Rational Design of a Metallic Functional Layer for High-Performance

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Rational Design of a Metallic Functional Layer for High-Performance Solid Oxide Fuel Cells Mingi Choi, Sangyeon Hwang, Seo Ju Kim, Jongseo Lee, Doyoung Byun, and Wonyoung Lee ACS Appl. Energy Mater., Just Accepted Manuscript • DOI: 10.1021/acsaem.9b00151 • Publication Date (Web): 06 May 2019 Downloaded from http://pubs.acs.org on May 7, 2019

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ACS Applied Energy Materials

Rational Design of a Metallic Functional Layer for High-Performance Solid Oxide Fuel Cells Mingi Choi§, Sangyeon Hwang§, Seo Ju Kim, Jongseo Lee, Doyoung Byun*, and Wonyoung Lee* Department of Mechanical Engineering, Sungkyunkwan University, 2066 Seobu-ro, Jangangu, Suwon-si, Kyunggi-do 16419, South Korea *Corresponding authors. E-mail: [email protected], [email protected] §These

authors contributed equally.

Abstract The rational design of the electrode–electrolyte interface plays a crucial role in expediting the oxygen reduction reaction (ORR) kinetics of intermediate temperature solid oxide fuel cells (IT–SOFCs). We employed metallic functional layers because of their high electrical conductivities

and

catalytic

activities

with

respect

to

ORR

kinetics.

Using

electrohydrodynamic (EHD) jet printing, we printed the metallic grid structure at the interface of Sm0.5Sr0.5CoO3-δ (SSC) and Gd0.1Ce0.9O2-δ (GDC) with Al, Ni, and Ag to systematically quantify the effects of the electrical conductivity and catalytic activity on ORR kinetics. Substantial improvements in interfacial properties were achieved with the metallic functional layers, manifested by reducing the polarization resistance to 12.5% of the bare SSC cathode. I–V characterization, electrochemical impedance spectroscopy (EIS) measurements, and distributed relaxation times (DRT) based on impedance fitting enabled the quantitative deconvolution and revealed that the enhanced electrical conductivity of the metallic functional layer was primarily responsible for the increased electrochemical performance compared to the enhanced catalytic activity. The SSC cathode with the Ag functional layer exhibited the highest peak power density of ~670 mW/cm2 at 650 °C, which was higher than that of the bare SSC cathode by ~1.8 times. Keywords: solid oxide fuel cell, electrohydrodynamic jet printing, metallic functional layer, interfacial properties, sheet resistance, charge distribution, high performance

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1. Introduction Retarded oxygen reduction reaction (ORR) kinetics at the cathode is one of the major challenges in the development of high-performance solid oxide fuel cells (SOFCs), particularly with regard to operation in the intermediate temperature (IT, 600–700 °С) and the lowtemperature regimes (LT, < 600 °С).1 In both temperature regimes, ORR kinetics have been reported to be the most active at the electrode/electrolyte interfaces, or at the triple-phase boundaries (TPBs), which are characterized by the co-presence of oxygen gas, electrode, and electrolyte.2 At such interfaces, the ionic conductivities, electronic conductivities, and catalytic activities of the electrode and electrolyte materials directly affect the ORR kinetics because oxygen ions and electrons are required at the TPBs in the ORRs.2 Thus, tailoring the interfacial properties to facilitate ORR kinetics is crucial to the enhancement of electrode performance, and also to the development of SOFCs that operate at reduced temperatures.2–4 Given the importance of the interfacial properties, a number of approaches have been investigated to develop the functional layer at the interfaces between the electrode and the electrolyte to expedite the reaction rates and to extend the reaction sites for ORR.3,5–11 In addition, the use of composite materials that capitalize on the synergetic advantages of the various constituent material properties of the cathode functional layer constitutes the basis of most of these extensively employed methods.7,12,13 A non-uniform electron distribution near the electrode/electrolyte interfaces due to the relatively low-electrical conductivities of the electrode materials and the insufficient interfacial contacts can also significantly impede the charge transfer kinetics.14,15 Fleig et al. reported the detrimental effect of the uneven electron distribution at the interfaces owing to the discrete contact of the porous electrode on the electrolyte that limited the reaction sites for ORRs and caused increases in ohmic resistance.14,15 This current constriction effect led to an increased charge transport resistance at high frequencies, and to the surface exchange resistance in medium frequencies owing to the insufficient electron pathway at the interface.16 Bertei et al. conducted 3–D simulations to verify the current constriction effect at the interfaces, and demonstrated the performance enhancement with a sufficient electron supply.17 Recently, the substantial performance enhancement due to the uniform electron distribution of the functional layers at the interfaces had been demonstrated, which facilitated a sufficient electron distribution at the interface, and resulted in the reductions of the polarization and ohmic resistances.16,18–20 Choi et al. reported a performance enhancement of ~16% at 650 °С (~1.4 W


