Dual relaxation behaviors driven by a homogeneous and stable dual-interface charge layer based on an EGaIn absorber

Abstract

Interface engineering, by modulating defect distribution and impedance at interfaces and inducing interfacial polarization, has proven to be an effective strategy for optimizing dielectric properties. However, the inherent incompatibility between heterogeneous phases presents a significant challenge in constructing multi-heterointerfaces and understanding how their distribution influences dielectric performance. Herein, we constructed an EGaIn@Ni/NiO/Ga₂O₃ composite structure by employing a low-intensity ultrasound-assisted galvanic replacement reaction followed by high-temperature annealing. The controlled addition of Ni salts allowed for the fine-tuning of Ni, NiO, and In concentrations and their spatial distribution within the interfacial architecture. Annealing treatment induced a transition from amorphous to crystalline phases, triggering dual relaxation behaviors between EGaIn/Ni and NiO/Ga₂O₃. Additionally, significant charge accumulation was observed at the NiO/Ga₂O₃ interface, likely due to the substantial work function difference between Ni and NiO, coupled with the low barrier height between EGaIn and Ni, which facilitates electron migration. Consequently, the optimized samples exhibited a maximum absorption bandwidth of 7.92 GHz, which is the highest among the EGaIn-based absorbers reported in the literature. This work not only elucidates the mechanism by which multi-heterogeneous interfacial distributions regulate the dielectric properties but also provides an effective approach for modulating the electromagnetic wave performance of liquid metals.

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New concepts

In this study, we propose a novel strategy centered on the regulation of multi-heterointerface distribution (the interfaces formed by diverse substances and their evolving distribution) to achieve multiple polarization effects. Previous studies have explored the construction of core–shell structures in liquid metals (LMs) through covalent substitution reactions; however, these efforts primarily focused on controlling surface components and only induced a single polarization phenomenon. In our work, utilizing a low-intensity ultrasound-assisted covalent substitution reaction, combined with high-temperature annealing and precise control over nickel salt addition, we successfully drove the self-assembly of heterogeneous components on the surface of the eutectic gallium–indium (EGaIn) alloy, resulting in the formation of an EGaIn@Ni/NiO/Ga₂O₃ multi-heterointerface structure. Dielectric performance characterization revealed that this multi-heterointerface configuration facilitated the formation of multiple interfacial charge accumulation zones, thereby triggering multiple polarization phenomena. The distinct charge accumulation at different interfaces significantly enhanced the EGaIn alloy's responsiveness to electromagnetic waves (EMWs) across various frequency bands. Overall, this innovative approach offers new perspectives and methodologies for regulating the dielectric properties of liquid metals within the field of materials science.

Introduction

The dielectric properties of materials are closely linked to their polarization behavior, energy dissipation, and stability under an electric field, which directly influence the material's response to external electric fields and its potential applications in fields such as information technology, energy storage, and smart devices.^1–4 This relationship is particularly critical in the domain of electromagnetic wave absorption, where the material's polarization behavior and carrier transport capacity are of paramount importance.^5,6 Recent advances have highlighted the manipulation of dielectric properties through compositional adjustments, defect engineering, and interface engineering, which have emerged as key areas of research. Among these, interface engineering has gained significant attention due to its ability to induce interfacial polarization, modulate defect distribution, and enhance interfacial impedance, thereby markedly improving dielectric performance.^7,8 Despite these advances, the construction of multiple heterogeneous structures remains a formidable challenge due to the distinct physical and chemical properties of heterogeneous phases.

Gallium-based LMs have recently been recognized as ideal platforms for designing complex interface structures, due to their unique properties, such as fluidity, stretchability, high electrical conductivity (∼34 000 S cm⁻¹), low melting point, and excellent interfacial compatibility.^9–13 For instance, leveraging the high reactivity of LMs, it is possible to achieve a broad range of metal or oxide shell compositions through galvanic replacement reactions, while maintaining an unchanged core composition.^14,15 However, existing studies have predominantly focused on compositional control of the interface, often neglecting the influence of multi-heterogeneous interfacial distributions. Theoretical studies indicate that variations in multi-heterogeneous interfacial distributions can lead to different electron transport behaviors and polarization effects, which in turn significantly impact the dielectric properties of materials.^16–18 Therefore, constructing multiple heterogeneous interfaces with controlled phase distributions using gallium-based LMs and investigating their influence on dielectric properties is of crucial importance for optimizing EMW absorption performance.

Gallium in EGaIn alloy exhibits a relatively low negative standard reduction potential (Ga(iii)/Ga(0) redox couple, E₀ = −0.549 V vs. SHE), facilitating its interaction with other metals and oxides through galvanic replacement reactions, thereby enabling the modulation of interfacial phases.^19,20 Additionally, the cavitation and mechanical effects of ultrasonication can disrupt the oxide layer on the surface of gallium-based LMs, thereby promoting galvanic replacement and grain growth. The intensity of ultrasonication is crucial in determining the nature of the interfacial phases, with lower intensity favoring the formation of amorphous phases by preventing crystallization.^21–23 Moreover, thermal annealing can further drive elemental diffusion and the transition from amorphous to crystalline phases. The synergistic application of these techniques shows promise for achieving precise control over the multi-heterogeneous interfacial distributions of gallium-based LMs.

