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[Small Structures--Bonding Heterogeneity and Nanoprecipitation on Substituting the Anionic Framework in Mg3Sb2 for p-type Zintl Thermoelectrics.pdf](https://mdr.nims.go.jp/filesets/9a49c972-f5fe-4fc1-8d13-4d143ebd6351/download)

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[Nagendra Singh Chauhan](https://orcid.org/0000-0003-2579-6642), [Takao Mori](https://orcid.org/0000-0003-2682-1846)

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[Bonding Heterogeneity and Nanoprecipitation on Substituting the Anionic Framework in Mg<sub>3</sub>Sb<sub>2</sub> for p‐Type Zintl Thermoelectrics](https://mdr.nims.go.jp/datasets/386e1465-66e0-4ab0-b155-e4cde95e2fac)

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Bonding Heterogeneity and Nanoprecipitation on Substituting the Anionic Framework in Mg3Sb2 for p‐Type Zintl ThermoelectricsBonding Heterogeneity and Nanoprecipitation onSubstituting the Anionic Framework in Mg3Sb2for p-Type Zintl ThermoelectricsNagendra Singh Chauhan and Takao Mori*1. IntroductionLow-grade waste heat recovery by thermoelectrics (TE) hasgarnered renewed attention in recent years, due to the out-standing n-type performance obtained in Mg3(Sb, Bi)2 solidsolutions.[1–6] To complement this exceptional performance atthe device level, ongoing efforts are aimedtoward developing their low-cost andhigh-performance p-type counterpart,besides further optimization of their n-typeTE figure of merit (zT ) near roomtemperature.[7–9] Intrinsically, undopedMg3Sb2 compound exhibits p-type con-duction with a low hole concentration(n)� 1017 cm�3 and a multivalley (Nv� 6)conduction band behavior near theFermi level.[10,11] Upon heavy Bi alloying(�25 at%) and subsequent doping by�1 at% with chalcogen group elements(i.e., Te, Se, S), a higher n� 1019 cm�3for n-type conduction was attained, leadingto maximization of zT values rangingbetween 1.5 and 2 at �723 K for the opti-mized compositions.[1–4] Such high zT iscomparable to, and even surpasses, manystate-of-the-art n-type TE materials, leadingto an impressive device performance with aconversion efficiency (η≥ 7%) under a300 K temperature difference (ΔT ).[12,13]Thus, the anionic framework constitutinga covalently bonded network formed bycomplex anions or metalloids plays a vital role in sustainingthe “electron–crystal” electronic structure in doped Mg3Sb2 crys-tals, which also exhibit an inherently low lattice thermal conduc-tivity (κL) �1Wm�1 K�1 near room temperature.[14–17]As a prominent and widely explored zintl phase, Mg3Sb2 ispart of the broader family of zintl antimonides, which encom-passes diverse crystal structures and compositions, though onlya small subset has been investigated for their TE properties.[18]On crystal structure basis, Mg3(Sb, Bi)2 and its mixed-anionderivatives are classified as 1–2–2 zintl family that exhibitsCaAl2Si2-type structure (Space Group: P3m1, 164)[18–22] with anasymmetric unit cell having complicated bonds. Interestingly,Mg atoms reside in both the octahedrally coordinated interlayerMg(I) site (i.e., the cation site) and the tetrahedrally coordinatedintralayer Mg(II) site (i.e., within the anionic [Mg2Sb2]2� frame-work). The chemical bonding analysis of α-Mg3Sb2 indicatescoordination flexibility of Mg─Sb bonds having both ionicand covalent character, wherein covalent bonding in the[Mg2Sb2]2� anionic framework enables high carrier mobility,while the ionic Mg2þ cations offer better dopability for carrierconcentration optimization.[10,11] The site preference and occu-pancy patterns during the alloying/doping indicate that Mg(I)sites tend to be preferentially occupied by highly electropositiveN. S. Chauhan, T. MoriResearch Center for Materials Nanoarchitectonics (MANA)National Institute for Materials Science (NIMS)Namiki 1-1, Tsukuba 305-0044, JapanE-mail: mori.takao@nims.go.jpT. MoriGraduate School of Pure and Applied SciencesUniversity of Tsukuba1-1-1 Tennodai, Tsukuba, Ibaraki 305-8573, JapanThe ORCID identification number(s) for the author(s) of this articlecan be found under https://doi.org/10.1002/sstr.202400632.© 2025 The Author(s). Small Structures published by Wiley-VCH GmbH.This is an open access article under the terms of the Creative CommonsAttribution License, which permits use, distribution and reproduction inany medium, provided the original work is properly cited.DOI: 10.1002/sstr.202400632Designing zintl compounds with complex crystal structures and large unit cellscontaining heavy elements may inherit bonding heterogeneity-induced latticeanharmonicity, leading to intrinsically low thermal conductivity. In this work,alloying-induced bonding heterogeneity in the extensively explored α-Mg3(Sb, Bi)2 phase due to local atomic ordering, site preferences, and prevailingheterogenous interfaces is evaluated in the p-type Mg3(Sb1�2xBixSnx)2-basedpolyanionic nanocomposites for different alloying concentrations. The inherentsusceptibility for partial phase transition (trigonal!monoclinic) is observedupon alloying, which is driven by alterations in bonding patterns, localizeddistortion, and secondary phase formation. At low alloying content (x≤ 0.05), atrigonal (Sb, Sn) phase is observed, while for higher alloying content (x≥ 0.1), acubic Mg2Sn nanophase emerges. A synergistic reduction in thermal conductivityand enhanced power factor maximize the zT� 0.25(�0.05) at 673 K in theoptimized p-type Mg3(Sb0.9Bi0.025Sn0.025)2 nanocomposites. The study highlightsbonding heterogeneity-induced structural transitions as an inherent challenge,where the dominant role of anionic sites becomes pivotal for determining/deriving favorable structural and functional properties in Mg3(Sb, Bi)2-based zintlcompounds.RESEARCH ARTICLEwww.