Charge Disproportionation Triggers Bipolar Doping in FeSb2

May 10, 2019 - (33) Remarkably, the XPS spectrum of Sb 3d5/2 shell electrons for samples with x ≥ 0.2 also showed a shoulder peak at lower binding e...
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Cite This: J. Am. Chem. Soc. 2019, 141, 9249−9261

Charge Disproportionation Triggers Bipolar Doping in FeSb2−xSnxSe4 Ferromagnetic Semiconductors, Enabling a Temperature-Induced Lifshitz Transition Honore Djieutedjeu,† Juan S. Lopez,† Ruiming Lu,† Brandon Buchanan,† Xiaoyuan Zhou,‡ Hang Chi,‡ Kulugammana G. S. Ranmohotti,†,§ Ctirad Uher,‡ and Pierre F. P. Poudeu*,† †

Laboratory for Emerging Energy and Electronic Materials (LE3M), Department of Materials Science and Engineering and Department of Physics, University of Michigan, Ann Arbor, Michigan 48109, United States § Division of Science, Mathematics and Technology, Governors State University, University Park, Illinois 60484, United States

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‡

S Supporting Information *

ABSTRACT: Ferromagnetic semiconductors (FMSs) featuring a high Curie transition temperature (Tc) and a strong correlation between itinerant carriers and localized magnetic moments are of tremendous importance for the development of practical spintronic devices. The realization of such materials hinges on the ability to generate and manipulate a high density of itinerant spin-polarized carriers and the understanding of their responses to external stimuli. In this study, we demonstrate the ability to tune magnetic ordering in the p-type FMS FeSb2−xSnxSe4 (0 ≤ x ≤ 0.20) through carrier density engineering. We found that the substitution of Sb by Sn FeSb2−xSnxSe4 increases the ordering of metal atoms within the selenium crystal lattice, leading to a large separation between magnetic centers. This results in a decrease in the Tc from 450 K for samples with x ≤ 0.05 to 325 K for samples with 0.05 < x ≤ 0.2. In addition, charge disproportionation arising from the substitution of Sb3+ by Sn2+ triggers the partial oxidation of Sb3+ to Sb5+, which is accompanied by the generation of both electrons and holes. This leads to a drastic decrease in the electrical resistivity and thermopower simultaneously with a large increase in the magnetic susceptibility and saturation magnetization upon increasing Sn content. The observed bipolar doping induces a very interesting temperature-induced quantum electronic transition (Lifshitz transition), which is manifested by the presence of an anomalous peak in the resistivity curve simultaneously with a reversal of the sign of a majority of the charge carriers from hole-like to electron-like at the temperature of maximum resistivity. This study suggests that while there is a strong correlation between the overall magnetic moment and free carrier spin in FeSb2−xSnxSe4 FMSs, the magnitude of the Curie temperature strongly depends on the spatial separation between localized magnetic centers rather than the concentration of magnetic atoms or the density of itinerant carriers.

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INTRODUCTION Manipulating magnetic ordering in various magnetic semiconductors by chemical (doping, solid solution) or physical (external stimuli such as applied bias voltage or magnetic field) means has been intensively investigated,1−3 with the goal of demonstrating the ability to use both the charge and spin degrees of freedom of free carriers in spintronics devices.1,4 Important milestones toward realizing this goal are the ability to control the density and transport of spin-polarized charge carriers as well as understanding their interactions with localized magnetic moments in magnetic semiconductors and particularly in high-temperature ferromagnetic semiconductors (FMSs).5−9 In this respect, significant progress has been made over the past three decades on dilute ferromagnetic semiconducting (FMS) material systems.4,10−13 For instance, nonmagnetic conventional semiconductors, such as III−V, II−VI, and IV−IV compounds, can be turned into magnetic semiconductors (MS) by introducing sizable concentrations of © 2019 American Chemical Society

magnetic ions, such as Mn, Fe, Ni, and Co into the crystal lattice. In addition, the electronic properties of such conventional semiconductors are modulated by the incorporation of the magnetic impurities. This approach has been used for several decades with III−V, II−VI, and IV−IV semiconductors.1,4,14,15 While III−V semiconductors such as GaAs are already being used in a large number of electronic devices, the integration of the magnetically doped materials (Ga1−xMnx)As into practical devices remains restricted by the low magnetic transition temperature (173 K) and low solubility of the magnetic ion in the host semiconductor.4,13 Nevertheless, such multifunctional materials could find application in devices where the magnetic properties can be controlled through optically activated charge carriers,16 such as the spin-based field effect transistors (FETs), spin-polarized Received: February 22, 2019 Published: May 10, 2019 9249

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compounds with low-symmetry crystal structure in which the substructure responsible for magnetic behavior is decoupled from the substructure responsible for electronic conduction. Such low-symmetry structure enables flexibility in independently tuning the carrier concentration through doping within the semiconducting substructure, as well as altering the concentration and the separation between magnetic atoms through solid solution substitution within the magnetic substructure. Along these lines, our research group has developed over the past decade several complex structures with tremendous compositional flexibilities.5,28−37 These include (M 2− x Pb x )Pb 2 Sb 4 Se 10 (M = Mn, Fe); 3 6 , 3 8 FePb4Sb6Se14;34 and homologous series M2Pn4M′N−3SeN+5 (N ≥ 3 (M = Mn, Fe; M′ = Pb, Sn; Pn = Sb, Bi).5,30,33,35−37 The fascinating class of ternary compounds MPn2Se4 (Figure 1) corresponds to the first member of the

lasers, light-emitting diodes (LEDs), nonvolatile memory, and quantum computing. In magnetically doped conventional semiconductors, extrinsic carriers (electrons or holes) generated within the valence band (VB) or the conduction band (CB) are coupled to the concentration of the magnetic impurity (Mn, Fe, etc.), which serves both as a doping agent (donor or acceptor) and as the magnetically active species. Fascinating collective carrierinduced or carrier-mediated ferromagnetic phenomena, which incorporate the effects of phonons, electrons, and spins on the overall magnetic and transport behavior, have been reported in magnetic semiconductors and diluted magnetic semiconductors.2,17−19 However, the mechanism of the exchange interaction between the spin of charge carriers and the localized moment of magnetic atoms remains controversial.10,13 In the case of carrier-induced ferromagnetism, the critical carrier concentration has been reported for the onset of ferromagnetism and for magnetic transitions such as AFM to FM ordering and FM to PM disorder behavior. For instance, it was found in [(PbTe)1−x(SnTe)x]1−y(MnTe)y that with a very small value of y the FM onset hole concentration was p ≈ 2 × 1020 cm−3, while the maximum Tc value was observed with samples exhibiting hole concentrations of p ≈ 7 × 1020 cm−3.20 Ferromagnetic ordering in this system was rationalized by considering that the number of effective carriers taking part in the RKKY (Ruderman-Kittel-Kasuya-Yosida) exchange interaction mechanism increases when the electrons from the L valence band are 4 times more mobile than those in the Σ band.21,22 In material systems with s−d and p−d orbital hybridization, such as p-type magnetic semiconductors with large hole densities (Mn-doped GaAs), it is widely accepted that the FM ordering of adjacent magnetic centers (Mn−Mn) within the host material is mediated by a very strong spin−spin interaction involving free carriers. However, the intrinsic contribution of individual material parameters (carrier density, carrier mobility, fraction of magnetic centers (concentration of Mn), and the distance between magnetic centers) to the change in the Curie temperature (Tc) remains unclear.10,13,23 In addition, the exact band picture resulting from the substitution of Ga by Mn in (Ga, Mn)As remains controversial.10,24 This limited understanding of the mechanism of ferromagnetism in Mn-doped FMSs is due to the fact that both the density of holes and the concentration of magnetic centers in FMSs are modified simultaneously when Mn ions substitute for Ga in (Mn, Ga)As. In addition, it is challenging to control the Mn−Mn separation within the host material as well as the exact concentration of Mn atoms contributing to the ferromagnetic order in Mn-doped FMSs. This is due to the fact that some of the Mn ions enter interstitial sites rather than occupying the crystallographic position of the substituted atoms.25−27 This limited understanding of the mechanism by which holes mediate ferromagnetic ordering and affect Tc in Mn-doped FMSs severely hampers our ability to use spin-dependent electronic transport to manipulate magnetic ordering and to develop efficient ferromagnetic semiconductors suitable for spin-based devices at practical temperatures. To achieve a greater level of understanding of key parameters (carrier concentration, concentration of magnetic atoms, distance between magnetic atoms, etc.) controlling the correlations between electronic transport and magnetic behavior in magnetic semiconductors, it is necessary to design

