induced by adding TiO2 is the key to enhancing the power\nfactor.\n12) If the method can be applied to the high ZT half Heusler compounds, this work will\ndefinitely deserve to be published in Nature Communications.", "reviewer_2": "Minor Revision. The authors introduced the enhancement in the thermoelectric performance\nof ZrNiSn-based alloys using atomic layer deposition technology, which effectively decouple\nthe interrelated thermoelectric parameters in ZrNiSn. I have carefully read the manuscript,\nand I found that this is an interesting work with promising results, and the structure and logic\nare good. I am happy to recommend this work to be published in this journal after some\nrevisions, since there are still unclear points in the current manuscript. My comments listed\nbelow may help the authors further improve their work:\n1) The diagram in Figure 1a should be redrawn, such that the second phase is tens to\nhundreds of nanometers in size and should not be represented by small balls. The thickness\nof the Ti-doped second phases coating the ZNSS matrix should not be less than the\ninterplanar spacing.\n2) It can be seen from the XPS spectra in Figure S3 and EDS results in S8 that there is a lot\nof amorphous TiO2 on the ZNSS surfaces, but no obvious changes in the diffraction peaks\ncaused by amorphous TiO2 as shown in Figure S4, whether it is due to the deduction of the\nXRD pattern background or the fact that the content of amorphous TiO2 is less than the\ndetection limit of XRD. The author needs further clarification.\n3) It is difficult to explain the interface reaction mechanism based on STEM and HAADF\nresults in Figure 2. This is because the HAADF image is essentially an atomic column\ncontrast image, and its contrast is affected by many factors. On page 8 of the manuscript,\nthe author's conclusion, “the residual Ti atoms occupy vacant Zr sites, which can create a Ti-\nZNSS layer…”, cannot be judged by the HAADF results alone. The STEM-EDS results\nshould be combined to explain the exact mechanism.\n4) On page 10 of the manuscript, if only Zr position elements change, the energy spectrum\nof other elements should be uniform. However, the distribution of elements of Ni and Sb is\nnot very uniform in Figure 3a, so is it appropriate to judge the reaction equation from the\nexperimental results in EPMA? The authors can add some energy spectrum analysis to\nverify the accuracy of the reaction formula.\n5) Figure 6d and Figure 6e are repeated, and Figure 6e should be the lattice thermal\nconductivity. In addition, the test temperature when the minimum lattice thermal conductivity\nis reached should be added.\n6) The point-scanned data in EPMA image of Table S1 and S2 show that the content of Ti in\nC20 and C40 samples is different with the increased of thickness of the deposited atomic\nlayer.\nThe authors need to add the point-scanned data of the other composition ZNSS/C (C=10,\n60, 80), because I believe there is a certain difference with the nominal chemical\ncomposition.\n7) The serial numbers of the references in Figure 8 do not correspond to each other. Please\ncheck them carefully.\n8) Error estimates for the different measured quantities can be included in experimental part\nor figures for better comprehension.\n9）Kindly check the formatting of references and image resolution in accordance to journal\nstandards. Such as there are some superfluous words (Adv. Energy Mater. 37/2021) in\nreference 17, the numbers in the material composition in references 5 and 17 should be\nsmall corner marks, and the resolution in Figure S1 is obviously poor. These mistakes\nshould not have happened, and the authors should confirm them carefully before\nresubmitting.\n10) The XPS fit peaks in Figure S3 are rough, especially the peaks near the 462 eV, so it is\nrecommended to fit them after removing the background of XPS spectra.\n11) For a good scientific paper, supporting documents and manuscript are equally important.\nHowever, the figures in the supporting files are rough, such as no picture numbers and ruler\nin Figure S9 Figure S11, and Figure S13.\n12) Some outlooks should be added before the conclusion such as the future direction for\npractical applications.", "reviewer_3": "Zhang and coworkers have explored thermoelectric ZrNiSn using a range of advanced\ntechniques, including aberration-corrected high-resolution microcopy. The authors have\nutilized an interesting idea to engineer the interfaces by inclusions of TiO2 layers through\natomically controlled processes. They have shown that scattering phonons against such\ninterfaces, while maintaining high carrier mobility and high effective mass (energy filtering),\nis a potent approach to enhance thermoelectric efficiency (ZT). Consequently, the ZT value\nreaches 1.3 at 900 K, which is an important step towards broader applications. Such a\ncoherent interface design is novel and deserves attention of the thermoelectric community\nand may have impact on other functional materials. The manuscript is well written. Hence,\nthis manuscript should be accepted for publication after addressing the points raised below.\n1) Since thermoelectric materials are often exposed to harsh conditions (atmosphere, high\ntemperatures) under cyclic loading, a coherent interface may also be relevant to counteract\nthese effects. Is it conceivable that further oxidation can be hindered by inclusion of TiO2?\nPlease discuss these issues in the manuscript. This would give a larger perspective.\n2) The introduction is too short regarding the physical (transport) properties of ZnNiSn and\nrelated open questions to be tackled in the manuscript. Please provide a broader review on\nZrNiSn, especially alternative pathways to enhance its thermoelectric performance (interface\nengineering with Al2O3 (Key Engineering Materials Vol. 249 (2003) pp 79-82) and\namorphous configurations (Computational Materials Science 230 (2023) 112530)). Another\nreference regarding Al2O3 has been mentioned in the results section, but the introduction\nleaves the impression that interface engineering has not been attempted. Furthermore, why\nis TiO2 a better choice than Al2O3? Could an amorphous interface (e.g., crystalline ZrNiSn\ninterfaced with amorphous ZrNiSn) have a similar performance? Furthermore, the addition of\nSb to ZrNiSn and selection of TiO2 should also be more motivated.\n3) ZrO2 nanoparticles are obviously important. Please also discuss their role in the abstract.\n4) Why is there only oxidation of Zr when TiO2 is in contact with ZrNiSn-Sb? Why are Sn,\nNi, and Sb inert? For instance, Sn can easily form oxides. Electronegativity, as mentioned on\npage 10, is only a partial cause. What about thermodynamics? Eq. (1) is not really\nsupported.\n5) How does the ZT value obtained in this work correlate with other Heusler systems?\nPlease discuss.\n6) Designating TiO2 layers as “C = 0, C = 10, C = 20, C = 40, C = 60 and C = 80” is\nmisleading. If somebody wants to reproduce these results (perhaps by using another\ntechnique), more general descriptors are needed (e.g., thickness rather than cycles).\nFor example, C = 40 means that the TiO2 thickness is about 3 nm (page 8). Why is this\ncomplicated designation (method oriented) necessary?\n7) It would be useful to designate the binding energy of both contributions in Fig. S3. Please\nalso compare the binding energy with literature.\n8) Does hydrogen have any role here? Having H2O as a precursor should not be ignored. In\nmany fields, H is instrumental. Please discuss.\n9) Fig. S4 has never been mentioned in the manuscript. If it’s redundant, please remove it.\nPlease check other figures in the supplementary materials as well.\n10) Is the compositional analysis as accurate as presented in Table S1? Are all digits\nstatistically reasonable to supply? What is the error bar? Why is O omitted in Table S1?\nSimilar goes for Table S2.\n11) The authors argue about strain in the interfaces. How is strain affecting the transport\nproperties?\n12) Please provide a reference for the procedure pertaining to the Pisarenko curve.\n13) There are some typos in the manuscript (e.g., “fromthe” on page 16 and “Mpa” on page\n23 (it should be MPa)). Please proofread the text.\n14) Band structure in Fig. S10 is given for ZrNiSn and TiNiSn as well as ZrO2 in Fig. S12,\nrather than the actual system investigated in the current work. Drawing conclusions from a\nvery idealized model system(s) can be ambiguous. Please elaborate.\n15) The statement that the “base pressure of ~ 1 Torr” was achieved in ALD experiments\ndoesn’t seem to be correct. This is probably the working pressure (1 Torr as a base pressure\nmay give rise to full oxidation). Please state the base pressure (pressure before\ndepositions).\n16) More details are needed for the characterization techniques (e.g., power setting for XRD,\npass energy for XPS, accelerating volage for microscopy, calibration procedure for DSC,\netc.)."}, "author_response": "Referee #1:\nGeneral comment: The manuscript submitted by Zhang et al. demonstrates a phase interface\nengineering approach to enhance the thermoelectric properties of ZrNiSn-based half-Heusler\ncompounds. The experimental design is quite adjective, and ALD-coated TiO as a source\n2\nreacted with ZrNiSnSb and formed ZrNiSnSb-ZrTiNiSnSb-ZrO nanocomposite. In this case, a\n2\nunique interface ZrNiSnSb-ZrTiNiSnSb-ZrO with a suitable interface potential barrier is built\n2\nup to scatter low-energy electrons, enhancing the Seebeck coefficient. A huge effort has been\ndevoted to characterizing the interface microstructures, which helps understand the structure-\nproperties relation and supports the argument of the energy filtering effect. It offers a new\ntechnical route to design the interface mircostructure. However, I have several concerns\nregarding technical aspects that need addressing before considering this manuscript for\npublication in Nature Communications.