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (1 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbHmailto:mori.takao@nims.go.jphttps://doi.org/10.1002/sstr.202400632http://creativecommons.org/licenses/by/4.0/http://creativecommons.org/licenses/by/4.0/http://www.small-structures.comalkaline earth metals and lanthanides (such as Ca, Sr, Ba, La, andYb), while Mg(II) sites exhibit a higher affinity for alloying withmore electronegative metals like Zn, Mn, and Cd.[21]Among its anionic solid solutions, Mg3Sb2�xBix and its deriv-atives exhibit an excellent bandgap tunability which has allowedoptimization of n-type properties by reducing the bandgap toenable efficient doping efficiency with Te/Se/S as dopants.[1–4]However, for p-type conduction, anionic doping strategies(i.e., substituting dopants on the Sb/Bi sites) in Mg3(Sb, Bi)2and its related solid solutions have remained largelyineffective,[23–26] primarily due to its weak valence band degen-eracy (Nv= 1).[27] Nevertheless, codoping approaches (Zn, Na,Yb, Ag, and/or Cu)[28–32] primarily targeting the cationic Mg(I)and/or anionic Mg(II) sites have remained effective in regulatingand optimizing holes concentration, reaching an optimaln� 1019 cm�3 to attain a maximum zT� 0.8 at around773 K.[28–33] From a practical perspective, enhancing the TE per-formance of p-type Mg3(Sb, Bi)2 to match its n-type counterpartstill remains critical and an ongoing challenge for realization ofthe α-Mg3Sb2-based device, with a maximum conversion effi-ciency (η) �5.5% at the hot-side temperature of 573 K, reportedtill date.[34]Complex multinary zintl compounds constituting a polya-nionic (ternary, quaternary, and higher) framework can offernew basis for deriving novel functionalities and understandingof the poor dopability of p-type Mg3(Sb, Bi)2. We reportMg3(Sb1�2xBixSnx)2-based polyanionic nanocomposites withbonding heterogeneity-induced locally distorted structures andnanoprecipitation, synthesized employing ball milling and sparkplasma sintering. The preferences for specific site occupancies of(Bi, Sn) substitutions in Mg3Sb2 were assessed employing a(3þ 1) dimensional superspace approach, which provided criti-cal insights into the atomic-scale interactions and local atomicordering in the constituting α-Mg3Sb2 phase. Partial phasetransition (trigonal!monoclinic), nanoprecipitation of trigonal(Sb, Sn, Bi) phase for low alloying content x� 0.05, and Mg2Snnanoprecipitation for increasing alloying content x> 0.05, wasobserved implying locally distorted structures. The present studyhighlights the inherent susceptibility of Mg3(Sb, Bi)2-based poly-anionic solid solutions for bonding heterogeneity, structuraltransitions, and secondary phase formations and their implica-tions on thermal and electrical transport properties. It also dem-onstrates how alloying strategies require vital consideration ofstructural and bonding modifications to effectively attain betterdopability and targeted functionalization of transport propertiesof zintl compounds.2. Results and Discussion2.1. Phase Compositions, Phase Segregation, and HierarchicalMicrostructureThe powder diffraction X-ray diffraction (PXRD) patterns of rep-resentative Mg3(Sb1�2xBixSnx)2 nanocomposite samples areshown in Figure 1a, which indicate major peaks correspondingto trigonal α-Mg3Sb2 phase (Space Group : P3m1, 164) in all thealloyed compositions. For low alloying content x� 0.05, addi-tional peaks corresponding to the trigonal (Sb, Sn, Bi) phase(Space Group: R3m, 166) were also indexed and indicated by star(*), while for higher alloying content x≥ 0.1, peaks correspond-ing to cubic Mg2Sn phase (Space Group: Fm3m, 225), as indi-cated by inverted triangle ( ), appear, implying α-Mg3Sb2 phaseinstability and secondary phase precipitation process duringsynthesis. With increasing equiatomic (Sn, Bi) content, that is,x, the major and overlapping (101) and (011) PXRD peaks shiftto lower angles (Figure 1b) when compared with synthesizedMg3Sb1.9Sn0.1, suggesting an expansion of the α-Mg3Sb2 lattice.In addition to lattice expansion, shifts in major overlapping peakscan be attributed to secondary-phase formation, potentially alter-ing the Sb/Bi ratio in the synthesized nanocomposites. Beyondx≥ 0.2, additional peaks corresponding to tetragonal Sn phase(Space Group: I41/amd 141) indicated by an empty inverted tri-angle ( ) were also observed. Samples with higher alloying con-tent are also more chemically stable, during unavoidableexposure to atmospheric conditions during XRD measurements,as indicated by absence of MgO impurity peaks, which is distinc-tively observed for x� 0.05 and indicated by filled circles. Thepresence of extra peaks corresponding to secondary phases evenin synthesized Mg3Sb1.9Sn0.1 is indicative of the discordantnature of Sn/Sb atoms in Mg─Sb─Sn-based compositions,wherein secondary-phase precipitation was also observed.[35–37]Notably, the synthesized Mg3(Sb1�2xBixSnx)2 compositions liewell within the solid solubility limits of Bi (i.e., x� 0.4)[23] forMg3Sb2�xBix and agree well with the previous structural obser-vations reported for Mg3Sb2�xSnx polycrystals.[36]To understand phase stability, potential phase transforma-tions, and optimal compositions in the synthesizedMg3(Sb1�2xBixSnx)2 nanocomposites, ternary phase diagram ofMg─Sb─Sn at room temperature (an isothermal section)[35] isoverlaid over the evaluated modified Gibbs free energy of mixingand is shown in Figure 1c. Previous investigation has revealedthe observed difficulty in synthesizing ternary Mg─Sb─Sncompounds,[35–38] wherein binary α-Mg3Sb2 compound was eval-uated[35] to coexist with Mg2Sn, Sn, and SnSb in phase equilibria.The Gibbs free energy of mixing (ΔGm) landscape influences theextent of solid solution formation, where the lowest ΔGm impliesthermodynamic stability, in regions corresponding to the synthe-sized compositions. The secondary-electron scanning electronmicroscopy images presented in Figure 1d display backscatteredelectron (BSE) micrograph of characteristic fracture surface mor-phology i a representative sample, indicative of transgranularfracture. This fracture mode is characterized by crack propaga-tion occurring through the grains rather than along grain bound-aries. The fracture surfaces reveal prominent cleavage facets withaligned lamellae patterns, which are characteristic of plasticdeformation within individual grains. The nanoprecipitationpromotes local ordering, contributing to the development oflayered structures that accommodate anionic substructures.Consequently, the material’s response to stress likely involvesboth grain-level plasticity and substructural adjustments, leadingto the observed layered morphology.