Figure 1. Schematic illustration of the crystal structure of FeSb2−xSnxSe4 projected along the b axis. The relevant semiconducting and magnetic substructural units are highlighted. The M1 and M2 sites (blue) have mixed occupancy by Sb (96%) and Fe (4%), whereas the M4 site (cyan) has mixed occupancy by Fe (84%) and Sb (16%). The M3 site (green) is fully occupied by Fe.

M2Pn4M′N−3SeN+5 series with N = 3. The monoclinic crystal structure (space group C2/m) contains four crystallographically independent metal positions [M1(4i), M2(4i), M3(2d), and M4(2a)] that are distributed within two distinct building units alternating along [001] denoted layer A and layer B (Figure 1). Layer A, which contains metal positions M1 and M3, consists of double rods with face-sharing monocapped trigonal prismatic coordination of Se atoms around the M1 site that alternate along the a axis with M(3)Se6 octahedra that shared edges to form a single chain parallel to the b axis. Adjacent layer A units are separated by the NaCl-type building unit, layer B, that consists of octahedrally coordinated M2 and M4 metal positions. The pnictogen atoms (Pn) predominantly occupy the M1 and M2 sites, whereas the transition-metal atoms (M) are located within the M3 and M4 sites. This preferential distribution of Pn and M atoms within the crystal lattice results in the formation of two important substructures. The first is the one-dimensional (1D) chain, [MnSe4n+2], of edge-sharing octahedra of Se atoms containing the magnetic transition-metal atoms (M3 and M4). This will be thereafter referred to as the magnetic subunit (Figure 1). The second substructure is the three-dimensional (3D) framework separating adjacent magnetic subunits and containing the nonmagnetic main group metals (M1 and M2) and selenium atoms. This nonmagnetic substructure will be referred to as the 9250

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Journal of the American Chemical Society semiconducting subunit. Our investigation of the MPn2Se4 phases led to the synthesis of several compositions and the study of their structure−properties relationship. This includes compositions MnSb2Se4,35 FeSb2Se4,37 Mn1−xSnxBi2Se4,5 FeBi2Se4,30 Fe1−xSnxSb2Se4,32 MnSb2−xSnxSe4,33 and FeSb2−xInxSe4.29 We find that the type of magnetic atoms (Mn, Fe) within the [MnSe4n+2] chains governs the nature of magnetic ordering in these phases. For instance, the FePn2Se4 phases display ferromagnetic behavior,30,37 whereas the Mn analogs (MnPn2Se4) are predominantly antiferromagnetic.5,35 In addition, the type of majority charge carriers in MPn2Se4 is controlled by the nature of the main group element (Bi, Sb) within the semiconducting unit. For example, the conduction type in MPn2Se4 switches from p type for MSb2Se4 to n type for MBi2Se4.5,30,35,37 This ability to independently control the magnetic ordering and carrier type within the same crystal lattice through chemical manipulation of the composition of various structural subunits makes it possible to examine the effect of free carriers on ferromagnetic ordering in spintronic materials. Here, we report on the FeSb 2−x Sn x Se 4 system, a ferromagnetic (FM) analogue of the MnSb2−xSnxSe4 antiferromagnetic (AFM) system,33 in which an interesting transition from AFM to FM behavior was observed for an optimum substitution of Sb with Sn (0.05 ≤ x ≤ 0.15).33 The mechanism for ferromagnetic behavior in Sn-doped MnSb2−xSnxSe4 samples can be well understood using the bound magnetic polarons model (BMP).33 The substitution of Sb3+ by Sn2+ in the FeSb2−xSnxSe4 system is anticipated to alter the carrier density while maintaining a constant concentration of the magnetic atoms. This enables a better assessment of the interactions of localized magnetic moments with charge carrier spin and the consequences of these interactions on the electronic and magnetic properties in ferromagnetic semiconductors. Polycrystalline powders of FeSb2−xSnxSe4(0 ≤ x ≤ 0.25) were synthesized by reacting stoichiometric mixtures of the elements at 775 K.33 The synthesized dark-gray polycrystalline powder of FeSb2−xSnxSe4 (0 ≤ x ≤ 0.25) was structurally characterized using X-ray powder diffraction and X-ray photoelectron spectroscopy (XPS). The phase purity and thermal stability of various samples were evaluated using differential scanning calorimetry (DSC), and their electronic and magnetic properties were investigated. Details of the synthesis procedure, structural characterization, and measurements of the electronic and magnetic properties were described elsewhere33 and can also be found in the Supporting Information.

Figure 2. Structural characterization of selected FeSb2−xSnxSe4 compositions. (A) Example of a Rietveld refinement result using the X-ray diffraction pattern of the sample with x = 0.01. (B) Refined lattice parameters as a function of Sn content. Distribution of Fe (C) and Sn (D) at M1, M2, M3, and M4 sites as a function of Sn content.

etry (DSC), which displayed a single peak of melting with an onset temperature of between 862 and 873 K (Figure 3A). The

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RESULTS AND DISCUSSION Crystal Structure. Polycrystalline powders of selected FeSb2−xSnxSe4 compositions were synthesized by the solidstate reaction of high-purity elements in the appropriate ratios. Careful comparison of the powder XRD patterns of all samples with the theoretical pattern calculated using single-crystal data of FeSb2Se4 (ICSD 421938)37 indicated an excellent match in peak positions (Figure 2A), suggesting that the synthesized FeSb2−xSnxSe4 compositions are isostructural with FeSb2Se4 and that all diffraction peaks can be indexed in the space group C2/m (no. 12). No obvious extra peak of an impurity phase could be detected on the XRD patterns of various compositions up to x = 0.2. The phase purity of all samples was further confirmed through differential scanning calorim-

Figure 3. (A) Differential scanning calorimetric heating curves of selected FeSb2−xSnxSe4 compositions revealing a single peak of melting between 862 and 873 K. X-ray photoelectron spectroscopy spectra of Sn 3d (B), Sb 3d (C), and Fe 2p (D) shell electrons for selected FeSb2−xSnxSe4 compositions.

difference between the melting-point onset temperature (∼10 K for the highest value observed) is within the tolerable range to conclude that substituting Sb with Sn in the FeSb2Se4 crystal lattice does not significantly affect the melting temperature (875 K)37 of the Fe1−xSb2SnxSe4 system for Sn concentrations of x ≤ 0.25. The effect of the substitution between Sb and Sn 9251