\nGeneral response: We appreciate the reviewer’s positive comments. Please find our point-by-\npoint responses below.\nDetailed comment 1: “[1] Many references are cited incorrectly. The authors need to double-\ncheck all references and cite them correctly. ”\nResponse: Thank you for bringing this to our attention. We apologize for any confusion caused\nand have rigorously reviewed all references to ensure their accuracy. Each citation has been\ncorrected to comply with the journal's formatting requirements, and we have verified the\naccuracy of each reference to uphold the integrity of our manuscript.\nDetailed comment 2: “[2] In the introduction part, page 4, lines 89-90, “these methods fall\nshort in precisely constructing coherent phase interface with multiple delta E”. Well, it is hard to\njudge this argument. At least, all those cited methods (ball milling, hydrothermal, self-\nprecipitation, and melt-spinning) are not aimed at creating the coherent phase interface with\nmultiple delta E. To prove the authors’ statement, they can try to ball mill ZrNiSnSb-\n(Zr Ti )SnSb-ZrO (same chemical composition as C=40, which I calculated according to\n0.67 0.33 2\nthe EMPA results in the supporting information). ”\nResponse: Thank you for your insightful suggestion. Following your advice, we conducted ball\nmilling of the mixture of ZrNi Sn Sb and ZrO to investigate the formation of a coherent\n1.03 0.99 0.01 2\ninterface. Based on the microstructural analysis presented in Fig. R1, it is evident that the ZrNiSn\nmatrix and the ZrO second phase do not exhibit a coherent interface. Additionally, experiments\n2\nsuch as those by X. Y. Huang et al. (Solid State Communications 130, 181–185, (2004))\ninvolving the incorporation of ZrO into ZrNiSn via ball milling have consistently shown a\n2\ndecrease in carrier mobility ( ), failing to achieve the ideal decoupling of electron and phonon\nH\ntransport. These results suggest that the introduction of ZrO through ball milling is insufficient\n2\nto create the desired coherent interface necessary for optimized electron and phonon transport\nproperties. The ball milling process may lack the precision required to control the interface\nchemistry and structure, which are critical for achieving the desired thermoelectric (TE)\nperformance.\nIn contrast, atomic layer deposition (ALD) has proven more effective in achieving a\ncoherent interface between the ZrNiSn matrix and the ZrO layer, as shown in Fig. 4. The ALD\n2\nprocess allows for precise control over the thickness and uniformity of the ZrO layer, which is\n2\ncrucial for forming a coherent interface that enhances the thermoelectric performance of the\nmaterial.\nWe appreciate the opportunity to discuss these alternative methods and clarify the\nadvantages of our chosen approach. We believe that our study provides valuable insights into the\nimportance of interface engineering in developing high-performance thermoelectric materials.\nFig. R1. a, b and c HRTEM image of ZrO nanoparticles composited under the ball milling\n2\nprocess and corresponding EDS energy spectra. e and f show the detail view of the interface\ncontact between ZrO and ZrNi Sn Sb in d.\n2 1.03 0.99 0.01\nDetailed comment 3: “[3] To me, it is more a modulation doping. If TiO can coat all matrix\n2\nparticles, then what you get after the SPS is a core-shell structure. How can you make sure TiO\n2\ncan coat all matrix particles? How long will it take to coat the 8 g matrix?”\nResponse: Thank you for your valuable comments. In this work, we utilized ALD to achieve a\nrelatively uniform, atomic-scale coating of TiO on the powder matrix materials prior to the\n2\n\nspark plasma sintering (SPS) process. ALD is especially effective for coating powders with\ncomplex morphologies due to its gas-phase reactions, which ensure uniform deposition across\neven the most irregular surfaces.\nDuring the ALD process, precursors in vapor form are introduced sequentially into a\nreaction chamber under vacuum. This approach allows for the deposition of shell materials with\natomic-scale precision on powder samples, regardless of their morphological complexities.\nIdeally, sintering the as-prepared powder with the core-shell structure results in polycrystalline\nmaterials with continuous heterogeneous grain boundaries.\nThe time required for the ALD process is determined by the number of cycles needed to\nachieve the desired coating thickness. For the deposition of the TiO layer, we used\n2\ncontaining TDMAT was heated to 40 °C, and the delivery line was maintained at 70 °C. The\nprecursor dose and purge times were set at 30 seconds and 60 seconds, respectively. Typically,\none ALD cycle takes approximately 90 seconds.\nThe use of ALD for phase interface engineering not only enhances the thermoelectric\nperformance of various materials but also demonstrates significant potential for industrial\nscalability and precision. This method is compatible with commercialization efforts due to its\ncontrolled and repeatable nature. We appreciate your inquiries and hope this explanation clarifies\nthe methodology and rigor employed in our research.\nFor a detailed delineation of the distinctions between this work and conventional modulation\ndoping, please refer to our response to detailed comment 11.\nDetailed comment 4: “[4] It is nice to see the authors show that the matrix is a Ni-rich phase.\nThe Ni-rich nature has been confirmed by neutron scattering experiments.”\nResponse: We appreciate the reviewers' acknowledgment of our demonstration that the matrix is\na Ni-rich phase, as confirmed. ZrNiSn, a half-Heusler compound, is particularly intriguing due to\nits distinctive electrical and thermal transport properties, as well as the notable discrepancies\nbetween theoretical predictions and experimental results.\nNeutron scattering proves invaluable for elucidating the macroscopic atomic arrangement\nwhile also detecting microscopic local disorders and defects. Previous studies (J. Mater. Chem. C\n3, 10534-10542 (2015), J. Appl. Phys. 116,163514 (2014)) utilizing neutron scattering have\nunderscored the essential feature of ZrNiSn: the excess of Ni atoms. This Ni excess introduces\nlocal chemical disorder, which subsequently affects the electronic structure and, consequently,\nthe thermoelectric performance of the material.\nThe respective content has been updated in the revised manuscript, as seen below: The\nexcess of Ni in both the matrix and the Ti-ZNSS layer is a result of two factors: firstly, it is a\nconsequence of our experimental design, and secondly, it reflects an inherent characteristic of\nZrNiSn itself.\nDetailed comment 5: “[5] If the authors measured DSC to get Cp, do they observe the Sn\nmelting peak in the DSC curve? Although the XRD does not show visible impurity, the\nSEM&EDS show the Sn/rich area (Fig. 3a). Also, Ni is not uniformly distributed in Fig. 3a.”\nResponse: Thank you for your insightful comment. We employed differential scanning\ncalorimetry (DSC) to determine the heat capacity (Cp) of our samples, using sapphire as the\nreference material. DSC is highly sensitive to detecting minute heat flow changes associated with\nthermal events, which is essential for characterizing both physical and chemical transformations\nin materials.\nWe carefully selected three samples, each weighing approximately 50 milligrams, to\nminimize any discrepancies that might arise from sample heterogeneity. In Figure S17b, none of\nthe samples exhibit a tin melting peak (around 504 K). In contrast to the Scanning Electron\nMicroscopy (SEM) and Energy-Dispersive X-ray Spectroscopy (EDS) analyses described in\nFigure 3a, where Sn-rich regions were observed, it is determined that these regions do not\nrepresent the presence of elemental tin but rather small areas enriched with Sn-containing\ncompounds. This enrichment is a result of slightly over adding tin during the preparation of the\nZrNi Sn Sb matrix using the suspension melting method to prevent tin evaporation.\n1.03 0.99 0.01\nBecause these areas are sufficiently small, their content is below the detection threshold of X-ray\ndiffraction (XRD), which is typically around 5%. This is the reason why they are not visible in\nthe XRD spectrum.\nFurthermore,In addition, as shown in Fig. 3a, the distribution of Ni is relatively more\nheterogeneous, which is mainly due to the over-sensitivity of Electron Probe Micro-Analysis\n(EPMA) to different local compositions. Combined with the data from our spot-scanning (Table\nS3), it can be seen that the Ni in the matrix is about 0.3% more than that in the Ti-ZNSS layer,\nand that the relative amount of Ni is probably due to the interfacial reaction, where the Ni\ncontent disappears completely in the corresponding ZrO nanoparticles. It is important to\n2\nhighlight that the melting point of Ni (1728 K) far exceeds the maximum testing temperature\nrange of our DSC (up to 873 K). Thus, no melting peak for Ni could be observed within our DSC\ntest conditions.\nAdditionally, we have provided logarithmic mode XRD spectra (Fig. S4b) to offer a clearer\ndepiction of the phases present, confirming the absence of distinct Sn or Ni peaks within the\ndetection limits of this technique.\nFig. S17. b Temperature dependent specific heat capacity Cp for ZNSS.