# Fileset

[MDR-article_Yb4Sb3_Sylvain.pdf](https://mdr.nims.go.jp/filesets/cf274a32-3f82-4508-83d1-e165b294b49c/download)

## Creator

[Sylvain Le Tonquesse](https://orcid.org/0000-0002-1939-7816), [Hugo Bouteiller](https://orcid.org/0009-0004-2132-2962), [Yoshitaka Matsushita](https://orcid.org/0000-0002-4968-8905), Araseli Cortez, Sabah K. Bux, Kazuki Imasato, Michihiro Ohta, [Jean-François Halet](https://orcid.org/0000-0002-2315-4200), [Takao Mori](https://orcid.org/0000-0003-2682-1846), Franck Gascoin, [David Berthebaud](https://orcid.org/0000-0002-2892-2125)

## Rights

This document is the unedited Author’s version of a Submitted Work that was subsequently accepted for publication in ACS Applied Energy Materials, copyright © 2023 American Chemical Society after peer review. To access the final edited and published work see https://doi.org/10.1021/acsaem.3c01693.[In Copyright](http://rightsstatements.org/vocab/InC/1.0/)

## Other metadata

[Enhanced High-Temperature Thermoelectric Performance of Yb<sub>4</sub>Sb<sub>3</sub> <i>via</i> Ce/Bi Co-doping and Metallic Contact Deposition for Device Integration](https://mdr.nims.go.jp/datasets/eb2e2cd5-a6c2-47a9-b133-d7c7f3ce4782)

## Fulltext

Enhanced high-temperature thermoelectric performanceof Yb4Sb3 via Ce/Bi double doping and metallic contactdeposition for device integrationSylvain Le Tonquesse,∗,† Hugo Bouteiller,† Yoshitaka Matsushita,‡ AraseliCortez,¶ Sabah K. Bux,¶ Michihiro Ohta,§ Kazuki Imasato,§ Jean-François Halet,∥Takao Mori,⊥ Franck Gascoin,† and David Berthebaud∗,∥,#†Laboratoire CRISMAT UMR 6508 CNRS ENSICAEN, 14050 Caen Cedex 04, Caen,France‡National Institute for Materials Science, 1-2-1 Sengen, Tsukuba, Ibaraki, 305-0047, Japan¶Thermal Energy Conversion Research and Advancement Group, Jet PropulsionLaboratory, California Institute of Technology, 4800 Oak Grove Drive, Pasadena, CA91109-8099, USA§Global Zero Emission Research Center, National Institute of Advanced Industrial Scienceand Technology (AIST), Umezono 1-1-1, Tsukuba, Ibaraki 305-8568, Japan∥CNRS-Saint-Gobain-NIMS, IRL 3629, Laboratory for Innovative Key Materials andStructures (LINK), National Institute for Materials Science, Tsukuba 305-0044, Japan⊥National Institute for Materials Science (NIMS), WPI-MANA, 1-1-1 Namiki, Tsukuba305-0044, Japan#Nantes Université, CNRS, Institut des Matériaux de Nantes Jean Rouxel, IMN, F-44000Nantes, FranceE-mail: sylvain.le-tonquesse@ensicaen.fr; david.berthebaud@cnrs.fr1AbstractThermoelectric (TE) for very high temperatures (> 800 K) have numerous poten-tial applications in the heavy industry and the space exploration. This article focuson the compound Yb4Sb3 which is a promising p-type counterpart to the structurallyrelated and high-performance n-type RE3Te4 (RE = Nb, La, Pr) for the fabrication ofhigh temperature TE modules. A quick and e�cient method for synthesizing pure andfully-dense Yb4Sb3 samples was developed and optimized using high-energy ball millingfollowed by reactive spark plasma sintering. The technique was utilized to produce anew series of doubly doped CexYb4−xBi0.2Sb2.8 compounds. X-ray di�raction and scan-ning electron microscopy (SEM) were employed to establish the solubility limit of Ce,which was determined to be x = 0.4. TE properties for Yb4Sb3 and Ce0.4Yb3.6Bi0.2Sb2.8were measured up to 1273 K, revealing that the doping strategy was e�ective in reduc-ing the charge carrier concentration and thermal conductivity. This led to a signi�cantincrease in the TE �gure-of-merit ZT from 0.2 to 0.4 at 1273 K. This study eventuallyconducted a preliminary examination of the realization of metallic contacts on Yb4Sb3through SPS. The results showed that two robust TE legs with Ni and Cu contacts weresuccessfully produced, and SEM analyses was used to study the di�usion that occurredat the interfaces. The measured electric contact resistances were very promising, withaverage values of 2 and 1 µΩ cm2 for Ni and Cu contacts, respectively.KeywordsThermoelectrics, Intermetallics, Electric contacts, Ball-milling, High-temperature1 IntroductionThermoelectric (TE) generators are reliable solid-state devices that take advantage of theSeebeck e�ect to produce electrical energy from a temperature gradient. Typically, TE2generators consist of two TE materials with opposite electrical conduction types, one is n-type and the other p-type, which are connected electrically in series but thermally in parallel.The conversion e�ciency of the device is directly related to the �gures-of-merit ZT of theindividual materials which are given by:ZT =α2Tρκ(1)with α the Seebeck coe�cient, ρ the electrical resistivity, κ the thermal conductivity and Tthe absolute temperature. These intrinsic properties of materials vary with temperature andare optimized di�erently for each materials. Thus, the choice of TE materials for a generatorprimarily depends on the operational temperature of the target application. Previous re-searches have primarily focused on materials suitable for near room-temperature applications(300 K - 400 K), such as Bi2Te3 1 or Mg3Bi2,2 or for mid-temperature applications (600 K- 800 K), such as PbTe,3,4 skutterudite CoSb3 5,6 and transition metal silicides.7,8 Nonethe-less, fewer studies have been devoted to high-temperature applications (1000 K - 1200 K)despite their signi�cant potential in heavy industry and space exploration. Si-Ge alloys havehistorically been the most investigated material for high-temperature applications,9 but newmaterials exhibiting promising properties have emerged, such as the Zintl phase Yb14MSb11(M = Mg, Mn)10,11 or borides.12 However, to improve the industrial potential of TE gen-erators in this