[39]The presence of nanoinclusions within the microstructure canbe attributed to localized compositional variations caused by thediscordant Sn atoms, which nucleate and segregate from thesupersaturated Mg3(Sb, Bi)2 solid solution as indicated by ele-mental mapping shown in Figure 1e. The energy-dispersivewww.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (2 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comX-ray (EDX) spectrum shown in Figure 1f, confirms the existenceof an extended solid solution comprising Mg2Sn, Bi, and SnSbwithin the Mg3Sb2 matrix, agreeing well with the previouslyreported phase equilibria studies in the Mg─Sb─Sn system.[35]Thus, a complex interplay between the different phases in theheterogeneous microstructure with nanoscale features of thesynthesized nanocomposites is indicative of the structural differ-ences between Mg2Sn and Mg3Sb2 phases that preclude the for-mation of a complete solid solution agreeing well with previousobservations.[37,38]2.2. Structural Parameters and Monoclinic Phase TransitionsSubstitutions at the anionic framework in zintl phases are highlysusceptible to causing structural distortions, which may result inlowering of symmetry, changes in bonding patterns, andformation of monoclinic variants.[18,21,22,40] These distortionsare often driven by electronic factors or size differences betweensubstituted atoms.[26] In the synthesized Mg3(Sb1�2xBixSnx)2nanocomposites, an extensive structural disorder with a partialloss of trigonal symmetry leading to monoclinic phase transi-tions (Space Group: C2/m, 12) was also indexed during refine-ment, implying a locally distorted configuration and chemicaldisordering in the synthesized polycrystals. It is noteworthy thatthe observance of monoclinic structure in the synthesized poly-crystals as a highly distorted variant owes structural similarity tothe high-pressure monoclinic structure for Mg3Sb2 and Mg3Bi2reported previously at critical transition at above 7.8 and 4.0 GPa,respectively.[41] However, the synthesized polycrystals exhibit anirreversible and co-existing monoclinic phase as shown inFigure 2a, in contrast to the pressure-induced displacive andreversible (trigonal ↔ monoclinic) phase transition, reportedpreviously.[41] The refinements of the PXRD evaluates the mono-clinic C2/m unit cell (formula unit= 6) with a volume performula unit of �500 Å3.The refined lattice parameter of the α-Mg3Sb2 phases, asshown in Figure 2b, obtained by Le Bail refinement, revealsan initial decrease in lattice parameters a and c for lower alloyingcontent, followed by a consistent increase with higher alloyingFigure 1. a) PXRD diffraction of synthesized Mg3(Sb1�2xBixSnx)2 polycrystals. Magnified view indicating shift in b) major 101 and 011 peaks of α-Mg3Sb2.c) Isothermal section at 573 K of Mg─Sn─(Sb, Bi) phase diagram and triangulation of modified Gibbs free energy indicating low ΔGm region for thesynthesized compositions. d) Fractured BSE micrograph for with x� 0.2, indicating transgranular fracture and embedded nanoprecipitation of (Sn, Sb)phase and Mg2Sn, distinctively identifiable in e) elemental mapping and f ) EDX spectrum showing elemental peaks.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (3 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comlevels. The refined XRD patterns are shown in Figure S1–S3,Supporting Information, which displays the observed, calculated,and difference profiles along with the Bragg peaks of the consti-tuting phases and refinement parameters are tabulated inTable 1. The Sn, Bi dissolution into the Mg3Sb2 matrix, particu-larly of Bi atoms within the α-Mg3Sb2 phase, implies accordancewith Vegard’s Law. However, even for low alloying content, thatis, x≤ 0.05, complex coordination and trigonal (Sb, Sn, Bi) phaseprecipitation may result in α-Mg3Sb2 phase lattice contraction.For the higher alloyed compositions, that is, x≥ 0.1, cubicMg2Sn phase precipitation and larger elemental volume of Biin comparison to Sb atoms led to larger and increasing latticeparameter a and c, implying dopant solubility in the constitutingphases. When Sn is introduced at Sb site, it occupies specificanionic sites (causing Sb atoms precipitation), leading to altera-tions in the local lattice environment and symmetry of α-Mg3Sb2phase. Thus, the average crystal structure of the synthesizedMg3(Sb1�2xBixSnx)2 polycrystals and stability of constitutingphases are related to the amount of dopant, besides the differ-ences in the atomic size and electronegativity of the host andsubstituted Sb anions. It is worth mentioning that the presenceof excess Mg at cationic sites can stabilize or destabilize certainphases, potentially promoting or preventing the monoclinicphase transitions. In contrast to the anisotropic structural col-lapse with relative compression of the a and c-axes,[41] the mono-clinic phase formation is accompanied by unit cell expansion ofα-Mg3Sb2 unit cell. The alloying content dependence of the c/aratio and unit cell volume of the major trigonal α-Mg3Sb2 phase,as shown in Figure 2c, reveals a more elongated unit cell forhigher alloying content, with varying c/a ratio. When comparedto