Improved Thermoelectric Properties of Re-Substituted Higher

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Functional Nanostructured Materials (including low-D carbon)

Improved thermoelectric properties of Re substituted HMS by inducing phonon scattering and energy filtering effect at grain boundary interfaces Swapnil Ghodke, Akio Yamamoto, Hsuan-Chun Hu, Shunshe Nishino, Takuya Matsunaga, Dogyun Byeon, Hiroshi Ikuta, and Tsunehiro Takeuchi ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b09397 • Publication Date (Web): 05 Aug 2019 Downloaded from pubs.acs.org on August 5, 2019

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ACS Applied Materials & Interfaces

Improved Thermoelectric Properties of Re Substituted HMS by Inducing Phonon Scattering and Energy Filtering Effect at Grain Boundary Interfaces Swapnil Ghodke§*, Akio Yamamoto§, Hsuan-Chun Hu§, Shunsuke Nishino§, Takuya Matsunaga§, Dogyun Byeon§, Hiroshi Ikuta†, Tsunehiro Takeuchi§

§Energy

Materials Laboratory, Toyota Technological Institute, 2-12-1 Hisakata, Tempaku

Ward, Nagoya, 468-0034, Japan

†Department

of Materials Physics, Nagoya University, B3-1 (622), Furo-Cho, Chikusa-Ku,

Nagoya 464-8603, Japan

Keywords: higher manganese silicide, phonon scattering, energy filtering effect, grain boundaries, liquid quenching, ribbon sample, thick films.

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Abstract

In this study, the effect of grain boundary density on transport properties of Re substituted higher manganese silicide Mn30.4Re6Si63.6 has been investigated. The efficiency of electrical energy conversion from waste heat, mainly in thermoelectric generators depends on how the thermal conduction is reduced while the charge carrier electrons/ holes contribute to possess a large magnitude both of electrical conductivity σ and Seebeck coefficient S to consequently lead to a large power factor PF= S2σ. In this work, we tried to obtain such a condition with a novel approach of merging the energy filtering effect at the grain boundaries to improve the power factor. The nano-structuring and heavy element substitution were also employed to greatly scatter the phonon conduction. As a result, the enhancement of power factor was observed in the diffused nanostructure of annealed ribbon samples and the enhancement was correlated to the formation of Schottky barriers at the grain boundary interface. Together with the reduction of thermal conductivity to very low magnitude 1.27 Wm-1K-1, we obtained a maximum ZT = 1.15 at 873 K for the annealed ribbon samples.


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INTRODUCTION Thermoelectric generators (TEG’s) are solid-state devices that can convert waste heat into useful electrical energy without any moving mechanical parts or fluids. The performance of a thermoelectric generator is measured by the dimensionless figure of merit (ZT = S2σT/ (κe+κl)). Here, S, σ, κe, κL, and T are Seebeck coefficient, electrical conductivity, electron thermal conductivity, lattice thermal conductivity, and absolute temperature, respectively1–3. S, σ, and κe are mainly dependent on the electronic structure near the chemical potential, while κL is the ability of phonons to conduct energy as collective lattice vibrations4,5. The state-of-art thermoelectric materials are Pb-Te6, Bi-Sb-Te7,8, Zn-Sb9, SiGe10,11, SnSe, Cu2Se12, etc., which possess ZT > 1 because of their electronic structure and low lattice thermal conductivity13,14. However, the drawbacks of these compounds are that the constituent elements that are either toxic or expensive, hence limiting its usage for practical application. The most prospective thermoelectric materials are silicides, such as Mg2Si, MnSi, βFeSi2, and CrSi2, because their constituent elements are abundant, inexpensive, non-toxic, with low cost of production, and have good oxidation resistance at high-temperature15,16. The