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of 0 ≤ x ≤ 0.25 confirmed the presence of all elements. The intensity of the Sn 3d5/2 peak at a binding energy of ∼486.4 eV, which is associated with Sn2+ in SnSe,42 consistently increases with increasing concentration of Sn in the FeSb2−xSnxSe4 structure (Figure 3B), confirming an effective incorporation of Sn atoms within the FeSb2Se4 crystal lattice for the 2+ oxidation state. The Fe 2p3/2 peaks are observed for all samples at a binding energy of 711.3 eV along with a broad satellite peak at 716.5 eV for compositions with x ≥ 0.15 (Figure 3D). This is consistent with binding energy associated with Fe2+ in FeO.43−45 The Sb 3d5/2 peak appears at ∼530.5 eV in the sample with x = 0 and shifts to higher values with increasing concentration of Sn. The observed binding energy for the Sb 3d5/2 peak in FeSb2Se4 is consistent with the reported values for MnSb2−xSnxSe4,33 Sb2Se3,5,33 and Sb2O3 (530.2 eV),46 implying a 3+ oxidation state for the Sb atom. Interestingly, the Sb 3d5/2 peak shape broadens with increasing concentration of Sn from x = 0.05 up to 0.25, and the binding energy shifts from ∼530.5 to ∼530.9 eV. Such peak shifting suggests that a fraction of the Sb3+ atoms have been oxidized to Sb5+ with increasing Sn concentration. The suggested partial oxidation is triggered by the local charge imbalance (excess negative charge) resulting from the substitution of Sb3+ by Sn2+ and the intermixing of these atoms at M1 and M2 positions in the crystal lattice. Such partial oxidation of the larger Sb3+ (IR: 76 pm) to the slightly smaller Sb5+ (IR: 60 pm) compensates for the anticipated increase in the lattice parameter due to the substitution of small Sb3+ (IR: 76 pm) by a larger Sn2+ (IR: 93 pm), leading to nearly constant lattice parameters for samples with x ≤ 0.15. For samples with higher Sn content (x ≥ 0.15), a significantly large fraction of Sb3+(4d105s25p0), which contains a stereoactive electron lone pair, is oxidized to Sb5+(4d105s05p0), which has no electron lone pair, diminishing the impact of the stereoactivity of the Sb lone pair on the lattice parameter. The active lone pair, which is oriented along the c axis, hampered lattice contraction along that direction for samples with x ≤ 0.15, and the removal of such stereoctivity through the oxidation of Sb3+ to Sb5+ well explains the observed sharp contraction of the c axis for samples with x ≥ 0.2. A similar effect was observed in MnSb2−xSnxSe4.33 Remarkably, the XPS spectrum of Sb 3d5/2 shell electrons for samples with x ≥ 0.2 also showed a shoulder peak at lower binding energy, 529 eV, which matches the binding energy of elemental Sb.47 This suggests that for samples with x > 0.15 a partial reduction of Sb3+ to elemental Sb0 also takes place simultaneously with the partial oxidation of Sb3+ to Sb5+. The effect of the substitution of Sb3+ by Sn2+ in FeSb2−xSnxSe4 and the resulting change in the chemical state of Sb from Sb3+ to Sb5+ and Sb0 (for samples with x ≥ 0.2) observed from XPS data on the electronic properties of the resulting substitution compound can be rationalized using eqs 1−5. The substitution of Sb3+ by Sn2+ in FeSb2−xSnxSe4 leads to the formation of acceptor states accordingly with eq 1.

on the lattice parameters, atomic positions, bond distances, and distribution of metal atoms (Fe, Sb, and Sn) within the crystal structure of FeSb2−xSnxSe4 compositions was investigated by Rietveld refinement on XRD patterns of the synthesized materials using FullProf.39,40 The fairly good agreement in both peak position and intensity between the observed and calculated XRD data for all compositions, as shown in Figure 2A as well as in Supporting Information Figure S1, attests to the good quality of the Rietveld refinement results, which is also supported by the relatively low values of the weighted-profile R factors and χ2 (Table S1). The refinement result revealed no significant change in the lattice parameters and unit cell volume (Figure 2B and Table S1) upon substituting Sb by Sn, which is surprising because the 22% larger effective ionic radius (IR) of Sn2+ (93 pm) in octahedral coordination compared to that of Sb3+ (76 pm)41 points to an increase in the lattice parameters and unit cell volumes. Astonishingly, the c parameter slightly contracts with increasing Sn content, and the largest drop (Δc/c = 7%) is observed for composition x = 0.2 (Figure 2B). In addition, an irregular trend in the local distortion of metal coordination environments can be noticed from the variation of both atomic positions and M−Se bond lengths within similar polyhedra in selected compositions (Tables S2 and S3). Although this unanticipated alteration of the structural parameters can be partially supported by the redistribution of metal atoms within various atomic sites upon increasing Sn content (Figure 2C,D), it also suggests a more complex variation in the chemical states (oxidation state) of elements within the crystal structure. The distribution of metal atoms at various positions was refined in the last step of the Rietveld refinement process assuming full occupancy of all atomic positions. The refinement results show that the M3(2d) site is exclusively occupied by Fe, regardless of the Sn content, which is consistent with the occupancy of the M3 site in the parent structure, FeSb2Se4.37 The remaining metal sites (M1(4i), M2(4i), and M4(2a)) show mixed occupancy with an interesting redistribution of Fe and Sn in various samples. For instance, starting with the parent structure (x = 0), Sb atoms are preferentially found at M1(Sb:96%/Fe:4%) and M2(Sb:96%/Fe:4%) sites, whereas Fe is predominantly at the M4(Sb:16%/Fe:84%) site.37 Upon substitution of Sb by Sn, the concentration of Fe gradually decreases at the M1 and M2 sites, whereas higher Fe content is found at the M4 sites (Figure 2C). This implies an increased ordering in the distribution of Fe atoms with increasing substitution of Sb by Sn. Full ordering of Fe at M3 and M4 sites is obtained for compositions with x ≥ 0.1. Simultaneously, Sn atoms preferentially occupied the M1 site for compositions with x ≤ 0.04 and started to fill both M1 and M2 sites in samples with x > 0.04 (Figure 2D). This variation in the distribution of Fe and Sn atoms within the crystal structure of samples with increasing Sn content can partially explain the irregular alterations of the structural parameters such as lattice constants and bond distances and is also anticipated to have a significantly impact on the magnetic and electronic properties of FeSb2−xSnxSe4 samples. To further investigate the origin of the unexpected changes observed in the unit cell parameters, we have probed the oxidation state of various cations (Sn, Fe, and Sb) in all FeSb2−xSnxSe4 compositions using X-ray photoelectron spectroscopy (XPS). The spectra (Figure 3B−D) for Sn 3d, Sb 3d, and Fe 2p shell electrons for various compositions in the range

xSb3 +

xSn 2 + ⎯⎯⎯⎯→ x(Sn 2 +)′Sb + x h.

