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Role of crystal transformation on the enhanced thermoelectric performance in Mn-doped Cu2SnS3
Accepted Manuscript Role of crystal transformation on the enhanced thermoelectric performance in Mndoped Cu2SnS3 Zhen Zhang, Huiwen Zhao, Yifeng Wang, Xiaohui Hu, Yinong Lyu, Changchun Cheng, Lin Pan, Chunhua Lu PII:
S0925-8388(18)34461-X
DOI:
https://doi.org/10.1016/j.jallcom.2018.11.329
Reference:
JALCOM 48555
To appear in:
Journal of Alloys and Compounds
Received Date: 15 July 2018 Revised Date:
20 November 2018
Accepted Date: 24 November 2018
Please cite this article as: Z. Zhang, H. Zhao, Y. Wang, X. Hu, Y. Lyu, C. Cheng, L. Pan, C. Lu, Role of crystal transformation on the enhanced thermoelectric performance in Mn-doped Cu2SnS3, Journal of Alloys and Compounds (2018), doi: https://doi.org/10.1016/j.jallcom.2018.11.329. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Role of crystal transformation on the enhanced thermoelectric performance in Mn-doped Cu2SnS3 Zhen Zhang,a Huiwen Zhao,a Yifeng Wang,a,b* Xiaohui Hu,a,b Yinong Lyu,a* Changchun Cheng,a Lin Pana,b, Chunhua Lua,b
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a. College of Materials Science and Engineering, Nanjing Tech University, Nanjing 210009, China. b. Jiangsu Collaborative Innovation Center for Advanced Inorganic Function Composites, Nanjing Tech University, Nanjing 210009, China.
Abstract
Mohite-type Cu2SnS3 exhibits a large degree of manipulation in both electronic and phonon
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transport properties by transition metal doping. In this work, the effects of Mn doping on the transport properties of Cu2SnS3 were investigated. A crystal transition occurred upon doping from cation-ordered monoclinic structure to cation-disordered cubic and tetragonal structures. Such a
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crystal transition caused a significant increase of DOS effective mass of carriers, m*, which should be caused by the improved degeneracy of the bands at the Fermi level due to the improved symmetry, and also by the participation (in the upper valence bands) of unfilled d-orbitals of Mn cations that share crystallographically equivalent sites with Cu cations. Meanwhile the lattice thermal conductivity (~0.4 W m-1K-1 at ~773K) was suppressed effectively, which should also be linked to the disorder structure induced by Mn-doping. Benefitting from the increased carrier
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density and much improved m*, a remarkable power factor of 0.92 Wm-1K-2 has been obtained at 723 K. The highest ZT ~0.68 at 723 K is almost 15 times increased as compared to that of pristine Cu2SnS3.
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Keywords: Cu2SnS3, d-unfilled dopant, crystal transition, thermoelectric performance
ACCEPTED MANUSCRIPT Introduction Thermoelectric (TE) materials have garnered extensive interests and shown promising prospects in applications of energy conversion and semiconductor cooling without using moving parts nor greenhouse gases emission [1,2]. The performance of TE materials is quantified by the dimensionless figure of merit ZT, defined as ZT = S2σT/κ = S2σT/(κe+κl), where S is the Seebeck
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coefficient, σ is the electrical conductivity, T is the absolute temperature, and κ is the thermal conductivity which comprises of the electronic thermal conductivity (κe) and lattice thermal conductivity (κl) [3,4]. High TE efficiency requires high ZT values of materials to become comparable with other energy conversion routes and to expand the application fields. For this purpose, a large power factor (PF = S2σ) and a low κ should be secured. However, S, σ and κ are
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strongly coupled through the carrier concentration (n) [5-7]. It is still one of the biggest challenges to independently optimize these three parameters to obtain high-performance TE materials with large ZT values.
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Recently, as an attractive earth abundant material, p-type ternary sulfide Cu2SnS3 (CTS) has been found excellent not only for absorber in solar cells [8, 9] but also as a potential eco-friendly TE material with phonon-glass-electron-crystal features, with peak ZTs around 0.6~0.85 at 723 K to date [10,11]. CTS has polymorphic crystal structures including monoclinic, tetragonal and cubic, all of which can be viewed as a diamond-like network of inter-connecting S-centered tetrahedra of SM4 (M = Cu and Sn) with varied arrangement and symmetry. The complex crystallographic
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network and soft bonding between the metal and S atoms with diverse and even random coordination endue CTS a very low κl [12-14]. Furthermore, in the monoclinic phase of CTS, there is a 3-dimensional conductive lattice network, with the upper valence bands (VB) mainly composed of hybrid Cu 3d orbitals and S 3p orbitals (while containing no obvious contribution from Sn [12]), which ensures the improvement of σ in CTS by way of doping to control the carrier
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density. However, S is inevitably reduced as σ increases, which still impedes the significant improvement of ZT value, since Sn-site acceptor doping seems to minimally affect the upper VB
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shape and only denotes holes in the monoclinic structure [12]. Nevertheless, when it transforms into a high symmetry tetragonal and cubic structure,
appropriate dopant atoms at Sn-site, which are crystallographically equivalent to Cu atoms, may have contributions to enhance the density-of-states (DOS) of the upper VB, while such a phase transformation can be readily realized by transition metal (TM, e.g. Zn, Ni, and Co) doping. Moreover, as revealed in our previous work, doping with 3d-unfilled TM, e.g. Co [10], seemed much more favorable to improve the TE performance than Zn-doping as a result of the enhanced DOS effective mass of carriers m* attributed to the participation of their additional 3d states in the VB. Effects that would be induced in terms of the more unfilled orbitals contained in Mn cations in the present study. Results revealed an order-to-disorder crystal transformation of CTS from monoclinic to cubic and tetragonal phases upon doping, a greatly enhanced DOS effective mass
ACCEPTED MANUSCRIPT m* together with an increase in n. While the κl was suppressed largely, due to the enhanced phonon scattering attributed to disordered cation arrangement and other defects related to the phase transformation or doping. Finally, a maximal ZT ~0.68 was obtained at 723 K, which is about 15 times higher than that of the pristine CTS. It strongly suggests the significant role of crystal transformation due to d-unfilled TM doping for CTS and the promising prospective of CTS as a
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good mid-temperature TE material.
Experimental
High purity powders of Cu, Mn, Sn, and S were first mixed thoroughly as raw materials in a
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molar ratio of 2: x: 1-x: 3 for Cu2MnxSn1-xS3 (x = 0, 0.05, 0.10, 0.15, 0.20) samples. The mixtures were sealed under vacuum in silica tubes that were then heated from room temperature to 1193 K at a rate of 5 K min-1 with a holding time of 8h, then rapidly cooled to 973 K and kept there for
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48h. Subsequently, the obtained ingots were naturally cooled to room temperature and ground into powders. The powders were finally sintered by Spark Plasma Sintering (SPS) in graphite dies at 773 K for 5 min under an axial pressure of 50 MPa. Phase composition and crystal structure were checked by X-ray diffraction (XRD) analysis with an ARL X’TRA diffractometer (SmartLab3, RIGAKU, Japan) using Cu Kα radiation. S and σ were measured in the radial direction of a bar-shaped specimen with dimensions of 10 mm × 3 mm × 3 mm by a conventional steady state method and a four-probe method, respectively, in a He atmosphere at 323~723 K with a
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commercial system (LSR-3, Linseis). The morphologies of the bulk samples were observed by field-emission scanning electron microscopy (FE-SEM, FEI Nova NanoSEM450). Thermal conductivity was calculated by κ = DdCp, where D was thermal diffusivity measured in the axial direction of a disk-shaped sample of Φ10 mm × 1 mm using a Netzsch laser flash diffusivity
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instrument (LFA457, Netzsch, Germany), Cp is the specific heat capacity determined by differential scanning calorimetry (DSC, 200 F3, TA instruments), and d is the mass density measured using the Archimedes method. Hall effect measurements for n and Hall mobility were
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conducted with a van der Pauw configuration under vacuum using a ResiTest8300 system (Toyo Tech. Co., Japan). Sound velocity was detected by an ultrasonic pulse-echo method (Model 5800 PR, Olympus) at room temperature.