temperature range, new materials demonstrating superior performance andstability need to be discovered.Refractory rare earth tellurides crystallizing in the cubic Th3P4 structure-type are an-other class of promising high temperature TE materials. La3−xTe4, a material that canaccommodate large vacancies concentration on the La site, has long been known to showZT ≈ 1.1 at 1273 K for an optimized vacancy concentration of x = 0.23.13 Recently, theisostructural compounds Nd2.78Te4 14 and Pr2.74Te4 15 have been reported to have even higherZT values of 1.2 and 1.7 at 1273 K, respectively. These materials exhibit good TE properties3due to (i) the relatively low κ caused by the presence of vacancies, (ii) the rare-earth 4f or-bitals forming a heavy band near the Fermi level in the electronic density of state (DOS) andresulting in relatively high α and (iii) the strong correlation between the vacancy concen-tration x and the charge carrier concentration n allowing for easy optimization of the powerfactor PF = α2/ρ. The relationship between vacancy concentration and electronic propertiesof these materials can be intuitively understood by considering the charge balanced formu-lation RE3+3−xTe2−4 e−1−3x where excess of electrons (e−) are necessary to compensate vacanciesformation on the rare earth (RE) network, which results in n-type electronic conduction. In-terestingly, no p-type counterpart with similar properties has been discovered, which limitstheir integration in thermoelectric generators.Yb4Sb3 crystallizes in the anti-Th3P4 structure type, in which the anionic and cationicsub-lattices are reversed compared to the Th3P4 structure type, is good candidate showing p-type electrical conduction. Unlike rare earth tellurides, Yb4Sb3 is stochiometric and cannotaccommodate large vacancy concentration on any of its sub-lattices. Yb valence �uctuationis known to occur in this material, with most studies indicating that the great majority ofYb are in a valence state of 2+ as determined by magnetic measurements and X-ray absorp-tion spectroscopy.16,17 The charge balanced formula can therefore be written as Yb2+4 Sb3−3 h+which correctly predict an excess of positive charge carrier (h+) and therefore the p-typeconduction of this materials. Yb4Sb3 was reported to exhibit low resistivity in the order of10 µΩ m at 1273 K due to its elevated hole concentration in the order of 1021 cm−3 and apositive Seebeck coe�cient from 500 K that reaches about 80 µV K−1.18 However, despiteits potential, the thermal conductivity and ZT of pristine Yb4Sb3 has not been reported inthe literature to date. Nevertheless, studies have investigated the e�ect of doping on Yb4Sb3,where Yb was substituted with La,19 Sm,20 Ce or Eu,21 and Sb with Bi18 or I.21 However,it should be noted that the thermoelectric properties of most of these doped compoundswere only partially reported, and Hall measurements were rarely performed. Based on these4partial results, doping with these elements seems to decrease the charge carrier concentra-tion, as a simultaneous increase of α and ρ were observed, resulting in an enhanced PF .18,20The substitution also signi�cantly reduces thermal conductivity due to the induced atomicdisorder on the crystallographic sites.19 As a direct consequence, a �gure-of-merit ZT ashigh as 0.75 at 1273 K was reported, for example, for the composition La0.5Yb3.5Sb3.19 Thepromising thermoelectric performance of singly doped Yb4Sb3 compounds could potentiallybe improved further by double doping, which involves simultaneous substitutions on bothYb and Sb sublattices. This approach aims to increase the concentration of electron donorelements in the material while simultaneously increasing the atomic disorder to reduce ther-mal conductivity. However, the synthesis of Yb4Sb3 reported in literature was conductedusing a conventional high-temperature process in expensive Nb sealed tubes, which requiredseveral annealing steps at temperatures as high as 1323 K over a total of two weeks. Toenable industrial applications for this compound, more cost-e�cient and scalable synthesisprocesses need to be developed.This article presents the synthesis of undoped Yb4Sb3 and doubly doped CexYb4−xBi0.2Sb2.8(x = 0.4, 0.5) compounds through a ball-milling step followed by a reactive SPS. The solu-bility limit of Ce in the series CexYb4−xSb3 has been previously reported as x = 0.4, withthe Ce oxidation state determined to be +3 based on magnetic measurements.21 Therefore,Ce shows promises as a dopant for decreasing the charge carrier concentration according tothe Ce+3x Yb2+4−xSb3−3 h+1−x formalism. The Bi concentration of y = 0.2 was selected because ityields the highest properties in the Yb4BiySb3−y series.20 Thermoelectric properties of thesynthesized single-phased and fully-densi�ed materials were measured up to 1273 K and arepresented and discussed. Additionally, the high-temperature thermal expansion and oxida-tion behavior of pristine Yb4Sb3 are discussed, as these properties are crucial for industrialapplications. Eventually, preliminary results for the fabrication of a TE leg are presented,involving the deposition of Ni and Cu metallic contacts on Yb4Sb3 using a direct SPS process.52 Experimental part2.1 SynthesisPristine Yb4Sb3 and substituted CexYb4−xBi0.2Sb2.8 (x = 0.4, 0.5) samples were synthesizedfrom pure elements, Yb (ingot, 99.9 %, Alfa Aesar), Ce (ribbon, 99.9 %, Alfa Aesar), Sb(shot, 99.999 %, Sigma-Aldrich) and Bi (shot, 99.999 %, Sigma-Aldrich), using a two-stepprocess combining high-energy ball-milling and reactive Spark Plasma Sintering (SPS). Therare earth ingots were �rst cleaned with a �le inside a glove box (O2 < 1 ppm, H2O < 1 