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cm-2) with a densely coated PrBa0.5Sr0.5Co1.5Fe0.5O5+δ (PBSCF) layer at the interfaces between the BaZr0.4Ce0.4Y0.1Yb0.1O3 electrolyte and the porous PBSCF cathode layer compared to nondeposited cells.18 Jang et al. showed a ~63% improved peak power density at 550 °С (0.57 W/cm2) compared to non-coated cells with the use of an LSCF nanoweb structure at the interfaces between LSCF and GDC.20 In this regard, the metallic functional layer could be a more promising candidate because of its higher electrical conductivity for extended reaction sites with uniform electron distribution.3,8,21–24 For example, Choi et al. fabricated an Ag-deposited Y0.08Zr0.92O2-δ (YSZ) interlayer between La0.65Sr0.3MnO3-δ (LSM) and YSZ, and showed a substantially reduced charge transport resistance at the interface, and a ~16-fold increase in the peak power density at 650 °C, compared to the interlayer that was not deposited with Ag.8 Choi and Hwang reported a ~3-fold reduction in the polarization resistance of the Ag grid structure at the LSCF/YSZ interfaces, and demonstrated that the ORR kinetics was significantly promoted. They also showed that the reaction sites were extended with the fast and uniform electron supply at the interfaces.3 Furthermore, the increased catalytic activity of the ORR at the metallic functional layer can expedite the ORR kinetics, such as the oxygen adsorption and dissociation reactions.25–30 For example, exsolved Ag nanoparticles from the Sr0.95Nb0.1Co0.9O3 cathode yielded significantly reduced activation energy equal to 94 kJ/mol compared to cathodes without Ag deposition (100 kJ/mol) owing to its higher catalytic activity. This resulted in a reduced polarization resistance of ~40%, which yielded a value of 0.641 Ω cm2, and an outstanding peak power density of ~2 W/cm2 at 600 °C.30 Other metals with relatively lower catalytic activities, such as Ni, Cu, and Zn, have also been reported to substantially enhance the electrochemical performance in the IT regime.24,31,32 Despite the notable advantages of the metallic functional layers on the electrochemical performance, systematic investigations are still lacking. In particular, the quantitative understanding of the reaction mechanism associated with the enhanced electrical conductivity and catalytic activity of the metallic functional layer at the electrode–electrolyte interfaces is essential to achieve a rational design for the metallic functional layer with an improved performance. In this study, we systematically investigated the specific role of the metallic functional layer on the electrochemical performance of IT–SOFCs. Electrohydrodynamic (EHD) printing was used to print the metallic functional layers at the electrode/electrolyte interfaces with Ag, Ni, and Al, which have different electrical conductivities and catalytic activities. Substantial



reductions of the polarization resistances were achieved with the use of the metallic functional layers in the order of Al, Ni, and Ag. The smallest polarization resistance was obtained with the Ag functional layer, which was 0.016 cm2 at 650 °С, and ~8-fold smaller than that of the bare SSC cathode. Tafel, electrochemical impedance spectroscopy (EIS), and quantitative impedance fitting with distributed relaxation time (DRT) analyses, were performed to quantitatively deconvolute the different contributions from the electrical conductivity and catalytic activity on the polarization resistance. Lastly, the SSC cathode with the Ag functional layer demonstrated a peak power density of ~670 mW/cm2 at 650 °C, which was ~1.8-fold higher than that of the bare SSC cathode. Our results can provide an in-depth understanding for the rational design of the metallic functional layer that has led to the considerable performance enhancement of IT–SOFCs.