\nFig. S4. b XRD pattern (log mode) of TiO = 3.2 nm sample after SPS process.\nResponse: We apologize for the oversight regarding the absence of Ti EDS mapping in Fig. 3b.\nWe have now included the Ti EDS map in the revised Fig. 2d to provide a comprehensive view\nof its distribution within the sample. Your attention to detail and valuable feedback are greatly\nappreciated.\nFig. 2 ALD Synthesis Technology Route. a Schematic diagram illustrating the synthesis of the\nTiO layer through ALD coating. b HRTEM images of ZNSS powders coated with 40 cycles of\n2\ncoating. d ABF images and corrsponding (Zr, O, Ti element) EDS mapping of the bulk ZNSS\nsample coated with 40 cycles of TiO (TiO = 3.2 nm).\n2 2\nDetailed comment 7: “[7] Temperature-dependent  is missing in Fig. 6. Figs 6d and 6e are\nL\nthe same.”\nResponse: Thank you for your meticulous review and valuable suggestions. We apologize for\nthe oversight in Fig. 6 regarding the missing temperature-dependent κ . We have addedthe data in\nL\nthe revised version. This revision enhances the clarity of the temperature dependence and overall\nunderstanding of the thermal properties discussed.\nFig. 6 Electrical and thermal transport properties. The electrical and thermal properties of the ZNSS\nsamples before and after different ALD cycles, where uncoated, TiO = 0.8 nm, TiO = 1.6 nm, TiO\n2 2 2\n= 3.2 nm, TiO = 4.8 nm, TiO = 6.4 nm.a, b, cPF, d, e and fzT.\n2 2 L\nDetailed comment 8: “[8] It is assumed that the density functional theory (DFT) calculations\nwere conducted for ZrNiSn, TiNiSn, and ZrO at 0 K. However, these data cannot accurately\n2\nrepresent the band structure at 300 K and 800 K.”\nResponse: Thank you for your insightful comment. It is true that band structures can evolve with\ntemperature changes, which is critical for accurately linking DFT results to experimental data at\nelevated temperatures like 300 K and 800 K.\nIn thermoelectrics, DFT calculations at 0 K are standard due to computational limitations\nand typically provide a baseline approximation of the electrical transport properties. To address\ntemperature effects on the band structure, we rely on variable-temperature X-ray diffraction\n(XRD) to observe any significant changes in lattice parameters, indicative of thermal expansion\nand its impact on electronic properties.\nOur variable-temperature XRD results indicate that the lattice parameters for our samples\nremain relatively invariant between 300 K and 800 K. This stability is likely due to the robust\ncovalent bonds in ZrNiSn, resulting in a very low thermal expansion coefficient. The monoclinic\nphase of ZrO similarly exhibits minimal lattice changes with temperature, as supported by\nK band structure calculations for our thermoelectric materials at 300 K and 800 K.\nFurthermore, we computed the density-of-states effective mass for ZrNiSn at these\ntemperatures and confirmed it remains consistent at 2.7 m , corroborating the high stability of the\ne\nband structures as observed in studies by Prof. Tiejun Zhu and his team (Sci. Rep. 4, 6888\n(2014)).\nThis consistency across temperatures reinforces the validity of using 0 K DFT calculations\nto estimate material behaviors at higher temperatures within the specific context of this work. We\nacknowledge the limitations of this approach and continue to explore more precise methods to\nincorporate temperature-dependent changes in future work.\nFig. S5. a Variable temperature XRD patterns of TiO = 3.2 nm sample and b the corresponding\n2\nlattice parameters\nDetailed comment 9: “[9] Besides, if we take the calculated band structures of ZrNiSn and\nTiNiSn, we can only conclude that the delta E between them is around 0.2. How can you\nconclude that the delta E between Zr Ti NiSn Sb and ZrNiSn Sb is in the range of\n0.67 0.33 0.99 0.01 0.99 0.01\n0-0.2?”\nResponse: Thank you for your critical observations and recommendations regarding our\ncalculations of the band structures for ZrNiSn and TiNiSn, and their implications for the alloys\nZr Ti NiSn Sb and ZrNiSn Sb . We appreciate the opportunity to clarify these\n0.67 0.33 0.99 0.01 0.99 0.01\npoints further.\nRegarding the E calculations, the energy difference between ZrNiSn and TiNiSn is\napproximately 0.2 eV, as noted. For the alloys Zr Ti NiSn Sb and ZrNiSn Sb , we\n0.67 0.33 0.99 0.01 0.99 0.01\ninfer the E in the range of 0 to 0.2 eV based on linear interpolation between the two end-\nmembers (pure phases). This inference is based on the assumption that the partial substitution of\nTi for Zr introduces a proportional but limited perturbation to the band structure compared to the\nend-member phases.\nIn response to your query, we have revised our analysis and supplemented it with an updated\nband structure diagram for Zr Ti NiSn. This new diagram more accurately depicts the\n0.67 0.33\nconduction band minimum, clearly showing that it lies between those of ZrNiSn and TiNiSn. The\ninclusion of Ti exhibits a subtle upward shift in the conduction band minimum (around 0.1 eV),\nwhich is critical for understanding the energy barrier modifications when designing materials for\nenhanced thermoelectric performance.\nThe primary effect of Sb doping is to optimize carrier concentration rather than significantly\naltering the band structure itself. Sb doping tends to raise the Fermi level to near or within the\nconduction band, enhancing electrical conductivity without a substantial shift in band structures.\nThis aspect has been corroborated by previous studies involving Sb-doped ZrNiSn, which have\nshown minimal changes in the band features, thus supporting our assumptions and calculations\n(Sci Rep. 4, 6888 (2014)).\nFig. S11. Band structure diagram of a ZrNiSn, TiNiSn and b Zr Ti NiSn.\n0.67 0.33\nDetailed comment 10: “[10] The conductivity data has a certain trend with increased c numbers.\nI guess the conductivity is a tradeoff of the Ti doping (increase the n) and interface scattering\n(decrease the mobility).”\nResponse: We appreciate the reviewer’s keen observation and valuable insights. Indeed, after\napplying over 40 cycles of ALD, there is a slight increase in the carrier concentration. This\nenhancement maybe due to the incorporation of additional Ti atoms into the ZNSS matrix during\nthe TiO deposition process, which mildly boosts the carrier concentration.\nreduce carrier scattering, an excessive number of such interfaces—particularly with increased\nALD cycles—can act as scattering centers. This exacerbates carrier scattering and negatively\nimpacts carrier mobility.\nDespite these observations, the overall effect on electrical conductivity remains modest,\nprimarily due to the balance between the benefits of Ti doping and the drawbacks of interface\nscattering. We have included a comprehensive discussion of these trade-offs in Fig. S19 in the\nrevised supporting information.\nDetailed comment 11: “[11] I think modulation doping induced by adding TiO is the key to\n2\nenhancing the power factor.”\nResponse: We appreciate the reviewer's insightful commentary. Modulation doping effectively\nfacilitates the migration of charge carriers from doped to undoped zones, thereby reducing\nionized impurity scattering and enhancing carrier mobility. This mechanism significantly\nimproves electrical conductivity, a key determinant in the enhancement of PF.\nNevertheless, our observations indicated that the C = 40 sample displayed slightly lower\nelectrical conductivity compared to the pristine C = 0 sample. This reveals that the increase in\nPF in this work primarily results from a substantial increase in the Seebeck coefficient across all\nmeasured temperatures, coupled with sustained high carrier mobility. The notable enhancement\nof the Seebeck coefficient and the preservation of carrier mobility are attributed to the presence\nof coherent dual-interface energy barriers, which represent a deviation from traditional\nmodulation doping effects.\nMoreover, while modulation doping is conventionally utilized to boost carrier mobility, its\napplication extends beyond this to include sophisticated interface engineering. This approach can\nsignificantly refine the quality of semiconductor heterostructure interfaces. With an\nunderstanding of the broader potential of this technique, we are committed to exploring the\nunderlying physical mechanisms of modulation doping and its extensive impact on\nthermoelectric performance in our future research endeavors.\nDetailed comment 12: “[12] If the method can be applied to the high ZT half Heusler\ncompounds, this work will definitely deserve to be published in Nature Communications.”\nResponse: We appreciate the reviewers' high regard for our work. We fully concur with your\nassessment, and indeed, this vision has propelled our innovative application of ALD technology\nwithin the field of thermoelectrics. Our chosen matrix material, ZrNi Sn Sb , currently\n1.03 0.99 0.01\nachieves a peak zT of approximately 1.1—one of the highest values reported for n-type ZrNiSn\nhalf-Heusler materials. Through the implementation of coherent composite phase interface\nengineering via ALD, we have successfully elevated the zT to a record 1.3 within this category.\nIn addressing thermal transport, ALD facilitates the formation of a consistent three-\ndimensional interface network, which effectively scatters phonons at grain boundaries.\nFurthermore, an aspect not previously underscored in our discussion is the ability of ALD to\ninhibit Ostwald ripening. This is crucial as it prevents grain growth and maintains the\nthermoelectric performance over extended periods. From an electronic transport perspective,\nALD allows for the meticulous adjustment of interface barriers, thus optimizing the PF.