Result and discussion
Powder X-ray diffraction (XRD) patterns of Cu2Sn1-xMnxS3 (x = 0-0.20) samples are shown in Fig. 1(a). For the x = 0, 0.05 samples, the peaks can be well indexed to monoclinic structure of CTS (PDF#27-0198) as shown in Fig. 1(b). Upon Mn-doping with x = 0.10, 0.15, the patterns turn to match well with cubic structure of CTS (PDF#89-2877) shown in Fig. 1(c), while with x increasing to 0.20, the XRD pattern was found to be dominated by the peaks of the tetragonal structure of CTS (PDF#89-4714) illustrated in Fig. 1(d). And as shown in the photograph in the
ACCEPTED MANUSCRIPT left insert, for doped samples, powder in dull blue on the surface of the as-synthesized product can be seen and was confirmed to be Cu-rich sulfide particles, while the undoped one was completely with a silver gray luster [10]. Actually, some small stray peaks for CuS (PDF#78-2121) can be seen in doped samples from the insert in Fig. 1(a). In order to check the phase composition, Cu2Sn1−xMnxS3 ceramic samples with x = 0, and 0.15 were selected for FE-SEM observations. As can be seen in Fig. 2(a) and (b), the samples, either undoped or doped, are all dense, and the grain
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size is coarse and uneven, ranging from several to dozens of microns. In particular, as shown in Fig. 2(c), numerous small-sized precipitates are formed at the surface of grains (Region a). Energy dispersive spectrometer (EDS) pattern in Fig. 3(d) confirmed the precipitates to be a Cu-rich secondary phase as compare to the matrix (Region b), which agrees with the presence of CuS diffraction peaks in the XRD result for doped samples. Here, a defect reaction is proposed for
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Mn-doping in CTS similar to Co-doping process [10] as below,
xMn Cu 2SnS3 → ( Cu1− x□x )2 ( Sn1− x Mn x ) ( S3− 2 x y □2 x y ) + ( 2 x y ) Cu yS + 2 x hole⋅ , 424 3 14243 14243 14 4244 3 1 4244 3 14 Sn-site
precipitate
S-site
carrier
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Cu-site
(1)
which shows that the generation of holes as carriers should be influenced little despite the formation of Cu-rich sulfide precipitate. Instead, enhanced phonon scattering at the interfaces should exert additional influences to help suppress the thermal transport.
It is worth noting here the crystallographic difference of cations arrangement between the 3 structures. For the monoclinic phase, Cu (3 types) and Sn atoms are ordered at different specific
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sites, respectively. While contrastingly, in the cubic and tetragonal phases, the metal atoms are disordered, e.g. at the 4a sites in the former and the 2b and 4d sites in the latter. This structural evolution of order-disorder transition have been found very common in the CTS system, especially when acceptor-doped [11,15-17]. The reasons can be related first to the local adjustment
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of cation coordination in the SM4 tetrahedra so as to minimize the lattice energy as stated in the octet rule (the sum of the valence electrons of the cations surrounding each anion should be equal to eight), which breaks the ordering of cations in monoclinic structure by randomizing their
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positioning and forms essentially the tetragonal and cubic structures. Second, the energy required for cations to exchange positions should be very small, as for Cu and Zn in (001) layers of Cu2ZnSnS4 (CZTS) which adapts a similar structure to CTS, so that such exchange should occur facilely under normal synthesis conditions to allow the coexistence of these phases [18]. Table 1. XRD Rietveld refinement results of phase fraction and lattice parameters for Cu2Sn1-xMnxS3 Lattice parameters (Å) x
Phase fraction
0
0.05
a
b
c
Cc (90.5%) F-43m (9.5%)
6.657 5.437
11.539 5.437
6.669 5.437
Cc (27.9%)
6.650
11.535
6.661
F-43m (21.6%)
5.432
5.432
5.432
Rwp (%)
3.62
5.52
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0.15
0.20
5.409
5.409
10.932
Cc (13.2%)
6.664
11.547
6.663
F-43m (30.7%)
5.444
5.444
5.444
I-42m (56.1%)
5.430
5.430
10.903
Cc (12.5%)
6.661
11.551
6.663
F-43m (27.5%)
5.442
5.442
5.442
I-42m (60.0%)
5.439
5.439
10.839
Cc (5.8%)
6.655
11.545
6.663
F-43m (31.7%)
5.438
5.438
5.438
I-42m (62.5%)
5.408
5.408
10.898
5.93
5.65
6.63
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0.10
I-42m (50.4%)
Very importantly, such structural transition should be very favorable for achieving a high TE performance for CTS. On one hand, the symmetry of SM4 tetrahedra in the cubic and
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tetragonal phases is high and hence can increase the degeneracy of the VBs that are not strictly degenerate (at the Γ point) in the monoclinic CTS. On the other hand, this kind of cation-disordered structure can significantly disrupt the phonon propagation, leading to an
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intensified phonon scattering at high temperatures, as has been found in many disordered
systems [19-23]. These two aspects will be discussed in the context on the transport properties of Mn-doped CTS.
As the standard diffraction profiles of cubic and tetragonal phases are highly overlapped
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in-between and by the monoclinic one as well, Rietveld refinement was performed for more information as shown in Fig. 3(a). The Rwp indexes were effectively lowered down to 5~7%
by taking the 3 phases into refinement, suggesting their coexistence in these samples. As has been reported in our previous work [16], the monoclinic phase took 96.1% of mole fraction and
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dominated in the undoped CTS, while upon Mn-doping, the fraction of monoclinic phase
reduces continuously, from 27.9% (x = 0.05) to 5.8% (x = 0.20), while the sum of the
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tetragonal and the cubic phases increases rapidly (Table 1), accounting for over 80~90% with x over 0.1. The tetragonal phase is always dominating (~60%) at x >0.10, which agrees well with the result that, the amount of tetragonal phase became the largest as doping improved, e.g. in the cases of Zn, Ni, and Co-doping. However, with the further increase of Mn-doping, the increment of the two phases becomes slow, which means the phase transition can be launched at a very low content of Mn and reaches an equilibrium with x more than 0.10.