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Higher Manganese Silicide (HMS, MnSiγ) (1.73 ≤ γ ≤ 1.75), has the potential to show a high thermoelectric performance among all the silicides because (a) the electronic structure has wide band gap ~0.8 eV with a sharp intense peak near the valance band edge in the density of states17,18, (b) the complex Nowotny chimney ladder (NCL) crystal structure has a high probability of phonon scattering19,20 resulting in lower lattice thermal conductivity compared to CrSi2/ β-FeSi2, and (c) MnSi has good oxidation stability compared to Mg2Si. Typical non-doped p-type HMS shows a large Seebeck coefficient ~ 230 μVK-1, low electrical resistivity ~ 2 ×10-3 Ωcm, and moderately low thermal conductivity ~ 2.5 WmK-1 to consequently possess ZT of 0.3-0.45 at around 800 K21–24. Unfortunately, the ZT value was not competitive with other state-of-art thermoelectric materials. A high-performance thermoelectric material should possess high power factor S2σ along with low lattice thermal conductivity. Theoretically, all the physical quantities are strongly coupled with each other, which make it more challenging to obtain a highperformance in the thermoelectric material. In order to improve the performance of HMS, Miyazaki et al., Ponnambalam et al., and Truong et al. have performed carrier concentration tuning by substitution of Cr, Fe, V, Co, Ru, W, Mo at Mn Site24–31 or Al / Ge at Si site 28,32,33 and achieved a maximum ZT value of 0.7. The lattice thermal conductivity of HMS was


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effectively reduced by W substitution34, Re substitution in HMS35,36 by inducing impurity scattering with large mass difference in heavier (W/Re) and lighter (Mn) atoms. It was also confirmed that the heavy element substitution does not create any impurity states in the electronic density of states near the chemical potential to affect the electron transport properties35. Nano-structuring has drawn our attention to investigate a novel technique that can improve the power factor and reduce the lattice thermal conductivity simultaneously. Here, the grain boundaries can act as a scattering center both for electrons and phonons. In the case of phonon scattering, it reduces the phonon mean free path and consequently reduces lattice thermal conductivity37–41. Likewise, the barriers introduced by the interfaces of inclusion or grain boundaries act as scattering centers for low-energy/high-energy carriers in n-type/ptype materials leading to energy filtering effect42–46, with which Seebeck coefficient was significantly increased. Though the scattering of the carrier increases electrical resistivity, the Seebeck is a squared term and the slight increase in S makes a significant contribution to power factor. Recently, Dong-Kyun Ko et al. reported that by making nano-composite with Pt nano-inclusion in Sb2Te3 matrix, the Pt nano-crystals (NC) acted as a barrier for low-energy



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carriers and the ZT was increased47. Soni et al. found out similar effect to argue the energy filtering at grain boundaries48. They studied the effect of sintering temperature on grain size, as the grain size was increased the effect of energy filtering on power factor was minimized. By considering these reports, we investigated, in the previous work, the effect of carrier filtering due to secondary phases (Si / MnSi) and grain boundaries in HMS49. As a result, we derived that the energy filtering effect at the grain boundaries might be due to the formation of Schottky barriers (Figure 1a) that can enhance the power factor by 20-30%.

Low-energy carriers Mid / long wavelength phonons

High-energy carriers Short wavelength phonons

Figure 1. (a) Schematic of a band diagram representing the mechanism of energy filtering effect. Formation of the Schottky barrier at the interface of grain boundary and matrix phase (p-type HMS scatters the low-energy holes and allows high-energy holes). (b) Schematic of microstructure which can induce energy filtering effect (grain boundaries) and phonon scattering (nanograins, heavy elements, atomic defects). Therefore, in this work, a novel strategy was implemented to enhance ZT of HMS by combining two effects simultaneously (1) employing the energy filtering effect at grain boundaries in HMS to increase the power factor, and (2) reducing the lattice thermal