\nThese findings affirm that ALD is not merely effective, but also exhibits significant\nversatility and scalability. This suggests promising potential for its application across a broader\nspectrum of materials beyond half-Heusler compounds. Encouraged by these outcomes, we\nintend to further explore these possibilities in our subsequent research initiatives.\nReferee #2:\nGeneral comment: “Minor Revision. The authors introduced the enhancement in the\nthermoelectric performance of ZrNiSn-based alloys using atomic layer deposition technology,\nwhich effectively decouple the interrelated thermoelectric parameters in ZrNiSn. I have carefully\nread the manuscript, and I found that this is an interesting work with promising results, and the\nstructure and logic are good. I am happy to recommend this work to be published in this journal\nafter some revisions, since there are still unclear points in the current manuscript. My comments\nlisted below may help the authors further improve their work:”\nGeneral response: We appreciate the reviewer’s positive comments. Please find our point-by-\npoint responses below.\nDetailed comment 1: “[1] The diagram in Fig. 1a should be redrawn, such that the second\nphase is tens to hundreds of nanometers in size and should not be represented by small balls. The\nthickness of the Ti-doped second phases coating the ZNSS matrix should not be less than the\ninterplanar spacing.”\nResponse: We sincerely appreciate your insightful suggestions concerning Fig. 1a. Following\nyour recommendations, we have revised the schematic diagram to more accurately depict the\nsize and morphology of the second phase, ensuring it is represented within the range of tens to\nhundreds of nanometers, consistent with dimensions observed experimentally. Furthermore, we\nhave adjusted the depiction of the Ti-doped layers to ensure that their thicknesses align with\nscientific accuracy, notably ensuring these thicknesses do not fall below the interplanar spacing.\nFig. 1 Schematic diagram of the decoupled thermoelectric properties of the working action\nof the coherent composite-phase. a Schematic diagram of electron and phonon transport in the\nmulti-scale phase interfaces. b Temperature-dependent zT for the TiO = 3.2 nm sample in this\n2\nwork, compared with those of other high-zT HH thermoelectric materials25,39-47.c Schematic\ndiagram of energy filtering effect. d Schematic diagram of electron and phonon transport in the\ncoherent phase interfaces.\nDetailed comment 2: “[2] It can be seen from the XPS spectra in Fig. S3 and EDS results in S8\nthat there is a lot of amorphous TiO on the ZNSS surfaces, but no obvious changes in the\n2\ndiffraction peaks caused by amorphous TiO as shown in Fig. S4, whether it is due to the\n2\ndeduction of the XRD pattern background or the fact that the content of amorphous TiO is less\n2\nthan the detection limit of XRD. The author needs further clarification.”\nResponse: We greatly value the reviewer's constructive critique. In response to your\nobservations concerning the XPS spectra in Fig. S3 and the EDS results in Fig. S8, we confirm\nthe presence of an amorphous TiO layer on the ZNSS surfaces. This observation is consistent\n2\nwith the absence of significant changes in the XRD diffraction peaks displayed in Fig. S4.\nThe lack of discernible diffraction peaks in the XRD patterns for the amorphous TiO can\n2\nbe attributed primarily to its non-crystalline nature, which inherently does not produce the sharp\nBragg peaks typical of crystalline materials. Additionally, the thickness of the amorphous TiO\n2\nlayer is less than 10 nm, a dimension that falls below the detection threshold of conventional\nXRD, which accounts for the absence of any discernible amorphous humps in the XRD patterns\n(Nanoscale 3, 1580-1588 (2020)).\nTo provide further clarity within our manuscript, we have expanded upon the discussion\nregarding the formation mechanism of the amorphous TiO layer and its detection limitations\n2\nwith XRD technology. A comprehensive explanation has been included in the main text to\nelucidate how the amorphous nature of TiO and its minimal thickness contribute to the absence\n2\nof detectable XRD signals.\nDetailed comment 3: “[3] It is difficult to explain the interface reaction mechanism based on\nSTEM and HAADF results in Fig. 2. This is because the HAADF image is essentially an atomic\ncolumn contrast image, and its contrast is affected by many factors. On page 8 of the manuscript,\nthe author's conclusion, “the residual Ti atoms occupy vacant Zr sites, which can create a Ti-\nZNSS layer…”, cannot be judged by the HAADF results alone. The STEM-EDS results should be\ncombined to explain the exact mechanism.”\nResponse: We appreciate the reviewer's insightful comments and have made significant revisions\nto our manuscript accordingly. We have enhanced our analysis by integrating high-angle annular\ndark-field (HAADF) imaging with energy-dispersive X-ray spectroscopy (EDS) results to offer a\nmore comprehensive explanation of the interface reaction mechanism.\nIn the revised Fig. 2d, we have replaced the HAADF image with detailed EDS elemental\nmappings, clearly depicting the distribution of Zr, O, and Ti. This approach provides more robust\nevidence of the interactions at the interface. The EDS mappings reveal that residual Ti atoms\nindeed occupy vacant Zr sites, forming a Ti-ZNSS layer. This observation not only corroborates\nour initial conclusion but also strengthens the scientific validity of our findings by presenting\ndirect and compelling evidence of the proposed mechanism.\nFig. 2 ALD Synthesis Technology Route. a Schematic diagram illustrating the synthesis of the TiO\n2\nlayer through ALD coating. b HRTEM images of ZNSS powders coated with 40 cycles of TiO . c\nimages and corrsponding (Zr, O, Ti element) EDS mapping of the bulk ZNSS sample coated with 40\ncycles of TiO (TiO = 3.2 nm).\n2 2\nDetailed comment 4: “[4] On page 10 of the manuscript, if only Zr position elements change,\nthe energy spectrum of other elements should be uniform. However, the distribution of elements\nof Ni and Sb is not very uniform in Fig. 3a, so is it appropriate to judge the reaction equation\nfrom the experimental results in EPMA? The authors can add some energy spectrum analysis to\nverify the accuracy of the reaction formula.”\nResponse: Thank you for raising this valuable question. We have provided further clarification\non the rationale behind the chemical equation and the preferential reaction of Zr with O when in\ncontact with TiO and ZrNi Sn Sb matrix, based on several factors:\n2 1.03 0.99 0.01\n1. Chemical affinity and reactivity: Zr exhibits a stronger chemical affinity for O compared\nto Ti, Sn, Ni, and Sb. This chemical affinity is largely dictated by the electronegativity\ndifferences between the metal atoms and oxygen. Specifically, Zr and O have an\nelectronegativity difference of 2.11, which is higher than those of Ti-O (1.90), Sn-O (1.48), Ni-O\n(1.53), and Sb-O (1.34). This greater difference enhances the likelihood of Zr bonding with O to\nform more stable compounds.\n2. Thermodynamic considerations: The formation energies of various oxides during high-\ntemperature processes are critical determinants of their stability and presence in the final product.\nZrO exhibits a formation energy of -3.801 eV, which is lower than that of TiO (-3.3 eV), SnO\n2 2 2\n(-2.123 eV), NiO (-1.761 eV), and Sb O (-1.728 eV). This indicates that ZrO is\n2 2 3 2\nthermodynamically the most favorable oxide under the experimental conditions, explaining the\npreferential formation of this phase.\n3. In-situ chemical dynamics: The ALD process promotes specific in-situ chemical reactions,\nparticularly between the TiO coating and the ZrNiSn matrix. This interaction facilitates the\n2\nformation of ZrO due to the reactivity of Zr with O, resulting in the creation of stable ZrO\n2 2\nnanoparticles. This dynamic is a key factor in the interface chemistry of our materials.\nRegarding the non-uniform distribution of Ni and Sn observed in Fig. 3a, this can be\nattributed to the intentional design of our matrix. The matrix is Ni-rich, a strategy based on prior\nresearch that suggests stability benefits from an excess of Ni, as evidenced by neutron scattering\nexperiments (J. Mater. Chem. C 3, 10534-10542 (2015)). Similarly, the enrichment of Sn is due\nto a small excess of Sn present in the samples, this is a result of slightly overadding Sn during the\npreparation of the ZrNi Sn Sb matrix using the suspension melting method, in order to\n1.03 0.99 0.01\nprevent the volatilization of Sn.\nDetailed comment 5: “[5] Fig. 6d and Fig. 6e are repeated, and Fig. 6e should be the lattice\nthermal conductivity. In addition, the test temperature when the minimum lattice thermal\nconductivity is reached should be added.”\nResponse: We sincerely appreciate the reviewer's meticulous and insightful feedback. We\napologize for the confusion caused by the duplication in Fig. 6d and 6e and the omission of\nspecific details regarding the temperature-dependent lattice thermal conductivity (κ ).\nL\nUpon reviewing the noted Fig. 6d and Fig. 6e, we have corrected Fig. 6e to properly display\nthe κ as a function of temperature in the revised manuscript. The revised Fig. 6e now includes\nL\nannotations that clearly indicate the temperature at which κ reaches its minimum value. These\nL\nmodifications have greatly enhanced the clarity and precision of our data presentation.