The lattice parameters (a, c, and M-S bond lengths) deriving from the Rietveld refinement are shown in Fig. 3(b) and (c) as a function of x. The lattice constants a of the cubic and tetragonal phase increase first and then decrease as c varies oppositely with the
ACCEPTED MANUSCRIPT doping amount, which is in contrast to the radii order of rMn2+ (66 pm at coordination number, CN = 4) > rSn4+ (55 pm at CN = 4) [21,24]. This would suggest a complicated mechanism similar to that in the Co-doping case [10] which would involve the formation of CuS and even Cu and S vacancies as discussed above. For electrical transport properties, the temperature dependence of the electrical conductivity (σ) and Seebeck coefficient (S) from 323 K to 723 K are shown in Fig. 4(a) and (b). It has been
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reported[14] for pristine CTS that there present Cu vacancies thermodynamically to create holes as the major carrier. However, the concentration is very low, and σ is only between 1~10 Scm-1, with a semiconductor and small polaron behavior. Upon Mn-doping, the σ increases almost linearly with the doping amount of Mn, which can be clearly seen in the insert of Fig. 4(a) , from ~2 Scm-1 for x = 0 and ~100 Scm-1 for x = 0.05 to a high maximum of 1235 Scm-1 for x = 0.20 at
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323 K, reflecting the acceptor role of Mn doping. Besides, except the pristine one, they decrease gradually with temperature as a result of intensified phonon scattering. To examine the influence of chemical composition on the electrical transport properties, energy dispersive X-ray
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spectroscopy (EDX) measurement was carried out on the polished sample pellets, and the data is provided in the Supporting Information. Generally, the content of Cu and the Cu:(Sn+Mn) atomic ratio (about 1.86 ± 0.08) are essentially independent of the Mn amount qualitatively, which signifies that the effective improvement of electrical transport properties with the increased Mn amount should be directly related to the Mn substitution for Sn as acceptor. S values for all samples are positive in the whole measured temperature range, and they
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decrease with the doping level, from ~290 µVK-1 for pristine one to ~60 µVK-1 for x = 0.20 at 323 K, as shown in the insert of Fig. 4(b). It should be noted that, when compared with Zn-doped CTS, the σ values are similar and, interestingly, the S values are higher at the same time, which means better electrical transport properties than Zn-doping [11]. Nevertheless, when compared with
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Co-doped CTS, the S values in Mn-doped CTS are basically slightly smaller, but the electrical properties should be comparable as the σ values are relatively larger than that of Co-doped
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ones[10].
In order to better understand the transport mechanism of Mn doped CTS, Hall effect
measurements were conducted, and based on the single parabolic band (SPB) model [25] with acoustic phonon scattering as the main carrier scattering mechanism, DOS effective mass (m*) of carriers were calculated and shown together with carrier concentration n in Fig. 4(c) as a function of doping level x. As can be seen, the n increased continuously with the doping level x, from about 1 × 1018 cm-3 for the pristine CTS to ~6×1021 cm-3 for x = 0.20. The obtained m* for Cu2Sn1-xMnxS3 with x = (0.05-0.20) are 3.0 ~ 9.0 mo at 323 K and increase with x almost linearly. These high m* should be favored by the electronic nature in the tetragonal and cubic CTS, and can be related to the deeper Fermi level in the VB as the doping level rises, which corresponds to a larger m* [10]. Furthermore, both n and m* are strikingly larger than those in Zn-doped CTS, and a
ACCEPTED MANUSCRIPT little higher than those in Co-doped counterparts with the same doping level. As shown in Fig.4(d) together, this can be explained by the fact that, compared with Zn (3d10) and Co (3d7), Mn (3d5) has more unfilled 3d orbitals contributable to the VB of Cu2Sn1-xMnxS3, which could enhance the n and m* at the same time. Similar results [26, 27] have been obtained by Co-doping in CTS and also in Cu2CoSnS4. On the other hand, however, the mobility µ (1~2 cm2V-1s-1 at 323 K) is much lower than that in Zn-doped CTS (2~4 cm2V-1s-1 at 323 K), while comparable with that for
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Co-doping. This can be reasonably attributed to the increased interaction between carriers when n increases, and to the enhanced m*, and as well to the defects introduced by Mn-doping.
The relative large increase of n and m* with Mn-doping in CTS can be easily seen when
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compared with those for Co- and Zn-doping in the Pisarenko plot as shown in Fig. 5(a). Because the n in these heavily doped CTS are generally very high (all above 1021 cm-3), the calculation
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based on SPB should stand reasonable [12]. The continuous growth of m* with n contributed to the large PF by heavy doping. As shown in Fig. 5 (b), the values of PF increase obviously with increasing doping content and all the doped samples exhibit higher PF than the undoped sample in the whole temperature range. A maximal value of 0.90-0.92 mWm-1 K-2 is achieved at 723 K in the x = 0.15 and 0.20 sample. The enhancement in PF for the doped samples can be attributed to
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the substantial improvement in σ and a slight increase of S at high temperatures, which is mainly caused by the greatly improved hole concentration and the enhancement of m*. By comparison, it is encouraging to see that the maximal PF of Mn-doped CTS in the whole temperature range is electrical transport property
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higher than that of Co-doped CTS, which means an improved
relative to that of Co-doping. However, the highest PF obtained at 723 K for Cu2Sn1-xMnxS3 (x = 0.15 or 0.20), is only comparable with that for Cu2Sn0.8Co0.2S3 (0.94 mWm-1K-2 at 723 K), despite
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about 10% larger than that for Zn-doping, which is probably due to the excessive increase of n. This essentially reflects the trade-off effects of n and m* on S, which can be given by S = [8π2kB2m*T(π/3n)2/3]/(3eh2) [28], where kB and h are the Boltzmann constant and the Plank constant.
The temperature dependence of the total thermal conductivity κ for Cu2Sn1-xMnxS3 (x = 0-0.20) samples is shown in Fig. 6(a). The κ data for all the samples decrease gradually with increasing temperature, indicative of phonon-phonon interaction dominating with temperature increasing. For pristine CTS, the κ ranging from 2.4 Wm-1K-1 at room temperature to 0.8 Wm-1K-1 at 723 K, is
ACCEPTED MANUSCRIPT well comparable with those reported in literatures [21,22]. For the doped samples, the κ values increased with x in the whole temperature range, indicative of considerable contribution from the electrical conduction.
The electronic thermal conductivity (κe) is estimated from the Wiedeman–Franz law as κe =
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LσT, for the insert of Fig.6(b), the Lorenz number L was obtained from reduced chemical potential
η which was calculated from the experimental S and n data as previously [29,30]. With the increase of Mn content , L increases from 1.6 for the undoped sample to 2.2 for the x = 0.2 sample, in accordance with the enhanced degenerate level of CTS due to Mn-doping. As shown in the Fig.
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6(b), κe are quite constant in the whole temperature range, while increase with x remarkably due to the increase of σ. For the undoped sample, κe is negligible to the total κ, indicating a primary role
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of phonon conduction in it. While for the doped CTS, the values increased from 0.06 Wm-1K-1 (x = 0.05) to 0.92 Wm-1K-1 (x = 0.20) at 323 K due to the high σ and L values in heavier doped samples. By subtracting κe from κ, the lattice thermal conductivity κl was calculated and is shown in the insert of Fig. 6(a) as a function of temperature. As can be seen, the κl for all the doped samples (x = 0.05-0.20) drop gradually from ~1.5 Wm-1K-1 at 323 K to 0.3 ~ 0.5 Wm-1K-1 at 723 K,
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a level close to the theoretical minimum of 0.3 Wm-1K-1 [12]. In fact, most sulfide TE materials have large κl (TiS2 [31, 32] ~ 1.3 Wm-1K-1, PbS [33] ~ 1 Wm-1K-1) due to the light mass of sulfur and strong chemical bonds between sulfur and the metal ions. However, for diamond-like sulfide
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TEs, low κl have been frequently reported [23, 34-36], which is linked primarily to the disordered occupation of cations in cubic and tetragonal structures, the formation of Cu-S precipitates, and
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even M and S vacancies. As has been reported, the liquid-like copper ions in the rigid S- sublattice in the binary Cu2-xS compounds play a significant role to cause an extraordinarily low phonon speed [37], which is to say, the presence of CuyS phase is contributable to suppress the lattice thermal conductivity. Nevertheless, it is still a challenging issue to quantitatively analyze its contribution at the present. In addition, we found that the reduction of κl corresponds well with the decrease of monoclinic phase fraction from 90.5% ( x=0 ) to 13.2% ( x=0.1 ), which suggests the strong effect of cation disordering on phonon scattering, while with x goes higher, since the disordered phases dominating in the samples keep almost constant around at 80-90%, that is to say, the phonon scattering due to cation disordering would be maintained at a constant level, thus the
ACCEPTED MANUSCRIPT lattice thermal conductivity remains almost unchanged with the Mn-doping.