2. Experimental section EHD jet printing process (1) Preparation of the metal nanoparticle suspension The Ag (NPK Co.), Ni, and Al (NURIVISTA Co.) nanoparticle suspensions were used to fabricate the metal grids. The nanoparticle concentrations were adjusted to achieve densely packed metal lines with 70 wt% compositions for Ag, Ni, and Al grids. The solvents used in each suspension included triethylene glycol ethyl ether for Ag and dihydroteripinyl acetate for Ni and Al. Sm0.5Sr0.5CoO3 nanopowder (Kceracell Co.) was purchased to prepare the viscous SSC nanoparticle suspension. SSC nanopowder (4 g) was dispersed in 50 mL of ethanol, and the mixture was sonicated using a horn type sonicator for 6 h to achieve well-dispersed particles. The mixture was dried at 100 °С for 2 days to eliminate the solvent. Dimethylformamide (5.4 g, Sigma Aldrich) was added to the dried powder, and the mixture was sonicated with the use of a bath-type sonicator for 30 min. Polyvinylpyrrolidone (1.4 g, with a M.W. of 1,300,000, Sigma Aldrich) was also added to the mixture with the use of vortexing. The mixture was stirred at 60 °С for 24 h to achieve a well dispersed SSC suspension. All of the suspensions were stirred with a magnetic stirrer for 12 h and underwent vortex mixing before printing. (2) Experimental setup The EHD jet printing was conducted with the use of an EHD jet printer (NP–200, Enjet Inc.).


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The motorized X–Y–Z stages were precisely controlled at the micrometer scale with the fixed printing speed of 30 cm/s. The ceramic nozzle tip with an inner diameter of 150 μm was placed at a distance of 2 mm above the bottom collector, and the stainless steel nozzle body was connected to the high-voltage supply unit. The nanoparticle suspensions were emitted from the nozzle tip, which was supplied by a microsyringe pump, while the flow rate was fixed at 150 nL/min. The voltage difference between the nozzle and the grounded collector was adjusted to within the range of 1.4 to 1.6 kV, while it maintained a stable electrified jet. A relative movement was applied between the nozzle and the collector to draw the line patterns with desired widths and grid line pitch. Preparation of the samples (1) Symmetric cell fabrication The Gd0.1Ce0.9O2-δ (GDC) powder was pressed at a pressure of 40 MPa and was sintered at 1500 °C for 5 h to form the dense GDC electrolyte with a thickness of ~1.3 mm. Following the EHD jet printing process with a symmetric configuration on the as-prepared GDC pellets, presintering was performed at 200 °C for 2 h to remove the polymers and solvents in the printed structures. SSC powder (Kceracell Co.) was mixed with the binder (VEH, Fuel Cell Materials) at a weight ratio of 1:2.5 to fabricate the porous cathode structures and was symmetrically screen printed on the as-printed GDC pellets at a thickness of ~10 µm. After screen printing, the sintering process was executed at 800 °C for 3 h to remove the binder and to form the microstructure of the electrode. (2) Single-cell fabrication To evaluate the exchange current density of each material, single cells were fabricated with the use of an electrolyte-supported configuration. The EHD jet printing process was conducted on the cathodic sides of the GDC pellets, and presintering was performed at 200 °C for 2 h. After presintering, Pt paste (Pt paste 5542, Advanced Materials Technologies) was printed on the anodic sides of the GDC pellets. Sintering was performed at 800 °C for 3 h to remove the binder and organic materials in the Pt paste. To assess the electrochemical performance, anode-supported cells were fabricated. The NiO (Fuel Cell Materials), YSZ (Tosoh Corp.), and poly(methyl methacrylate) (PMMA) pore former (~5 μm) were mixed by a ball-milling process at a weight ratio of 6:4:1.5 with dispersant (HypermerTM KD-1, Croda) and with a polyvinyl butyral (PVB) binder. As-prepared anodic