ppm)to removed the oxide layer at the surface. Stoichiometric amounts of metal precursors wereweighed, except for Yb which was added with a 2 % excess due to its tendency to adhere tothe milling jar's walls. About 3.5 g of metal precursor mixtures were loaded inside a 65 mLhardened steel grinding jar with two Ø = 12.6 mm hardened steel grinding balls (weighing7 g in total) under Ar atmosphere. The mixtures were ball-milled for 8 h using a SPEX8000M Mixer apparatus. The resulting powders, black in color and highly sensitive to air,were recovered within the glove box. The powders underwent reactive sintering using a Dr.Sinter Lab. SPS-322Lx apparatus in a Ø = 10 mm graphite die. The sintering processwas carried out at 1573 K for 10 min with a uniaxial pressure of 50 MPa under dynamicvacuum. Finally, the densi�ed pellets were wrapped in 0.05 mm thick Ta foil and annealedin evacuated silica tubes (10−3 mbar) for 24 h at 1273 K.Ni and Cu powders were utilized for the electrical contact fabrication on pristine Yb4Sb3.The Ni powder (particle size below 45 µm, 99.5 %, Sigma-Aldrich) and Cu powder (particlesize of 150 µm, 99.9 %, Sigma-Aldrich) were arranged in a layered succession with the ball-milled Yb-Sb powder mixture inside a Ø = 10 mm graphite die. The metal powder (250mg for Ni, 200 mg for Cu) was placed at the bottom and on top of the Yb-Sb mixture layer(1.5 g) before being cold-pressed at 50 MPa inside a glove-box. The whole set-up was thensintered under an uniaxial pressure of 50 MPa and under dynamic vacuum at 1373 K for620 min for Ni and 1173 K for 20 min for Cu. The sintering temperature was lowered inthe case of Cu to avoid melting the metal contact. Finally, the obtained Ni/Yb4Sb3/Ni andCu/Yb4Sb3/Cu pellets with thickness of about 2 mm were polished to remove the coveringgraphite foil and cut into 5×5 mm2 square-shaped pellets using a wire saw.2.2 Material characterizationsX-ray di�raction (XRD) analyses were conducted on powders obtained by grinding parts ofthe annealed SPS-sintered pellets using a mortar and pestle inside the glove-box. Room-temperature XRD patterns were carried out using a Rigaku SmartLab di�ractometer (Curadiation, λKα1 = 1.54056 Å and λKα2 = 1.5444 Å with an intensity ratio λKα2/λKα1 of 0.5)with scanning step size of 0.02◦. The high-temperature XRD measurement was performedwith a Rigaku SmartLab di�ractometer (Cu rotating target, monochromatized incident beamλKα1 = 1.54056 Å) equipped with a D/teX Ultra 250 detector. The patterns were measuredfrom 300 K to 900 K every 50 K under Ar �ow (1 L min−1, 5N purity) with a scanning stepsize of 0.02◦ and using variable slits. The powder was only exposed a brief time to air (<1 min) during the positioning of the Al2O3 sample holder inside the di�ractometer furnacechamber, which was then purged three times with Ar before the measurement. The X-raypowder patterns were �tted using the FullProf software.22 The peak shapes were modeledusing Thompson-Cox-Hastings pseudo-Voigt functions.23 Berar's factors were applied to theestimated standard deviation to obtain more realistic values.24 The scanning electron micro-scope (SEM) back-scattered electron images of the polished pellet surfaces were taken usinga Hitachi Tabletop Microscope TM3000. The SEM analyses of the TE leg cross-sections wereperformed using a Jeol 7200LV apparatus equipped with a Bruker X-Flash60 EDS detector.2.3 Thermoelectric properties measurementsThermal di�usivity (D) measurements were performed under N2 atmosphere (99.999 %)on Ø = 10 mm and 2 mm thick pellets coated with graphite using the laser �ash analysis7(LFA) apparatus Netzsch 467 HyperFlash. The thermal conductivity (κtot) was calculatedusing the expression κ = D Cp d with d the density and Cp the speci�c heat. The thermaldependence of Cp was estimated using a Netzsch Pyroceram reference during the thermaldi�usivity measurements. It was found to increase approximately linearly from 0.17 to0.20 J g−1 K−1 from 298 to 1270 K, respectively. d was determined by the Archimede methodin absolute ethanol. The lattice thermal conductivity κL was estimated by subtractingthe electronic contribution κE, calculated using the Wiedemann-Franz relationship κE =L T/ρ with L = 2.44 10−8 W Ω K−2, to κtot. The electrical resistivity (ρ) and the Hallcoe�cient (RH) were measured using the Van der Pauw method on a custom-built systemunder dynamic vacuum (10−5 mbar).25 The measurements were carried out on a Ø = 10 mmpellet shaped sample with thickness of about 1 mm. The charge carrier concentration (n)and mobility (µ) were calculated from the RH measurements realized under a magnetic �eldof 0.75 T using the relations (2) and (3), respectively.n =1RHe(2) µ =RHρ(3)with e the elemental charge of the electron. The Seebeck coe�cient (α) was measured with acustom-built instrument using the light pulse method under dynamic vacuum (10−5 mbar).26Scanning resistance measurement across the Yb4Sb3 leg with Cu and Ni metal contacts wereperformed at room temperature in air using a home-made apparatus described elsewhere.273 Results and discussionThe compound Yb4Sb3 was reported to melt congruently at 1813 K.28 Its synthesis usingconventional fusion-solidi�cation techniques is impractical due to the high vapor pressure ofSb at such elevated temperatures. In this study, an alternative approach using high-energyball milling was employed, o�ering the advantages of signi�cantly reducing the thermal treat-ment duration and being more scalable. The �nely divided powder mixture obtained after8the ball milling step exhibited high reactivity, allowing for simultaneous reaction and sinter-ing at a relative density exceeding 99 % in only 10 min at 1573 