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conductivity by heavy element substitution & nano-structuring (Figure 1b). Re substituted HMS were selected to study the effect of controlled modification in microstructure on the thermoelectric properties by employing liquid quenching technique, subsequent annealing, and spark plasma sintering. EXPERIMENTAL PROCEDURES High purity Manganese (99.99%), Silicon (99.99%), and Rhenium (99.99%) powders were weighed in a stoichiometric ratio of Mn30.64Re6Si63.36. The powders were mixed thoroughly by mortar and pestle in air, then cold-pressed into a pellet. The density of the prepared pellet was good enough to hold the shape to be melted in the arc-melting furnace. In the arcfurnace, the ingot was repeatedly melted over 3 to 4 times under argon gas atmosphere to obtain a homogenously mixed compound. The obtained ingot was re-melted in a silica tube using induction heating and rapidly quenched by being injected on a copper wheel rotating at 4500 rpm in the pressurized argon atmosphere. The melt-spinning parameters were selected as follows: - Nozzle diameter was  = 0.4 mm, temperature ~1473 K, and cooling rate of 105106 K/s. The liquid quenching process produced a ribbon-shaped samples (thick films) with 10 μm ~ 15 μm in thickness, 1 mm – 1.5 mm in width, and 10 mm-15 mm in length. Some of the ribbons were annealed at 1200 K for 5 hours under vacuum condition and some others



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were sintered by spark plasma sintering (SPS) at 1200 K, 60 MPa uniaxial pressure, 50 °C/min heating rate, and 5 min soaking. The phases formed in the non-annealed ribbons, annealed ribbons, and sintered samples were identified on grounded powders by conventional powder X-ray diffraction using BRUKER D8 Advance with Cu-Kα radiation (𝜆 = 0.15418 nm). The evolution of microstructure was analyzed by scanning electron microscope (SEM) JEOL JSM-6330F. The Seebeck coefficient along the in-plane direction was measured by steady-state method using the Seebeck Measurement System developed by Micro-miniature refrigerator (MMR) technologies Inc. in the temperature range of 300 K – 700 K. We also employed a handmade setup50, with which ΔV vs. ΔT plots were used to determine S, at 300 K – 973 K. Electrical resistivity was measured by standard four-probe technique under a vacuum atmosphere ~ 10–2 Pa along the in-plane direction. The uncertainty in resistivity measurement for the nonannealed ribbons was tabulated in Table S1. Thermal conductivity of the ribbon samples was calculated using κ = D*ρ*Cp, the thermal diffusivity (D) along the in-plane direction measured by AC calorimetry method51,52 (Supplementary Figure S1 & S2), the density ρ measured with Archimedes-method (ρnon-Ann = 4.30(50) g/cm3, ρANN = 5.32 (50) g/cm3, ρSPS = 5.90(5) g/cm3, and ρcalc.Re6 = 6.266(5) g/cm3), and specific heat was determined by the


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Differential Scanning Calorimetry (DSC) (Supplementary Figure S3). The thermal diffusivity (D) of the sintered pellet (out-of-plane direction) was also measured in the temperature range of 300 – 973 K using Laser flash method (LFA457, NETZSCH), specific heat (Cp) was extracted from the reference sample (Pyroceram9606). The error bars or the uncertainty for the Seebeck coefficient (7%) was obtained from several measurements and instrumental inaccuracy. Similarly, for electrical resistivity 10% uncertainty was obtained from the dimensional inaccuracy such as thickness, width, the distance between the voltage probes. The error in the thermal conductivity (5-7%) is due to the thickness (t), density (ρ), Cp, and thermal diffusivity (D), which are measured multiple times. Therefore the error in ZT was about 20%. RESULTS The measured X-ray diffraction (XRD) patterns were plotted (Figure 2) for the non-annealed ribbons, annealed ribbons, and sintered Mn30.36Re6.0Si63.64 sample. All the XRD patterns show no sign of MnSi (manganese mono-silicide), which was often reported to appear as an impurity phase53,54. The absence of MnSi would be closely related to the rapid quenching process. The XRD pattern for the non-annealed ribbons showed broadening in all diffraction peaks.