(1)

The local excess negative charges resulting from the substitution of Sb3+ by Sn2+ triggers the partial oxidation of Sb3+ to Sb5+ (for compositions with x ≤ 0.15) or the simultaneous oxidation and reduction of Sb3+ to Sb5+ and elemental Sb0 for samples with higher Sn content. These oxidation and reduction reactions lead to the generation of 9252

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Journal of the American Chemical Society various concentrations of electrons according to the chemical reactions in eqs 2 and 3. y(Sb3 + → Sb5 + + 2e−)

(2)

z(Sb3 + + 3e− → Sb0 )

(3)

The global equation for this oxidation/reduction reaction is given by eq 4. (y + z)Sb3 + → ySb5 + + zSb0 + (2y − 3z)e−

(4)

The complete chemical equation of the substitution and the subsequent oxidation and reduction processes taking place within the FeSb2−xSnxSe4 crystal lattice is given by eq 5. Fe2 + + (2 − x)Sb3 + + xSn 2 + + 4Se 2 − → Fe2 + + (1 − y − z)(2 − x)Sb3 + + y(2 − x)Sb5 + + z(2 − x)Sb0 + x(Sn 2 +)′Sb + x h. + (2y − 3z) (2 − x)e−

(5)

with y = x + [3z(2 − x)]/2(2 − x). Therefore, the final chemical formula of the substituted composition is (Fe2 +)(Sb3 +)[(4 − 3x)/2 − 5z(2 − x)/2](Sb5 +)[x /2 + 3z(2 − x)/2]

Figure 4. Temperature dependence of the electronic properties of selected FeSb2−xSnxSe4 samples. (A and B) Electrical resistivity highlighting the electronic transition at 62 K from p-type semiconducting behavior to n-type metallic behavior for the sample with x = 0.15 (inset of 4B). (C and D) Thermopower showing a sign reversal at 62 K for the composition with x = 0.15. (E) Thermal conductivity. (F) Figure of merit.

−

. (VB) (Sb0 )z(2 − x)(Sn 2 +)x (Se 2 )4 (e−)[x − 3z(2 − x)/2](CB)(h) x

One can derive from the above equations that the substitution of x moles of Sb3+ by an equivalent amount of Sn2+ within the structure of FeSb2−xSnxSe4 induces the formation of [x/2 + 3z(2 − x)/2] moles of Sb5+ and z(2 − x) moles of elemental Sb0 along with the generation of both [x − 3z(2 − x)/2] moles of electrons and x moles of holes within the FeSb2−xSnxSe4 crystal lattice. Therefore, bipolar electronic conduction is anticipated for various FeSb2−xSnxSe4 compositions along with strong changes in the magnetic properties. Electronic Transport. To fully characterize the impact of the substitution of Sb by Sn on the electronic transport properties of FeSb2−xSnxSe4 samples, electrical resistivity (Figure 4A,B) and thermopower (Figure 4C,D) data were measured in the temperature range from 300 to 10 K. Regardless of the temperature, the electrical resistivity and thermopower of various samples decrease drastically with increasing Sn content, reaching the lowest values for x = 0.15 and slightly increasing with further increases in Sn content (x > 0.15). For example, the electrical resistivity of the sample with x = 0 is ∼200 Ω·m at 300 K and decreases drastically to 8 Ω·m for the sample with x = 0.02. The resistivity further drops to 0.1 Ω·m for samples with x ≥ 0.1 (Figure 4A,B). Likewise, the thermopower at 300 K is 217 μV·K−1 for FeSb2Se4 (x = 0) and then decreases to ∼150 μV·K−1 for x = 0.1. Increasing the Sn content to x = 0.15 leads to a drastic drop in the thermopower to 36 μV·K−1, and a further increase in Sn content to x = 0.2 induced a slight increase in the thermopower to 50 μV·K−1 (Figure 4C,D). The observed drastic drop in the electrical resistivity and thermopower of various FeSb2−xSnxSe4 samples with a small increase in the Sn content can be associated with the large increase in the carrier density (both electrons and holes) arising from the substitution of Sb3+ by Sn2+ and the accompanying oxidation−reduction processes as discussed above. Indeed, chemical equation 5 indicates the generation of equal numbers of holes and electrons, for compositions with

x ≤ 0.15, and an increase in the concentrations of both carrier types with Sn content (x value). For instance, the generation of ∼8.4 × 1020 electrons per cm3 and an equivalent density of acceptor states are anticipated for the composition with x = 0.15. However, for compositions with higher x values (x > 0.15), the presence of a Sb 3d5/2 peak at 529 eV in the XPS data, which is associated with elemental Sb0, suggests the generation of a slightly lower density of electrons ([x − 3z(2 − x)/2] moles, where z(2 − x) is the mole fraction of Sb0 formed) compared to that of holes (x moles). This analysis is in agreement with the alteration of the electrical resistivity and thermopower, where a decrease in the resistivity and thermopower with increasing Sn content up to x = 0.15 and the slight increase in both electrical resistivity and thermpower for the sample with x = 0.2 are observed. The large density of both electrons and holes generated in various FeSb2−xSnxSe4 samples and the observed sharp drop in the measured thermopower with a small increase in Sn content imply a bipolar conduction mechanism. Therefore, the overall electrical resistivity and thermopower measured correspond to the contributions of both holes and electrons to the electronic transport. Here, the total electrical conductivity that is measured is the sum of electron and hole contributions (σ = σ+ + σ−), whereas the thermopower measured at a given temperature is the weighted sum of the positive thermopower (S+) values from hole transport and the negative thermopower (S−) values from electron transport (ST = [σ+S+ + σ−S−]/(σ+ + σ−)). All samples show positive overall values of the thermopower within the temperature range from 300 to 60 K, suggesting holes as the overall majority carriers. The 9253

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maximum of 0.41 Ω·m at peak temperature Tp ≈ 62 K and sharply decreases to 0.22 Ω·m with further cooling of the sample to 6.7 K. This anomalous switching of the temperaturedependent electrical resistivity from the semiconducting behavior from 300 to 62 K to the metallic behavior in the temperature range below 62 K is a signature of an electronic topological transition known as the Lifshitz transition (LT), which is associated with the change in the Fermi surface (FS) topology due to the variation of electronic band structure or the change in the Fermi energy.48 A similar electronic transition was also observed at Tp = 19 K for the sample with x = 0.02 that was successfully measured to very low temperatures (Figure S2). A slightly higher electrical resistivity was observed at low temperatures upon increasing the Sn content to x = 0.2. For this sample, the electrical resistivity gradually increases with decreasing temperature and plateaus at ∼700 Ω·m at 20 K without a full reversal to metallic behavior. Further confirmation of the alteration in the electronic structure of FeSb2−xSnxSe4 samples with x ≥ 0.15 is obtained through the temperature-dependent measurement of the thermopower. Both samples (x = 0.15 and 0.2) initially show a slow decrease in thermopower with decreasing temperature up to a critical temperature (Tc) (from 36 μV·K−1 at 300 K to 30 μV·K−1 at 150 K for x = 0.15 and from 50 μV·K−1 at 300 K to 40 μV·K−1 at 90 K for x = 0.2), beyond which a sharp drop is observed (Figure 4D). This is a surprising trend given that the electrical resistivity, within the same temperature range, increases with decreasing temperature. Interestingly, the sample with x = 0.15 shows a sign reversal of the thermopower at 62 K, which also corresponds to the temperature of the maximum electrical resistivity curve (Figure 4B). Beyond this temperature, negative values of the thermopower (n type) were observed (−2 μV·K−1 at 40 K), which implies a switching of the majority carrier type from hole to electron. Astonishingly, the change in the majority carrier type vanishes upon increasing the Sn content to x = 0.2, as the thermopower remains positive down to 10 K (10.7 μV·K−1). This further points to the high sensitivity of the electronic structure of FeSb2−xSnxSe4 samples to doping. The observed larger electrical resistivity and thermopower for the sample with x = 0.2 compared to the samples with x = 0.15 is consistent with chemical equation 5 that predicts the generation of fewer electrons as a result of the formation of elemental Sb0 according to the XPS data. Origin of the Lifshitz Transition. The simultaneous change in the conduction behavior from semiconducting to metallic along with the rapid decrease in the thermopower with decreasing temperature and the reversal of the sign of majority charge carrier type from hole to electron observed at 62 K for the sample with x = 0.15 and at 19 K for the sample with x = 0.02 (Figure S2) strongly points to a rapid change in the Fermi surface topology with doping as well as temperature, which is a manifestation of the Lifshitz transition.48 Similar types of unusual transport properties (strong resistivity peak along with a sign reversal of majority carriers) were reported in iron pnictides,49 transition-metal pentatellurides (ZrTe5, HfTe5, etc.),28,50 high-temperature superconductors,51,52 and strongly correlated electron systems such as NaxCoO2.53 However, the precise origin of such anomalous transport properties remains a subject of longstanding debate. Several mechanisms have been proposed for the rational understanding of the observed transport anomalies. These include bipolar conduction,54 structural phase transition,55 the formation of a charge density