powder was uniaxially pressed until it achieved a thickness of ~1 mm, and was subsequently presintered at 1000 °C. The pressed anode support was overprinted with a ~10-μm-thick anode functional layer (AFL) comprising a mixture of binder, NiO, and YSZ powder. The YSZ slurry contained ethanol (as a solvent), dispersant (Triton X–100), plasticizer (polyethylene glycol, Alfa Aesar), and binder (PVB), and was coated using a custom-made air-assisted electrospraying apparatus to obtain a dense and uniform electrolyte layer. The AFL and electrolyte were co-sintered at 1400 °C. The GDC diffusion barrier was electrosprayed at a thickness of ~5.5 μm followed by the sintering at 1200 °C for 3 h. Subsequently, the cathode layer, which included the metallic functional layer, was fabricated using the same way as that described in the symmetric cell preparation. Characterization methods A scanning electron microscope (SEM, JSM 7000F, JEOL) was employed to investigate the microstructure of the metallic functional layers and fabricated cells. Atomic maps were used to characterize the embedded functional layer at the interface and were analyzed by the energy dispersive spectrometer (EDS, JSM 7000F, JEOL). EIS measurements were conducted to compare the electrochemical behaviors of each SSC cathode with respect to the metallic functional layers with symmetric cell configurations. For these comparisons, a custom-made test station was used in ambient air equipped with an impedance analyzer (GAMRY Reference 600, GAMRY Inc.) in the frequency range of 0.01 to 106 Hz. For symmetric cell operations, both the cathodic and anodic sides of the electrodes were exposed to ambient air. Impedance fitting was conducted by an equivalent circuit model based on the calculation of the distribution of relaxation times (DRT) to identify the specific frequency range of the electrochemical reactions. The mathematical calculation was conducted with the software program FTIKREG. For single-cell operations, 100 sccm of H2 was supplied to the anode and cathode which were exposed to ambient air. Electrochemical I–V behaviors were measured with a cyclic voltammetry analyzer (GAMRY Reference 600, GAMRY Inc.). The patterned functional layers were observed with an optical microscope (Nikon Eclipse LV100ND, Nikon Co.). A four-point probe measurement system (MST–4000A, MS Tech.) was used to measure the inplane sheet resistances of bare and functional layer-coated SSC cathodes. 3. Results and discussion


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Figure 1. (a) Illustration of the overall fabrication process and (b) optical images of the printed metallic functional layer deposited using various metals. Figure 1(a) shows a schematic of the overall fabrication processes of the structured metallic functional layers of the grid. EHD jet printing was employed to fabricate precisely the metallic functional layers on the surfaces of the GDC electrolyte. The pitch of the grid structure was set to 50 µm for all metallic functional layers. Thereafter, the SSC powder was screen printed and achieved a thickness of ~10 µm. Detailed characterizations involving images at highmagnifications of the printed grid, and cross-sectional images of the as-prepared sample and the embedded Ag grid at the electrolyte and cathode using SEM, are shown in Figure S1. Specifically, EDS mapping in Figure S2 clearly shows the embedded Ag functional layers at the interface. Three different metals (Ag, Ni, and Al) were used to systematically investigate the effects of the metallic functional layers on the interfacial properties. Ag was reported to be the one of the most active materials in ORRs with a reactivity that is almost comparable to Pd and Pt, which are known to be the best catalysts for ORRs.25 Ni is less active in ORR than Ag, but one of the most conventional catalysts.9,31,33,34 To the best of our knowledge, Al has not been reported as a catalyst in ORRs. We selected Al as the least active catalyst. Figure 1(b) shows optical images of patterned metallic functional layers of Ag, Ni, and Al, on the GDC pellet. There are no discernible differences among metals, yielding a pitch of ~50 µm, widths in the range of 8–10 µm, and heights in the order of 1–2 µm. Such identical structures allow the direct investigation of the effects of the metallic functional layers from a material perspective.



Figure 2. I–V characterization of metallic functional layers in the single-cell configuration at 650 °С. (a) I–V curves and (b) Tafel plots based on the I–V curves. Before the detailed evaluation of the effects of the SSC cathode with the metallic functional layer on the electrochemical performance, comparisons of the electrical conductivities and catalytic activities of the metals were performed. The electrical conductivities of the metals in previous reports are shown in Figure S3. For example, at 600 °С, the electrical conductivity was the highest in Ag (9.1 × 105 S/cm) followed by Al (5.0 × 105 S/cm) and Ni (1.8 × 105 S/cm). The same trend was identified in all the corresponding temperature ranges. In contrast, SSC—the cathode material used in this work—yielded an electrical conductivity (1.5 × 103 S/cm) which was lower by approximately two orders of magnitude compared to all the studied metals. Even though there were minor differences among the metals, Ag, Al, and Ni, yielded considerably higher electrical conductivities than SSC. Catalytic activities of each metal were quantitatively assessed using Tafel plots. Based on the