doped Mg3Sb1.9Sn0.1, codoping Sn and Bi atoms causes areduction in the c-axis length relative to the a-axis, which canbe ascribed to instability and partial loss of symmetry of majorα-Mg3Sb2 phase. As the volume increase is significant, it is indic-ative of dopant solubility and a growing potential for precipitationof secondary phases at higher alloying contents, discussedsubsequently. The partial loss of trigonal symmetry in the syn-thesized Mg3(Sb1�2xBixSnx)2 by phase transition occurs at theanionic [Mg2Sb2]2� framework and can be ascribed to the discor-dant Sn atoms, preferentially occupying specific lattice sites inthe alloyed polycrystal. The expanding unit cell volume for trigo-nal phase is indicative of a displacive phase transition whereinthe distorted coordination polyhedral within the anionic frame-work may constitute bonding heterogeneity.2.3. Bonding Heterogeneity and Low Thermal ConductivityMixed-anion compounds, which contain polyanions as separatebuilding blocks, may result in locally distorted structureswhere more than one anion bonds to a cation. The bondingenvironment in Mg defect complex, particularly the coordinationand geometric arrangement of Mg and Sb atoms, contributesto its overall stability or functionality. In the synthesizedMg3(Sb1�2xBixSnx)2 nanocomposites, the formation of MgSb6octahedra and MgSb4 tetrahedra as well as their interactionsthrough shared corners and edges to prevailing nanoclusterswas understood by evaluating bond lengths and occupanciesin the trigonal α-Mg3Sb2 phase, a polar intermetallic phase.The Mg(I)─Mg(II) distances in Mg3Sb2 are �3.736 Å, whilethe Mg─Sb distances range from about 2.8 to 3.2 Å.[11]Particularly, Mg(I)─Sb is slightly longer than Mg(II)─Sb, sug-gesting that both Mg─Mg and Mg─Sb interactions are signifi-cant in α-Mg3Sb2. The Mg(I) atoms are coordinated by six Sbatoms in a trigonal prismatic arrangement, while Mg(II) atomsare tetrahedrally coordinated by four Sb atoms as shown inFigure 3a. The bonding in Mg3Sb2 is predominantly ionic withsome covalent character in the Mg─Sb interactions,[42] whereinthe Mg(I)─Mg(II) interactions, while present, are relativelyweaker than the Mg─Sb bonds.To evaluate bonding heterogeneity, positional modulation ofthe atomic sites was introduced, after several refinement cycles,considering isotropic displacement parameters up to the fourthorder of the Fourier terms An and Bn (n= 1–4) using both cosineand sine components. The anion 4b Wyckoff site is occupied bythe Sb and Mg(II) atoms whereas the cation 4a Wyckoff site isoccupied by Mg(I) atoms, as shown in Figure 3b. The presence ofthree anions (i.e., Sb, Sn, and Bi) at the 4a Wyckoff site in(Mg2Sb2�2xBixSnx)2� motifs may result in an enhanced degreeof anionic disorder in the trigonal α-Mg3Sb2 phase leading tolocally distorted configurations with partially broken symmetry.The altering bonding distances of octahedral Mg(I) and tetrahe-dral Mg(II) environments, ranging from 2.87 to 3.02 Å, for lower(x≤ 0.05) and higher (x≥ 0.1) substitution, suggest compress-ibility and expansion, respectively. Softer and longerFigure 2. a) Unit cells of α-Mg3(Sb, Bi)2 and its partially transformedmonoclinic phase, with Mg(I) in gray, Mg(II) in red, and (Sb, Bi) in green.b) Refined lattice parameters a (square) and c(triangle), along with c). c=aratio (diamond) and volume (circle) for unit cell of major α-Mg3(Sb, Bi)2phase in Mg3(Sb1�2xBixSnx)2 nanocomposites. Corresponding structuralparameters for synthesized Mg3Sb1.9Sn0.1, shown with empty symbols.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (4 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comMg(I)─(Sb, Bi, Sn) bonds reveal an initial decrease and subse-quent increase in bond length, corroborating well with thedecreasing and increasing lattice parameter c, respectively.Remarkably, a similar trend for the shorter Mg(II)─(Sb, Bi, Sn)bonds constituting the polyanionic framework was evaluated,wherein an altering bonding patterns was also evaluatedfor higher (x≥ 0.1) substitution in the synthesized Mg3(Sb1�2xBixSnx)2 nanocomposites. Interestingly, atomic displace-ment evaluated for trigonal α-Mg3Sb2 and shown in Figure S7,Supporting Information, indicates a significant variation forMg(II) and Sb atoms constituting the anionic framework for vary-ing alloying content. Thus, synthesized Mg3(Sb1�2xBixSnx)2 con-stituting polyanionic framework is bound to have bondingheterogeneity-induced localized distortion and secondary phaseformation with altered bonding lengths. A comprehensiveunderstanding of bonding heterogeneity and structural transi-tions requires integrating advanced techniques such as atomprobe tomography,[43–45] nuclear magnetic resonance,[43] X-rayphotoelectron spectroscopy,[46] high-resolution transmissionelectron microscopy,[45–48] and Raman spectroscopy,[46,47] whichshould be the focus of future research to explore local atomic andelectronic environments. Moreover, interfacial lattice matchinganalysis[49] can be used in conjunction with crystal orbitalHamilton population evaluations[50] to study the compatibilityand interactions between different phases at their interfaces.The constituting locally distorted structures with a diversebonding environment in mixed-anion Mg3(Sb1�2xBixSnx)2-basednanocomposites are anticipated to result in significant variationsin interatomic force constants, splitting of phonon frequencies,and enhanced anharmonicity when compared to single-anionMg3Sb2.