\nFig. 6 Electrical and thermal transport properties. The electrical and thermal properties of the ZNSS\nsamples before and after different ALD cycles, where uncoated, TiO = 0.8 nm, TiO = 1.6 nm, TiO\n2 2 2\n= 3.2 nm, TiO = 4.8 nm, TiO = 6.4 nm. a, b, c PF, d, e and f zT. (The test temperature at\n2 2 L\nwhich the minimum  is observed is 873 K)\nL\nDetailed comment 6: “[6] The point-scanned data in EPMA image of Table S1 and S2 show that\nthe content of Ti in C20 and C40 samples is different with the increased of thickness of the\ndeposited atomic layer.\nThe authors need to add the point-scanned data of the other composition ZNSS/C (C=10, 60, 80),\nbecause I believe there is a certain difference with the nominal chemical composition.”\nResponse: We thank the reviewer for their insightful observations regarding the variation in Ti\ncontent across our samples. In response to your constructive suggestion, we have expanded our\ndataset to include point-scanned data for ZNSS/C samples with C = 10 and C = 60. These\nadditional data are now comprehensively presented in the revised Supplementary Information\n(Tables S1 - S4).\nThe EPMA point scanning analysis at points 1 and 2 indicates that the Ti content at the Ti-\nZNSS interface increases progressively with the number of ALD cycles. This trend substantiates\nthe effective integration of Ti into the ZNSS matrix. At point 3, the matrix composition remains\nunaffected by the interface, displaying the original composition (ZrNi Sn Sb ) without Ti.\n1.03 0.99 0.01\nThis consistency across multiple scans underscores the statistical robustness of our data.\nTable S1. Point-scanned data in EPMA image of TiO2 = 0.8 nm sample (1, 2 and 3 points).\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 30.375 33.755 32.895 2.602 0.373\n2 30.672 33.304 32.693 2.969 0.362\n3 32.157 34.237 33.221 0.000 0.381\nTable S2. Point-scanned data in EPMA image of TiO = 1.6 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 26.643 33.741 33.104 6.202 0.310\n2 26.150 33.955 33.308 6.269 0.318\n3 32.157 34.237 33.221 0.000 0.385\nTable S3. Point-scanned data in EPMA image of TiO = 3.2 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 23.770 33.141 32.928 9.874 0.287\n2 23.526 33.315 33.897 9.969 0.293\n3 33.154 33.554 32.971 0.000 0.321\nTable S4. Point-scanned data in EPMA image of TiO = 4.8 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 22.577 33.774 32.111 11.289 0.249\n2 22.103 33.367 32.617 11.662 0.251\n3 32.654 34.084 32.986 0.000 0.276\nDetailed comment 7: “[7] The serial numbers of the references in Fig. 8 do not correspond to\neach other. Please check them carefully.”\nResponse: We sincerely apologize for the oversight regarding the mismatch of reference serial\nnumbers in Fig. 8 and express our gratitude to the reviewer for identifying this issue. We have\nthoroughly reviewed and corrected the serial numbers of the references in the Fig. to ensure they\nare accurately aligned with the corresponding citations throughout the revised manuscript.\nDetailed comment 8: “[8] Error estimates for the different measured quantities can be included\nin experimental part or Fig.s for better comprehension.”\nResponse: Good advice. After thorough consideration, we have made the following\nenhancements to the revised manuscript to improve the clarity and completeness of the data\npresentation:\nIn the main text, we have focused on presenting data for the optimally prepared sample\nZNSS (C = 40). For this sample, we have now incorporated error bars into Fig. 6 to clearly\ndelineate the uncertainty and variability in our measurements.\nTo ensure the clarity of all Fig.s while providing comprehensive error data, we have\nincluded detailed experimental data lines with error estimates for all other samples in Fig. S19.\nThis allows us to present a full set of error estimates without overcrowding the main Fig.s,\nthereby improving the readability and accessibility of our data.\nFig. 6 Electrical and thermal transport properties with error lines. The electrical and thermal\nproperties of the ZNSS samples before and after different ALD cycles, where uncoated, TiO = 0.8 nm,\n2 2 2 2 L\nFig. S19 Electrical and thermal transport properties with error lines. The electrical and thermal\nproperties of the ZNSS samples before and after different ALD cycles, where uncoated, TiO = 0.8 nm,\n2 2 2 2 L\nDetailed comment 9: “[9] Kindly check the formatting of references and image resolution in\naccordance to journal standards. Such as there are some superfluous words (Adv. Energy Mater.\n37/2021) in reference 17, the numbers in the material composition in references 5 and 17 should\nbe small corner marks, and the resolution in Fig. S1 is obviously poor. These mistakes should\nnot have happened, and the authors should confirm them carefully before resubmitting.”\nResponse: We sincerely apologize for the oversight and have taken comprehensive measures to\naddress these issues to enhance the quality of our submission:\nReference Formatting: We have thoroughly reviewed and corrected all references to ensure\nstrict adherence to the journal's formatting requirements. Specifically, we have removed the\nsuperfluous words in reference 17 and adjusted the material composition numbers in references 5\nand 17 to small corner marks as required.\nImage Resolution and Fig. Quality: We acknowledge the poor resolution of Fig. S1 and\nhave replaced it with a high-resolution version that meets the journal’s standards. Additionally,\nwe have carefully examined all other Fig.s in the revised manuscript and supplementary\ninformation to ensure they exhibit optimal clarity and consistency. Each Fig. has been\nmeticulously revised to include clear scales, comprehensive descriptive text, and to maintain\nuniformity in Fig. numbering and layout.\nFig. S1. XRD pattern of the as-prepared ZrNi Sn Sb (ZNSS) powders.\n1.03 0.99 0.01\nDetailed comment 10: “[10] The XPS fit peaks in Fig. S3 are rough, especially the peaks near\nthe 462 eV, so it is recommended to fit them after removing the background of XPS spectra.”\nResponse: Thank you for your valuable feedback. In response, we have addressed the concerns\nby removing the background noise and refining the peak fitting process, particularly around 462\neV. This adjustment has substantially improved the clarity and accuracy of our spectral analysis,\nleading to a more precise interpretation of the XPS data.\nFig. S3. XPS spectra of Ti 2p and Ti 2p core level for TiO = 3.2 nm sample.\n1/2 1/3 2\nDetailed comment 11: “[11] For a good scientific paper, supporting documents and manuscript\nare equally important. However, the Fig.s in the supporting files are rough, such as no picture\nnumbers and ruler in Fig. S9 Fig. S11, and Fig. S13.”\nResponse: We appreciate the reviewer's attention to detail and valuable feedback. In response,\nwe have enhanced the quality of the Fig.s in our revised supplementary materials. Specifically,\nwe have updated Fig.s S9, S11, and S13 by adding Fig. numbers and scales, ensuring clarity and\nconsistency. These improvements significantly increase the readability and interpretability of the\nFig.s, aligning with the high standards expected for scientific documentation.\nFig. S10.a HADDF, b ABF and c FFT images of matrix ZNSS.\nFig. S11. a HADDF and b ABF images of ZrO .\nfuture direction for practical applications.”\nResponse: We appreciate the reviewer's suggestion regarding the inclusion of future research\ndirections. In response, we have expanded our discussion to underscore the potential practical\napplications of the ZrNiSn-based thermoelectric materials explored in this work.\nALD enables the creation of a uniform three-dimensional interface network, which\nsignificantly scatters phonons at interfaces and allows for the modulation of interface barriers.\nAn important feature, previously not emphasized, is the capability of coated interface layers to\nprevent Ostwald ripening. This property is critical as it inhibits grain growth, thereby preserving\nthe thermoelectric performance over extended durations. Given their robust thermoelectric\nproperties and superior oxidation resistance, these materials are highly suitable for demanding\nenvironments encountered in applications such as high-temperature waste heat recovery and\npower generation. Their suitability for such extreme conditions is bolstered by their consistent\nperformance under high temperature and oxidative stress.\nReferee #3:\nGeneral comment: “Zhang and coworkers have explored thermoelectric ZrNiSn using a range\nof advanced techniques, including aberration-corrected high-resolution microcopy. The authors\nhave utilized an interesting idea to engineer the interfaces by inclusions of TiO layers through\n2\natomically controlled processes. They have shown that scattering phonons against such\ninterfaces, while maintaining high carrier mobility and high effective mass (energy filtering), is a\npotent approach to enhance thermoelectric efficiency (ZT). Consequently, the ZT value reaches\n1.3 at 900 K, which is an important step towards broader applications. Such a coherent interface\ndesign is novel and deserves attention of the thermoelectric community and may have impact on\nother functional materials. The manuscript is well written. Hence, this manuscript should be\naccepted for publication after addressing the points raised below.”\nGeneral response: We appreciate the reviewer’s positive comments. Please find our point-by-\npoint responses below.\nDetailed comment 1: “[1] Since thermoelectric materials are often exposed to harsh conditions\n(atmosphere, high temperatures) under cyclic loading, a coherent interface may also be relevant\nto counteract these effects. Is it conceivable that further oxidation can be hindered by inclusion\nof TiO ? Please discuss these issues in the manuscript. This would give a larger perspective.”