Besides, the
outermost electron dissatisfaction would make the chemical bond between Mn and S less compact and lead to softened chemical bonds, which is beneficial to the suppression of κl [10, 39].
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It is interesting to see in Fig. 6(c) that, for the doped Cu2Sn1-xMnxS3 samples, the κl values are very close despite their different doping levels, furthermore, they are almost at the same level with the data of Zn-doped CTS in the whole temperature range. However, the κl of Co-doped samples are different from them at 323 K. Generally, κl can be given as κl = 1/3vavgCvl, where vavg, Cv, and l are the average speed of sound, heat capacity, and phonon mean free path, respectively. In order to
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explicate the reduction of κl, the average sound velocities of longitudinal and transverse branches, as an important parameter determining the κl, are measured and listed in Table 2. The vl and vt of Cu2Sn1-xMnxS3 decrease gradually with the doping amount at 300 K, which can be traced back to
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the weak chemical bonds due to the crystal structure transformation. Comparing to Co-doping samples, the vt of Cu2Sn1-xMnxS3 are a little larger which may result from the larger relative density because of different SPS temperature (773K and 823K for Co- and Mn-doped CTS, respectively). The vavg of both Co- and Mn-doped samples do not differ from each other and both are lower than the pristine CTS for a certain extent. Clearly, the small vavg is one reason for the abnormally low κl of CTS doped samples. The Grüneisen parameter γ can be calculated from the
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measured vl and vt by γ = 3 (1+ν P ) 2 ( 2-3ν P ) , where Poisson ratio vp can be derived from the vl and vt by v p = 1 − 2 ( vt vl )2 2 − 2 ( vt vl )2 . The calculated γ of the Mn-doped samples ranges from
1.71 to 1.88, nearly the same as that of Co-doping samples (1.71~1.92). This similarity indicates that the distinction of anharmonic lattice vibrations is not obvious for the doped sample, which is
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to say, the relatively larger κl in the low temperature range in Mn-doped CTS can be mainly attributed to the higher density than that of Co-doped CTS.
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Table 2. Parameters of density ρ (g·cm-3), relative density K (%), longitudinal νl (m·s-1), transverse νt (m·s-1), average speed of sound νavg (m·s-1), Debye temperature θD (K) and Grüneisen parameter γ of bulk materials Cu2Sn1-xMxS3 (x = 0-0.20; M = Mn, Co) at 300 K. Dopant
Mn
Co
x
ρ
K
νl
νt
vavg
θD
γ
0 0.05 0.10 0.15 0.20 0.05 0.10 0.15 0.20
4.58 4.57 4.56 4.54 4.56 4.52 4.46 4.43 4.53
96.42 96.61 96.40 96.18 96.81 95.16 93.89 93.46 95.56
4648 3975 3879 3634 3860 3687 3940 3540 3872
2373 2057 2043 1980 2075 1968 2015 1924 1988
2658 2303 2284 2209 2317 2198 2257 2146 2227
287.14 249.39 248.00 240.18 253.13 264.70 264.01 264.90 264.35
1.93 1.88 1.82 1.71 1.75 1.78 1.92 1.71 1.91
Fig. 6(d) shows the figure of merit ZT calculated from the data above. The ZT values increase
ACCEPTED MANUSCRIPT sharply with rising temperature in doped samples with x = 0.10~0.20. And the highest ZT of ~0.68 is achieved in the x = 0.15 sample due to the largest PF and the relatively low κ. However, this value is a little smaller than the optimal ZT in Co doped samples, it can be considered to suffer from primarily its larger thermal conductivity which is caused by both the increased κe and κl.
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Conclusions In this work, Mn is demonstrated as an effective d-unfilled dopant for tuning carrier concentration and effective mass of carriers for CTS as a TE material. It is shown that d-unfilled TM (such as Co, Mn) doping induced an order-to-disorder crystal transformation from monoclinic to tetragonal and cubic structures. Such a transformation is crucial for both electrical and thermal
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benefits for TE performance. On one hand, it causes a high carrier concentration and enables both the improved degeneracy of up VBs and the participation of d-orbitals to achieve an enhanced DOS effective mass, thus leading to a high PF. On the other hand, a low κl even close to
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theoretical minimum can be obtained in the tetragonal and cubic dominated bulks, due primarily to the high degree disorder cation occupation and bond softening, and other accompanying defects (precipitation and vacancies) as well. These findings point out the significance of phase transformation of CTS by way of suitable TM doping and would be helpful to get a systematical understanding of the fundamental parameters determining the transport properties of CTS.
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Acknowledgements
Reference
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This research was supported by the National Natural Science Foundation of China under Grant No. 51272103, 51672127, No. 11604047, the Natural Science Foundation of Jiangsu Province (No. BK20160694), the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD).
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non-toxic earth-abundant copper sulfide, Adv. Mater. 26 (2014) 3974-3978. [38] S. Wang, Y. Sun, J. Yang, B. Duan, L. Wu, W. Zhang, J. Yang, High thermoelectric performance in Te-free (Bi,Sb)2Se3via structural transition induced band convergence and chemical bond softening, Energy Environ. Sci. 9 (2016) 3436-3447.
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Fig.1 (a) Powder X-ray diffraction patterns of Cu2Sn1-xMnxS3 (x = 0-0.20), together with photographs of the as-synthesized powder of undoped and doped CTS in insert 1, and a local zoom demonstrating the impurity peaks of CuS in insert 2. Crystal structures of CTS in monoclinic (b), cubic (c) and tetragonal (d), with ordered cation occupation in (b), disordered (at the 4a site) in (c), and (at the 4d and 2b sites) in (d).
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Fig. 2 SEM images of fractured surface for samples with (a) x = 0 and (b) x = 0.15, and a magnified local view (c) for the x = 0.15 sample showing precipitates on the grain boundaries (Region a) and the interior (Region b). (d) shows the corresponding EDS patterns for Region a and b in (c).
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Fig.3 Rietveld refinements results (a) for Powder XRD patterns of Cu2Sn1-xMnxS3 (x = 0.05-0.20), and derived lattice parameters of a and c (b), and the bond lengths of M-S and Cu-S (c) as functions of the doping level.
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Fig. 4 Thermoelectric parameters of (a) electrical conductivity σ, (b) Seebeck coefficient S in the measured temperature range, together with inserts showing their dependence on the Mn-doping level, respectively. Variation of carrier concentration n and Hall mobility µHall (c) and DOS effective mass of carriers m*(d) with doping level of x in comparison with cited date for Co-[10] and Zn-doped[11] CTS at 300 K.
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Fig.5 (a) Pisarenko plot of S v.s. carrier concentration n in comparison with cited data of Co-[10] and Zndoped[11] CTS, where the lines represent the calculated relation based on SPB model, and (b) PF = S2σ for the Cu2Sn1-xMnxS3 (x = 0-0.20) bulks at different temperatures, with cited data for 20%Co- and Zndoped CTS for comparison.
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Fig. 6 Thermal transport properties of (a) total thermal conductivity κ and lattice thermal conductivity κl (insert), (b) electronic thermal conductivity κe and Lorenz number L (insert) and (c) its variation with doping level at 323 and 723 K, with cited data of Co-[10] and Zn-doped[11] CTS for comparison. (d) shows the calculated dimensionless figure of merit ZT for Cu2Sn1-xMnxS3 (x = 0-0.20) samples, where the maximal ZT~0.68 was obtained in the 15%Mn-doped CTS at 723 K.