observed irregular trend in the variation of the electrical conductivity and thermopower of various FeSb2−xSnxSe4 with increasing Sn content can be attributed to the alteration in the distribution of Fe and Sn atoms in the crystal structure of various samples. One interesting feature of the bipolar electronic transport in various FeSb2−xSnxSe4 samples is the observation of multiple discontinuities and transitions on the temperature-dependent electrical resistivity and thermopower curves. For example, Figure 4A,B revealed a striking variation in the magnitude of the electrical resistivity of various samples upon cooling from 300 to 50 K. At 300 K, an electrical resistivity of ∼200 Ω·m was measured for the sample with x = 0. The electrical resistivity gradually increases to ∼400 Ω·m at 200 K and then suddenly increases by several orders of magnitude, exceeding ∼1.4 × 106 Ω·m at 50 K. Temperature-dependent electrical resistivity curves of Sn-substituted samples with x = 0.02 and 0.1 showed a similar trend. However, the magnitude of the electrical resistivity at a given temperature and the temperature at which the sharp increase in the electrical resistivity is observed gradually decreases with increasing Sn content. For instance, a close examination of the temperature-dependent electrical resistivity curve of samples with x = 0.02 and 0.1 revealed discontinuities at 180 and 60 K, respectively, where a small drop in the electrical resistivity is observed. Likewise, a large increase in the overall thermopower was observed for samples with x = 0, 0.02, and 0.1 upon cooling from 300 to 60 K. The thermopower for the sample with x = 0 increases from ∼217 μV·K−1 at 300 K to ∼400 μV·K−1 at 60 K. Likewise, the thermopower of Sn-substituted samples (x = 0.02 and 0.10) rapidly increases upon cooling, reaching a maximum of 380 μV·K−1 at 90 K for the sample with x = 0.02 and 300 μV·K−1 at 80 K for the sample with x = 0.1. The observed sudden increase in both the electrical resistivity and thermopower of the samples with x = 0, 0.02, and 0.1 could be associated with the large localization of charge carriers upon cooling. However, this picture is not consistent with the drop in the thermopower upon cooling to below 80 K. Therefore, the observed temperature dependence of the electronic transport in FeSb2−xSnxSe4 samples may also be associated with the onset of an interesting topological electronic transition as discussed below. The electrical resistivity and thermopower curves for the samples with x ≥ 0.15 surprisingly showed a completely different temperature dependence and an abrupt change with increasing Sn content compared to samples with x ≤ 0.1. For example, the electrical resistivity and thermopower at 100 K drastically drop from 49 Ω·m and 265 μV·K−1 for the sample with x = 0.1 to 0.27 Ω·m and 17.15 μV·K−1 for the sample with x = 0.15, respectively. In addition, the sample with x = 0.15 showed only a marginal increase in the electrical resistivity upon cooling from 300 K (0.08 Ω·m) to 100 K (0.27 Ω·m), whereas the electrical resistivity of the sample with x = 0.1 increases by nearly 600% within the same temperature range (Figure 4A). The abrupt drop in both electrical resistivity and thermopower over a wide temperature range with a small increase in the Sn content from x = 0.1 to 0.15 suggests a significant change in the underlying electronic structure of FeSb2−xSnxSe4 samples with x ≥ 0.15. Such an alteration of the electronic structure can be probed through measurements of macroscopic physical properties such as temperature-dependent electrical resistivity and thermopower. As can be seen from Figure 4B, the electrical resistivity initially increases with decreasing temperature from 0.08 Ω·m at 300 K up to the 9254

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Journal of the American Chemical Society wave,56 polaronic behavior,57 a semimetal−semiconductor transition,58 or a temperature-induced band shift.59−61 Building on our current understanding of the Lifshitz transition in semiconductors, the anomalous transport properties observed for various FeSb2−xSnxSe4 samples can be rationalized within the context of bipolar conduction coupled with a downshift of the overall electronic structure with decreasing temperature as detailed below. The substitution of Sb3+ by Sn2+ leads to the creation of acceptor states (extra holes), which are located within the band gap and just above the valence band (VB). The density of these acceptor states increases with increasing Sn content (x values). At sufficiently high concentrations, x > 0.1, these acceptor states form a holelike sub-band, hereafter denoted the lower band (LB) within the band gap (Figure 5). Simultaneously, the local charge

However, the UB downshifts gradually upon cooling, touching the Fermi energy level at 62 K for the sample with x = 0.15 (Figure 5B). At this point, the topology of the Fermi surface transitions from hole-like to electron-like, which is manifested in the macroscopic properties by a peak on the resistivity curve and the transition from p-type semiconducting to n-type metallic behavior. Further cooling to below 62 K results in the UB crossing the Fermi energy level, leading to the observed ntype metallic conductivity at low temperatures. Although the proposed band picture explains very well the anomalies observed in the temperature-dependent electrical resistivity and thermopower of FeSb2−xSnxSe4 samples, a direct visualization of changes in their electronic structure with doping level and temperature using high-resolution angle-resolved photoemission spectroscopy (ARPES)62 is highly desired to confirm the proposed band structure picture. Temperaturedependent evolution of the overall electronic band structure was observed in WTe260 and ZrTe5.50 High-resolution angleresolved photoemission spectroscopy (ARPES) study of WTe2 and ZrTe5 showed that two branches of bands (upper band (UB) and lower band (LB)) with linear dispersion that are separated by an energy gap dominate the electronic conductions in both compounds. It was found that the overall electronic structure shifts downward (lower energy) by as much as 70 meV upon cooling from 255 to 2 K. The energy gap between the upper band (UB) and lower band (LB) decreases but persist down to very low temperature.50 By analogy, one can assimilate the UB and LB bands observed in both WTe 2 and ZrTe 5 to the sub-bands formed in FeSb2−xSnxSe4 samples by donor and acceptor states, respectively. This implies that the Lifshitz transition can potentially be switched “on” and “off” on demand in any narrow band gap semiconductor through the engineering of the electronic band structure near the Fermi level via doping so as to create both hole-like (LB) and electron-like (UB) subbands within the band gap. Thermal Transport and ZT. To fully evaluate the impact of Sb substitution by Sn on the thermoelectric performance of FeSb2−xSnxSe4 samples, thermal conductivity data were collected on hot pressed pellets of selected compositions within the temperature range from 300 to 2 K. The total thermal conductivity of the pristine sample (x = 0) is ∼2 W· K−1·m−1 at 300 K and rapidly increases with decreasing temperature reaching a maximum of ∼3.4 W·K−1·m−1 at 180 K (Figure 4E). A further decrease in temperature resulted in a drastic drop in the thermal conductivity to ∼0.5 W·K−1·m−1 at 90 K, which remains constant down to 2 K. Interestingly, the magnitude of the total thermal conductivity and the temperature at which the maximum total thermal conductivity is observed decrease rapidly with increasing Sn content (Figure 4E). For instance, the total thermal conductivity at 300 K drops to 1.5 W·K−1·m−1 for the samples with x = 0.02 and further decreases to 1 W·K−1·m−1 for the samples with x = 0.1. The total thermal conductivity of the sample with x = 0.02 remains nearly constant at 1.5 W·K−1·m−1 upon cooling from 300 to 225 K and then rapidly increases with further cooling to a maximum value of 2.5 W·K−1·m−1 at 150 K and finally drops to 0.5 W·K−1·m−1 with further cooling to 50 K. Further increasing the Sn content to x = 0.1 resulted in a drop in the peak thermal conductivity to 1.5 W·K−1·m−1 at 125 K. The overall thermoelectric figure of merit (ZT) of FeSb2−xSnxSe4 samples is very low (Figure 4F) because of the extremely large electrical resistivity of the pristine sample and the bipolar