K. It should be noted thatafter sintering, the pellets displayed weak mechanical stability and are prone to breakageduring cutting or high-temperature measurements. To enhance their mechanical integrity,an additional annealing step at 1273 K for 24 h was implemented. The precise e�ect of thisannealing step is not yet fully understood, as the density, grain size, and chemical composi-tion of the samples were found to be similar before and after annealing. Another observationis the slow oxidation of the sintered samples when stored in air, indicated by the progressivedarkening of their initially shiny metallic surface within a few hours. To minimize oxidation,the samples discussed in this study were stored in a glove box, with exposure to air mini-mized as much as possible.Figure 1: SEM back-scattered electrons images of the polished surface of (a) Yb4Sb3, (b)Ce0.4Yb3.6Bi0.2Sb2.8 and (c) Ce0.5Yb3.5Bi0.2Sb2.8, (d) picture of a sintered pellet of Yb4Sb3after the annealing step showing a shiny metallic aspect9Table 1: Nominal compositions, experimental compositions determined by SEM-EDS,absolute and relative densities of the sintered samples Yb4Sb3, Ce0.4Yb3.6Bi0.2Sb2.8 andCe0.5Yb3.5Bi0.2Sb2.8Nominal composition SEM-EDS analyses Density (g cm−3) Relative density (%)Yb4Sb3 Yb4.10Sb2.90 8.64 99.6Ce0.4Yb3.6Bi0.2Sb2.8 Ce0.41Yb3.68Bi0.20Sb2.71 8.43 99.0Ce0.5Yb3.5Bi0.2Sb2.8 Ce0.46Yb3.64Bi0.18Sb2.72 8.28 98.2Figures 1 and Table 1 present SEM back-scattered electron images and SEM-EDS analy-ses conducted on the polished surface of the annealed sintered pellets, respectively. The ab-sence of chemical contrast on the surface of Yb4Sb3 and Ce0.4Yb3.6Bi0.2Sb2.8 indicates a goodchemical homogeneity. Only a low amount of porosity is visible on the surface which alignswith the high relative densities of the samples reported in Table 1. Most importantly, thechemical compositions determined by EDS analyses are in good agreement with the targetednominal compositions. However, the back-scattered electron image of Ce0.5Yb3.5Bi0.2Sb2.8reveals the presence of an impurity (dark gray area) that was identi�ed as CeSb by EDSanalysis. The formation of less dense CeSb (6.37 g cm−3) compared to Yb4Sb3 (8.67 g cm−3)can explain the lower relative density of this sample.Yb4Sb3 adopts the anti-Th3P4 structure type (I 4̄3d, space group no. 220) as illustratedin Figure 2(a). The unit cell is composed of 28 atoms, with Yb occupying the 16c Wycko�site and Sb occupying the 12a Wycko� site. Figures 2(b,c,d) shows the Rietveld re�nedXRD powder patterns measured on crushed sintered pellet after annealing and Table 2presents the obtained re�ned structural parameters. Consistent with the SEM analyses, thedi�raction patterns of Yb4Sb3 and Ce0.4Yb3.6Bi0.2Sb2.8 are single-phased with all di�ractionpeaks successfully indexed with the Th3P4 structure type. However, the di�raction patternof Ce0.5Yb3.5Bi0.2Sb2.8 shows additional small peaks corresponding to the CeSb impurity(Fm3̄m, space group no. 225), corroborating the �ndings from the SEM analyses. The re-�ned lattice parameter of a = 9.3228(2) Å for Yb4Sb3 agrees relatively well with a previous10Figure 2: (a) crystal structure of Yb4Sb3 and Rietveld re�ned XRD patterns of (b) Yb4Sb3,(c) Ce0.4Yb3.6Bi0.2Sb2.8 and (d) Ce0.5Yb3.5Bi0.2Sb2.8. The experimental data are plotted inred symbols, the calculated pattern with a black line and the di�erence with a blue line. Thevertical ticks indicate the theoretical Bragg positions of Yb4Sb3 (black) and CeSb impurity(green).study reporting a = 9.321(5) Å.18 Upon doping, the re�ned lattice parameter increased toa = 9.3997(2) Å and 9.4106(1) Å for Ce0.4Yb3.6Bi0.2Sb2.8 and Ce0.5Yb3.5Bi0.2Sb2.8, respec-tively. This increase is consistent with the larger atomic radii of the substituting elements,Ce (185 pm) and Bi (160 pm), compared to Yb (175 pm) and Sb (145 pm), respectively.29The presence of CeSb impurity in Ce0.5Yb3.5Bi0.2Sb2.8, as con�rmed by both XRD andSEM, and the limited increase in lattice parameter upon further Ce addition in Ce0.4Yb3.6Bi0.2Sb2.8suggests that the solubility limit of Ce in this system is slightly above x = 0.4. This �nding11Table 2: Structural parameters and R-factors obtained by Rietveld re�nement of Yb4Sb3,Ce0.4Yb3.6Bi0.2Sb2.8 and Ce0.5Yb3.5Bi0.2Sb2.8 XRD patterns. Single (*) and double stars (**)indicate �xed (non-re�ned) and constrained occupancy parameters (Occ. Yb = 1 - Occ. Ce,Occ. Sb = 1 - Occ. Bi), respectively.Yb4Sb3 Ce0.4Yb3.6Bi0.2Sb2.8 Ce0.5Yb3.5Bi0.2Sb2.8a (Å) 9.3228(2) 9.3997(2) 9.4106(1)Wycko� site 16c (x,x,x )x 0.0732(3) 0.0728(5) 0.0728(3)Occ. Ce - 0.1(1) 0.12(9)Occ. Yb 1* 0.9** 0.88**Biso (Å2) 0.5(1) 1.0(1) 0.6(1)Wycko� site 12a (3/8, 0, 1/4)Occ. Bi - 0.10(5) 0.11(3)Occ. Sb 1* 0.90** 0.89**Biso (Å2) 0.5(1) 0.9(1) 0.6(1)R-factorsχ2 4.45 5.67 4.62RBragg 2.02 2.85 2.15is consistent with previously reported values for the CexYb4−xSb3 series synthesized throughconventional high-temperature synthesis methods. It suggests that the substitution with Bion the Sb site has minimal e�ect on the solubility limit of Ce.21 It is important to notethat the anti-Th3P4 crystal structure features large cavities located at the 12b Wycko� site(7/8,0,1/4), which can accommodate interstitial atoms like Au in Au3Y3Sb4.30 However, inthis study, the structural model used, where all Ce and Bi atoms substitute Yb and Sbon their respective sites, shows good agreement with experimental data. The main sourceof discrepancy between the measured and calculated data (blue line on Figure 2(b,c,d)) isprimarily attributed to the asymmetric shape of the more intense peaks, which could not befully captured in the re�nements. The free atomic parameters of