ACCEPTED MANUSCRIPT The main highlights of the paper includes (1) Mn-doping in Cu2SnS3 caused transitions from cation-ordered to disordered phases. (2) Low lattice thermal conductivity was obtained in cation-disordered structures. (3) High carrier effective mass and large PF caused by contribution of Mn 3d states.
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(4) Maximal ZT ~0.68 was achieved at 723 K by optimal Mn-doping in Cu2SnS3.
S0925-8388(18)34461-X
DOI:
https://doi.org/10.1016/j.jallcom.2018.11.329
Reference:
JALCOM 48555
To appear in:
Journal of Alloys and Compounds
Received Date: 15 July 2018 Revised Date:
20 November 2018
Accepted Date: 24 November 2018
Please cite this article as: Z. Zhang, H. Zhao, Y. Wang, X. Hu, Y. Lyu, C. Cheng, L. Pan, C. Lu, Role of crystal transformation on the enhanced thermoelectric performance in Mn-doped Cu2SnS3, Journal of Alloys and Compounds (2018), doi: https://doi.org/10.1016/j.jallcom.2018.11.329. This is a PDF file of an unedited manuscript that has been accepted for publication. As a service to our customers we are providing this early version of the manuscript. The manuscript will undergo copyediting, typesetting, and review of the resulting proof before it is published in its final form. Please note that during the production process errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain.
ACCEPTED MANUSCRIPT Role of crystal transformation on the enhanced thermoelectric performance in Mn-doped Cu2SnS3 Zhen Zhang,a Huiwen Zhao,a Yifeng Wang,a,b* Xiaohui Hu,a,b Yinong Lyu,a* Changchun Cheng,a Lin Pana,b, Chunhua Lua,b
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a. College of Materials Science and Engineering, Nanjing Tech University, Nanjing 210009, China. b. Jiangsu Collaborative Innovation Center for Advanced Inorganic Function Composites, Nanjing Tech University, Nanjing 210009, China.
Abstract
Mohite-type Cu2SnS3 exhibits a large degree of manipulation in both electronic and phonon
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transport properties by transition metal doping. In this work, the effects of Mn doping on the transport properties of Cu2SnS3 were investigated. A crystal transition occurred upon doping from cation-ordered monoclinic structure to cation-disordered cubic and tetragonal structures. Such a
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crystal transition caused a significant increase of DOS effective mass of carriers, m*, which should be caused by the improved degeneracy of the bands at the Fermi level due to the improved symmetry, and also by the participation (in the upper valence bands) of unfilled d-orbitals of Mn cations that share crystallographically equivalent sites with Cu cations. Meanwhile the lattice thermal conductivity (~0.4 W m-1K-1 at ~773K) was suppressed effectively, which should also be linked to the disorder structure induced by Mn-doping. Benefitting from the increased carrier
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density and much improved m*, a remarkable power factor of 0.92 Wm-1K-2 has been obtained at 723 K. The highest ZT ~0.68 at 723 K is almost 15 times increased as compared to that of pristine Cu2SnS3.
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Keywords: Cu2SnS3, d-unfilled dopant, crystal transition, thermoelectric performance
ACCEPTED MANUSCRIPT Introduction Thermoelectric (TE) materials have garnered extensive interests and shown promising prospects in applications of energy conversion and semiconductor cooling without using moving parts nor greenhouse gases emission [1,2]. The performance of TE materials is quantified by the dimensionless figure of merit ZT, defined as ZT = S2σT/κ = S2σT/(κe+κl), where S is the Seebeck
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coefficient, σ is the electrical conductivity, T is the absolute temperature, and κ is the thermal conductivity which comprises of the electronic thermal conductivity (κe) and lattice thermal conductivity (κl) [3,4]. High TE efficiency requires high ZT values of materials to become comparable with other energy conversion routes and to expand the application fields. For this purpose, a large power factor (PF = S2σ) and a low κ should be secured. However, S, σ and κ are
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strongly coupled through the carrier concentration (n) [5-7]. It is still one of the biggest challenges to independently optimize these three parameters to obtain high-performance TE materials with large ZT values.
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Recently, as an attractive earth abundant material, p-type ternary sulfide Cu2SnS3 (CTS) has been found excellent not only for absorber in solar cells [8, 9] but also as a potential eco-friendly TE material with phonon-glass-electron-crystal features, with peak ZTs around 0.6~0.85 at 723 K to date [10,11]. CTS has polymorphic crystal structures including monoclinic, tetragonal and cubic, all of which can be viewed as a diamond-like network of inter-connecting S-centered tetrahedra of SM4 (M = Cu and Sn) with varied arrangement and symmetry. The complex crystallographic
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network and soft bonding between the metal and S atoms with diverse and even random coordination endue CTS a very low κl [12-14]. Furthermore, in the monoclinic phase of CTS, there is a 3-dimensional conductive lattice network, with the upper valence bands (VB) mainly composed of hybrid Cu 3d orbitals and S 3p orbitals (while containing no obvious contribution from Sn [12]), which ensures the improvement of σ in CTS by way of doping to control the carrier
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density. However, S is inevitably reduced as σ increases, which still impedes the significant improvement of ZT value, since Sn-site acceptor doping seems to minimally affect the upper VB
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shape and only denotes holes in the monoclinic structure [12]. Nevertheless, when it transforms into a high symmetry tetragonal and cubic structure,
appropriate dopant atoms at Sn-site, which are crystallographically equivalent to Cu atoms, may have contributions to enhance the density-of-states (DOS) of the upper VB, while such a phase transformation can be readily realized by transition metal (TM, e.g. Zn, Ni, and Co) doping. Moreover, as revealed in our previous work, doping with 3d-unfilled TM, e.g. Co [10], seemed much more favorable to improve the TE performance than Zn-doping as a result of the enhanced DOS effective mass of carriers m* attributed to the participation of their additional 3d states in the VB. Effects that would be induced in terms of the more unfilled orbitals contained in Mn cations in the present study. Results revealed an order-to-disorder crystal transformation of CTS from monoclinic to cubic and tetragonal phases upon doping, a greatly enhanced DOS effective mass
ACCEPTED MANUSCRIPT m* together with an increase in n. While the κl was suppressed largely, due to the enhanced phonon scattering attributed to disordered cation arrangement and other defects related to the phase transformation or doping. Finally, a maximal ZT ~0.68 was obtained at 723 K, which is about 15 times higher than that of the pristine CTS. It strongly suggests the significant role of crystal transformation due to d-unfilled TM doping for CTS and the promising prospective of CTS as a
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good mid-temperature TE material.