[51] The locally distorted structures resulting from alteredcoordination numbers (CN) in mixed-anion materials introducebonding heterogeneity, which effectively contributes to achievinglow κL as indicated in Figure 3c. The Sn atoms preferentiallyoccupy lattice sites with distorted square prismatic coordination(CN � 8, Fm-3m), distinct from the monocapped trigonal pris-matic sites (CN � 7, P-3m1) typically occupied by the substitutedhost Sb atoms in the anionic framework, leading to bond hetero-geneity. Moreover, higher dopant solubility at the polyanionicframework enhances scattering of phonons by mass/strain fluc-tuations particularly for lower substitution.Despite Sn limited solubility, a pronounced phonon scat-tering was attained, thus lowering the κL to >30% in comparisonto Bi-[23] and Sn-doped Mg3Sb2 counterparts in the measuredtemperature range. In Mg3(Sb1�2xBixSnx)2-based nanocompo-sites, κL decreases from �1.35Wm�1 K�1 (Mg3Sb1.9Sn0.1)to �0.9Wm�1 K�1 (x= 0.05) at 298 K, implying enhancedscattering of lower-frequency phonons. The limited dopantsolubility and coexisting secondary phases in the synthesizedMg3(Sb1�2xBixSnx)2 nanocomposites have contrasting effectson κL: wherein heterogeneous interfaces (presumably Mg3Sb2─Mg2Sn)[37,38] instead of enhancing phonon scattering apparentlydecrease interfacial thermal resistance, resulting in higher κL.Furthermore, higher κL is evaluated for higher alloying contentx� 0.2, where undesirable (Sn/SnSb) phases evolve, violatingthe T�1 temperature-dependent κL more drastically, thereby com-pensating the phonon scattering occurring at heterogeneousinterfaces.[17] Thus, thermal transport indicates critical role ofTable 1. Constituting phases in synthesized Mg3(Sb1�2xBixSnx)2 nanocomposites.Compositiona) Space group Lattice parameters Volume R-factorsb)GOF Rp RwpMg3Sb1.9Sn0.1 P3m1 a= b= 4.5724(1), c= 7.2505(2), α= β= 90°, γ= 120° 130.56(1) 1.59 3.76 5.12C2/m a= 14.493(1), b= 4.396(1), c= 7.908(1), α= γ= 90°; β= 90.858(3) 503.74(1)R3m a= b= 4.303(1), c= 11.335(1), α= β= 90°, γ= 120° 181.80(1)x� 0.025 P3m1 a= b= 4.5691(1), c= 7.2322(2), α= β= 90°, γ= 120° 130.76(1) 1.63 3.65 5.05C2/m a= 14.413(1), b= 4.348(1), c= 7.873(1), α= γ= 90°; β= 90.907(6) 493.38(1)R3m a= b= 4.299(1), c= 11.300(3), α= β= 90°, γ= 120° 180.92(1)x� 0.05 P3m1 a= b= 4.5673(1), c= 7.2314(1), α= β= 90°, γ= 120° 130.64(1) 1.55 3.38 4.80C2/m a= 14.452(1), b= 4.340(1), c= 7.866(1), α= γ= 90°; β= 91.013(5) 493.35(1)R3m a= b= 4.317(1), c= 11.307(2), α= β= 90°, γ= 120° 182.51(1)x� 0.1 P3m1 a= b= 4.5828(1), c= 7.2502(2), α= β= 90°, γ= 120° 131.87(1) 1.72 4.02 5.30C2/m a= 14.517(1), b= 4.463(1), c= 7.935(1), α= γ= 90°; β= 90.606(6) 503.74(1)Fm3m a= b= c= 6.6795(2), α= β= γ= 90° 298.01(1)x� 0.2 P3m1 a= b= 4.5909(1), c= 7.2799(3), α= β= 90°, γ= 120° 132.88(1) 1.34 3.00 3.91C2/m a= 14.581(1), b= 4.496(1), c= 7.998(1), α= γ= 90°; β= 90.847(5) 524.29(2)Fm3m a= b= c= 6.7477(2), α= β= γ= 90° 307.24(1)I41/amd a= b= 5.8287(2), c= 3.1798(2), α= β= γ= 90° 108.03(1)a)Space group, lattice parameters, unit cell volume (Å2), and refinement parameters (weighted profile (Rwp); profile (Rp); and expected (Rexp) R-factor). Standard deviations aregiven in parentheses. b)Rwp ¼ffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiPwi yobsi � ycalci� �2Pwi yobsi� �2s;Rp ¼P jyobsi � ycalci jPyobsi; GOF ¼ RwpRexp.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (5 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comstructural inhomogeneity, bonding heterogeneity, and phaseseparation in increasing the lattice anharmonicity, which iseffective in reducing κL of the synthesized Mg3(Sb1�2xBixSnx)2nanocomposites.2.4. Electronic Band Structure, Ionized Impurity ScatteringMechanism, and zT EnhancementThe Mg3(Sb, Bi)2-based zintl compounds exhibit a complex bandstructure with a challenging large and asymmetric unit cell.[11,52]Similar to state-of-the-art (Sb, Bi)2Te3 alloys, electronic bandstructure of Mg3(Sb, Bi)2 constitutes low-symmetry conductionband minimum (CBM) with high band degeneracy (Nv� 6).However, low band degeneracy (i.e., Nv� 1) of valence bandmaximum (VBM) limits the optimization of power factor forp-type conduction.[27] Notably, transitioning from semiconduc-tor-to-metallic behavior occurs on varying the Sb contentMg3Sb2�xBix with minimum κ measured for x� 0.5.[14] To sim-ulate and analyze the effects of disorder on the electronic bandstructure of the synthesized Mg3(Sb1�2xBixSnx)2 nanocompo-sites, Korringa-Kohn-Rostoker (KKR) -coherent potential approx-imation (CPA) approach is employed which replaces the actualdisordered crystal structure with an effective medium Green’sfunction and a self-energy.[53–55] Even at low doping levels, theKKR-CPA approach provides a realistic representation of thedisordered doped system without necessitating largesupercells,[56–59] making it particularly relevant for our evaluationof (Sn, Bi) codoped Mg3Sb2. The electronic band structurewith estimated spectral weight variations corresponding toBloch spectral functions (BSFs) in arbitrary units (a.u.) ofMg3Sb1.9Sn0.1 and Mg3Sb1.8Sn0.1Bi0.1 is shown in Figure 4a,b,respectively.The band dispersion shows notable changes upon Sn doping,with the VBM at the Γ point and the CBM at the K point, agreeingwell with the evaluation made previously employing supercellapproach.[36] Since the VBM and CBM are located at differentpoints in k-space (Γ and K, respectively) with coinciding energies,the resulting zero indirect bandgap allows carriers to undergotransition from the valence to the conduction band without achange in momentum, but via an intermediate state. Diffusedspectral weight and overlapping bands at the Γ and A pointsin the valence band suggest that electronic states are spread overa range of energies implying disorder, impurities, or strongelectron–electron interactions. The upward shift of the VBMpoints to an enhanced p-type conductivity, indicating metallicor semimetallic behaviour due to the zero indirect bandgapand band overlap. Previous calculations have also suggestedthe VBM of Mg3Sb2 at the Brillouin zone center (Γ) being domi-nated by the p orbitals of Sb anions that are composed of in-planepx,y and out-of-plane pz orbital.