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Figure 2. XRD patterns of Re free sample HMS (calculated Mn15Si26) and Re substituted HMS samples (non-annealed ribbons, annealed ribbons, and sintered/ SPS sample)

The peak broadening, could be mainly caused by nano-grains, some contribution from inhomogeneous composition and/or heterogeneous strain in the sample. Notably, the peak broadening was diminished after annealing at 1200 K for 5 hours, the main peak at 41.6° (2110) transforms into two peaks. The sharpened peaks were very similar to those of the bulk sample sintered at 1200 K. The XRD patterns indicate the crystallinity of sample changes with the heat treatment process. The XRD of as-quenched ribbons also showed some preferred orientation, that was confirmed at (1010) and (1120) peaks associated with the Mn subsystem. However, it turned out to disappear from the annealed ribbons and sintered XRD patterns. Such preferred 10 ACS Paragon Plus Environment

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orientation might affect the transport properties as reported by Sadia et al.55. However, the anisotropic TE properties of polycrystalline ribbons have not been fully investigated yet and will be reported in the future.

Figure 3. SEM images of HMS (a) Re free HMS with an average grain size of 2 m and (b) Re substituted HMS with an average grain size of 250 nm.

Figure 3a and 3b show the microstructure of a liquid quenched ribbon sample with/without the Re substitution. The average grain size observed for Re-free sample was ~2 m while 6 at. % Re substitution drastically reduced the grain size to ~250 nm, presumably due to the increased rate of nucleation around Re atoms in the rapid quenching process. Considering the microstructure with very fine equiaxed grains in the order of nanometers, we expected a large energy filtering effect as we observed for Re free HMS49 and a strong phonon scattering at the grain boundaries38. Modification of the fine grain structure (Figure 4a) was observed in annealed ribbons (Figure 4b) which show a diffused microstructure with an average grain



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size of ~2 m. In the SPS process, the grain size increases drastically to 8 m (Figure 4c) due to rapid grain boundary migrations caused by the Joule heating at the resistive portion, i.e. grain boundaries. The grain boundary density (GBD) was effectively decreased with the sintering process and also with the annealing. Thus, we strongly expected that the change in GBD would influence the energy filtering effect and phonon scattering.

Figure 4. SEM images showing the effect of the thermal processing on the microstructure of 6 at. % Re-substituted HMS. (a) non-annealed ribbon sample with fine grain structure, (b) annealed ribbon with diffused microstructure, and (c) microstructure of cracked surface of SPS pellet shows a drastic grain growth. The Seebeck coefficient, electrical resistivity, power factor, and thermal diffusivity was measured for non-annealed ribbons, annealed ribbons and SPS samples at room temperature and were plotted as a function of grain size in Figure 5. It was evident that the fine microstructured ribbon sample possessed the largest Seebeck coefficient, which decreases with an increase in grain size (Figure 5a). The reduction in electrical resistivity (Figure 5b) with increasing grain size could be correlated to (a) decrease in porosity and (b) decrease in low-energy carrier scattering at the grain boundaries. 12 ACS Paragon Plus Environment

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The grain size dependence of PF was calculated and shown in Figure 5c. Among three samples of different preparation conditions, the annealed ribbon sample possessed the largest power factor both at 300 K and 700 K. Improvement in the PF was 30% in temperature range of 300K - 400 K and 11% at temperatures above 600 K. The absolute value of thermal diffusivity (Figure 5d) increased with the grain growth. The increase in thermal diffusivity must be due to suppression of phonon scattering at the grain boundaries. Since we realized that the non-annealed sample showed a very low PF presumably due to the macroscopic voids and/or cracks, we eliminated it for further temperature-dependent analysis.

Figure 5. Grain size dependence of (a) Seebeck coefficient (300 K), (b) electrical resistivity (300 K), (c) power factor at 300 K & 700 K, and (d) thermal diffusivity (300 K) for nonannealed ribbon, annealed ribbon, and Sintered/ SPS sample.