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Butler–Volmer equation, the exchange current density (𝑗0) can be extrapolated by fitting the current density as a function of overpotential in the Tafel plots.35 Because the exchange current density represents the reaction kinetics on the surface at the cathode, it has been used as an indicator of the catalytic activity of the materials.6,35,36 In the low-current density regime, activation loss—which arises from the charge transfer rate at the cathode—can be simplified according to Eq. (1), as listed below. Fitting of voltage–exchange current density plots in Figure 2(a) followed by extrapolation allows one to obtain y-intercepts at zero current density and hence, calculate ηact,cathode (Figure 2(b)) according to Eq. (1). 𝜂𝑎𝑐𝑡,𝑐𝑎𝑡ℎ𝑜𝑑𝑒 =

RT (αnF )ln (𝑗𝑗 ), 0

(1)

where 𝜂 is the overpotential, R is the ideal gas constant, T is the temperature, α is the chargetransfer coefficient, F is Faraday’s constant, and n is the number of the moles of the electrons transferred. The metallic functional layers with different metals were printed on the cathodic side using EHD jet printing with a pitch size of 50 μm, and Pt paste was then screen printed on the anodic side, as illustrated in Figure 2(a). Figure 2(a) shows the I–V behaviors measured at 650 °С. The limiting current density which represents the current values of the x-intercept in the I–V behaviors is the highest in Ag (0.035 mA/cm2) followed by Ni (0.0056 mA/cm2) and Al (4.6 × 10-6 mA/cm2). Figure 2(b) shows the Tafel plots quantified at 650 °С. Similar with the trend of the limiting current density, the extracted exchange current density is the highest in the Ag (3.06 × 107 A/cm2) functional layer, followed by that for the Ni (1.12 × 107 A/cm2) and Al (0.0018 × 107 A/cm2) functional layers. The specific order for the exchange current densities of the Ag and Ni functional layers are the same as the previously reported catalytic activities of ORRs.25 Conversely, the Al functional layer yielded an exchange current density which was approximately three orders of magnitude lower compared to those for the Ag and Ni functional layers. This indicated a significantly lower catalytic activity for ORR compared to the Ag and Ni functional layers. Thus, the significantly retarded catalytic activity of the Al functional layer in ORRs compared to the activities of Ni and Ag functional layers can be ascribed not only to the lower exchange current density observed in the former case, but also to the lower limiting current density of this (Al) layer.



Figure 3. EIS results of the SSC cathode with the metallic functional layer. (a) Nyquist plots at 650 °С, and (b) Arrhenius plots in the temperature range of 650–500 °С. To assess the effect of the SSC cathode with the metallic functional layer on the electrochemical performances at the cathode/electrolyte interface, EIS measurements were performed in a symmetric cell configuration in the temperature range of 650–500 °С in ambient air, as shown in Figures 3(a)–(b). Figure 3(a) shows a representative Nyquist plot at 650 °С. The ohmic resistance which is due to the ionic transport through the electrolyte was eliminated for the direct comparison of the polarization resistance (Rp). This resistance is attributed to the electrode reactions that involve complicated ORR, such as gas diffusion, adsorption, dissociation, surface diffusion, and incorporation. The SSC powder electrode which was not deposited with the metallic functional layer (bare SSC cathode) yielded an Rp of 0.13  cm2 at 650 °С, which well matched the previously reported values (0.1–0.15  cm2).4 With the metallic functional layer at the cathode/electrolyte interfaces, the Rp was substantially reduced. The SSC cathode with the Al, Ni, and Ag functional layers, respectively yielded ~4-fold (0.035