Figure 5. Schematic illustration of the electronic band structure of FeSb2−xSnxSe4 samples. (A) The substitution of Sb3+ by Sn2+ leads to the formation of two subbands within the band gap, a lower band (LB) of acceptor states, and an upper band (UB) of donor states. (B) Proposed downshift, upon cooling, of the overall band structure that explains the change in the Fermi surface topology from hole-like to electron-like (Lifshitz transition).

imbalance resulting from the substitution of Sb3+ by Sn2+ triggers the partial oxidation of Sb3+ to Sb5+, which is accompanied by the generation of donor states (extra electrons) that are located within the band gap. These donor states, for sufficiently high x values (x > 0.1), also form a subband, denoted the upper band (UB), just below the conduction band (CB). The resulting electronic band structure for FeSb2−xSnxSe4 samples with x ≥ 0.15 contains two subbands (UB (electron-like) and LB (hole-like)) that are separated by a small energy gap (Figure 5) within the narrow band gap (Eg ≈ 0.2 eV) of the FeSb2Se4 phase.37 Therefore, the topology of the Fermi surface in the resulting band structure and hence the transport behavior are dictated by the relative position of the sub-bands (UB and LB) with respect to the Fermi energy (EF). It should be noted that the relative position of the sub-bands with respect to the Fermi energy level depends on the doping level (the width of the sub-band increases with x values) but most importantly on the temperature, assuming a temperature-dependent evolution of the overall band structure. Within this proposed band picture, the semiconducting behavior of the temperature-dependent resistivity as well as the gradual slow decrease of the thermopower observed from 300 to 62 K for samples with x ≥ 0.15 implies that at 300 K the Fermi energy is located within the energy gap (Figure 5A) between the UB and LB but closer to the LB (hole-like Fermi surface), which leads to bipolar conduction and an overall p-type semiconducting behavior. 9255

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magnetic atom (Fe) is essentially the same at all compositions, which should lead to comparable magnetic susceptibility values. Therefore, one can speculate that the observed large change in the measured magnetic susceptibility of various samples takes into account contributions from free carriers as well as the change in the atomic ordering of Fe atoms in various samples (Figure 2C). Indeed, the increase in the magnitude of the susceptibility observed on various samples with increasing Sn content is in agreement with the observed drop in the electrical resistivity. This further confirms the effect of increasing carrier concentration on the overall magnetic susceptibility. Further discussion of a possible mechanism leading to the observed magnetic property is provided below. The ZFC susceptibility for the samples with x ≤ 0.1 gradually increases with decreasing temperature up to 130 K, sharply drops within a narrow temperature range between 130 and 110 K, and then continues to decrease slightly down to 40 K, beyond which another sharp drop is observed with further cooling (Figure 6B). The observed transition in the susceptibility at 130 K for both samples (x = 0.05 and 0.1) is similar to the one previously reported in FeSb2Se437 and FeSb2−xInxSe429 phases and is in agreement with the similarity in the separation between magnetic centers (Fe atom) within the crystal lattice of FeSb 2−x Sn x Se 4 (x ≤ 0.1) and FeSb2−xInxSe4 phases. It is believed that such a magnetic transition at 130 K is related to the anisotropic distortion of the [FeSe6] octahedra upon cooling,37 whereas the sharp drop in the susceptibility below 40 K suggests the presence of AFM (antiferromagnetic) ordering within the [FenSe4n+2] magnetic chain that competes with the dominating FM exchange interaction between adjacent magnetic chains in the bulk crystal. Interestingly, the shape of the temperature dependence of the ZFC susceptibility curve drastically changes for samples with Sn content x > 0.1. The drop in the ZFC susceptibility below the magnetic transition at 130 K is weak for the composition with x = 0.15 and nearly fades away for the sample with x = 0.2 (Figure 6A). This is a very significant result that further supports our hypothesis that the stereoactivity of the electron lone pair on the Sb atom contributes to the sharp transition in the magnetic susceptibility at 130 K.37 Indeed, as observed from the XPS data (Figure 3B), the fraction of such lone pairs decreases with increasing Sn content due to the oxidation of a large fraction of Sb3+ to Sb5+. This severely diminishes the effect of the lone pair, which is manifested by a contraction of the c axis and a slight increase in the a axis for samples with x ≥ 0.15. This change in the lattice parameters leads to a more regular shape of the [FeSe6] octahedra. Therefore, the observed fading of the magnetic transition at 130 K for FeSb2−xSnxSe4 samples with x > 0.1 implies that the regular shape of the [FeSe6] octahedra remains upon cooling below 130 K, which inhibits the redistribution of spin within the 3d6 orbitals. The analysis of FC susceptibility curves revealed a large difference in both the shape and magnitude compared to the ZFC curve from the same sample. For example, the FC magnetic susceptibility for compositions with x ≤ 0.10 increases with decreasing temperature down to 130 K, where the magnetic transition is observed. While the samples with composition x = 0 and 0.05 showed a drop in the susceptibility as anticipated, the magnitude of such a change in the susceptibility is significantly smaller than that of the ZFC. Interestingly, the sample with x = 0.1, instead of the anticipated decrease in the magnetic susceptibility, exhibits a

conduction resulting from the substitution of Sb by Sn. The very high electrical resistivity of the pristine sample (x = 0) leads to a very low ZT value of ∼7.0 × 10−7 at 300 K despite the large thermopower value (∼217 μV·K−1) and low total thermal conductivity (2 W·K−1·m−1). However, partial substitution of Sb atoms by Sn atoms leads to a significant drop in the electrical resistivity and total thermal conductivity and a moderate reduction of the thermopower owing to bipolar doping. This results in a marginal improvement in the overall figure of merit of FeSb2−xSnxSe4 samples with increasing Sn content. Magnetism. The magnetic properties of magnetic semiconductors and correlated electronic materials are very sensitive to the alteration of carrier density as well as the separation between magnetic atoms within the crystal. Therefore, changes in the carrier density of FeSb2−xSnxSe4 samples upon substituting Sb by Sn are anticipated to play a significant role in the observed overall magnetic behavior. To investigate the impact of Sb3+ to Sn2+ substitution on the magnetism in an FeSb2−xSnxSe4 solid solution, we have measured the field-cooled (FC) and zero-field-cooled (ZFC) temperature-dependent magnetic susceptibility of selected compositions under a 100 Oe applied magnetic field (Figures 6 and S3) in the temperature range from 2 to 300 K.

Figure 6. Variation of the magnetic susceptibility of selected FeSb2−xSnxSe4 samples with temperature under zero-field-cooled (ZFC) (A and B) and field-cooled (FC) (C and D) conditions using an external applied field of 100 Oe.