the Yb atoms on the 16csite, regardless of the chemical composition, were found to be similar and consistent with apreviously reported value of x = 0.074(1).28 The re�ned chemical compositions for the doped12compounds, Ce0.40(4)Yb3.60Bi0.3(1)Sb2.7 and Ce0.5(3)Yb3.50Bi0.3(1)Sb2.7, which were obtained byconstraining the 16c and 12a sites to full occupancy, are in good agreement with both thenominal compositions and the compositions determined by SEM-EDS analysesFigure 3: (a) High temperature XRD patterns of Yb3Sb4 measured up to 900 K and (b) thethermal evolution of the re�ned lattice parameter up to 600 KYb3Sb4 is aimed to be used for ultra high-temperature TE applications, making the char-acterization of its structure and stability under operating condition of great interest. Figure3(a) shows the high temperature XRD patterns of Yb3Sb4 collected at intervals of 50 K from300 to 900 K. The initial pattern at 300 K mostly show of Yb3Sb4 phase with contribution of13the Al2O3 sample support. Additional peaks in the 31-33◦ 2θ region appeared from 550 K,which are attributed to the low temperature phase of Yb5Sb3 (P63/mcm, space group no.193). From 650 K onwards, the decomposition of Yb4Sb3 accelerates with a signi�cant de-crease in peak intensities and the appearance of broad peaks associated with elemental Sb(R3̄m, space group no. 166) and Yb2O3 (Ia3̄, space group no. 206). From 800 K, the Yb4Sb3and Yb5Sb3 intermetallic phases completely disappeared and only Sb and Yb2O3 remain.The progressive decomposition of the Yb4Sb3 phase is attributed to oxidation caused byresidual oxygen absorbed at the surface of the sample particles or in experimental set-up.It demonstrates that this phase is extremely sensible to oxidation above 500 K even in anenvironment where air was thoroughly purged. While this extreme oxidation behavior poseschallenges for many terrestrial applications, it is not a primary concern for space applicationswhere oxygen is absent and high operating temperatures are required. Le Bail re�nementswere performed on the XRD patterns from 300 to 600 K, and the re�ned lattice param-eters are presented in Figure 3(b). The thermal evolution follows a linear trend allowingthe calculation of the linear thermal expansion parameter of 12.7(6)×10−6 K−1 at 300 K.The determination of this parameter holds great signi�cance for the realization of metalliccontacts, which will be discussed further below.The TE properties of the single-phased samples with nominal compositions Yb4Sb3 andCe0.4Yb3.6Bi0.2Sb2.8 were investigated from 298 to 1273 K. Figures 4(a) and 4(b) shows theelectrical resistivity and Seebeck coe�cient measurements, respectively. The pristine sampleexhibits a metallic behavior, with the electrical resistivity increasing from 2 to 12 µΩ m asthe temperature rises from 298 K to 1273 K. The Seebeck coe�cient displays a consistentupward trend, starting from a negative value of -20 µV K−1 at 298 K, indicating an n-typeelectrical behavior. It transitions to p-type behavior at around 500 K, eventually reaching amaximum of 85 µV K−1 at 1273 K. Such change of the electrical behavior from n to p-typewas already reported multiple times in the literature at similar temperatures.19,31 From Hall14Figure 4: High-temperature thermoelectric properties of Yb4Sb3 (black) andCe0.4Yb3.6Bi0.2Sb2.8 (red): (a) electrical resistivity, (b) Seebeck coe�cient, (c) powerfactor, (d) total (symbols) and lattice (solid lines) thermal conductivities, and (e) �gure-of-merit ZT.measurements conducted at 298 K, a charge carrier concentration of 3.0×1021 cm−3 and amobility of 9.5 cm2 V−1 cm−1 were determined. These values are in good agreement with aprevious study and are consistent with the observed metallic behavior of this material.16 TheSeebeck coe�cient of the present sample is signi�cantly higher at high temperatures than thepreviously reported value of 65 µV K−1 at 1273 K for pristine Yb4Sb3 synthesized throughconventional high-temperature synthesis.20 However, the electrical resistivity remains in ex-15cellent agreement across the entire temperature range. The exact reason for this discrepancyis not yet fully understood, but it could potentially be attributed to di�erences in microstruc-ture or undetectable contamination resulting from the utilization of di�erent synthesis meth-ods. Doping with Ce and Bi leads to an increase in both the electrical resistivity and theSeebeck coe�cient, resulting in maximum values of 15 mΩ m (+ 30 %) and 100 µV K−1 (+18 %) at 1273 K, respectively. This increase can be directly attributed to the decrease incharge carrier concentration to 1.8×1021 cm−3 as a result of Yb2+ substitution with Ce3+.The doping also has a slight e�ect on the charge carrier mobility, which decreases to 6.4 cm2V−1 cm−1. Despite the successful doping, the resulting power factors (PF ) shown in Figure4(c) remain comparable for the two samples with maximum value of 0.65 mW m−1 K−2 at1273 K. In comparison, the stoichiometric n-type counterparts materials La3Te4, Pr3Te4 andNb3Te4 only show slightly superior PF in the range 0.7 to 1.0 mW m−1 K−2 at the sametemperature.13�15 However, in these cases, the decrease in charge carrier concentration canbe successfully achieved by creating vacancies on the rare earth site, resulting in stronglyenhanced PF values greater than 1.5 mW m−1 K−2.15 For instance, a comparable reductionin charge carrier concentration from 4.0 to 2.0×1021 cm−3 by going from the compositionsNb2.9Te4 to Nb2.86Te4 leads to a signi�cant 40 % increase of the PF .14 The optimization ofcharge carrier concentration is essential to achieve substantial boost of the TE properties ofthe n-type rare-earth tellurides but