\ndeposition (ALD) to apply a uniform TiO coating on the ZrNiSn matrix, which significantly\n2\nenhanced the oxidation resistance of the ZrNiSn powders, as illustrated in Fig. 2b.\nDuring the high-temperature sintering process, this TiO coating reacted in situ with the\n2 2\nrenowned for its stability, further bolsters the material's resistance to oxidation. Consequently,\nsamples with the TiO coating exhibited superior oxidation resistance at elevated temperatures\n2\ncompared to the uncoated ZrNiSn matrix.\nFollowing your suggestion, we have expanded our discussion in the revised manuscript to\nfurther explore the potential of the TiO layer to enhance the material’s durability and inhibit\n2\nfurther oxidation, particularly under severe conditions and cyclic loading. This enhancement\nbroadens the scope of these thermoelectric (TE) materials' applications in challenging\nenvironments.\nDetailed comment 2: “[2] The introduction is too short regarding the physical (transport)\nproperties of ZnNiSn and related open questions to be tackled in the manuscript. Please provide\na broader review on ZrNiSn, especially alternative pathways to enhance its thermoelectric\nperformance (interface engineering with Al O (Key Engineering Materials Vol. 249 (2003) pp\n2 3\n79-82) and amorphous configurations (Computational Materials Science 230 (2023) 112530)).\nAnother reference regarding Al O has been mentioned in the results section, but the\n2 3\nintroduction leaves the impression that interface engineering has not been attempted.\nFurthermore, why is TiO a better choice than Al O ? Could an amorphous interface (e.g.,\n2 2 3\ncrystalline ZrNiSn interfaced with amorphous ZrNiSn) have a similar performance?\nFurthermore, the addition of Sb to ZrNiSn and selection of TiO should also be more motivated.”\nthem together as follows.\nZrNiSn, an n-type thermoelectric material, exhibits a substantial density-of-states effective\nmass (m* ~ 2.7 m ) and a low deformation potential. These characteristics enable it to maintain a\ne\nhigh Seebeck coefficient while retaining decent carrier mobility ( ), resulting in a high power\nH\nfactor. However, the high lattice thermal conductivity () limits its thermoelectric performance.\nL\nTo mitigate this, alloying has proven effective in enhanciing the zT by reducing . Additionally,\nL\nnanostructuring strategy, such as grain refinement and increased grain boundary scattering, have\nbeen employed to scatter phonons more effectively. Interface engineering, including\nincorporating Al O through ball milling, has successfully reduced  and improved zT (Key\n2 3 L\nEngineering Materials 249, 79-82 (2003)). Furthermore, introducing amorphous ZrNiSn has\nachieved ultralow  (Computational Materials Science 230, 112530 (2023)). Despite these\nL\nadvancements, these approach did not enhance the electrical PF due to decreased  and,\nH\nconsequently, electrical conductivity. Thus, the coupled nature of electron and phonon transport\nproperties limits the extent of zT improvement.\nIn this work, we adopted a composite coherent interface engineering strategy aimed at\neffectively scattering phonons while minimizing the reduction in  . By leveraging the energy\nH\nfiltering effect, we aimed to enhance the density-of-states effective mass and achieve synergistic\noptimization of electron and phonon transport properties, significantly improving the zT of\nZrNiSn-based thermoelectric materials.\nRegarding the choice of Al O and TiO as coating materials, Al O was initially considered\n2 3 2 2 3\ndue to its high melting point and chemical stability. However, high-temperature phase interface\nreactions revealed that the Al element's stability is inferior to that of the Ti element under similar\nconditions. XPS analysis indicated that after sintering at 1100 °C, the Al O coating often\n2 3\nexhibited Al element deficiency, likely due to Al volatilization or complex reactions with the\nmatrix (J. Mater. Chem. A 7, 26053–26061 (2019)). Conversely, TiO demonstrated better\n2\nstability during high-temperature sintering. The Ti element is less volatile and more likely to\nform a stable solid solution with the ZrNiSn matrix. We observed Ti doping into the interface\nlayer, enhancing interface stability and optimizing the band structure, thereby improving\nthermoelectric performance.\nRegarding Sb doping, our primary goal was to optimize the carrier concentration (n ) of the\nH\nZrNiSn matrix. As a donor dopant, Sb effectively increases the n from approximately 1019 cm-3\nH\nto 1020 cm-3, which is crucial for improving the PF. Building on Sb doping, we further facilitated\nthe co-optimization of both carrier and phonon transport properties, substantially enhancing the\nzT through composite coherent interface engineering. This strategy ensures a balanced\nimprovement in thermoelectric performance by effectively managing the interplay between\nelectrical and thermal transport.\nThese discussions have been incorporated into the revised introduction Page 3 - 4 (70-73\nlines).\nThe respective content has been updated in the revised manuscript, as seen below: For\nexample, the introduction of Al O into ZrNiSn or the creation of amorphous ZrNiSn have both\n2 3\nsuccessfully led to a significant reduction in . However, these approaches did not effectively\nL\nenhance PF due to the decrease in . Consequently, the coupling between electronic and phonon\nH\ntransport properties has limited the potential for improving the figure of merit zT.\nDetailed comment 3: “[3] ZrO nanoparticles are obviously important. Please also discuss their\n2\nrole in the abstract.”\nResponse: We appreciate the reviewer's insight regarding the critical role of ZrO nanoparticles.\n2\nthe presence of numerous twin boundaries within the ZrO , which significantly enhances phonon\n2\nscattering and thereby effectively reduces .\nL\nMoreover, we have engineered a coherent interface between the ZrO nanoparticles and the\nfurther diminishes the  but also mitigates potential negative impacts on carrier transport.\nL\nAcknowledging the constraints on word count in abstracts, we have succinctly outlined the\nrole of ZrO in the revised abstract to underscore its significance. A more detailed discussion on\n2\nthe implications of ZrO for enhancing thermoelectric performance is comprehensively presented\n2\nin the introduction and related work sections of the manuscript Page 5 (117-119 lines).\nThe respective content has been updated in the revised manuscript, as seen below:\nFurthermore, we are pleasantly surprised to discover that there are a significant number of twins\nwithin the ZrO nanoparticles. This twin structure provides an additional scattering center for\n2\nphonon transport at the nanoscale.\nDetailed comment 4: “[4] Why is there only oxidation of Zr when TiO is in contact with\nElectronegativity, as mentioned on page 10, is only a partial cause. What about thermodynamics?\nEq. (1) is not really supported.”\nResponse: We thank the reviewer for this insightful comment. The preferential oxidation of Zr,\ndespite the presence of other metals like Ti, Sn, Ni, and Sb, can be attributed to a combination of\nfactors influencing the chemical reactivity and thermodynamic stability during the synthesis\nprocess:\n1. Chemical affinity: Zr exhibits a stronger chemical affinity for O compared to Ti, Sn, Ni,\nand Sb. This affinity is largely dictated by the electronegativity differences between the metal\natoms and oxygen. Specifically, Zr and O have an electronegativity difference of 2.11, which is\nhigher than those of Ti-O (1.90), Sn-O (1.48), Ni-O (1.53), and Sb-O (1.34). This greater\ndifference enhances the likelihood of Zr bonding with O to form more stable compounds.\n2. Thermodynamic stability: From a thermodynamic perspective, ZrO is more stable than\n2\nthe oxides of the other metals involved. During the high-temperature sintering process, the\nformation energies of the oxides are crucial in determining their formation likelihood. ZrO has a\n2\nformation energy of -3.801 eV, which is notably lower than those of TiO (-3.3 eV), SnO (-\n2 2\n2.123 eV), NiO (-1.761 eV), and Sb O (-1.728 eV). The lower formation energy of ZrO\n2 2 3 2\nindicates a higher thermodynamic propensity for its formation.\n3. Interface energy considerations: The synthesis process fosters a coherent interface\nbetween ZrO and the ZrNiSn matrix. This interface, being energetically favorable, further\n2\nfacilitates the selective formation of ZrO nanoparticles at the interface due to reduced interface\n2\nenergy.\nThese factors collectively explain why Zr reacts preferentially with O during the synthesis,\nleading to the predominant formation of ZrO , while the other elements remain relatively inert\n2\nunder the same conditions.\nDetailed comment 5: “[5] How does the zT value obtained in this work correlate with other\nHeusler systems? Please discuss.”\nResponse: Thank you for your insightful comments. The zT value obtained in this study\nrepresents both an evolutionary step and a substantial advancement from traditional Heusler half-\nHeusler (HH) systems. Our approach builds upon an optimized ZrNiSn matrix, derived from\nexisting literature (Adv. Energy Mater. 5, 1500588 (2015)), where excess Ni reduces , and Sb\nL\ndoping optimizes n . This foundational strategy enabled our non-ALD coated samples to achieve\nH\na zT value of 1.1 at 873 K, one of the highest for n-type ZrNiSn systems.