[2,36,60] Furthermore, the reductionin σ(T ) upon Bi codoping could be well associated with a pro-nounced chemical disorder which enhances carrier scatteringin the (Mg2Sb2�2xBixSnx)2� motifs. The negative bandgap,accompanied by an upward shift of the valence bands, indicatesenhanced interactions between electronic states, while anincreased orbital overlap may spatially confine electrons, reduc-ing the density of free charge carriers in the system.To understand carrier characteristics, Hall measurementresults at room temperature are shown in Figure 4c. TheMg3(Sb1�2xBixSnx)2 nanocomposites for x� 0.025 reveal a sig-nificant increase in both carrier mobility (μH) and carrier concen-tration (nH), in comparison to synthesized Mg3Sb1.9Sn0.1, whichcan be ascribed primarily to the upshift of VBM. However, withincreasing substitution, decreasing nH and deteriorating μH weremeasured, attributed to an enhanced carrier scattering. Theeffects of Sn and Bi substitutions on temperature-dependentelectrical transport behavior of the synthesized Mg3Sb1.9Sn0.1and Mg3(Sb1�2xBixSnx)2 nanocomposites are shown inFigure 4d–f, alongside the previously reported Mg3Sb1.8Bi0.2for comparison.[23] The temperature-dependent Seebeck coeffi-cient S(T ), shown in Figure 4d, implies p-type conduction, as alsocorroborated from the Hall measurement data. At room temper-ature, S for Mg3Sb1.9Sn0.1 (�143 μV K�1) increases linearly withtemperature, reaching a value of �253 μV K�1 at 673 K. As thecosubstitution of Sn and Bi increases, the Seebeck coefficientat room temperature shows a marginal decrease for x� 0.25,likely due to changes in the carrier concentration or scatteringmechanisms. However, for phase-separated sample x� 0.1 hav-ing heterogeneous interfaces (presumably Mg3Sb2─Mg2Sn), Sincreases to �227 μV K�1 at 300 K and decreases thereafter,which can be ascribed primarily to carrier filtering.Figure 3. a) Schematic showing two nonequivalent Mg sites withtetrahedral and octahedral polyhedra and b) evaluated bond distancesfor Mg(I)─Sb and Mg(II)─Sb atoms, in trigonal α-Mg3(Sb, Bi)2 phase.c) Temperature-dependent lattice thermal conductivity (κL) of the synthe-sized Mg3(Sb1�2xBixSnx)2 nanocomposites shown along with previouslyreported[23] values of Mg3Sb1.8Bi0.2 polycrystals.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (6 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comThe temperature-dependent electrical conductivity σ(T ) withinthe temperature range of 300–673 K shown in Figure 4e indicatesa significant increase in σ(T ) for x� 0.25 to a measured value of�8.2� 103 Sm�1 at room temperature, corresponding to �1.6-fold enhancement over Mg3Sb1.9Sn0.1 (�3.1� 103 Sm�1), whichagrees well with the values reported previously.[36] This increaseis primarily due to the introduction of additional charge carriersby Sn substitution. However, with (Sn, Bi) cosubstitution,σ(T ) decreases as a result of additional scattering centers (reduc-ing carrier mobility), primarily due to increasing phase fractionsof secondary phases, which reduce carrier mobility. Additionally,modifications in the band structure by Bi substitution are likelycontribute to this decline in electrical conductivity, as observedpreviously in Mg3Sb1�xBix.[23] The σ(T ) of phase-separated nano-composites at higher alloying content aligns with the σ(T ) ofMg3Sb1�xSnx reported previously and implies the detrimentaleffects of phase separation on electrical conductivity.[36] This sug-gests that the presence of segregated phases leads to increasedscattering and reduced μH ultimately compromising the electri-cal transport for higher substitutions.The weighted mobility (μw)[61] derived S(T ) and σ(T ) measure-ments, that closely mirror μH, are also shown in Figure 4f as itreflects electron mobility weighted by the density of electronicstates, providing insights into the electronic structure and scat-tering mechanisms. The observed increase in μw with tempera-ture suggests ionized impurity scattering as the primarymechanism of carrier scattering. At lower temperatures, slowercarriers allow more interactions with charged impurities, whileas temperature rises, carriers move faster, reducing scatteringfrom ionized impurities and thus increasing μw. However, athigher temperatures, mobility typically decreases due toincreased lattice vibrations (phonon scattering), indicating thecoexistence of multiple scattering mechanisms, each varyingin influence with temperature and alloying content.The temperature-dependent thermal conductivity κ(T )shown in Figure 5a also indicates significant reduction inMg3(Sb1�2xBixSnx)2 nanocomposites, in comparison to Bi-[23]and Sn-doped Mg3Sb2 counterparts ascribed primarily to latticeanharmonicity, phase separation, and dynamic structural disor-der. κ(T ) is significantly lowered even for low alloying content asshown in Figure 5b, when compared to Mg3Sb1.9Sn0.1 both at low(�300 K) and high (�673 K) temperature. The increase in κL forx> 0.05 can be attributed to the higher κ of the Mg2Sn phase[62]and Sn segregation, which presumably enhances the phononFigure 4. Calculated electronic band structure of a) Mg3Sb1.9Sn0.1 and b) Mg3Sb1.8Sn0.1Bi0.1. Normalized scale reflects the BSFs in arbitrary units. c) Hallcarrier concentration and mobility at 300 K. Temperature-dependent d) Seebeck coefficient, e) electrical conductivity, and f ) weighted mobility ofMg3(Sb1�2xBixSnx)2 nanocomposites shown along with previously reported[23] values of Mg3Sb1.8Bi0.2 polycrystals.