it does not appear to hold for the present p-type Yb4Sb3.Figure 4(b) presents the measured total (κtot) and the estimated lattice (κL) thermal con-ductivities for the two compositions. The undoped sample Yb4Sb3, which was measured forthe �rst time to the best of our knowledge, exhibits κtot values going from 4.2 W m−1 K−1 at298 K to 3.8 W m−1 K−1 and 1273 K. Similar weak temperature dependence is observed forstoichiometric n-type compounds Pr3Te4 and Nb3Te4, with slightly lower κtot values around3.5 W m−1 K−1.14,15 The estimated charge carrier contribution (κE) was found to representabout 70 % of κtot which is consistent with the high charge carrier concentration of this com-16pound. The estimated κL remains constant at approximately 1.2 W m−1 K−1 throughout thetemperature range. This relatively low κL values can be attributed to the complex crystalstructure and the presence of heavy constituting elements in Yb4Sb3. Upon doping, κtot issigni�cantly reduced by more than 40 %, reaching values of about 2.2 W m−1 K−1 acrossthe entire temperature range. The decrease in κtot is partially attributed to the reduction ofκE consecutive to the reduce charge carrier concentration in the doped compound. Further-more, κL is signi�cantly reduced by about 50 % to reach approximately 0.6 W m−1 K−1 overthe entire temperature range. This reduction of κL is directly linked to the introduction ofatomic disorder in the structure due to partial substitution on both crystallographic sites.Such a signi�cant reduction in κL is generally advantageous for improving the thermoelectricperformance of materials since it is the only parameter in the �gure of merit ZT that is notcorrelated with the charge carrier concentration.Finally, Figure 4(e) displays the �gure-of-merit ZT values for both the pristine and dopedsamples. At temperatures below 800 K, ZT is small for both samples due to the low absoluteSeebeck coe�cients. However, ZT increases rapidly at higher temperatures. For the �rsttime to our knowledge, a maximum ZT of 0.21 is achieved for the pristine sample at 1273 K.As a direct consequence of the induced structural disorder, Ce0.4Yb3.6Bi0.2Sb2.8 exhibits asigni�cantly improved ZT of 0.39 (+85 %), primarily due to the substantial reduction inthermal conductivity. The measured ZT values obtained for the investigated compositions,although signi�cant, remained lower than the reported values of approximately 0.75 forLa0.5Yb3.5Sb3 and 0.65 for Sm0.4Yb3.6Sb3.19 This di�erence in properties are attributed tothe higher Seebeck coe�cient observed in these compositions, which reached 120 µV K−1 at1273 K. Further investigations are required to gain a better understanding of the in�uenceof the various doping elements and synthesis processes on the Seebeck coe�cient of thesematerials. It is worth noting that during the TE properties measurements at the highesttemperatures, the surface of the samples undergoes signi�cant oxidation. This oxidation is17evident in Figure 4(f), which shows the initial shiny metallic surface turning black after themeasurement. The oxidation process poses challenges in accurately measuring the transportproperties due to the deterioration of electrical contact quality. This is believed to be thecause of the slight deviations observed in the electrical resistivity data of the doped sampleabove 1100 K in Figure 4(a). However, it should be emphasized that the oxidized layer iscon�ned to the surface of the pellet and can be e�ectively removed by polishing after themeasurement, restoring the sample to its original condition.Figure 5: SEM back-scattered electron images of the contact interface and correspondingEDS mappings. The images (a,b) and (c,d) correspond to the legs with Ni and Cu contacts,respectively. EDS compositions of the phases indicated with numbers in (a) and (c) aregiven in Table 3The deposition of metallic contacts on TE materials is crucial for their successful inte-gration into functional devices. This task presents several challenges as the contacts needto possess strong adhesion, low electrical and thermal interface resistances, and compatiblethermal expansion properties with the TE material to prevent crack formation during synthe-18Table 3: SEM-EDS analyses of the di�erent phases found at the contact interface of the Niand Cu legs. The analyses numbers correspond to the areas indicated in Figure 5(a,c)Analysis number Ni (at.%) Cu (at.%) Yb (at.%) Sb (at.%) Attributed phase(1) 98 - 1 1 Ni(2) 69 - 21 10 Ni2Yb1−xSbx(3) 39 - 33 28 NiYbSb(4) 22 - 45 33 Ni1−xYbSb(5) < 1 - 60 40 Yb4Sb3(6) - 94 3 3 Cu(7) - < 1 60 40 Yb4Sb3sis and operation at high temperatures. To address these requirements, Ni (13.4×10−6 K−1)and Cu (16.5×10−6 K−1) were chosen due to their relatively close match with the thermalexpansion parameter of Yb4Sb3 (12.7×10−6 K−1) determined earlier by XRD. Figures 5(a,b)and 5(c,d) display representative SEM back-scattered electron images and correspondingEDS mappings of the contact interfaces for Ni and Cu, respectively.The elemental composition of the interface phases were determined by EDS analyses, andthe results are presented in Table 3. In both cases, the contact interfaces exhibited no cracksor excessive porosity which is consistent with the good mechanical properties of the legs. Forthe Ni contacts shown in Figure 5(a), two distinct phases (2) and (3) were observed at theinterface between Ni and Yb4Sb3, which correspond to the phases (1) and (5), respectively.The EDS mapping revealed that both phases are composed of the three elements, with adecrease in Ni concentration from the metal contact to the TE phase. Composition of thephase (2) suggests that it may be the Laves phase