\nFurthe, we employed ALD to introduce composite phase interface engineering. This method\nfurther simultaneously tunes the electron and phonon transport in ZrNiSn-based alloys, elevating\nthe zT value to 1.3 at 873 K, the highest recorded to date in this context.\nThe precision of the ALD technique, which allows for precisely control over specific\ninterfaces, is critical. This level of control, challenging to attain through traditional alloying and\ndoping techniques, enhances the zT values significantly. Moreover, the application of coherent\ncomposite phase interface engineering via ALD, as developed in this work, not only benefits the\nZrNiSn-based system but also holds promise for enhancing other HH and thermoelectric systems.\nDetailed comment 6: “[6] Designating TiO layers as “C = 0, C = 10, C = 20, C = 40, C = 60\n2\nand C = 80” is misleading. If somebody wants to reproduce these results (perhaps by using\nanother technique), more general descriptors are needed (e.g., thickness rather than cycles).\nFor example, C = 40 means that the TiO thickness is about 3 nm (page 8). Why is this\n2\ncomplicated designation (method oriented) necessary?”\nResponse: Thank you for your valuable comments on the nomenclature of the TiO layer.Our\n2\ninitial rationale for using the number of ALD cycles as descriptors was to provide a direct and\nstraightforward method to control and report the deposition of TiO layers. We recognize that\n2\nwhile specifying ALD cycles is simple and unambiguous for those familiar with the technique, it\nmay not be universally applicable or intuitive for researchers using different deposition methods.\nYou rightly pointed out that using thickness as a descriptor could offer a more general and\neasily reproducible metric, though it is subject to variations based on the specific deposition\ntechnique, materials used, and process parameters. To bridge this gap, we have included detailed\ncorrelations between the number of ALD cycles and the actual thicknesses of TiO layer in Table\nTiO layer approximately 3 nm thick.\nemphasizes physical dimensions like thickness rather than process-specific metrics such as\ncycles. This adjustment will undoubtedly enhance the clarity, applicability, and reproducibility of\nour findings across different research settings.\nDetailed comment 7: “[7] It would be useful to designate the binding energy of both\ncontributions in Fig.. S3. Please also compare the binding energy with literature.”\nResponse: Thank you for your valuable suggestion. We have marked the binding energies of\nTi’s 2p orbitals in the Fig. S3. Specifically, our fitting results indicate that the binding energies\nfor the Ti 2p and Ti 2p orbitals are approximately 464.5 eV and 458.8 eV, respectively.\n1/2 3/2\nFurthermore, we have compared these values with the binding energy data from similar\nXPS studies reported in the literature on functional materials. Our findings align closely with\nthese reports (Nat. Commun. 4, 2214-2222 (2013), Nat. Commun. 11, 4613-4621 (2020).) which\nconsistently indicate a Ti valence state of +4 and corroborate the formation of amorphous TiO .\n(SI) to provide a more comprehensive understanding and verification of our results.\nFig. S3. XPS spectra of Ti 2p and Ti 2p core level for TiO = 3.2 nm sample.\n1/2 1/3 2\nDetailed comment 8: “[8] Does hydrogen have any role here? Having H O as a precursor\n2\nshould not be ignored. In many fields, H is instrumental. Please discuss.”\nResponse: Thank you for your question about the role of hydrogen in this study. The ALD cycle\nincludes multiple steps: exposure to TDMAT, followed by an Ar (g) purge, exposure to H O\n2\nvapor, and a final Ar (g) purge. The TDMAT precursor and the substrate are heated at\ntemperatures of 40 °C and 150 °C, respectively, while the H O vapor is kept at room temperature.\n2\nthe reaction, are then purged from the reaction chamber with argon gas. Additionally, the\nrelatively low temperatures used in our process are designed to prevent the diffusion of elements,\nincluding hydrogen.\nBased on the analysis and consistent with other studiesusing H O as an ALD precursor (Nat.\ndeposited TiO thin films and the final sample is negligible. This has been empirically confirmed\n2\nand is considered to have no significant impact on the properties of the films produced by the\ndescribed ALD technique.\nDetailed comment 9: “Fig.. S4 has never been mentioned in the manuscript. If it’s redundant,\nplease remove it. Please check other Fig.s in the supplementary materials as well.”\nResponse: We apologize for the oversight of not mentioning Fig. S4 in the initial manuscript. It\nhas been noted that Fig.. S2 and S4, while appearing similar, each depict critical stages of the\nexperimental process. Fig.. S2 illustrates the XRD pattern of the powder post-ALD, highlighting\nan amorphous coating that is sufficiently thin—less than 10 nm—rendering it below the\nresolution capabilities of XRD for detecting secondary phase peaks.\nIn contrast, Fig. S4 represents the XRD diffraction pattern following SPS, which similarly\ndoes not display distinct secondary phase peaks. This absence is attributed to the composite\nphase content being below the detection threshold of XRD (5%). This phenomenon aligns with\nfindings reported in related ALD studies, such as those documented in Nano Energy, 49, 257–\n266, (2018).\nAn extensive review of all Fig.s in the supplementary materials was conducted, ensuring\nthat each is accurately mentioned within the revised manuscript to clarify their relevance and\nimportance to the study.\nDetailed comment 10: “[10] Is the compositional analysis as accurate as presented in Table S1?\nAre all digits statistically reasonable to supply? What is the error bar? Why is O omitted in\nTable S1? Similar goes for Table S2.”\nResponse: We appreciate your attention to detail and constructive feedback. Here is a refined\nexplanation addressing your concerns about the data presented in Tables S1 and S2:\n1. Accuracy and error margin of EPMA analysis: EPMA is renowned for its high precision\nand quantitative accuracy, with a detection limit typically between 0.01 and 0.05 weight percent.\nThe typical error margin for EPMA analysis is within ±3%, ensuring that the presented data are\nwithin a reasonable uncertainty range. As a result, we assert that the compositional data in Tables\nS1 - S4 are both accurate and reliable.\n2. Statistical validity: The EPMA point scanning analysis at points 1 and 2 indicates the Ti\ncontent at the Ti-ZNSS interface increases progressively with the number of ALD cycles. This\ntrend substantiates the effective integration of Ti into the ZNSS matrix. At point 3, the matrix\ncomposition remains unaffected by the interface, displaying the original composition\n(ZrNi Sn Sb ) without Ti.\n1.03 0.99 0.01\nThis consistency across multiple scans underscores the\nstatistical robustness of our data.\n3. Omission of oxygen: The focus of our point scanning data collection was primarily on the\nTi-ZNSS layer and the matrix, not on the ZrO secondary phase. Consequently, oxygen was\n2\nomitted in Tables S1 to S4. However, the EPMA backscattered electron image in Fig. 3a of the\nmain text clearly shows the presence of oxygen and confirms the formation of ZrO .\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 30.375 33.755 32.895 2.602 0.373\n2 30.672 33.304 32.693 2.969 0.362\n3 32.157 34.237 33.221 0.000 0.381\nTable S2. Point-scanned data in EPMA image of TiO = 1.6 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 26.643 33.741 33.104 6.202 0.310\n2 26.150 33.955 33.308 6.269 0.318\n3 32.157 34.237 33.221 0.000 0.385\nTable S3. Point-scanned data in EPMA image of TiO = 3.2 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 23.770 33.141 32.928 9.874 0.287\n2 23.526 33.315 33.897 9.969 0.293\n3 33.154 33.554 32.971 0.000 0.321\nTable S4. Point-scanned data in EPMA image of TiO = 4.8 nm sample (1, 2, 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n1 22.577 33.774 32.111 11.289 0.249\n2 22.103 33.367 32.617 11.662 0.251\n3 32.654 34.084 32.986 0.000 0.276\nDetailed comment 11: “[11] The authors argue about strain in the interfaces. How is strain\naffecting the transport properties?”\nResponse: Thanks for your valuable comment. In thermoelectrics, strain distribution crucially\ninfluences both carrier and phonon transport. In particular, high-entropy alloys exhibit significant\nstrain fields that intensify phonon scattering, effectively decreasing the , as corroborated by\nL\nrecent research (Science 384, 81–86 (2024)).\nEmploying geometric phase analysis (GPA) technology, we precisely delineated local strain\ndistributions. Contrary to traditional nanocomposites, where strain is predominantly localized at\nphase boundaries, our findings demonstrate a strain distribution throughout the entire particle.\nThis strain distribution is particularly noteworthy along the ε direction at the twin boundaries\nxx\nwithin the ZrO structure, where it markedly impedes phonon propagation. This mechanism\n2\nsignificantly contributes to the observed reduction in , highlighting the critical influence of\nL\nstrain on thermal transport in thermoelectric materials.\nDetailed comment 12: “[12] Please provide a reference for the procedure pertaining to the\nPisarenko curve.”\nResponse: We appreciate the reviewer's request for a reference detailing the procedure for the\nPisarenko plot. For an in-depth description of this method, please refer to the study published in\nAdv. Funct. Mater. 23, 5123–5130 (2013), which provides a comprehensive analysis of the\nPisarenko curve and its applications in thermoelectric materials.