In this study, we synthesized EGaIn@Ni/NiO/Ga₂O₃ composites using a combination of ultrasound-assisted galvanic replacement reactions and high-temperature annealing. By varying the Ni salt content, we achieved an evolution in the content and distribution of amorphous Ni, NiO, and In at the EGaIn interface, resulting in a single-polarization phenomenon. Subsequent annealing induced a phase transition from amorphous to crystalline at the interface, altering the charge carrier mobility and triggering dual relaxation effects between EGaIn and Ni, as well as between NiO and Ga₂O₃. The substantial work function difference between Ni and NiO, coupled with the low barrier height between EGaIn and Ni facilitated electron migration from Ni to NiO, further enhancing the polarization intensity between NiO and Ga₂O₃. As a result, the GaInNi-2-700 composites (EGaIn alloy synthesized with 2 mmol Ni salt and annealed at 700 °C) exhibited exceptional EMW absorption performance, with an effective absorption bandwidth (EAB) of 7.92 GHz (10.08–18 GHz) at a matching thickness of 2.3 mm, significantly surpassing the 0 GHz EAB of untreated EGaIn alloy. This study not only provides an effective strategy for optimizing the EMW absorption performance of LMs but also reveals the crucial role of multi-heterogeneous interfacial distributions in regulating dielectric properties in gallium-based LMs.

Results

Evolution of multi-heterogeneous interfacial distributions for EGaIn alloys

As illustrated in Fig. 1a, the evolution of multi-heterogeneous interfacial distributions on the surface of EGaIn alloys encompasses two principal stages: ultrasound-assisted galvanic replacement and annealing. In the initial stage, ultrasonic-induced shear forces rapidly disrupt the Ga₂O₃ shells surrounding the EGaIn alloy, exposing the oxide-free EGaIn alloy. Subsequently, driven by the substantial surface tension, small droplets of the EGaIn alloy separate from the oxide-free surface. Given the lower negative standard reduction potential of the Ga(iii)/Ga(0) redox couple (E₀ = −0.549 V vs. SHE) relative to the In(iii)/In(0) redox couple (E₀ = −0.338 V vs. SHE), Ga is more prone to galvanic replacement when the EGaIn alloy is exposed to Ni salt solutions.⁹ Upon the addition of Ni salts, the reduction-generated Ni disperses randomly across the surface of the EGaIn alloy, with a portion of the Ni subsequently reacting with dissolved oxygen to form NiO. As the concentration of Ni salt increases, the surface content of both Ni and NiO correspondingly rises. Moreover, excessive oxidation of Ga results in the precipitation of amorphous In. Concurrently, ultrasonically fragmented Ga₂O₃, along with Ga₂O₃ formed by Ga oxidation through ultrasonic overflow, undergo dissolution and dehydration, forming GaOOH. These GaOOH preferentially adsorb onto the surface of the EGaIn alloy or onto the Ni surface when a significant portion of EGaIn is encapsulated by Ni (Fig. 1a). For clarity, samples obtained via ultrasound-assisted galvanic replacement with varying Ni salt concentrations are denoted as EGaInNi-x (x = 1, 2, 3, 4, and 5) (Table S1, ESI†).

Fig. 1 (a) Schematic representation of the multi-heterogeneous interfacial distribution evolution for EGaIn alloys; (b) schematic representation of interfacial compositions; (c) schematic diagrams of state transitions; (d) XRD patterns of EGaIn alloy samples after annealing treatment; (e) high-resolution spectra of Ga 2p for EGaIn alloy samples after annealing treatment; (f) SEM images of EGaInNi-1; (g) EDS spectrum images of EGaInNi-1.

During the subsequent annealing stage, calcination promotes the crystallization of Ni, NiO, and precipitated In particles. Concurrently, the elevated temperatures drive the dehydration of amorphous GaOOH, leading to the formation of Ga₂O₃. This Ga₂O₃ subsequently undergoes temperature-induced atomic rearrangements, transitioning into its crystalline phase (Fig. 1b and c). The annealed samples are correspondingly labeled as EGaInNi-x-700 (x = 1, 2, 3, 4, and 5) (Table S1, ESI†).

Characterization of multi-heterogeneous interfacial distributions for EGaIn alloys

To characterize the variations in the interfacial phase composition of EGaIn alloys, X-ray diffraction (XRD) analysis was employed. As illustrated in Fig. S1, ESI,† all samples prepared via ultrasound-assisted galvanic replacement reactions exhibited a characteristic peak of amorphous EGaIn alloys, devoid of any crystalline phases. This indicates that low-power ultrasound was insufficient to provide the necessary energy to induce crystallization of GaOOH, Ni, and NiO. Upon annealing, the XRD patterns revealed diffraction peaks corresponding to the (111) and (002) planes of Ga₂O₃, as well as the (101) plane of NiO (Fig. 1d). These findings indicate that the annealing process facilitated the dehydration and crystallization of GaOOH at the surface of the EGaIn alloy, along with atomic rearrangement of amorphous NiO into its crystalline form. Additionally, XRD patterns of the EGaInNi-4-700 and EGaInNi-5-700 samples detected peaks corresponding to In. This observation is likely due to the excessive addition of Ni salt, which may have led to the depletion of Ga and the subsequent precipitation of In. Notably, the annealing process induced a transformation of In from an amorphous to a crystalline state.