Regardless of the temperature, the FC and ZFC susceptibilities increase in magnitude with increasing Sn content (x value) from 5 to 20%. For instance, the FC susceptibility at 300 K is 0.36 emu·mol−1·Oe−1 for the sample with x = 0 and increases to 0.67, 0.70, 1.01, and 1.42 emu·mol−1·Oe−1 for compositions with x = 0.05, 0.1, 0.15, and 0.2, respectively. While the change in the FC susceptibility with increasing Sn content is marginal at 300 K, a strong divergence of the FC susceptibility curves for various compositions is observed upon cooling. For example, the FC susceptibility at 60 K is 0.51 emu· mol−1·Oe−1 for the pristine sample (x = 0) and increases to 0.87, 1.70, 2.57, and 4.64 emu·mol−1·Oe−1 for compositions with x = 0.05, 0.1, 0.15, and 0.2, respectively. This result is quite surprising considering that the concentration of the 9256

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magnetic field between −10 kOe and +10 kOe. The S shape of the magnetization versus applied field (M vs H) curves for various samples at 2 and 300 K suggests the persistence of ferromagnetic character in the samples up to room temperature. Increasing the Sn content induces a marginal drop in the coercivity simultaneously with a remarkable increase in the saturation magnetization leading to steeper magnetization curves for compositions with x ≥ 0.15. For instance, the saturation magnetization at 2 K gradually increases with increasing Sn content, starting from 555 emu·mol−1 for the pristine sample (x = 0), reaching a maximum of ∼1400 emu· mol−1 for the sample with x = 0.15, and decreasing to ∼900 emu·mol−1 with further increases in the Sn content to x = 0.2. A similar trend was obtained for the magnetization curves at 300 K. Simultaneously, the coercive field (Hc) at 2 K marginally decreases with increasing Sn content from 180 Oe for x = 0 to 160 Oe for x = 0.15 and drops further to 66 Oe for the sample with x = 0.2. This implies that the large increase in the density of carriers (holes and electrons) arising from the substitution of Sb by Sn gradually transforms the FeSb2−xSnxSe4 samples from “semi-hard” magnetic materials to “soft” magnetic materials, which is characterized by a relatively facile spontaneous magnetization and spin reversal upon application of a weak external magnetic field. The observed trend in the alteration of the saturation magnetization and the coercivity with increasing Sn content is consistent with the change in the electrical resistivity with increasing Sn content, as discussed above, where the minimum resistivity is achieved for x = 0.15. This also indicates a strong correlation between electronic transport and magnetic ordering in FeSb2−xSnxSe4 samples. Another interesting result in this study is the surprisingly larger values of the coercivity at 300 K for various FeSb2−xSnxSe4 samples compared to the coercivity at 2 K (Figure 6B,D). For instance, the sample with x = 0 shows a coercivity of Hc ≈ 255 Oe at 300 K compared to 180 Oe at 2 K. Similarly, the coercivity for the sample with x = 0.15 increases to 300 Oe at 300 K from 160 Oe at 2 K. This anomalous increase in the coercivity with rising temperature suggests that the strength of magnetic coupling is mediated by the interactions of localized magnetic moments with free spin carriers in FeSb2−xSnxSe4 samples, which inhibits the drop in coercivity with rising temperature as is normally anticipated from thermal spin fluctuations. To probe the effect of Sb substitution by Sn on the Curie transition temperature of FeSb2−xSnxSe4 samples, high-temperature magnetic susceptibility data were collected under fieldcooled (FC) and zero-field-cooled (ZFC) conditions. ZFC data were collected after thermal treatment consisting of heating and cooling the sample between 300 and 600 K without any applied external field. The susceptibility measurement was performed during heating from 300 to 600 K under an applied field of 100 Oe. Then, the sample was cooled to 300 K under an external applied field of 100 Oe, and the FC susceptibility data was measured on heating with the same external applied field. The resulting FC and ZFC data for selected samples are shown in Figure 8. It can be noted from the FC data that all samples maintain ferromagnetic-like character well above 300 K. For example, the FC susceptibility curve for the sample with x = 0 shows a rapid drop with rising temperature from 300 to 325 K followed by a weaker drop in the susceptibility with a further increase in temperature up to 450 K, above which a constant value of the susceptibility is measured up to 600 K. This general shape of the FC

surprisingly large increase in susceptibility within a narrow temperature range from 130 to 110 K and then plateaus down to 40 K, below which a rapid drop in the susceptibility is observed (Figure 6D). The drop in the susceptibility below 40 K for the sample with x = 0.1 suggests a competition between the AFM intrachain exchange interaction (FenSe4n+2 magnetic chains) and the dominant interchain FM exchange interaction in FeSb2−xSnxSe4. Remarkably, the magnetic transition below 130 K completely vanishes from the FC susceptibility upon increasing the Sn content to x ≥ 0.15. The FC susceptibility curves of these samples show a continuous increase in the susceptibility with decreasing temperature from 300 K down to 2 K (Figure 6C). This is typical behavior for ferromagnetic semiconductors and suggests a dominant interchain FM exchange interaction in FeSb2−xSnxSe4. In addition, the larger magnitude of the FC susceptibility compared to the ZFC data at the same temperature further suggests a strong contribution from spin-polarized free carriers, which act through the formation of bound magnetic polarons (BMPs) with localized magnetic moments on FenSe4n+2 magnetic chains. The observed alteration in the nature of the magnetic transition below 130 K could be related to the variation in the distribution of the magnetic atom (Fe) within the crystal structure. As shown in Figure 2C, an ordering of Fe atoms exclusively within the M3 and M4 sites is observed for samples with x ≥ 0.1, whereas a small amount of Fe is also found at M1 and M2 sites for composition with x < 0.1. It originates from the above discussion that the magnetic transition at 130 K is related to the distribution of Fe atoms within the crystal lattice in addition to the contribution from the stereoactivity of the Sb3+ electron lone pair, whereas the magnitude of the susceptibility is controlled by the Sn content, which in turn determines the concentration of holes and electrons generated within the crystal lattice. Additional characterization of the nature of ferromagnetism in the synthesized FeSb2−xSnxSe4 samples was achieved through isothermal field-dependent magnetization measurements at 2 and 300 K (Figure 7) spanning the applied external

Figure 7. Isothermal field-dependent magnetization of selected FeSb2−xSnxSe4 samples collected at 2 K (A and B) and 300 K (C and D) showing ferromagnetism up to 300 K. Note the strong dependence of magnetic parameters, such as the coercivity (Hc) and saturation magnetization (Ms), on the Sn content. 9257