\nThe mathematical formulation for the room-temperature m* in the SPB model (Phys. Rev. B.\n80, 125205 (2009)) includes several components, with key equations provided below:\n(1)\n(2)\n(3)\nwhere e is the electron charge, k is the Boltzmann constant, ћ is the reduced Planck\nB\nconstant,  is the reduced Fermi level and r is the scattering factor, which is taken as r = -1/2 for\nboth the acoustic phonon scattering and alloy scattering, F() is the jth order Fermi integral.\nj\nDetailed comment 13: “[13] There are some typos in the manuscript (e.g., “fromthe” on page\n16 and “Mpa” on page 23 (it should be MPa)). Please proofread the text.”\nResponse: We appreciate the reviewer's attention to detail in identifying the typos in our\nmanuscript. We sincerely apologize for these oversights and have thoroughly proofread the\nmanuscript to correct these and any other errors.\nSpecifically, we have corrected \"fromthe\" to \"from the\" on page 16 and \"Mpa\" to \"MPa\" on\npage 23, ensuring all units and spacing are consistent with scientific standards.\nThank you for helping us improve the clarity and accuracy of our manuscript.\nDetailed comment 14: “[14] Band structure in Fig.. S10 is given for ZrNiSn and TiNiSn as well\nas ZrO in Fig.. S12, rather than the actual system investigated in the current work. Drawing\n2\nconclusions from a very idealized model system(s) can be ambiguous. Please elaborate.”\nResponse: In view of the complex material composition studied herein, the reviewer raised a\nvery important point. We address the reviewer’s question as follows:\nAddressing the concern regarding the use of idealized model systems, it is important to\nhighlight the inherent trade-offs between the reliability (accuracy and precision) and the cost of\nstructural modeling. Computational studies frequently utilize idealized models due to the\ncomplexity and variability inherent in experimental setups. To mitigate any potential ambiguities\narising from our modeling approach, we clarify that our coated samples are systematically\ndesigned and processed based on the matrix samples, with a specific focus on the thermoelectric\nperformance influenced by the proposed composite phase interface engineering.\nThe composite phase interfaces are bifurcated into two main components: the Ti-ZNSS\nlayer and the ZrO nanoparticles. For the Ti-ZNSS layer, additional band structure calculations\n2\nwere conducted for Zr Ti NiSn. This is corroborated by EPMA point scanning data (Tables\n0.67 0.33\nS1 - S4), which demonstrate a proportional increase in Ti content in the Ti-ZNSS layer as the\nthickness of the TiO coating layer increases. This increase corresponds to an upward shift in the\n2\nconduction band position as depicted in our band structure calculations (Fig.. S12, ZrNiSn-\nZr Ti NiSn-TiNiSn). Considering that a minor excess of Ni is primarily employed to\n0.67 0.33\ndecrease lattice thermal conductivity, while a small amount of Sb doping is utilized to optimize\ncarrier concentration, studies have reported that the influence of these modifications on the band\nstructure is negligible (Sci. Rep. 4, 6888 (2014), ACS Appl. Mater. Interfaces 11, 47830–47836\n(2019)). Thus, it is justifiable that band calculations omit the effects of both elements. Regarding\nZrO , its robust chemical bonds significantly restrict the incorporation of additional elements\n2\ninto its matrix.\nFurthermore, considering that the calculation of band structure is usually at 0 K, we have\nalso conducted variable-temperature XRD tests, which show that the lattice constants of the\nZrNi Sn Sb sample remain unchanged from room temperature to 873 K, indicating that\n1.03 0.99 0.01\nthe band structure calculated at the 0 K point is also applicable at high temperatures (Fig. S5).\nBesides, the monoclinic phase of ZrO also exhibits extremely small lattice temperature\n2\nvariations, which has been confirmed in the Phys. Rev. B.70, 245116 (2004).\nFig. S12. Band structure diagram of a ZrNiSn, TiNiSn and b Zr Ti NiSn.\n0.67 0.33\nFig. S5. a Variable temperature XRD patterns of TiO = 3.2 nm samples and b the corresponding\n2\nlattice parameters\nDetailed comment 15: “[15] The statement that the “base pressure of ~ 1 Torr” was achieved in\nALD experiments doesn’t seem to be correct. This is probably the working pressure (1 Torr as a\nbase pressure may give rise to full oxidation). Please state the base pressure (pressure before\ndepositions).”\nResponse: We appreciate the reviewer's astute observation and apologize for any confusion\nregarding the pressure terminology in our ALD experiments. In this work, the TiO layers were\n2\ndeposited on the matrix powder surface using a homemade continuous-flow ALD reactor. Ar gas\nwas employed both as a carrier and a purging gas, with the working pressure maintained at\napproximately 1 Torr.\nTo clarify, the mentioned pressure of ~1 Torr refers to the working pressure during the\ndeposition process, not the base pressure. The actual base pressure, measured before the start of\nthe depositions, was maintained at approximately 5 × 10-6 Torr, ensuring a clean and controlled\nenvironment to prevent full oxidation.\nDetailed comment 16: “[16] More details are needed for the characterization techniques (e.g.,\npower setting for XRD, pass energy for XPS, accelerating volage for microscopy, calibration\nprocedure for DSC, etc.).”\nResponse: We appreciate the reviewer's request for additional details regarding our\ncharacterization techniques. To enhance transparency and ensure reproducibility, we have\nexpanded the experimental section with the following specific details:\nXRD testing conditions: For phase structure analysis using the Rigaku Smartlab 9kW X-ray\ndiffractometer, we specified a tube voltage of 45 kV and a tube current of 200 mA.\nXPS testing details: In our X-ray photoelectron spectroscopy (XPS) analysis, the pass\nenergy was set to 20 eV, essential for accurate surface electronic state characterization.\nEPMA and HRTEM parameters: The acceleration voltages for electron probe microanalysis\n(EPMA) and high-resolution transmission electron microscopy (HRTEM) ARM300 were set at\n15 kV and 300 kV, respectively.\nDSC calibration procedure: We detailed the calibration procedure for differential scanning\ncalorimetry (DSC), using sapphire as a standard. This procedure includes blank tests, sapphire\nstandard tests, and sample tests, ensuring precise baseline correction for each DSC curve.\nREVIEWER COMMENTS\nReviewer #1 (Remarks to the Author):\nThe authors have adequately addressed all the comments that I raised. Therefore, I have no\nfurther comment, and in my opinion, it is ready for publication in Nature Communications.\nReviewer #2 (Remarks to the Author): attached\nReviewer #3 (Remarks to the Author):\nThe authors have improved the manuscript. The work is novel and of very high quality\nrequired for this journal. It should be provisionally accepted if the authors reduce the number\nof supplied digits in Table S1 – S4 to statistically relevant ones. In the rebuttal, they have\nclaimed that “the typical error margin for EPMA analysis is within ±3%”. Depending if this\nerror is taken as absolute or relative, the Zr content in Table S1 (position 1) should be either\n30 at.% or 30.4 at.% but certainly not 30.375 at.% (currently supplied value). For instance,\nreporting the composition with a “precision” higher than 1 at.% can often be very\nquestionable (see, e.g., Surface & Coatings Technology 473 (2023) 130020). Surely the\nauthors of the current manuscript don’t have the precision of 0.001 at.% by EPMA.\nThe author answered and resolved the questions raised by the reviewer very carefully,\nand therefore agreed to publish.\nReferee #3:\nGeneral comment: “The work is novel and of very high quality required for this journal. It\nshould be provisionally accepted if the authors reduce the number of supplied digits in Table S1\n– S4 to statistically relevant ones. In the rebuttal, they have claimed that “the typical error\nmargin for EPMA analysis is within ±3%”. Depending if this error is taken as absolute or\nrelative, the Zr content in Table S1 (position 1) should be either 30 at.% or 30.4 at.% but\ncertainly not 30.375 at.% (currently supplied value). For instance, reporting the composition\nwith a “precision” higher than 1 at.% can often be very questionable (see, e.g., Surface &\nCoatings Technology 473 (2023) 130020). Surely the authors of the current manuscript don’t\nhave the precision of 0.001 at.% by EPMA.”\nGeneral response: Thank you for your positive evaluation and the constructive comments on our\nmanuscript. We have taken your suggestions about the precision of the EPMA data in Tables S1-\nS4 seriously and have made the necessary adjustments to our supplementary tables accordingly.\nTable S1. Point-scanned data in EPMA image of TiO = 0.8 nm sample (1, 2 and 3 points).\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n30.4 33.8 32.9 2.6 0.4\n1\n30.7 33.3 32.7 3.0 0.4\n2\n32.2 34.2 33.2 0.0 0.4\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n26.6 33.7 33.1 6.2 0.3\n1\n26.2 34.0 33.3 6.3 0.3\n2\n32.2 34.2 33.2 0.0 0.4\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n23.8 33.1 32.9 9.9 0.3\n1\n23.5 33.3 33.9 10.0 0.3\n2\n33.2 33.6 33.0 0.0 0.3\n2\nposition Zr at% Ni at% Sn at% Ti at% Sb at%\n22.6 33.8 32.1 11.3 0.2\n1\n22.1 33.4 32.6 11.7 0.3\n2\n32.7 34.1 33.0 0.0 0.3\nReviewer #3 (Remarks to the Author):\nThe authors have modified the manuscript in a satisfactory manner. It’s acceptable for\npublication in the current form."}]

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