X-ray photoelectron spectroscopy (XPS) was employed to elucidate the surface elemental states of the samples. As shown in Fig. 1e and Fig. S2, ESI,† the high-resolution Ga 2p XPS spectra reveal peaks corresponding to Ga⁰ and Ga³⁺, indicating the presence of EGaIn and GaOOH following ultrasound-assisted galvanic replacement reactions.^24,25 Additionally, the characteristic peaks at 444.36 eV and 451.94 eV confirmed the presence of In in the form of In⁰, verifying the preferential oxidation of Ga (Fig. S2 and S3, ESI†).^26,27 The high-resolution O 1s spectra exhibited three characteristic peaks at 532.58 eV, 531.48 eV, and 530.12 eV, corresponding to surface adsorbed water, oxygen vacancies, and lattice oxygen, respectively (Fig. S2 and Tables S2, S3, ESI†).^28,29 Notably, a shift of approximately 1 eV towards higher binding energy in the lattice oxygen peaks was observed after annealing, proving the occurrence of the transition from amorphous to crystalline state during the annealing process (Fig. S2 and Tables S2, S3, ESI†). Post-annealing treatment also led to a significant reduction in the percentage of surface adsorbed water peaks, from 32% to 8%, further confirming the dehydration of GaOOH. Moreover, the oxygen vacancy concentration remained consistent across all samples, indicating that defect polarization due to oxygen vacancies is not the primary cause of EM losses (Table S2, ESI†).^30,31

Scanning electron microscopy (SEM) was employed to further elucidate the morphology of EGaIn alloys. As depicted in Fig. 1f and Fig. S4, ESI,† the EGaIn alloys subjected to ultrasound-assisted galvanic replacement reactions exhibit a typical spherical structure with a diameter of approximately 2 μm. Energy dispersive spectroscopy (EDS) mapping provided additional insights into the elemental composition of the interfacial phases. In the EDS mapping (Fig. 1g and Fig. S5–S13, ESI†), Ga, In, Ni, and O elements are clearly detected in the selected regions, with Ga and In concentrations being more pronounced in the central region, confirming that the core is composed of EGaIn alloys. Additionally, the distribution of Ga and O elements is slightly more extensive than that of Ni elements, suggesting that the shells of EGaIn alloy are mainly composed of GaOOH or Ga₂O₃ (Fig. S5–S13, ESI†).

To elucidate the evolution of the multi-heterogeneous interfacial distributions in EGaIn alloys, transmission electron microscopy (TEM) analysis was conducted on all samples (Fig. 2 and Fig. S14–S20, ESI†). As shown in Fig. 2b, the EGaIn alloy samples subjected to ultrasound-assisted galvanic replacement exhibit a core–shell structure, with a shell thickness of approximately 50 nm. High-resolution TEM (HR-TEM) images of these samples reveal no discernible lattice fringes, indicating an amorphous shell composition, consistent with the XRD results (Fig. S14 and S15, ESI†). Post-annealing, the shell region displays lattice fringes corresponding to Ga₂O₃ (111), (002), and (−311) crystal planes, as well as Ni (002) and NiO (200) crystal planes, confirming the coexistence of Ni, NiO, and Ga₂O₃ phases within the shell (Fig. S17–S20, ESI†). Additionally, as depicted in Fig. 2(c2), both Ga₂O₃ (400) and Ni (002) lattice fringes are observed, with Ga₂O₃ lattice fringes distributed throughout the region, indicating that Ga₂O₃ is situated on the outer layer of Ni. As the Ni content increases, there is a corresponding increase in NiO lattice fringes and a decrease in Ni lattice fringes, suggesting that at higher Ni salt concentrations, Ni reacts with oxygen in the solution under ultrasonic conditions to form more NiO (Fig. S17–S20, ESI†). Furthermore, post-annealing samples show only low-density defects and minimal lattice distortion in the regions corresponding to Ni, NiO, and Ga₂O₃, indicating high lattice integrity (Fig. S17–S20, ESI†). This low defect density and high lattice integrity facilitate carrier transport within the grains, thereby enhancing the material's response to electromagnetic fields. To further corroborate the low defect density within the grains of the annealed samples, geometric phase analysis (GPA) was performed based on the HR-TEM images (the observed color transitions from white to orange and from white to blue are associated with tensile and compressive strain, respectively). As depicted in Fig. 2e–h, no significant strain variation along the ε_xx direction was observed within the grains of Ga₂O₃ and Ni, providing further evidence of the low defect density in the annealed samples.³²

Fig. 2 (a) Schematic representation of the structural evolution of EGaIn alloy samples subjected to ultrasound-assisted galvanic replacement reactions followed by high-temperature annealing; (b) TEM and corresponding selective electron diffraction images for EGaInNi-2; (c) HR-TEM images of EGaInNi-2-700; (d) HR-TEM images of EGaInNi-3-700; (e) HR-TEM image of the Ga₂O₃ crystal phase for EGaInNi-2-700; (f) GPA analysis for Ga₂O₃ based on the HR-TEM image of EGaInNi-2-700; (g) HR-TEM image of the Ni crystal phase for EGaInNi-2-700; (h) GPA analysis for Ni based on the HR-TEM image of EGaInNi-2-700.