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Figure 6. Temperature dependence of transport properties, (a) Seebeck coefficient, (b) electrical resistivity, (c) Power factor, and (d) thermal conductivity for 6 at% Re HMS annealed ribbons and SPS sample. The temperature dependence of Seebeck coefficient, plotted in Figure 6a, showed higher values for annealed ribbons throughout the temperature range of measurements with the maximum value of 235 μVK-1 at 873 K. A typical metallic-like electrical conduction was observed for ribbons and sintered sample (Figure 6b). The PF for annealed ribbons was


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higher over the whole temperature range, except at 973 K, where it becomes almost equivalent to sintered sample (Figure 6c). Maximum PF of 2.29 mWm-1K-2 at 873 K was obtained for annealed ribbon sample. The estimated thermal conductivity for annealed ribbons was 1.27 Wm-1K-1, which is 34% smaller than that of the sintered bulk sample (1.93 Wm-1K-1). The Wiedemann-Franz law κe = L0σT was used to estimate the electron thermal conductivity, and the lattice thermal conductivity was derived from κl = κ – κe. The κl = 0.855 Wm-1K-1 at 300 K was obtained for the ribbon sample and the temperature-dependent κ for ribbons was calculated from κe(T) + κl (0.855 Wm-1K-1), assuming a very weak temperature dependence in κl (Figure 6d). The major difference between two samples was in the lattice thermal conductivity that would be related to a strong phonon scattering of phonons having mid/long-range wavelength compared to the nano-grained diffused microstructure of the annealed ribbons. As a result, we observed a ZT = 1.15 at 873 K for annealed ribbons (Figure 7). The enhancement in the ZT value was about 30% in comparison to the SPS sample (ZT = 0.8). Notably, this value, ZT = 1.15, is the largest ever reported for the HMS-based thermoelectric materials. This extraordinary value of ZT was successfully obtained by the unique strategy in



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which energy filtering effect, nano-structuring, and heavy element substitution are simultaneously employed.

Figure 7. Temperature dependence of ZT for (a) Re substituted annealed ribbons, (b) Re substituted SPS sample, and (c) Re free HMS24.

DISCUSSION To understand the observed energy filtering effect in terms of electronic structure, we derived Seebeck coefficient S(T) and electrical conductivity σ(T) from spectral conductivity σ(ε, T) in the context of linear response theory4,56 (Figure 8: Equation (1) - (4)). Here, S(T) and σ(T) are the functions of spectral conductivity σ(ε,T), Fermi-Dirac distribution function fFD, electronic density of states N(ε), mean group velocity vG(ε) of an electron, relaxation time τ (ε, T), and the chemical potential μ (T). The σ(T) is determined by W0 = (-∂fFD/∂ε). W0 shows


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a peak exactly at the chemical potential with a full-width-half-maximum of about 3.5 kBT. The S(T) is strongly dominated by the product of W1 = (ε-μ)(-∂fFD/∂ε) and σ(ε,T). Here, W1 has negative and positive maxima at ε -μ =  1.7 kBT. Generally, the energy dependence of relaxation time is ignored, but if we purposefully introduce some energy dependence by increasing scattering probability for low-energy carriers the S(T) and σ(T) would be modified. For example, if the scattering of carriers occurs below the chemical potential by introducing barrier by grains or inclusions, the negative contribution of S(T) below the chemical potential can be reduced (W1) (Figure 8). Thus the numerator part of Seebeck is increased. The W0 is also affected by the scattering and reduction of carriers that reduce the σ(T). This means the denominator term for S (T) is reduced.

Figure 8. Schematic of the energy dependence of W0, W1, and spectral conductivity with increased scattering probability (hatched) below the chemical potential.