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field of 100 Oe start from negative values at 300 K (−1.3 emu· mol−1·Oe−1 for x = 0 and ∼−0.1 emu·mol−1·Oe−1 for x = 0.1) and rapidly increase with temperature up to 325 K where a change in the slope is observed. The magnitude of the susceptibility continues to increase with further increases in temperature, crosses over to positive values at Tcross ≈ 325 K (for x = 0.1) and 410 K (for x = 0), and plateaus at 0.27 emu· mol−1·Oe−1 (x = 0) or 0.15 emu·mol−1·Oe−1(x = 0.1) in the temperature range from 450 to 600 K. One plausible origin of the observed negative magnetic susceptibility is the formation of a trapped internal magnetic field (TIMF) in the sample during cooling under zero applied field from 600 K. It appears that upon heating the samples to 600 K under no external magnetic field the randomization of magnetic moments from BMPs is achieved. During cooling under no external magnetic field, the alignment of magnetic moments is dictated by the strength of the local magnetic field on the individual magnetic sublattice, layers A and B (Figure 1). Because the concentration of the magnetic atom (Fe) within the M3 site in layer A is larger than that of the M4 site in layer B, for FeSb2−xSnxSe4 compositions with x ≤ 0.1 (Figure 4) a larger fraction of randomized magnetic moments will align along the direction dictated by the localized moment on M3 sites (layer A). Therefore, the TIMF could originate from the nonzero net magnetic moment resulting from the antiferromagnetic exchange interaction between the two ferromagnetic sublattices corresponding to layers A and B (Figure 1) as observed in our previous investigation of the FeSb2−xInxSe4 phases.29 The magnitude and orientation of such TIMF could significantly impact the sign and magnitude of the zero-fieldcooled magnetic susceptibility data under various applied external magnetic fields. For example, a large negative magnetic susceptibility will be measured if the magnitude of the external applied field is smaller than that of the TIMF and its orientation is opposite to that of the TIMF. This analysis is consistent with the observed drop in the magnitude of the negative ZFC magnetic susceptibility upon increasing the Sn content from −1.3 emu·mol−1·Oe1− at 300 K (for x = 0) to (−) to −0.1 emu·mol−1·Oe1− at 300 K (for x = 0.1). The TIMF vanishes for FeSb2−xSnxSe4 compositions with x > 0.1 because the full occupancy by Fe of both M3 and M4 sites in layers A and B, respectively, leads to similar magnetic moments for both sublattices.

Figure 8. High-temperature zero-field-cooled (ZFC) and field-cooled (FC) magnetic susceptibility of selected FeSb2−xSnxSe4 samples measured under an external field of 100 Oe. Zero-field-cooled (ZFC) data were collected after heating the as-synthesized samples from RT to 600 K under a zero applied external field.

susceptibility curve with the two magnetic transitions at 325 and 450 K is also observed for the sample with x = 0.1 (Figure 8A). However, the magnetic transition at 450 K vanishes with further increases in the Sn content to x ≥ 0.15, and only the transition at 325 K is observed on the FC susceptibility curves. This strong alteration in the general shape and magnitude of the temperature-dependent susceptibility can be correlated to the distribution of the magnetic metal atom (Fe) at various crystallographic sites within the crystal structure of FeSb2Se437 as well as the magnetic interaction of localized magnetic moments with charge carrier spin. In a previous study on Fe1−ySnySb2Se4 phases, we showed that a large separation between magnetic centers, achieved through the substitution of Fe by Sn, resulted in a large drop in the Curie temperature, Tc, from 450 K for the composition with y = 0 to 325 K for the structurally ordered sample with y = 0.13.32 In the present work, we found that a similar manipulation of the separation between magnetic centers within the structure of FeSb2−xSnxSe4 samples can be achieved through the preferential distribution of Sn at M1 and M2 sites whereas Fe atoms are ordered within M3 and M4 sites (Figure 2C). Therefore, the two magnetic transitions on the temperature-dependent FC susceptibility of samples with x ≤ 0.1 can be associated with local magnetic ordering within individual magnetic chains at 325 K through the interaction with charge carriers and the formation of magnetic polarons and the subsequent long-range collective ordering at 450 K of overlapping adjacent magnetic polarons (BMPs).32 For samples with x ≥ 0.15, the long-range ordering vanishes as a result of the larger separation between adjacent magnetic polarons. Therefore, only the intrachain magnetic ordering at 325 K (Figure 8B) is observed. Another surprising magnetic behavior in FeSb2−xSnxSe4 samples is the observation of negative values of the ZFC magnetic susceptibility for compositions with x ≤ 0.1 (Figure 8C,D). ZFC susceptibility data were collected after a thermal treatment, under zero applied field, that consisted of heating the samples to 600 K followed by cooling to 300 K. ZFC susceptibility data measured during heating under an applied

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CONCLUSIONS We have systematically investigated the impact of the partial substitution of Sb by Sn on the magnetic and electronic behavior of p-type ferromagnetic semiconductor FeSb2−xSnxSe4. We found that the charge imbalance arising from the substitution of Sb3+ by Sn2+ triggers the partial oxidation of Sb3+ to Sb5+, leading to simultaneous hole and electron doping of the FeSb2−xSnxSe4 matrix. This presumably leads to the formation of two subbands within the band gap of FeSb2−xSnxSe4: the upper band formed by donor states and the lower subband formed by acceptor impurities. Such bipolar doping results in a very interesting temperature-induced quantum electronic transition, which is manifested by the presence of an anomalous peak on the temperature-dependent resistivity curve at 62 K simultaneously with a reversal of the sign of the majority charge carriers from holes (above 62 K) to electrons (below 62 K) for the FeSb2−xSnxSe4 sample with x = 0.15. Such an electronic transition is reminiscent of the Lifshitz transition, which is associated with the change in the Fermi 9258

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surface topology from hole-like to electron-like upon cooling from 300 to 2 K. It was also found that increasing the Sn content leads to a preferential ordering of Fe atoms at both M3 and M4 metal positions for compositions with x > 0.1. This results in an increase in the separation between magnetic centers, leading to a decrease in the Curie temperature from 450 K for samples with x ≤ 0.1 to 325 K for compositions with x > 0.1. This result is significant because it unambiguously settles the debate on the correlation of the Curie transition temperature, Tc, with the concentration of magnetic centers on one hand and the separation between magnetic centers on the other hand. It is found that the separation between magnetic centers rather than the concentration of magnetic atoms dictates the magnitude of the Curie temperature. Interestingly, the difference in the fraction of the magnetic atom, Fe, within the sublattices, layers A and B, leads to the observation of negative ZFC magnetic susceptibility at high temperatures in FeSb2−xSnxSe4 samples with x ≤ 0.1.

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S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/jacs.9b01884.

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Article

Experimental section, results from the Rietveld refinement of the X-ray powder patterns of FeSb2−xSnxSe4 samples in the monoclinic space group, distribution of atoms at various metal positions for selected FeSb2−xSnxSe4 compositions, interatomic distances for selected FeSb2−xSnxSe4 samples from Rietveld refinement, Rietveld refinement results for selected FeSb2−xSnxSe4 samples indicating excellent agreement between the measured and calculated X-ray diffraction patterns, temperature-dependent electrical resistivity and thermopower for FeSb0.98Sn0.02Se4 and FeSb0.85Sn0.15Se4, and temperature-dependent magnetic susceptibility of selected FeSb2−xSnxSe4 samples under zero-field-cooled) and field-cooled conditions using an external applied field (PDF)

AUTHOR INFORMATION

Corresponding Author

*[email protected] ORCID

Xiaoyuan Zhou: 0000-0002-0930-1278 Pierre F. P. Poudeu: 0000-0002-2422-9550 Notes

The authors declare no competing financial interest.

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ACKNOWLEDGMENTS This work was supported by the National Science Foundation under award no. DMR-1561008. P.F.P.P. and C.U. gratefully acknowledge the financial support from Department of Energy, Office of Basic Energy Science under award no. DE-SC0018941 for electronic and thermal transport measurements. J.S.L. acknowledges financial support from a National Science Foundation Graduate Research Fellowship under award no. DGE 1256260. Magnetic data was measured on a SQUID magnetometer at the University of Michigan, which was purchased using MRI grants from the NSF (DMR-1428226 and CHE-104008). 9259

DOI: 10.1021/jacs.9b01884 J. Am. Chem. Soc. 2019, 141, 9249−9261

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