Modulation of EMW absorption properties and EM parameters through multi-heterogeneous interfacial distributions in EGaIn alloys

The EMW absorption properties of EGaIn alloys were systematically evaluated to assess the impact of multi-heterogeneous interfacial distributions. As displayed in Fig. S21, ESI,† the pristine EGaIn alloy exhibited poor EMW absorption, with an EAB of 0 GHz, likely due to the insulating effect of the Ga₂O₃ layer, which impedes electron transport. Upon galvanic replacement reactions, the EMW absorption properties initially improve and subsequently decline (Fig. 3a–d and Fig. S22, ESI†). This trend can be attributed to the introduction of small amounts of Ni salts, which create scattered Ni distributions on the EGaIn alloy surface, acting as charge transport bridges and inducing interfacial polarization between EGaIn and Ni.³³ However, as the Ni salt content further increases, Ni becomes more uniformly distributed on the EGaIn surface and gradually reacts with dissolved oxygen under ultrasonic conditions to form NiO, which impedes electron transfer and diminishes EMW absorption performance. Post-annealing, the samples display a similar trend in EMW absorption performance with varying Ni content, initially increasing and then decreasing (Fig. 3e–h and Fig. S22, ESI†). This behavior is likely associated with the crystallization of Ni, NiO, and Ga₂O₃, which significantly enhances carrier transport while providing additional regions for interfacial charge accumulation. Notably, the EGaInNi-2-700 sample demonstrated the optimum EMW absorption performance, achieving an EAB of 7.92 GHz, significantly surpassing the highest values reported for LM based composites in the literature (Fig. 3i and j).^{14,15,34–39} To further evaluate the practical application potential of the samples, computer simulation technology (CST) simulations were conducted. As shown in Fig. 3k and Fig. S23, ESI,† compared to the perfect electric conductor (PEC) samples, the EGaInNi-2-700 and EGaInNi-1-700 samples can effectively attenuate EMWs within the range of ±90°, theoretically demonstrating their potential for real-world applications.

Fig. 3 (a)–(d) 2D RL values of EGaInNi-1 to EGaInNi-4 and (e)–(h) EGaInNi-1-700 to EGaInNi-4-700; (i) histogram distribution of RLmin and EAB of EGaIn-based composites; (j) comparison of EAB and thickness of EGaInNi-2-700 composites with LM-based EMW absorbing materials; (k) RCS date of PEC and EGaInNI-1-700 and EGaInNi-2-700.

The mapping of electromagnetic parameters (EMPs) for all samples further elucidates the impact of multi-heterogeneous interfacial distributions (Fig. 4 and Fig. S22, S24, ESI†). As displayed in Fig. S21, ESI,† the pristine EGaIn alloy exhibits a real and imaginary part of the complex permittivity being 3.5 and 0, respectively, indicating negligible dielectric loss. This behavior is likely due to the insulating effect of the spontaneously formed Ga₂O₃ layer on the alloy surface. Following low-intensity ultrasound-assisted galvanic replacement, all samples except EGaInNi-2 displayed distinct single relaxation peaks, with their intensity first increasing and then decreasing as the Ni substitution amount increased (Fig. 4a, c and Fig. S24, ESI†). At trace levels of Ni salt, Ni is sparsely distributed across the EGaIn surface, forming localized inhomogeneous charge regions that induce interfacial polarization with EGaIn. However, the low Ni content limits the formation of a strong interfacial polarization effect, as illustrated in Fig. 4a. With a moderate increase in Ni salt concentration, a homogeneous and stable layer arrangement gradually forms on the surface of EGaIn alloys, resulting in a progressive enhancement in the intensity of the relaxation peak. As the Ni salt content further increases, amorphous NiO phases form on the uniformly distributed Ni, impeding electron transport and suppressing interfacial polarization, as shown in Fig. S24, ESI† (ref. 40–42). Notably, when the Ni salt concentration reaches 5 mmol, despite the EGaIn alloy surface being almost entirely encapsulated by the electrically insulating amorphous GaOOH, the precipitation of In particles induces an asymmetric charge distribution between In and EGaIn, resulting in a weak polarization effect (Fig. S24, ESI†).

Fig. 4 (a)–(d) The EM parameters of EGaInNi-1, EGaInNi-1-700, EGaInNi-2 and EGaInNi-2-700; (e)–(h) Cole–Cole plots of EGaInNi-1, EGaInNi-1-700, EGaInNi-2 and EGaInNi-2-700; (i) schematic diagram of the interface structure transformation before and after annealing; (j) EMW absorption performance of EGaInNi-2; (k) EGaInNi-2-700; (l) HR-TEM images of EGaInNi-2-700; (m) off-axis electron holography images of EGaInNi-2-700; (n) charge density distribution map of EGaInNi-2-700.

Following annealing, distinct dual relaxation peaks were observed in both EGaInNi-1-700 and EGaInNi-2-700 samples. This phenomenon is primarily attributed to the transition of Ni, NiO, and GaOOH from amorphous to crystalline phases, which significantly enhances carrier transport capabilities. Consequently, this transition induces asymmetric charge distribution at the EGaIn/Ni and NiO/Ga₂O₃ interfaces, as illustrated in Fig. 4b and d. With further increases in Ni salt content, the insulating characteristics of NiO and the outermost Ga₂O₃ layer began to dominate, resulting in the complex permittivity returning to levels like those of the initial EGaIn alloy (Fig. S24, ESI†). In the EGaInNi-5-700 samples, the increased precipitation of In particles promoted asymmetric charge distribution between In and EGaIn, giving rise to pronounced polarization effects. Notably, the relaxation behavior in the EGaInNi-5-700 sample was more pronounced than in the unannealed EGaInNi-5, suggesting that the enhanced carrier transport further amplified the polarization phenomenon (Fig. S24, ESI†).