Ni2Yb (MgCu2 structure-type, Fd3̄m,space group no. 227), potentially with some Sb substitution on the Yb site. However, tothe best of our knowledge, such substitution has not been reported in the literature, and noNi2Sb phase is reported in the Ni-Sb phase diagram. The phase (3) is identi�ed as the half-Heusler NiYbSb (F 4̄3m, space group no. 216), which is the only known ternary compoundin this system. Additionally, the presence of darker spots with sizes below approximately195 µm, visible within the Yb4Sb3 matrix, corresponds to phase (5). This phase, composed ofthe three elements, exhibits a relatively lower Ni content compared to phase (4). It may beattributed to Ni-de�cient Ni1−xYbSb as already reported by R. Mishra et al.32 However, itshould be noted that the EDS measurement accuracy for this phase may be compromiseddue to the relatively small size of the analyzed area, which can include some part of theYb4Sb3 matrix owing to the interaction volume of the electron beam (a few µm). In the caseof Cu contacts shown in 5(c,d), no formation of new phases was observed at the interfacebetween Cu and Yb4Sb3, which corresponds to phases (6) and (7), respectively. The EDSmapping clearly indicates a sharp separation of the elements at the interphase, indicatingthe absence of signi�cant interdi�usion. The formation of several new phases at the interfacewith Ni, which are absent for Cu, are expected to contribute to increase the electrical andthermal contact resistances of the TE legs.Figure 6: Room temperature scanning electrical resistance measurements of the (a) Ni and(b) Cu legs. The pictures in inset shows the samples used for the measurements.The electrical contact resistivities measured at room temperature for the Ni and Cu legsare shown in Figures 6(a) and 6(b), respectively. The contact resistivity was estimated fromthe jump of the electrical resistance occurring between the metal contact and TE materials.For the Ni leg, contact resistivities of 2.8 and 1.4 µΩ cm2 were measured, while for theCu leg, contact resistivities of 1.1 and 0.5 µΩ cm2 were obtained. These values, ranging20in the order of a few µΩ cm2, indicate good electrical contact quality. The higher contactresistivities observed for the Ni leg compared to the Cu leg are consistent with the formationof secondary phases observed in the SEM analyses. The electrical resistivity of Yb4Sb3 inthe legs was estimated to be 2.2 and 2.5 µΩ m based on the slope of the linear section.These values are of the same order of magnitude as the room temperature value of 2.4 µΩshown in Figure 4(a). The slight deviation is attributed to the density change resulting fromdi�erent sintering temperatures used for the leg synthesis. This con�rms that there is nosigni�cant di�usion of the metal contact into the TE material, which would otherwise a�ectits electronic properties.334 ConclusionSingle phased and fully densi�ed Yb4Sb3 materials were successfully synthesized par ballmilling followed by reactive SPS. However, the obtained samples encountered two chal-lenges. Firstly, they exhibited mechanical fragility, necessitating annealing at 1273 K for24 hours. Secondly, they demonstrated high sensitivity to oxidation above 550 K, even ina purged protective atmosphere, which posed di�culties in measuring the high-temperatureTE properties. A new series of doubly substituted CexYb4−xBi0.2Sb2.8 samples were syn-thesized. Despite the substitution with Bi on the Sb site, XRD and SEM analyses revealedan unchanged Ce solubility limit of x = 0.4, consistent with the previously reported singlydoped CexYb4−xSb3 series. The substitution of Yb2+ with Ce3+ e�ectively reduced thecharge carrier concentration of Ce0.4Yb3.6Bi0.2Sb2.8 resulting in an slightly increased Seebeckcoe�cient compared to pristine Yb4Sb3. However, this improvement was accompanied byan increase in electrical resistivity, leading to comparable maximum PF of approximately0.4 W m−1 K−2 at 1273 K for both samples. Nonetheless, the introduction of atomic disor-der through the double substitution signi�cantly decreased the lattice thermal conductivity,resulting in an improved ZT of 0.4 (+ 50 %) at 1273 K. Based on the determined linear21thermal expansion coe�cient of 12.7×10−6 K−1 from high-temperature XRD measurements,Cu and Ni were selected as contact metals for fabricating TE legs using SPS. Low averagedelectric contact resistances of 2 and 1 µΩ cm2 were measured for Ni and Cu, respectively.These promising results highlight the potential of these materials for fabricating TE devicesintended for high-temperature applications.AcknowledgementThe authors thank the Agence Nationale de la Recherche (ANR) for the funding ANR-18-CE05-0037. Xavier Larose (CRISMAT laboratory) is acknowledged for SEM images andEDS analyses performed on the thermoelectric legs.References(1) El Oualid, S.; Kosior, F.; Dauscher, A.; Candol�, C.; Span, G.; Mehmedovic, E.;Paris, J.; Lenoir, B. Innovative design of bismuth-telluride-based thermoelectric micro-generators with high output power. Energy Environ. Sci. 2020, 13, 3579�3591.(2) Liu, Z.; Gao, W.; Oshima, H.; Nagase, K.; Lee, C.-H.; Mori, T. Maximizing the perfor-mance of n-type Mg3Bi2 based materials for room-temperature power generation andthermoelectric cooling. Nat. Commun. 2022, 13, 1120.(3) Jood, P.; Male, J. P.; Anand, S.; Matsushita, Y.; Takagiwa, Y.; Kanatzidis, M. G.;Snyder, G. J.; Ohta, M. Na Doping in PbTe: Solubility, Band Convergence, PhaseBoundary Mapping, and Thermoelectric Properties. J. Am. Chem. Soc. 2020, 142,15464�15475.(4) Wang, X.; Veremchuk, I.; Burkhardt, U.; Bobnar, M.; Böttner, H.; Kuo, C.-Y.;Chen, C.-T.; Chang, C.-F.; Zhao, J.