Experimental
High purity powders of Cu, Mn, Sn, and S were first mixed thoroughly as raw materials in a
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molar ratio of 2: x: 1-x: 3 for Cu2MnxSn1-xS3 (x = 0, 0.05, 0.10, 0.15, 0.20) samples. The mixtures were sealed under vacuum in silica tubes that were then heated from room temperature to 1193 K at a rate of 5 K min-1 with a holding time of 8h, then rapidly cooled to 973 K and kept there for
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48h. Subsequently, the obtained ingots were naturally cooled to room temperature and ground into powders. The powders were finally sintered by Spark Plasma Sintering (SPS) in graphite dies at 773 K for 5 min under an axial pressure of 50 MPa. Phase composition and crystal structure were checked by X-ray diffraction (XRD) analysis with an ARL X’TRA diffractometer (SmartLab3, RIGAKU, Japan) using Cu Kα radiation. S and σ were measured in the radial direction of a bar-shaped specimen with dimensions of 10 mm × 3 mm × 3 mm by a conventional steady state method and a four-probe method, respectively, in a He atmosphere at 323~723 K with a
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commercial system (LSR-3, Linseis). The morphologies of the bulk samples were observed by field-emission scanning electron microscopy (FE-SEM, FEI Nova NanoSEM450). Thermal conductivity was calculated by κ = DdCp, where D was thermal diffusivity measured in the axial direction of a disk-shaped sample of Φ10 mm × 1 mm using a Netzsch laser flash diffusivity
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instrument (LFA457, Netzsch, Germany), Cp is the specific heat capacity determined by differential scanning calorimetry (DSC, 200 F3, TA instruments), and d is the mass density measured using the Archimedes method. Hall effect measurements for n and Hall mobility were
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conducted with a van der Pauw configuration under vacuum using a ResiTest8300 system (Toyo Tech. Co., Japan). Sound velocity was detected by an ultrasonic pulse-echo method (Model 5800 PR, Olympus) at room temperature.
Result and discussion
Powder X-ray diffraction (XRD) patterns of Cu2Sn1-xMnxS3 (x = 0-0.20) samples are shown in Fig. 1(a). For the x = 0, 0.05 samples, the peaks can be well indexed to monoclinic structure of CTS (PDF#27-0198) as shown in Fig. 1(b). Upon Mn-doping with x = 0.10, 0.15, the patterns turn to match well with cubic structure of CTS (PDF#89-2877) shown in Fig. 1(c), while with x increasing to 0.20, the XRD pattern was found to be dominated by the peaks of the tetragonal structure of CTS (PDF#89-4714) illustrated in Fig. 1(d). And as shown in the photograph in the
ACCEPTED MANUSCRIPT left insert, for doped samples, powder in dull blue on the surface of the as-synthesized product can be seen and was confirmed to be Cu-rich sulfide particles, while the undoped one was completely with a silver gray luster [10]. Actually, some small stray peaks for CuS (PDF#78-2121) can be seen in doped samples from the insert in Fig. 1(a). In order to check the phase composition, Cu2Sn1−xMnxS3 ceramic samples with x = 0, and 0.15 were selected for FE-SEM observations. As can be seen in Fig. 2(a) and (b), the samples, either undoped or doped, are all dense, and the grain
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size is coarse and uneven, ranging from several to dozens of microns. In particular, as shown in Fig. 2(c), numerous small-sized precipitates are formed at the surface of grains (Region a). Energy dispersive spectrometer (EDS) pattern in Fig. 3(d) confirmed the precipitates to be a Cu-rich secondary phase as compare to the matrix (Region b), which agrees with the presence of CuS diffraction peaks in the XRD result for doped samples. Here, a defect reaction is proposed for
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Mn-doping in CTS similar to Co-doping process [10] as below,
xMn Cu 2SnS3 → ( Cu1− x□x )2 ( Sn1− x Mn x ) ( S3− 2 x y □2 x y ) + ( 2 x y ) Cu yS + 2 x hole⋅ , 424 3 14243 14243 14 4244 3 1 4244 3 14 Sn-site
precipitate
S-site
carrier
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Cu-site
(1)
which shows that the generation of holes as carriers should be influenced little despite the formation of Cu-rich sulfide precipitate. Instead, enhanced phonon scattering at the interfaces should exert additional influences to help suppress the thermal transport.
It is worth noting here the crystallographic difference of cations arrangement between the 3 structures. For the monoclinic phase, Cu (3 types) and Sn atoms are ordered at different specific
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sites, respectively. While contrastingly, in the cubic and tetragonal phases, the metal atoms are disordered, e.g. at the 4a sites in the former and the 2b and 4d sites in the latter. This structural evolution of order-disorder transition have been found very common in the CTS system, especially when acceptor-doped [11,15-17]. The reasons can be related first to the local adjustment
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of cation coordination in the SM4 tetrahedra so as to minimize the lattice energy as stated in the octet rule (the sum of the valence electrons of the cations surrounding each anion should be equal to eight), which breaks the ordering of cations in monoclinic structure by randomizing their
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positioning and forms essentially the tetragonal and cubic structures. Second, the energy required for cations to exchange positions should be very small, as for Cu and Zn in (001) layers of Cu2ZnSnS4 (CZTS) which adapts a similar structure to CTS, so that such exchange should occur facilely under normal synthesis conditions to allow the coexistence of these phases [18]. Table 1. XRD Rietveld refinement results of phase fraction and lattice parameters for Cu2Sn1-xMnxS3 Lattice parameters (Å) x
Phase fraction
0
0.05
a
b
c
Cc (90.5%) F-43m (9.5%)
6.657 5.437
11.539 5.437
6.669 5.437
Cc (27.9%)
6.650
11.535
6.661
F-43m (21.6%)
5.432
5.432
5.432
Rwp (%)
3.62
5.52
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0.15
0.20
5.409
5.409
10.932
Cc (13.2%)
6.664
11.547
6.663
F-43m (30.7%)
5.444
5.444
5.444
I-42m (56.1%)
5.430
5.430
10.903
Cc (12.5%)
6.661
11.551
6.663
F-43m (27.5%)
5.442
5.442
5.442
I-42m (60.0%)
5.439
5.439
10.839
Cc (5.8%)
6.655
11.545
6.663
F-43m (31.7%)
5.438
5.438
5.438
I-42m (62.5%)
5.408
5.408
10.898
5.93
5.65
6.63
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0.10
I-42m (50.4%)
Very importantly, such structural transition should be very favorable for achieving a high TE performance for CTS. On one hand, the symmetry of SM4 tetrahedra in the cubic and
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tetragonal phases is high and hence can increase the degeneracy of the VBs that are not strictly degenerate (at the Γ point) in the monoclinic CTS. On the other hand, this kind of cation-disordered structure can significantly disrupt the phonon propagation, leading to an
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intensified phonon scattering at high temperatures, as has been found in many disordered
systems [19-23]. These two aspects will be discussed in the context on the transport properties of Mn-doped CTS.
As the standard diffraction profiles of cubic and tetragonal phases are highly overlapped
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in-between and by the monoclinic one as well, Rietveld refinement was performed for more information as shown in Fig. 3(a). The Rwp indexes were effectively lowered down to 5~7%
by taking the 3 phases into refinement, suggesting their coexistence in these samples. As has been reported in our previous work [16], the monoclinic phase took 96.1% of mole fraction and
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dominated in the undoped CTS, while upon Mn-doping, the fraction of monoclinic phase
reduces continuously, from 27.9% (x = 0.05) to 5.8% (x = 0.20), while the sum of the
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tetragonal and the cubic phases increases rapidly (Table 1), accounting for over 80~90% with x over 0.1. The tetragonal phase is always dominating (~60%) at x >0.10, which agrees well with the result that, the amount of tetragonal phase became the largest as doping improved, e.g. in the cases of Zn, Ni, and Co-doping. However, with the further increase of Mn-doping, the increment of the two phases becomes slow, which means the phase transition can be launched at a very low content of Mn and reaches an equilibrium with x more than 0.10.