To further elucidate the impact of multi-heterogeneous interfacial distributions on the dielectric polarization, Cole–Cole plots of the samples were drawn based on Debye relaxation theory (Fig. 4e–h and Fig. S25, ESI†). As depicted in the Fig. 4e, g and Fig. S25, ESI,† the Cole–Cole plots of samples subjected to ultrasound-assisted galvanic replacement reactions, except for EGaInNi-2, exhibit a distinct semicircle, indicative of a single relaxation process, which is consistent with the EMP results.^43–45 Following annealing treatment, the Cole–Cole plots of both EGaInNi-1-700 and EGaInNi-2-700 reveal the emergence of two distinct semicircles, signifying a transition to dual dielectric polarization processes (Fig. 4f and h). This shift from single to dual relaxation is likely attributable to the crystallization of Ni, NiO, and Ga₂O₃ from their amorphous states (Fig. 4i). This structural transformation enhances carrier transport within and between grains, facilitating polarization at the EGaIn/Ni and NiO/Ga₂O₃ interfaces. The asymmetric charge distribution at these interfaces may form over different timescales due to the varying electron-binding capacities of EGaIn, Ni, NiO, and Ga₂O₃. This temporal variation allows the material to respond to EMWs across different frequency bands, thereby broadening the frequency range over which EMW attenuation occurs. As illustrated in Fig. 4j and k, the EGaInNi-2 sample, which exhibits only single polarization behavior, has an EAB of 6.56 GHz. In contrast, the annealed EGaInNi-2-700 sample demonstrates dual relaxation behavior, with an EAB increased to 7.92 GHz, approximately 1.21 times that of the original.

To elucidate the polarization phenomena at the EGaIn/Ni and NiO/Ga₂O₃ interfaces, off-axis electron holography was employed to visualize the charge distribution at the EGaInNi-2-700 interface. As depicted in Fig. 4l and m, distinct regions of charge accumulation were identified at the EGaIn/Ni and NiO/Ga₂O₃ interfaces, respectively. This observation indicates the presence of dual relaxation regions at the EGaInNi-2-700 interface, consistent with the analysis of EMP and Cole–Cole plots. To further quantify the polarization intensity, a charge distribution map was generated. As shown in Fig. 4n, the charge density at the NiO/Ga₂O₃ interface reaches a peak value of 0.060 e nm⁻³, approximately twice that observed at the EGaIn/Ni interface. This phenomenon may be attributed to the substantial work function difference between Ni and NiO, combined with the lower barrier height between EGaIn and Ni, facilitating more efficient electron migration from Ni to the NiO/Ga₂O₃ interface.

Modulation mechanism of multi-heterogeneous interfacial distributions on EGaIn alloys

Based on the above analysis, the modulation mechanisms of dielectric properties and the corresponding EMW absorption performance of the EGaIn alloy by multi-heterogeneous interfacial distributions are summarized as follows (Fig. 5).

Fig. 5 (a) The EIS spectra of annealed samples; (b) work functions of amorphous EGaIn, Ga₂O₃, and NiO; (c) schematic diagram of the regulation mechanism of multi-heterogeneous interfacial distributions on the EMW absorption performance of EGaIn alloy.

Samples obtained through ultrasound-assisted galvanic replacement, except for EGaInNi-2, exhibit single polarization behavior, whereas EGaInNi-1-700 and EGaInNi-2-700 demonstrate dual relaxation behavior. According to dielectric loss theory, polarization behavior in the microwave frequency range originates from dipole polarization, defect-induced polarization, and interfacial polarization. The absence of significant functional groups on the surface of EGaIn alloys suggests that dipole polarization is not the cause of the observed polarization peaks. Additionally, the lack of notable changes in oxygen vacancy concentrations and defect contents, as evidenced by XPS analysis, indicates that defect-induced polarization is not the primary mechanism behind the dual relaxation (Fig. S2 and Tables S2, S3, ESI†). Therefore, the observed bipolar behavior in the samples is primarily attributed to interfacial polarization.

In the ultrasonically treated samples, the interfacial phases primarily consist of amorphous Ni, NiO, and GaOOH. The presence of the amorphous phase typically increases electron scattering during transport, leading to a reduction in carrier mobility (Fig. S26, ESI†).^46,47 In the EGaInNi-1 sample, characterized by a low Ni salt content, both EGaIn/Ni and NiO/GaOOH interfaces are observed, but the polarization behavior is primarily dictated by the single polarization phenomenon at the EGaIn/Ni interface (Fig. 4a and e). This is mainly attributed to the suppression of carrier transport by the amorphous phase and the low NiO content, which hampers the formation of significant charge asymmetry at the NiO/GaOOH interface, thereby limiting the manifestation of interfacial polarization effects. Furthermore, both EGaIn and Ni are amorphous metals, resulting in good interfacial compatibility and a relatively low barrier height, facilitating carrier transport across the interface under an external electric field. Consequently, the EGaIn/Ni interface exhibits a weaker polarization effect. According to Debye relaxation theory, the position of the relaxation peak is inversely proportional to the relaxation time.⁴⁸ Due to the low barrier and high carrier mobility at the EGaIn/Ni interface, this interface is more sensitive to external electromagnetic fields, leading to a relaxation peak near 13.1 GHz (Fig. 4a and e). Upon annealing, NiO and GaOOH gradually transition to crystalline phases, enhancing carrier mobility within NiO and Ga₂O₃, which in turn leads to significant charge accumulation at the NiO/Ga₂O₃ interface (Fig. 5a). However, this interface, characterized by a higher barrier and relatively weaker carrier transport, responds more slowly to external electromagnetic fields, resulting in a relaxation peak at the lower frequency of 7.8 GHz (Fig. 4b and f).