-T.; Grin, Y. Thermoelectric stability of Eu- andNa-substituted PbTe. J. Mater. Chem. C 2018, 6, 9482�9493.22(5) El Oualid, S.; Kogut, I.; Benyahia, M.; Geczi, E.; Kruck, U.; Kosior, F.; Masschelein, P.;Candol�, C.; Dauscher, A.; Koenig, J. D.; Jacquot, A.; Caillat, T.; Alleno, E.; Lenoir, B.High Power Density Thermoelectric Generators with Skutterudites. Adv. Energ. Mater.2021, 11, 2100580.(6) Le Tonquesse, S.; Alleno, E.; Demange, V.; Prestipino, C.; Rouleau, O.; Pasturel, M.Reaction mechanism and thermoelectric properties of In022Co4Sb12 prepared by mag-nesiothermy. Mater. Today Chem. 2020, 16, 100223.(7) Le Tonquesse, S.; Joanny, L.; Guo, Q.; Elkaim, E.; Demange, V.; Berthebaud, D.;Mori, T.; Pasturel, M.; Prestipino, C. In�uence of Stoichiometry and Aging at Operat-ing Temperature on Thermoelectric Higher Manganese Silicides. Chem. Mater. 2020,32, 10601�10609.(8) Le Tonquesse, S.; Zhang, W.; Srinivasan, B.; Fontaine, B.; Hiroto, T.; Mori, T.;Halet, J.-F.; Berthebaud, D.; Suzuki, T. S. Improvement of Thermoelectric Proper-ties via Texturation Using a Magnetic Slip Casting Process-The Illustrative Case ofCrSi2. Chemistry of Materials 2022, 34, 1143�1156.(9) Scoville, N.; Bajgar, C.; Rolfe, J.; Fleurial, J.-P.; Vandersande, J. Thermal conductivityreduction in SiGe alloys by the addition of nanophase particles. Nanostruc. Mater.1995, 5, 207�223.(10) Brown, S. R.; Kauzlarich, S. M.; Gascoin, F.; Snyder, G. J. Yb14MnSb11: New High Ef-�ciency Thermoelectric Material for Power Generation. Chem. Mater. 2006, 18, 1873�1877.(11) Perez, C. J.; Cerretti, G.; Wille, E. L. K.; Devlin, K. P.; Grewal, N. S.; Justl, A. P.;Wood, M.; Bux, S. K.; Kauzlarich, S. M. Evolution of Thermoelectric Properties in theTriple Cation Zintl Phase: Yb13˘xCaxBaMgSb11 (x = 1-6). Chem. Mater. 2021, 33,8059�8069.23(12) Sauerschnig, P.; Tsuchiya, K.; Tanaka, T.; Michiue, Y.; Sologub, O.; Yin, S.;Yoshikawa, A.; Shishido, T.; Mori, T. On the thermoelectric and magnetic properties,hardness, and crystal structure of the higher boride YbB66. J. Alloys Compd. 2020,813, 152182.(13) May, A. F.; Fleurial, J.-P.; Snyder, G. J. Thermoelectric performance of lanthanumtelluride produced via mechanical alloying. Phys. Rev. B 2008, 78, 125205.(14) Gomez, J., Steven; Cheikh, D.; Vo, T.; Von Allmen, P.; Lee, K.; Wood, M.; Sny-der, G. J.; Dunn, B. S.; Fleurial, J.-P.; Bux, S. K. Synthesis and Characterization ofVacancy-Doped Neodymium Telluride for Thermoelectric Applications. Chem. Mater.2019, 31, 4460�4468.(15) Cheikh, D.; Hogan, B. E.; Vo, T.; Von Allmen, P.; Lee, K.; Smiadak, D. M.; Ze-valkink, A.; Dunn, B. S.; Fleurial, J.-P.; Bux, S. K. Praseodymium Telluride: A High-Temperature, High-ZT Thermoelectric Material. Joule 2018, 2, 698�709.(16) Ochiai, A.; Nakai, S.; Oyamada, A.; Suzuki, T.; Kasuya, T. Valence �uctuation in anti-Th3P4 structure Yb4Sb3 and Yb4Bi3. J. Magn. Magn. Mater. 1985, 47-48, 570�572.(17) Aoki, H.; Ochiai, A.; Suzuki, T.; Helfrich, R.; Steglich, F. Systematic study in mixedvalence system Yb4(As1−xSbx)3. Physica B 1997, 230-232, 698�701.(18) Chamoire, A.; Viennois, R.; Tedenac, J.-C.; Koza, M. M.; Gascoin, F. Antimony-BasedCompounds with the Anti-Th3P4 Structure as Potential High-Temperature Thermo-electric Materials. J. Electron. Mater. 2011, 40, 1171�1175.(19) Chamoire, A.; Gascoin, F.; Estournès, C.; Caillat, T.; Tédenac, J.-C. High-temperaturetransport properties of complex antimonides with Anti-Th3P4 structure. Dalton Trans.2010, 39, 1118�1123.24(20) Chamoire, A.; Gascoin, F.; Estournès, C.; Caillat, T.; Tédenac, J.-C. High-TemperatureTransport Properties of Yb4−xSmxSb3. J. Electron. Mater. 2010, 39, 1579�1582.(21) Chamoire, A. Antimoniures Complexes de Type Th3P4 et Propriétés Thermoélec-triques. Université Montpellier II - Sciences et Techniques du Languedoc 2009, (PhDthesis written in french).(22) Rodriguez-Carvajal, J. Recent advances in magnetic-structure determination by neu-tron powder di�raction. Physica B 1993, 192, 55�69.(23) Thompson, P.; Cox, D. E.; Hastings, J. B. Rietveld re�nement of Debye-Scherrer syn-chrotron X-ray data from Al2O3. J. Appl. Cryst. 1987, 20, 79�83.(24) Bérar, J.-F.; Lelann, P. E.s.d.'s and estimated probable error obtained in Rietveldre�nements with local correlations. J. Appl. Cryst. 1991, 24, 1�5.(25) Borup, K. A.; Toberer, E. S.; Zoltan, L. D.; Nakatsukasa, G.; Errico, M.; Fleurial, J.-P.; Iversen, B. B.; Snyder, G. J. Measurement of the electrical resistivity and Hallcoe�cient at high temperatures. Rev. Sci. Instrum. 2012, 83, 123902.(26) Wood, C.; Zoltan, D.; Stapfer, G. Measurement of Seebeck coe�cient using a lightpulse. Rev. Sci. Instrum. 1985, 56, 719�722.(27) Mengali, O.; Seiler, M. Contact resistance studies on thermoelectric materials. Adv.Energy Convers. 1962, 2, 59�68.(28) Bodnar, R. E.; Stein�nk, H. The Phase Equilibria and Crystal Chemistry of the Inter-mediate Phases in the Ytterbium-Antimony System. Inorg. Chem. 1967, 6, 327�330.(29) Slater, J. C. Atomic Radii in Crystals. J. Chem. Phys. 1964, 41, 3199�3205.(30) Dwight, A. E. Yttrium-gold-antimony Y3Au3Sb4. Acta Cryst. 1977, B33, 1579�1581.25(31) Bucher, E.; Cooper, A.; Jaccard, D.; Sierro, J. Valence Instabilities and Related NarrowBand Gap Phenomena. P.D. Park (Plenum Press, New York) 1977, 529.(32) Mishraa, R.; Pöttgen, R.; Ho�mann, R.-D.; Fickenscher, T.; Eschen, M.; Trill, H.;Mosel, B. D. Ternary Antimonides YbTSb (T = Ni, Pd, Pt, Cu, Ag, Au) � Synthesis,Structure, Homogeneity Ranges, and 121Sb Mössbauer Spectroscopy. Z. Naturforsch. B2002, 57, 1215�1223.(33) Le Tonquesse, S.; Matsushita, Y.; Jood, P.; Ohta, M.; Mori, T.; Berthebaud, D. Fabri-cation and Evaluation of Low-Cost CrSi2 Thermoelectric Legs. Crystals 2021, 11 .26TOC Graphic27