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So the carrier scattering increases the S(T) and reduces σ(T), thus the power factor (Equation-4) increases since the S(T) is a squared term. Figure 9 shows the Schottky barrier formation at boundary and strong energy dependence of electron scatterings mechanism. The schematic shows a band structure of grain boundary interface (metallic) with p-type HMS matrix at T = 0 and T = T1 >>0. The chemical potential (0) of both materials stay at the same energy in the valance band (VB) at absolute zero kelvin. Since materials possessing a large Seebeck coefficient show a significant temperature dependence of S(T), the chemical potential of HMS is supposed to move into the band gap with an increase in temperature57 (T1), while that of grain boundary is kept almost unchanged because of the metallic nature. The difference of (T1) between two materials naturally causes band bending at the interface, thus forming a Schottky barrier. This barrier acts as a scattering center for the low-energy holes creating an energy filtering effect to improve the power factor. In this scenario, metallic grain boundaries formed between semiconducting materials of high Seebeck coefficient always produce the Schottky barrier regardless of carrier types. One of the typical example is Bi-Sb-Te prepared by using partial melting process, where metallic phase covered the whole surface of all Bi-Sb-Te grains inducing energy filtering effect8.


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Figure 9. Schematic of a band diagram representing (a) the interface of HMS matrix with grain boundary with chemical potential at absolute zero Kelvin, (b) the formation of the Schottky barrier at the interface at T = T1 K. Before concluding, we have to mention on the contradictory fact that the as-quenched ribbon samples (non-annealed samples) showed smaller values for PF even with the smallest grain size. The as-quenched samples are characterized not only by the small grains but also by large porosity or voids, which were presumably introduced by the shrinkage of each grain at the rapid-cooling process. The SEM images proving the presence of such voids are shown in Figure 4a. Although the value of the Seebeck coefficient is almost independent of the voids, the electrical resistivity is strongly increased by these voids. Therefore, non-annealed ribbons showed no improvement in the PF due to degradation in electrical resistivity caused by the porosity.



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Nevertheless, this consideration led to a prospect of further increase of PF with eliminating voids in the as-quenched ribbons. It would be possible to obtain a higher ZT in the bulk samples by precisely modifying sintering parameters to obtain densely packed nanostructure, for example, high-pressure & low-temperature sintering58 or liquid phase sintering8. It would be interesting to implement such a novel strategy for n-type material as well. CONCLUSION The significant reduction of grain size (~250 nm) was observed for the Re-substituted HMS, which was noticed under the combined effect of rapid quenching and solute substitution. We also observed improvement in PF most likely by the implementing energy filtering effect at the grain boundaries. The annealed liquid quenched samples showed ~30% improvement in PF, especially at low-temperature. We succeeded in implementing a novel strategy of simultaneously reducing the lattice thermal conductivity to 1.27 Wm-1K-1 (34%) by partial substitution of Re for Mn and nano-structuring, as well as improving PF (2.29 mWm-1K-2 at 873 K) by energy filtering at the grain boundaries. Thus we succeeded in obtaining a highest ZT = 1.15 at 873 K for HMS based alloy.


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ASSOCIATED CONTENT Supporting Information The Supporting Information is available free of charge on the ACS Publications website. The detailed description of AC calorimetry method/ setup with the schematic, frequency dependence of thermal diffusivity measured data, and working principle to measure the thermal diffusivity of ribbon samples has been explained. We have also added a comparison between specific heat data measurement by DSC and Laser flash method for ribbon and sintered samples. Table for room temperature resistivity measurement of the ribbon sample and the uncertainties in dimensions and absolute values has been added. AUTHOR INFORMATION

Corresponding Author

*E-mail: [email protected] ACKNOWLEDGMENT This work was conducted under the financial support of JST CREST and Japan Society for the Promotion of Science (JSPS) KAKENHI, Grant Nos. 18H01695, 18K18961. The authors would like to thank Dr. Saurabh Singh, Dr. Seongho. Choi, Dr. Muthusamy Omprakash, and Ashwini Pachkor for their fruitful discussions. 21 ACS Paragon Plus Environment


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