As the Ni salt content increases, Ni gradually encapsulates the EGaIn alloy, and the NiO content also rises. Consequently, in the EGaInNi-2 sample, significant charge accumulation begins to appear at the NiO/Ga₂O₃ interface. However, the low barrier at the EGaIn/Ni interface leads to the large reduction in charge accumulation at the EGaIn/Ni interface and the corresponding high-frequency relaxation peak declines (Fig. 4c and g). Given the work functions of amorphous EGaIn (6.12 eV) and crystalline Ni, NiO, and Ga₂O₃ (4.41 eV, 5.39 eV, 5.09 eV, respectively), electrons are more prone to accumulate in EGaIn and NiO, resulting in more pronounced dual charge accumulation zones at the EGaIn/Ni and NiO/Ga₂O₃ interfaces (Fig. 5b, c and Fig. S27, ESI†). Notably, the large work function difference between Ni and NiO, along with the low barrier at the EGaIn/Ni interface, facilitates electron migration from Ni to NiO, resulting in a more pronounced charge accumulation at the NiO/Ga₂O₃ interface (Fig. 4l–n). The differential polarization dynamics at the EGaIn/Ni and NiO/Ga₂O₃ interfaces enable these dual charge accumulation zones to respond to distinct frequency ranges of EMW, thereby inducing a dual relaxation phenomenon (Fig. 5c). This dual relaxation behavior contributes to the EGaInNi-2-700 sample achieving an EAB of 7.92 GHz, significantly surpassing the optimum performances of LM-based absorbers reported in the literature.

In summary, the pure EGaIn alloy exhibits a low complex permittivity, making it ineffective for significant electromagnetic loss. During the multi-heterogeneous interfacial distribution modulation of EGaIn alloy, it primarily acts as a template that drives the evolution of multi-heterogeneous interfacial distributions, thereby inducing dual relaxation behavior and facilitating EMW attenuation. While the EGaIn alloy itself contributes minimally to EMW absorption, it serves as an auxiliary by promoting the formation of other phases, leading to enhanced electromagnetic loss. This study presents a novel approach to tuning the EMW absorption properties of materials, leveraging EGaIn alloy to drive microstructural design and optimize EMW absorption performance (Fig. 5(c3)).

Conclusions

In this study, EGaIn alloys with diverse multi-heterogeneous interfacial distributions were synthesized using ultrasound-assisted galvanic replacement reactions followed by annealing processes. The application of low-intensity ultrasound induced an amorphous structure at the interface of the EGaIn alloy, which inhibited electron migration, leading to the emergence of single polarization behavior. Post-annealing, the interfacial phase transitioned from amorphous to crystalline, significantly enhancing carrier mobility. Furthermore, the introduction of an appropriate amount of Ni salt resulted in the formation of abundant EGaIn/Ni and NiO/Ga₂O₃ interfacial distributions, which induced dual relaxation behavior. The pronounced work function difference between Ni and NiO, along with the low barrier height between EGaIn and Ni, promoted electron migration from Ni to NiO, thereby further amplifying the polarization intensity between NiO and Ga₂O₃. This dual relaxation behavior markedly improved the EMW attenuation performance, enabling EGaInNi-2-700 to achieve an optimal EAB of 7.92 GHz at a thickness of 2.3 mm, substantially surpassing the initial EGaIn alloy's 0 GHz. This work not only elucidates the mechanisms by which multi-heterogeneous interfacial distribution evolution regulates the dielectric properties of EGaIn alloys but also proposes an effective strategy for designing high-performance LM-based EMW absorption materials.

Experimental

Raw materials

Nickel nitrate hexahydrate (Ni(NO₃)₂·6H₂O), anhydrous ethanol, and deionized water were purchased from Aladdin Industrial Corporation (Shanghai, China). Gallium (Ga 99.99%) and indium (In 99.99%) were purchased from Macklin (Shanghai, China). All chemicals were of commercial grade and were utilized as received, without any additional purification.

Synthesis of eutectic gallium indium (EGaIn) liquid metals

Initially, Ga and In were weighed in a mass ratio of 75.5% to 24.5% and placed in a beaker. The beaker was then positioned on a hot plate, where it was heated to 160 °C while being stirred for 0.5 h, followed by natural cooling to room temperature. Upon reaching room temperature, the resulting material displayed a shiny silver metallic appearance in its liquid state.

Synthesis of EGaInNi-1 and EGaInNi-1-700 materials

Initially, 1–5 mmol of Ni(NO₃)₂·6H₂O was weighed and dissolved in a beaker containing 100 mL of deionized water to form a solution. Subsequently, 1 g EGaIn alloy was weighed and added to the prepared solution. The resulting mixture was subjected to ultrasonication in a water bath for 30 minutes, followed by stirring with a magnetic stirrer for 1 h. The reaction products were then alternately washed three times with deionized water and anhydrous ethanol. Finally, the obtained products were dried in a vacuum oven at 60 °C for 24 h to yield the final products, which were designated as EGaInNi-1 through EGaInNi-5 based on the concentration of Ni(NO₃)₂·6H₂O.

For the preparation of EGaInNi-700-x (where x = 1, 2, 3, 4, and 5), the synthesized products were placed in a tube furnace and heated to 700 °C at a rate of 2 °C min⁻¹ under an Ar atmosphere, maintaining this temperature for 2 h. After natural cooling to room temperature, the EGaInNi-700-x products were obtained and designated according to the amount of Ni salt added, labeled as EGaInNi-700-1 through EGaInNi-700-5.

Author contributions

G. Chen, writing – original draft, conceptualization, and visualization; T. Zhang, writing – original draft, visualization; L. M. Zhang, K. Tao, and Q. Chen, conceptualization, and supervision; H. J. Wu, conceptualization, supervision, writing – review & editing.