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (7 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.compropagation by creating additional thermally conductive path-ways. The power factor (Figure 5c,d) shows a marginal deterio-ration compared to Mg3Sb1.9Sn0.1 for low alloying content(x≤ 0.05) and decreases significantly thereafter with increasingalloying content, primarily to reduction in both S(T ) and σ(T ) dueto secondary phase formations. However, the enhancementis more significant in relation to Mg3Sb1.8Bi0.2,[23] for lowertemperatures (<600 K), achieving a maximum power factor of�2.5 μWcm�1 K�2 at 573 K, thus, implying that the cosubstitu-tion of Sn and Bi sustains the power factor for low alloyingcontent (x≤ 0.05) and remains detrimental for increasing substi-tution (x≥ 0.1) due to Mg2Sn phase precipitation. The zT shownin Figure 5e for all the synthesized Mg3(Sb1�2xBixSnx)2 nano-composites increases with temperature, wherein Mg3Sb1.9Sn0.1exhibits significantly higher zT� 0.2 (� 0.05), which increasesmarginally for x� 0.25–0.5 (as shown in Figure 5f ), primarilydue to a notable κ reduction, thus, demonstrating that (Sn, Bi)substitutions effectively improve the zT, but inherent distortionand phase separation prove detrimental to further enhance thezT, which optimizes only for lower content (x≤ 0.05) in synthe-sized Mg3(Sb1�2xBixSnx)2 nanocomposites.Thus, a limited understanding of structural changes that occurduring doping or alloying hinders insights into their prominenteffects on conduction/valence bands and defect chemistry. Whilechallenges such as limited solubility and poor doping efficiencypersist, improving carrier mobility and power factor in p-typeMg3Sb2 necessitates exploring codoping or alloying, especiallywithin the anionic framework along with effective cationicdopants such as Na, Yb, Ag, Zn, and/or Cu.[28–32] Thus, complexmultinary p-type Mg3(Sb,Bi)2-based zintl nanocomposites,with polyanionic frameworks, offer potential for realizing all-Mg3Sb2-based TE devices with enhanced TE conversion efficiency.However, addressing intrinsic monoclinic phase transition, nano-precipitation susceptibility, and bonding heterogeneity remainscritical for optimizing transport properties.3. ConclusionAnionic solid solutions constituting a polyanionic substructurecan offer a versatile approach to favorably tune p-type conductionin Mg3(Sb, Bi)2 by regulating the composition of the anionicFigure 5. Temperature dependence and alloying content dependence of a,b) total thermal conductivity, c,d) power factor, and e,f ) TE figure of merit (zT )of nominal Mg3(Sb1�2xBixSnx)2 nanocomposites. Corresponding TE transport parameters for synthesized Mg3Sb1.9Sn0.1, shown with empty symbols.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (8 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.comframework. The synthesized (Sn, Bi) cosubstituted Mg3Sb2-based polyanionic solid solutions exhibit bonding heterogeneityin dominant α-Mg3(Sb, Bi)2 phase, intrinsic partial monoclinicphase transitions, and nanoprecipitation susceptibility with vary-ing alloying content. Coexistence of different crystal structuresand phases relays poor dopability and limited solubility,highlighting the critical role of anionic framework for phase sta-bility. Despite low content (x≤ 0.05), structural changes inducedby (Sn, Bi) substitution reduce κ and synergistically enhance car-rier mobility and concentrations, leading to improved power fac-tor and zT. Phase separation of cubic Mg2Sn for higher alloyingcontent (x≥ 0.1) enhances κL and deteriorates carrier transport,thereby drastically reducing the zT. Thus, alloying/substitutionstrategies must carefully account for structural and bondingmodifications to enhance dopability and tailor the transportproperties of zintl compounds.Supporting InformationSupporting Information is available from the Wiley Online Library or fromthe author.AcknowledgementsThis research was supported by JST Mirai Program, Japan (grant no.JPMJMI19A1).Conflict of InterestThe authors declare no conflict of interest.Author ContributionsNagendra Singh Chauhan: conceptualization (equal); formal analysis(lead); investigation: (lead); methodology (lead); writing—original draft(lead). Takao Mori: conceptualization (equal); funding acquisition (lead);project administration (lead); resources (lead); supervision (lead); writing—review and editing (lead).Data Availability StatementThe data that support the findings of this study are available from thecorresponding author upon reasonable request.Keywordsbonding heterogeneities, nanocomposites, phase separations,thermoelectrics, zintlReceived: December 2, 2024Revised: February 13, 2025Published online: March 5, 2025[1] H. Tamaki, H. K. Sato, T. Kanno, Adv. Mater. 2016, 28, 10182.[2] J. Zhang, L. Song, S. H. Pedersen, H. Yin, L. T. Hung, B. B. Iversen,Nat. Commun. 2017, 8, 13901.[3] J. Zhang, L. Song, A. Mamakhel, M. R. V. Jørgensen, B. B. Iversen,Chem. Mater. 2017, 29, 5371.[4] J. Zhang, L. Song, K. A. Borup, M. R. V. Jørgensen, B. B. Iversen, Adv.Energy Mater. 2018, 8, 1702776.[5] S. Bano, R. Chetty, J. Babu, T. Mori, Device 2024, 2, 100408.[6] J. Mao, H. Zhu, Z. Ding, Z. Liu, G. A. Gamage, G. Chen, Z. Ren,Science 2019, 365, 495.[7] L. Wang, N. Sato, Y. Peng, R. Chetty, N. Kawamoto, D. H. Nguyen,T. Mori, Adv. Energy Mater. 2023, 13, 2301667.[8] L. Wang, W. Zhang, S. Y. Back, N. Kawamoto, D. H. Nguyen, T. Mori,Nat. Commun. 2024, 15, 6800.[9] H. Cho, S. Y. Back, N. Sato, Z. Liu, W. Gao, L. Wang, H. D. Nguyen,N. Kawamoto, T. Mori, Adv. Funct. Mater. 2024, 34, 2407017.[10] H. Shang, Z. Liang, C. Xu, J. Mao, H. Gu, F. Ding, Z. Ren, Research2020, 2020, 1219461.[11] A. Li, C. Fu, X. Zhao, T. Zhu, Research 2020, 2020, 1934848.[12] Z. Liu, N. Sato, W. Gao, K. Yubuta, N. Kawamoto, M. Mitome,K. Kurashima, Y. Owada, K. Nagase, C.-H. Lee, J. Yi, K. Tsuchiya,T. Mori, Joule 2021, 5, 1196.[13] Y. Sun, Y. Zhu, H. Wu, N. Qu, L. Xie, J. Zhu, Z. Liu, Q. Zhang, W. Cai,F. Guo, J. Sui, Energy Environ. Sci. 2024, 17, 738.[14] K. Imasato, M. Wood, S. Anand, J. J. Kuo, G. J. Snyder, Adv. EnergySustainability Res. 2022, 3, 2100208.[15] T. Kanno, H. Tamaki, M. Yoshiya, H. Uchiyama, S. Maki, M. Takata,Y. Miyazaki, Adv. Funct. Mater. 2021, 31, 2008469.[16] J. Hu, J. Zhu, X. Dong, M. Guo, Y. Sun, W. Shi, Y. Zhu, H. Wu, F. Guo,Y. Zhang, Z. Ge, Q. Zhang, Z. Liu, W. Cai, J. Sui, Small 2023, 19,2301382.