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cm2), ~5-fold (0.025 cm2), and ~8-fold (0.016 cm2) decreases in Rp at 650 °С compared to the bare SSC cathode. These results confirmed that the electrochemical properties, especially at the interface, were significantly improved with the use of the metallic functional layers, and that these improvements were depended on the materials. Note that Rp of the SSC cathode with the SSC functional layer was also reduced ~2-fold (0.06 cm2) at 650 °С compared to the bare SSC cathode. The grid-structured SSC functional layer was also printed (Figure S4) to directly verify the effects of the grid structure on the electrochemical performance in the absence of catalytic activity. Because the SSC functional layer had the same material properties as those of the bare SSC cathode, and because it resulted in an identical catalytic activity, the Rp reduction of the SSC cathode with the SSC functional layer was entirely attributed to the electrical conductivity of the grid structure. According to previous work, the sheet resistance was reduced at the interfaces when the grid-structured metallic functional layer was used. This was attributed to the even distribution of electrons, which resulted in the substantial decrease in the Rp.3 In Figure S5, the SSC cathode with the SSC functional layer (40.9 ± 8 Ω/sq) shows a ~4-fold lower in-plane sheet resistance compared to that for the bare SSC cathode (162.2 ± 16 Ω/sq). This indicates that the enhanced electrical conductivity was attributed to the grid structure despite the fact that the same material was used. It can be stated that the printed SSC powders in the SSC functional layer are more densely packed owing to their higher powder concentrations (~40 wt%) compared to the screen printed SSC powder (~28 wt%). Consequently, a lower Rp can be achieved with the SSC functional layer at the interfaces owing to the higher electrical conductivity of the grid structure. Figure 3(b) shows the Arrhenius plots with the metallic functional layer in the temperature range of 650–500 °С. The slope of the Arrhenius plot denotes the activation energy (Ea), and provides valuable information about the RDS among the complicated electrochemical reactions.37,38 The bare SSC cathode shows that the Ea is 1.36±0.1 eV, which is similar to the previously reported values in the range of 1.3–1.4 eV. This value indicates that the charge transfer reactions at the interface are the slowest.37–39 Both the SSC cathode with the SSC functional layer and the SSC cathode with the Al functional layer show similar Ea values in the range of 1.35–1.37 eV, thus suggesting that the RDS remains the same as the bare SSC cathode. However, the absolute value of the Rp at 650 °С was reduced ~2-fold and ~4-fold in the cases of the SSC cathode with the SSC functional layer and the SSC cathode with the Al functional layer, respectively. The parallel downshift in the Arrhenius plot, or the reduction of the absolute



value of Rp, maintained the same Ea value and indicated the expansion of the reaction sites with no changes in the RDS toward the ORR.3,4 Therefore, the reduced Rp and the maintenance of the RDS can be attributed to the TPB enlargements mainly owing to the enhanced electrical conductivity with the SSC functional layer and Al functional layers, and not owing to the catalytic activity which changes the RDS.3 Conversely, the SSC cathodes with the Ni and Ag functional layers yielded reduced Ea values that were equal to 1.12±0.3 eV and 0.96±0.04 eV, respectively, and reduced Rp values. Such significant changes in Ea indicate that the RDS was changed in the presence of the Ni and Ag functional layers. In particular, Ea represents the RDS of the oxygen surface exchange and its value was in the range of 0.9–1.2 eV.11,37,38 Therefore, the reduced polarization resistances observed for the SSC cathode with Ni and Ag functional layers were explained not only in terms of electrical conductivity but also considering the catalytic activity of the employed materials, and the corresponding RDS values changed from 1.36 to 1.12 and 0.96 eV, respectively.21,26 The reaction barrier for the charge transfer was reduced when the catalysts Ag and Ni were used. This resulted in the promotion of the reaction kinetics and the reduction of the Rp values.26,27,40 Furthermore, the trends of the Ea and Rp in the case of the SSC cathode with Al, Ni, and the Ag functional layers, appear to be the same as those of the exchange current density in Figure 2(b) and the previously reported ORR activities, thus substantiating the specific catalyst role of the metallic functional layers in ORRs. Detailed values of the results of the EIS analysis are listed in Table S1.


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Figure 4. Representative DRT analyses and Gaussian fittings at 650 °С. (a) DRT analysis of the bare SSC cathode, (b) Gaussian fitting of the bare SSC cathode, (c) DRT analysis of the SSC cathode with the Ag functional layer, and (d) Gaussian fitting of SSC cathode with the Ag functional layer. To achieve quantitative deconvolution of the contributions from the electrical conductivity and the catalytic activity in the metallic functional layers, impedance spectra were fitted using equivalent circuit models based on the DRT analyses.41–43 Complicated electrochemical reactions, including those associated with the gas diffusion and ionic conduction in the electrolyte, are reflected in the relevant frequency ranges of the impedance spectra.9,22,38,39,44– 46

Adler et al. suggested three distinct frequency ranges, including the high- (HF, >103 Hz),

medium- (MF, 10–103 Hz), and low- (LF,

Rational Design of a Metallic Functional Layer for High-Performance

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