The lattice parameters (a, c, and M-S bond lengths) deriving from the Rietveld refinement are shown in Fig. 3(b) and (c) as a function of x. The lattice constants a of the cubic and tetragonal phase increase first and then decrease as c varies oppositely with the
ACCEPTED MANUSCRIPT doping amount, which is in contrast to the radii order of rMn2+ (66 pm at coordination number, CN = 4) > rSn4+ (55 pm at CN = 4) [21,24]. This would suggest a complicated mechanism similar to that in the Co-doping case [10] which would involve the formation of CuS and even Cu and S vacancies as discussed above. For electrical transport properties, the temperature dependence of the electrical conductivity (σ) and Seebeck coefficient (S) from 323 K to 723 K are shown in Fig. 4(a) and (b). It has been
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reported[14] for pristine CTS that there present Cu vacancies thermodynamically to create holes as the major carrier. However, the concentration is very low, and σ is only between 1~10 Scm-1, with a semiconductor and small polaron behavior. Upon Mn-doping, the σ increases almost linearly with the doping amount of Mn, which can be clearly seen in the insert of Fig. 4(a) , from ~2 Scm-1 for x = 0 and ~100 Scm-1 for x = 0.05 to a high maximum of 1235 Scm-1 for x = 0.20 at
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323 K, reflecting the acceptor role of Mn doping. Besides, except the pristine one, they decrease gradually with temperature as a result of intensified phonon scattering. To examine the influence of chemical composition on the electrical transport properties, energy dispersive X-ray
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spectroscopy (EDX) measurement was carried out on the polished sample pellets, and the data is provided in the Supporting Information. Generally, the content of Cu and the Cu:(Sn+Mn) atomic ratio (about 1.86 ± 0.08) are essentially independent of the Mn amount qualitatively, which signifies that the effective improvement of electrical transport properties with the increased Mn amount should be directly related to the Mn substitution for Sn as acceptor. S values for all samples are positive in the whole measured temperature range, and they
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decrease with the doping level, from ~290 µVK-1 for pristine one to ~60 µVK-1 for x = 0.20 at 323 K, as shown in the insert of Fig. 4(b). It should be noted that, when compared with Zn-doped CTS, the σ values are similar and, interestingly, the S values are higher at the same time, which means better electrical transport properties than Zn-doping [11]. Nevertheless, when compared with
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Co-doped CTS, the S values in Mn-doped CTS are basically slightly smaller, but the electrical properties should be comparable as the σ values are relatively larger than that of Co-doped
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ones[10].
In order to better understand the transport mechanism of Mn doped CTS, Hall effect
measurements were conducted, and based on the single parabolic band (SPB) model [25] with acoustic phonon scattering as the main carrier scattering mechanism, DOS effective mass (m*) of carriers were calculated and shown together with carrier concentration n in Fig. 4(c) as a function of doping level x. As can be seen, the n increased continuously with the doping level x, from about 1 × 1018 cm-3 for the pristine CTS to ~6×1021 cm-3 for x = 0.20. The obtained m* for Cu2Sn1-xMnxS3 with x = (0.05-0.20) are 3.0 ~ 9.0 mo at 323 K and increase with x almost linearly. These high m* should be favored by the electronic nature in the tetragonal and cubic CTS, and can be related to the deeper Fermi level in the VB as the doping level rises, which corresponds to a larger m* [10]. Furthermore, both n and m* are strikingly larger than those in Zn-doped CTS, and a
ACCEPTED MANUSCRIPT little higher than those in Co-doped counterparts with the same doping level. As shown in Fig.4(d) together, this can be explained by the fact that, compared with Zn (3d10) and Co (3d7), Mn (3d5) has more unfilled 3d orbitals contributable to the VB of Cu2Sn1-xMnxS3, which could enhance the n and m* at the same time. Similar results [26, 27] have been obtained by Co-doping in CTS and also in Cu2CoSnS4. On the other hand, however, the mobility µ (1~2 cm2V-1s-1 at 323 K) is much lower than that in Zn-doped CTS (2~4 cm2V-1s-1 at 323 K), while comparable with that for
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Co-doping. This can be reasonably attributed to the increased interaction between carriers when n increases, and to the enhanced m*, and as well to the defects introduced by Mn-doping.
The relative large increase of n and m* with Mn-doping in CTS can be easily seen when
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compared with those for Co- and Zn-doping in the Pisarenko plot as shown in Fig. 5(a). Because the n in these heavily doped CTS are generally very high (all above 1021 cm-3), the calculation
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based on SPB should stand reasonable [12]. The continuous growth of m* with n contributed to the large PF by heavy doping. As shown in Fig. 5 (b), the values of PF increase obviously with increasing doping content and all the doped samples exhibit higher PF than the undoped sample in the whole temperature range. A maximal value of 0.90-0.92 mWm-1 K-2 is achieved at 723 K in the x = 0.15 and 0.20 sample. The enhancement in PF for the doped samples can be attributed to
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the substantial improvement in σ and a slight increase of S at high temperatures, which is mainly caused by the greatly improved hole concentration and the enhancement of m*. By comparison, it is encouraging to see that the maximal PF of Mn-doped CTS in the whole temperature range is electrical transport property
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higher than that of Co-doped CTS, which means an improved
relative to that of Co-doping. However, the highest PF obtained at 723 K for Cu2Sn1-xMnxS3 (x = 0.15 or 0.20), is only comparable with that for Cu2Sn0.8Co0.2S3 (0.94 mWm-1K-2 at 723 K), despite
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about 10% larger than that for Zn-doping, which is probably due to the excessive increase of n. This essentially reflects the trade-off effects of n and m* on S, which can be given by S = [8π2kB2m*T(π/3n)2/3]/(3eh2) [28], where kB and h are the Boltzmann constant and the Plank constant.
The temperature dependence of the total thermal conductivity κ for Cu2Sn1-xMnxS3 (x = 0-0.20) samples is shown in Fig. 6(a). The κ data for all the samples decrease gradually with increasing temperature, indicative of phonon-phonon interaction dominating with temperature increasing. For pristine CTS, the κ ranging from 2.4 Wm-1K-1 at room temperature to 0.8 Wm-1K-1 at 723 K, is
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The electronic thermal conductivity (κe) is estimated from the Wiedeman–Franz law as κe =
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LσT, for the insert of Fig.6(b), the Lorenz number L was obtained from reduced chemical potential
η which was calculated from the experimental S and n data as previously [29,30]. With the increase of Mn content , L increases from 1.6 for the undoped sample to 2.2 for the x = 0.2 sample, in accordance with the enhanced degenerate level of CTS due to Mn-doping. As shown in the Fig.
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6(b), κe are quite constant in the whole temperature range, while increase with x remarkably due to the increase of σ. For the undoped sample, κe is negligible to the total κ, indicating a primary role
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of phonon conduction in it. While for the doped CTS, the values increased from 0.06 Wm-1K-1 (x = 0.05) to 0.92 Wm-1K-1 (x = 0.20) at 323 K due to the high σ and L values in heavier doped samples. By subtracting κe from κ, the lattice thermal conductivity κl was calculated and is shown in the insert of Fig. 6(a) as a function of temperature. As can be seen, the κl for all the doped samples (x = 0.05-0.20) drop gradually from ~1.5 Wm-1K-1 at 323 K to 0.3 ~ 0.5 Wm-1K-1 at 723 K,
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a level close to the theoretical minimum of 0.3 Wm-1K-1 [12]. In fact, most sulfide TE materials have large κl (TiS2 [31, 32] ~ 1.3 Wm-1K-1, PbS [33] ~ 1 Wm-1K-1) due to the light mass of sulfur and strong chemical bonds between sulfur and the metal ions. However, for diamond-like sulfide
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TEs, low κl have been frequently reported [23, 34-36], which is linked primarily to the disordered occupation of cations in cubic and tetragonal structures, the formation of Cu-S precipitates, and
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even M and S vacancies. As has been reported, the liquid-like copper ions in the rigid S- sublattice in the binary Cu2-xS compounds play a significant role to cause an extraordinarily low phonon speed [37], which is to say, the presence of CuyS phase is contributable to suppress the lattice thermal conductivity. Nevertheless, it is still a challenging issue to quantitatively analyze its contribution at the present. In addition, we found that the reduction of κl corresponds well with the decrease of monoclinic phase fraction from 90.5% ( x=0 ) to 13.2% ( x=0.1 ), which suggests the strong effect of cation disordering on phonon scattering, while with x goes higher, since the disordered phases dominating in the samples keep almost constant around at 80-90%, that is to say, the phonon scattering due to cation disordering would be maintained at a constant level, thus the
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Besides, the
outermost electron dissatisfaction would make the chemical bond between Mn and S less compact and lead to softened chemical bonds, which is beneficial to the suppression of κl [10, 39].