[17] Y. Zhu, Y. Xia, Y. Wang, Y. Sheng, J. Yang, C. Fu, A. Li, T. Zhu, J. Luo,C. Wolverton, G. J. Snyder, J. Liu, W. Zhang, Research 2020, 2020,4589786.[18] N. Kazem, S. M. Kauzlarich,Handbook on the Physics and Chemistry ofRare Earths, Elsevier, Amsterdam 2016, pp. 177–208.[19] M. Martinez-Ripoll, A. Haase, G. Brauer, Acta Crystallogr. B 1974, 30,2006.[20] C. Zheng, R. Hoffmann, R. Nesper, H. G. Von Schnering, J. Am.Chem. Soc. 1986, 108, 1876.[21] J. Wang, J. Mark, K. E. Woo, J. Voyles, K. Kovnir, Chem. Mater. 2019,31, 8286.[22] S. M. Kauzlarich, K. P. Devlin, C. J. Perez, Thermoelectric EnergyConversion, Elsevier, Amsterdam 2021, pp. 157–182.[23] A. Bhardwaj, N. S. Chauhan, D. Misra, J. Mater. Chem. A 2015, 3,10777.[24] A. Bhardwaj, D. Misra, RSC Adv. 2014, 4, 34552.[25] J. Tani, H. Ishikawa, Phys. B 2020, 588, 412173.[26] N. S. Chauhan, T. Mori, J. Mater. Chem. A 2024, 12, 32703.[27] P. Gorai, B. R. Ortiz, E. S. Toberer, V. Stevanović, J. Mater. Chem. A2018, 6, 13806.[28] L. Huang, T. Liu, X. Mo, G. Yuan, R. Wang, H. Liu, X. Lei, Q. Zhang,Z. Ren, Mater. Today Phys. 2021, 21, 100564.[29] Z. Ren, J. Shuai, J. Mao, Q. Zhu, S. Song, Y. Ni, S. Chen, Acta Mater.2018, 143, 265.[30] M. Tiadi, V. Trivedi, S. Kumar, P. K. Jain, S. K. Yadav, R. Gopalan,D. K. Satapathy, M. Battabyal, ACS Appl. Mater. Interfaces 2023,15, 20175.[31] V. P. Kannan, V. Lourdhusamy, I. Paulraj, S. Madanagurusamy,C.-J. Liu, ACS Appl. Mater. Interfaces 2024, 16, 58677.[32] M. Tiadi, M. Battabyal, P. Jain, A. Chauhan, D. K. Satapathy,R. Gopalan, Sustainable Energy Fuels 2021, 5, 4104.[33] A. Bhardwaj, N. S. Chauhan, S. Goel, V. Singh, J. Pulikkotil,T. Senguttuvan, D. Misra, Phys. Chem. Chem. Phys. 2016, 18, 6191.[34] Z. Liang, C. Xu, S. Song, X. Shi, W. Ren, Z. Ren, Adv. Funct. Mater.2023, 33, 2210016.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (9 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.com[35] J. Zhang, J. Wang, Y. Yuan, Calphad 2021, 75, 102361.[36] Y. Wang, X. Zhang, Y. Liu, Y. Wang, J. Zhang, M. Yue, Vacuum 2020,177, 109388.[37] E. Dong, Y. Zhu, S. Tan, W. Liu, J. Wang, W. Zhang, Acta Mater. 2022,236, 118139.[38] Y. Zhu, R. Shu, F. Jiang, A. Elsukova, P. O. Å. Persson, P. Eklund,W. Liu, ACS Appl. Nano Mater. 2023, 6, 6133.[39] P. Zhao, W. Xue, Y. Zhang, S. Zhi, X. Ma, J. Qiu, T. Zhang, S. Ye,H. Mu, J. Cheng, X. Wang, S. Hou, L. Zhao, G. Xie, F. Cao,X. Liu, J. Mao, Y. Fu, Y. Wang, Q. Zhang, Nature 2024, 631, 777.[40] S. Baranets, A. Ovchinnikov, S. Bobev, Handbook on the Physics andChemistry of Rare Earths, Elsevier, Amsterdam 2021, pp. 227–324.[41] M. Calderón-Cueva, W. Peng, S. M. Clarke, J. Ding, B. L. Brugman,G. Levental, A. Balodhi, M. Rylko, O. Delaire, J. P. S. Walsh,S. M. Dorfman, A. Zevalkink, Chem. Mater. 2021, 33, 567.[42] J. Zhang, L. Song, M. Sist, K. Tolborg, B. B. Iversen, Nat. Commun.2018, 9, 4716.[43] T. Luo, J. J. Kuo, K. J. Griffith, K. Imasato, O. Cojocaru-Mirédin,M. Wuttig, B. Gault, Y. Yu, G. J. Snyder, Adv. Funct. Mater. 2021,31, 2100258.[44] J. J. Kuo, Y. Yu, S. D. Kang, O. Cojocaru-Mirédin, M. Wuttig,G. J. Snyder, Adv. Mater. Interfaces 2019, 6, 1900429.[45] R. Shu, Z. Han, A. Elsukova, Y. Zhu, P. Qin, F. Jiang, J. Lu,P. O. Persson, J. Palisaitis, A. Le Febvrier, Adv. Sci. 2022, 9, 2202594.[46] M. Tiadi, S. Kumar, P. K. Jain, D. K. Satapathy, M. Battabyal, ACS Appl.Mater. Interfaces 2025, 17, 4996.[47] M. Tiadi, D. K. Satapathy, M. Battabyal, Phys. Rev. B 2024, 109,195201.[48] P. Nan, A. Li, L. Cheng, K. Wu, Z. Liang, F. Lin, C. Fu, T. Zhu, B. Ge,Mater. Today Phys. 2021, 21, 100524.[49] Q. Zhang, J. Li, N. S. Chauhan, L. Wang, Z. Huang, W. Fan,K. Hayashi, S. Chen, J. Fan, Y. Miyazaki, J. Mater. Chem. A 2023,11, 18811.[50] Q. Zhang, H. Li, N. S. Chauhan, L. Wang, W. Fan, S. Chen, J. Fan,Y. Miyazaki, Mater. Today Energy 2024, 44, 101656.[51] N. Sato, N. Kuroda, S. Nakamura, Y. Katsura, I. Kanazawa, K. Kimura,T. Mori, J. Mater. Chem. A 2021, 9, 22660.[52] Z. Li, P. Graziosi, N. Neophytou, NPJ Comput. Mater. 2024, 10, 8.[53] H. Akai, P. Dederichs, J. Phys. C: Solid State Phys. 1985, 18, 2455.[54] T. Kotani, H. Akai, Phys. Rev. B 1996, 54, 16502.[55] H. Akai, M. Ogura, N. H. Long, J. Phys.: Condens. Matter 2009, 21,064241.[56] A. Gonis, G. Stocks, W. Butler, H. Winter, Phys. Rev. B 1984, 29, 555.[57] N. S. Chauhan, Y. Miyazaki, Mater. Adv. 2022, 3, 8498.[58] N. S. Chauhan, Y. Miyazaki, J. Alloys Compd. 2022, 908, 164623.[59] S. Misra, B. Wiendlocha, S. El Oualid, A. Dauscher, B. Lenoir,C. Candolfi, J. Mater. Chem. A 2024, 12, 1166.[60] W. Zhang, J.-F. Halet, T. Mori, J. Mater. Chem. A 2023, 11, 24228.[61] G. J. Snyder, A. H. Snyder, M. Wood, R. Gurunathan, B. H. Snyder,C. Niu, Adv. Mater. 2020, 32, 2001537.[62] L. Zhang, J. Li, H. Chen, J. Feng, R. Liu, J. Mater. Chem. C 2024, 12,8935.www.advancedsciencenews.com www.small-structures.comSmall Struct. 2025, 6, 2400632 2400632 (10 of 10) © 2025 The Author(s). Small Structures published by Wiley-VCH GmbH 26884062, 2025, 4, Downloaded from https://onlinelibrary.wiley.com/doi/10.1002/sstr.202400632 by National Institute For, Wiley Online Library on [03/04/2025]. See the Terms and Conditions (https://onlinelibrary.wiley.com/terms-and-conditions) on Wiley Online Library for rules of use; OA articles are governed by the applicable Creative Commons Licensehttp://www.advancedsciencenews.comhttp://www.small-structures.com Bonding Heterogeneity and Nanoprecipitation on Substituting the Anionic Framework in Mg3Sb2 for p-Type Zintl Thermoelectrics 1. Introduction 2. Results and Discussion 2.1. Phase Compositions, Phase Segregation, and Hierarchical Microstructure 2.2. Structural Parameters and Monoclinic Phase Transitions 2.3. Bonding Heterogeneity and Low Thermal Conductivity 2.4. Electronic Band Structure, Ionized Impurity Scattering Mechanism, and zT Enhancement 3. Conclusion