Data availability

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

Conflicts of interest

There are no conflicts to declare.

Acknowledgements

Financial support was provided by the National Science Foundation of China (Grants no. 51872238, 52074227, and 21806129), and the Innovation Funds for Graduate Students of Northwestern Polytechnical University (CX2023094). Thanks Prof. Renchao Che for his support with the off-axis electron holography images.

References

1. J. H. Ye and J. C. Tan, Nano Energy, 2023, 114, 108687.
2. K. H. Kim, I. Karpov, R. H. Olsson III and D. Jariwala, Nat. Nanotechnol., 2023, 18, 422–441.
3. B. B. Yang, Q. H. Zhang, H. B. Huang, H. Pan, W. X. Zhu, F. Q. Meng, S. Lan, Y. Q. Liu, B. Wei, Y. Q. Liu, L. T. Yang, L. Gu, L. Q. Chen, C. W. Nan and Y. H. Lin, Nat. Energy, 2023, 8, 956–964.
4. F. Liu and Z. X. Fan, Chem. Soc. Rev., 2023, 52, 1723–1772.
5. Z. H. Zhao, Y. C. Qing, L. Kong, H. L. Xu, X. M. Fan, J. J. Yun, L. M. Zhang and H. J. Wu, Adv. Mater., 2024, 36, 2304182.
6. G. Chen, Z. J. Li, L. M. Zhang, Q. Chang, X. J. Chen, X. M. Fan, Q. Chen and H. J. Wu, Cell Rep. Phys. Sci., 2024, 5, 102097.
7. S. Sovizi, S. Angizi, S. Alem, R. Goodarzi, M. Boyuk, H. Ghanbari, R. Szoszkiewicz, A. Simchi and P. Kruse, Chem. Rev., 2023, 123, 13869–13951.
8. E. Gabriel, C. Ma, K. Graff, A. Conrado, D. Hou and H. Xiong, eScience, 2023, 3, 100139.
9. L. Ren, N. Y. Cheng, X. K. Man, D. C. Qi, Y. D. Liu, G. B. Xu, D. D. Cui, N. N. Liu, J. X. Zhong, G. Peleckis, X. Xu, S. X. Dou and Y. Du, Adv. Mater., 2021, 33, 2008024.
10. M. K. Zhang, P. J. Zhang, Q. Wang, L. Li, S. J. Dong, J. Liu and W. Rao, J. Mater. Chem. C, 2019, 7, 10331–10337.
11. D. H. Yu, Y. Liao, Y. C. Song, S. L. Wang, H. Y. Wan, Y. H. Zeng, T. Yin, W. H. Yang and Z. Z. He, Adv. Sci., 2020, 7, 2000177.
12. J. Y. Gao, J. Ye, S. Chen, J. H. Gong, Q. Wang and J. Liu, ACS Appl. Mater. Interfaces, 2021, 13, 17093–17103.
13. M. K. Zhang, P. J. Zhang, C. L. Zhang, Y. S. Wang, H. Chang and W. Rao, Appl. Mater. Today., 2020, 19, 100612.
14. B. Zhao, Y. Q. Du, H. L. Lv, Z. K. Yan, H. Jian, G. Y. Chen, Y. Y. Wu, B. B. Fan, J. C. Zhang, L. M. Wu, D. W. Zhang and R. C. Che, Adv. Funct. Mater., 2023, 33, 2302172.
15. B. Zhao, Z. K. Yan, L. L. Liu, Y. Y. Zhang, L. Guan, X. Q. Guo, R. S. Li, R. C. Che and R. Zhang, Adv. Funct. Mater., 2024, 34, 2314008.
16. Z. G. Gao, Z. H. Ma, D. Lan, Z. H. Zhao, L. M. Zhang, H. J. Wu and Y.-L. Hou, Adv. Funct. Mater., 2022, 32, 2112294.
17. S. C. Hui, X. Zhou, L. M. Zhang and H. J. Wu, Adv. Sci., 2024, 11, 2307649.
18. M. H. Li, W. J. Zhu, X. Li, H. L. Xu, X. M. Fan, H. J. Wu, F. Ye, J. M. Xue, X. Q. Li, L. F. Cheng and L. T. Zhang, Adv. Sci., 2022, 9, 2201118.
19. Y. Y. Zhang, Q. M. He, H. Yang, Z. S. Li, H. Jiang, Y. Zhang, X. Luo and Y. Zheng, Nano Lett., 2024, 24, 6247–6254.
20. S. S. Zhao, J. Q. Zhang and L. Fu, Adv. Mater., 2021, 33, 2005544.
21. Y. Q. Guo, X. G. Yang, G. Li, J. Yang, L. Liu, L. M. Chen and B. Li, Chem. Eng. J., 2021, 405, 126914.
22. G. Chen, H. S. Liang, J. J. Yun, L. M. Zhang, H. J. Wu and J. Y. Wang, Adv. Mater., 2023, 35, 2305586.
23. Y. Y. Zhao, F. Qian, W. F. Shen, C. L. Zhao, J. G. Wang, C. X. Xie, F. L. Zhou, C. T. Chang and Y. J. Li, J. Mater. Sci. Technol., 2021, 82, 144–152.
24. Y. F. Wang, M. Mayyas, J. Yang, J. B. Tang, M. B. Ghasemian, J. L. Han, A. Elbourne, T. Daeneke, R. B. Kaner and K. Kalantar-Zadeh, Adv. Funct. Mater., 2021, 31, 2005866.