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It is interesting to see in Fig. 6(c) that, for the doped Cu2Sn1-xMnxS3 samples, the κl values are very close despite their different doping levels, furthermore, they are almost at the same level with the data of Zn-doped CTS in the whole temperature range. However, the κl of Co-doped samples are different from them at 323 K. Generally, κl can be given as κl = 1/3vavgCvl, where vavg, Cv, and l are the average speed of sound, heat capacity, and phonon mean free path, respectively. In order to
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explicate the reduction of κl, the average sound velocities of longitudinal and transverse branches, as an important parameter determining the κl, are measured and listed in Table 2. The vl and vt of Cu2Sn1-xMnxS3 decrease gradually with the doping amount at 300 K, which can be traced back to
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the weak chemical bonds due to the crystal structure transformation. Comparing to Co-doping samples, the vt of Cu2Sn1-xMnxS3 are a little larger which may result from the larger relative density because of different SPS temperature (773K and 823K for Co- and Mn-doped CTS, respectively). The vavg of both Co- and Mn-doped samples do not differ from each other and both are lower than the pristine CTS for a certain extent. Clearly, the small vavg is one reason for the abnormally low κl of CTS doped samples. The Grüneisen parameter γ can be calculated from the
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measured vl and vt by γ = 3 (1+ν P ) 2 ( 2-3ν P ) , where Poisson ratio vp can be derived from the vl and vt by v p = 1 − 2 ( vt vl )2 2 − 2 ( vt vl )2 . The calculated γ of the Mn-doped samples ranges from
1.71 to 1.88, nearly the same as that of Co-doping samples (1.71~1.92). This similarity indicates that the distinction of anharmonic lattice vibrations is not obvious for the doped sample, which is
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to say, the relatively larger κl in the low temperature range in Mn-doped CTS can be mainly attributed to the higher density than that of Co-doped CTS.
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Table 2. Parameters of density ρ (g·cm-3), relative density K (%), longitudinal νl (m·s-1), transverse νt (m·s-1), average speed of sound νavg (m·s-1), Debye temperature θD (K) and Grüneisen parameter γ of bulk materials Cu2Sn1-xMxS3 (x = 0-0.20; M = Mn, Co) at 300 K. Dopant
Mn
Co
x
ρ
K
νl
νt
vavg
θD
γ
0 0.05 0.10 0.15 0.20 0.05 0.10 0.15 0.20
4.58 4.57 4.56 4.54 4.56 4.52 4.46 4.43 4.53
96.42 96.61 96.40 96.18 96.81 95.16 93.89 93.46 95.56
4648 3975 3879 3634 3860 3687 3940 3540 3872
2373 2057 2043 1980 2075 1968 2015 1924 1988
2658 2303 2284 2209 2317 2198 2257 2146 2227
287.14 249.39 248.00 240.18 253.13 264.70 264.01 264.90 264.35
1.93 1.88 1.82 1.71 1.75 1.78 1.92 1.71 1.91
Fig. 6(d) shows the figure of merit ZT calculated from the data above. The ZT values increase
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Conclusions In this work, Mn is demonstrated as an effective d-unfilled dopant for tuning carrier concentration and effective mass of carriers for CTS as a TE material. It is shown that d-unfilled TM (such as Co, Mn) doping induced an order-to-disorder crystal transformation from monoclinic to tetragonal and cubic structures. Such a transformation is crucial for both electrical and thermal
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benefits for TE performance. On one hand, it causes a high carrier concentration and enables both the improved degeneracy of up VBs and the participation of d-orbitals to achieve an enhanced DOS effective mass, thus leading to a high PF. On the other hand, a low κl even close to
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theoretical minimum can be obtained in the tetragonal and cubic dominated bulks, due primarily to the high degree disorder cation occupation and bond softening, and other accompanying defects (precipitation and vacancies) as well. These findings point out the significance of phase transformation of CTS by way of suitable TM doping and would be helpful to get a systematical understanding of the fundamental parameters determining the transport properties of CTS.
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Acknowledgements
Reference
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This research was supported by the National Natural Science Foundation of China under Grant No. 51272103, 51672127, No. 11604047, the Natural Science Foundation of Jiangsu Province (No. BK20160694), the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD).
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Fig.1 (a) Powder X-ray diffraction patterns of Cu2Sn1-xMnxS3 (x = 0-0.20), together with photographs of the as-synthesized powder of undoped and doped CTS in insert 1, and a local zoom demonstrating the impurity peaks of CuS in insert 2. Crystal structures of CTS in monoclinic (b), cubic (c) and tetragonal (d), with ordered cation occupation in (b), disordered (at the 4a site) in (c), and (at the 4d and 2b sites) in (d).
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Fig. 2 SEM images of fractured surface for samples with (a) x = 0 and (b) x = 0.15, and a magnified local view (c) for the x = 0.15 sample showing precipitates on the grain boundaries (Region a) and the interior (Region b). (d) shows the corresponding EDS patterns for Region a and b in (c).
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Fig.3 Rietveld refinements results (a) for Powder XRD patterns of Cu2Sn1-xMnxS3 (x = 0.05-0.20), and derived lattice parameters of a and c (b), and the bond lengths of M-S and Cu-S (c) as functions of the doping level.
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Fig. 4 Thermoelectric parameters of (a) electrical conductivity σ, (b) Seebeck coefficient S in the measured temperature range, together with inserts showing their dependence on the Mn-doping level, respectively. Variation of carrier concentration n and Hall mobility µHall (c) and DOS effective mass of carriers m*(d) with doping level of x in comparison with cited date for Co-[10] and Zn-doped[11] CTS at 300 K.
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Fig.5 (a) Pisarenko plot of S v.s. carrier concentration n in comparison with cited data of Co-[10] and Zndoped[11] CTS, where the lines represent the calculated relation based on SPB model, and (b) PF = S2σ for the Cu2Sn1-xMnxS3 (x = 0-0.20) bulks at different temperatures, with cited data for 20%Co- and Zndoped CTS for comparison.
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Fig. 6 Thermal transport properties of (a) total thermal conductivity κ and lattice thermal conductivity κl (insert), (b) electronic thermal conductivity κe and Lorenz number L (insert) and (c) its variation with doping level at 323 and 723 K, with cited data of Co-[10] and Zn-doped[11] CTS for comparison. (d) shows the calculated dimensionless figure of merit ZT for Cu2Sn1-xMnxS3 (x = 0-0.20) samples, where the maximal ZT~0.68 was obtained in the 15%Mn-doped CTS at 723 K.
ACCEPTED MANUSCRIPT The main highlights of the paper includes (1) Mn-doping in Cu2SnS3 caused transitions from cation-ordered to disordered phases. (2) Low lattice thermal conductivity was obtained in cation-disordered structures. (3) High carrier effective mass and large PF caused by contribution of Mn 3d states.
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(4) Maximal ZT ~0.68 was achieved at 723 K by optimal Mn-doping in Cu2SnS3.