25. W. Zhang, J. Z. Ou, S. Y. Tang, V. Sivan, D. D. Yao, K. Latham, K. Khoshmanesh, A. Mitchell, A. P. O'Mullane and K. Kalantar-zadeh, Adv. Funct. Mater., 2014, 24, 3799–3807.
26. M. G. Saborio, S. X. Cai, J. B. Tang, M. B. Ghasemian, M. Mayyas, J. L. Han, M. J. Christoe, S. H. Peng, P. Koshy, D. Esrafilzadeh, R. Jalili, C. H. Wang and K. Kalantar-Zadeh, Small, 2020, 16, 1903753.
27. M. Rahim, F. Centurion, J. L. Han, R. Abbasi, M. Mayyas, J. Sun, M. J. Christoe, D. Esrafilzadeh, F. M. Allioux, M. B. Ghasemian, J. Yang, J. B. Tang, T. Daeneke, S. Mettu, J. Zhang, M. H. Uddin, R. Jalili and K. Kalantar-Zadeh, Adv. Funct. Mater., 2021, 31, 2007336.
28. Y. R. Han, Y. F. Wang, S. H. Fu, J. G. Ma, H. Y. Xu, B. S. Li and Y. C. Liu, Small, 2023, 19, 2206664.
29. M. Qin, L. M. Zhang, X. R. Zhao and H.-J. Wu, Adv. Sci., 2021, 8, 2004640.
30. J. L. Liu, L. M. Zhang and H.-J. Wu, Adv. Funct. Mater., 2022, 32, 2200544.
31. Y. F. Yan, K. L. Zhang, G. Y. Qin, B. S. Gao, T. Zhang, X. X. Huang and Y. Zhou, Adv. Funct. Mater., 2024, 34, 2316338.
32. B. Zhao, Z. K. Zhao, Y. Q. Du, L. J. Rao, G. Y. Chen, Y. Y. Wu, L. T. Yang, J. C. Zhang, L. M. Wu, D. W. Zhang and R. C. Che, Adv. Mater., 2023, 35, 2210243.
33. Z. H. Zhao, K. C. Kou, L. M. Zhang and H. J. Wu, Carbon, 2022, 186, 323–332.
34. Y. Wang, Y. N. Gao, T. N. Yue, X. D. Chen, R. C. Che and M. Wang, J. Colloids Interface Sci., 2022, 607, 210–218.
35. K. Y. Zhao, C. L. Luo, C. Sun, M. L. Huang and M. Wang, Composites, Part A, 2023, 173, 107640.
36. L. C. Wang, L. Huang, Y. B. Li and Y. Yuan, J. Appl. Phys., 2022, 132, 194101.
37. X. F. Zhang, H. Jiang, L. L. Xu, K. Zuraiqi, T. Daeneke, J. T. Zhou, G. K. Li and A. Zavabeti, Ceram. Int., 2022, 48, 10066–10078.
38. F. Pan, L. Cai, Y. Y. Shi, Y. Y. Dong, X. J. Zhu, J. Cheng, H. J. Jiang, X. Wang, S. Zhong and W. Lu, Chem. Eng. J., 2022, 435, 135009.
39. X. F. Zhang, L. L. Xu, J. T. Zhou, W. J. Zheng, H. Jiang, K. Zuraiqi, G. K. Li, J. Liu and A. Zavabeti, ACS Appl. Nano Mater., 2021, 4, 9200–9212.
40. M. Qin, L. M. Zhang, X. R. Zhao and H.-J. Wu, Adv. Funct. Mater., 2021, 31, 2103436.
41. J. X. Xiao, B. B. Zhan, M. K. He, X. S. Qi, X. Gong, J. L. Yang, Y. P. Qu, J. F. Ding, W. Zhong and J. W. Gu, Adv. Funct. Mater., 2024 DOI:10.1002/adfm.202316722.
42. M. M. Liu, Z. C. Wu, L. T. Yang, X. W. Lv, R. X. Zhang, H. B. Zhang, X. Y. Wang, G. Y. Chen, J. C. Zhang, H. L. Lv and R. C. Che, Adv. Funct. Mater., 2023, 33, 2307943.
43. T. Gao, R. Z. Zhao, Y. X. Li, Z. Y. Zhu, C. L. Hu, L. Z. Ji, J. Zhang and X. F. Zhang, Adv. Funct. Mater., 2022, 32, 2204370.
44. H. S. Liang, S. C. Hui, G. Chen, H. Shen, J. J. Yun, L. M. Zhang, W. Lu and H. J. Wu, Small Methods, 2024, 8, 2301600.
45. H. S. Liang, G. Chen, D. Liu, Z. J. Li, S. C. Hui, J. J. Yun, L. M. Zhang and H. J. Wu, Adv. Funct. Mater., 2023, 33, 2212604.
46. H. P. Lv, C. Wu, J. Tang, H. F. Du, F. X. Qin, H. X. Peng and M. Yan, Chem. Eng. J., 2021, 411, 128445.
47. Z. G. Gao, D. Lan, L. M. Zhang and H. J. Wu, Adv. Funct. Mater., 2021, 31, 2106677.
48. S. L. Patil, R. K. Sun, S. N. Cheng and S. W. Cheng, Phys. Rev. Lett., 2023, 130, 098201.

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Footnotes

† Electronic supplementary information (ESI) available. See DOI: https://doi.org/10.1039/d4mh01564a

‡ G. Chen and T. Zhang contributed equally to this work.

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