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Enhanced thermoelectric and mechanical properties in hierarchical tubular porous cuprous selenide
Scripta Materialia 176 (2020) 104–107
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Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Enhanced thermoelectric and mechanical properties in hierarchical tubular porous cuprous selenide Jixing Liu a,b,1, Meng Li c,1,∗, Shenghui Yang d,∗, Shengnan Zhang a, Jianqing Feng a, Chengshan Li a,b, Pingxiang Zhang a,b, Lian Zhou a,b a
Superconducting Materials Research Center, Northwest Institute for Nonferrous Metal Research, Xi’an, 710016, China School of Materials Science and Engineering, Northeastern University, Shenyang, 110819, China School of Mechanical and Mining, University of Queensland, Brisbane, QLD 4067, Australia d School of Materials Science and Engineering, Xi’an University of Technology, Xi’an, 710000, China b c
a r t i c l e
i n f o
Article history: Received 18 June 2019 Revised 27 August 2019 Accepted 6 September 2019 Available online 11 October 2019 Keywords: Thermoelectric materials Cu2 Se Structure engineering
a b s t r a c t Superionic thermoelectric Cu2 Se bulks are successfully acquired via an identical gas reinforced eutectic transformation (gasar) process based on its melting-congruent nature. The gasar synthesis process introduces tubular pores to Cu2 Se matrix, which have hierarchical size and corrugated surface, and are distributed homogeneously and isotopically. The Cu2 Se bulk with such a unique porous structure shows enhanced thermoelectric figure-of-merit up to ∼2.1 at 873 K owing to the conspicuously reduced thermal conductivity, which can be attributed to the increased phonon scattering. Moreover, the porous Cu2 Se bulk also possesses higher mechanical robustness compared with traditional porous-structured material because of the corrugated surface, which is promising for making thermoelectric modules. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Thermoelectric (TE) technology provides a promising solution to the current global energy crisis by converting waste heat to electricity or vice versa [1,2]. However, the inferior intrinsic energy conversion efficiency of TE materials, primarily represented by the dimensionless figure-of-merit (zT), restricts TE technology from wide application. The zT of a given material is defined as zT = S2 σ T /κ , where S, σ and κ are Seebeck coefficient, electrical and thermal conductivity, respectively [3–5]. Due to the competing correlations of charge/phonon transport and the complicated entanglements between real/reciprocal space in TE materials, it is experimentally difficult to decouple these parameters and improve zT [6,7]. In recent decades, owing to the development in semiconductor theory [8], the breakthrough on nano-engineering [9] and the application of quantum concept [10], methods to enhance zT have flourishingly emerged, such as band engineering [11], defects scattering [12], bond anharmonicity [13], phonon spectra modifying, nano structuring [14], phase hybridization [15], etc. Recent studies reported high zT in porous TE materials because that the existence of voids in matrix can effectively suppress the phonon mean free path (MFP) to reduce thermal conductivity without sacrificing the power factor (P F = S2 σ ) [16–19]. Besides,
∗
1
Corresponding authors. E-mail addresses: [email protected] (M. Li), [email protected] (S. Yang). These authors contribute equally to this work.
https://doi.org/10.1016/j.scriptamat.2019.09.009 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the economical cost on pristine elements and the improved portability make porous TE materials more promising for making TE modules. In contrast to the aforementioned advantages, however, several issues of porous TE materials need to be settled before being considered for further application. (1) Lack of fast and clean (namely no residual impurity of void foaming medium) synthesis of porous TE materials. (2) Lack of precise control of the size, shape and distribution of pores. (3) Deterioration of chemical stability and mechanical robustness. In all, it should be imperative to develop porous TE materials with both high performance and good stability, while facile and cheap to synthesize. Herein, gas reinforced eutectic transformation (gasar) process was firstly utilized to form porous structure in superionic thermoelectric Cu2 Se due to its melting-congruent nature and comprehensively understood porosity properties [20–24]. The formed pores were in tubular shape scoping from nanometric to micrometric hierarchically, in which the corrugated inner surface was observed. The tubular porous Cu2 Se showed slight fluctuation of Seebeck coefficient and electrical conductivity, but conspicuously reduced thermal conductivity, leading to an enhanced zT up to ∼2.1 at 873 K. It also possessed high mechanical robustness, namely hardness, thermal expansion and fracture toughness, compared with reported porous Cu2 Se. The facile fabrication, optimized TE performance as well as the ameliorated mechanical properties indicate that the gasar process is promising for synthesizing porous TE materials.
J. Liu, M. Li and S. Yang et al. / Scripta Materialia 176 (2020) 104–107
Fig. 1. Schematic diagram illustrating the (a) pristine Cu2 Se with void foaming agent, (b) formation of tubular porous Cu2 Se via gasar process and (c) suppressed phonon transport by interfacial scattering and heat radiation.
Porous Cu2 Se samples are synthesized via a typical gasar process demonstrated in Fig. 1(a) and (b). Carbon nanoparticles (CNPs) pre-treated with acetone and the melting-congruent Cu2 Se powder are hand ground and sealed in an evacuated silica tube, then melted at 1200 °C. In the melted liquid, the CNPs tend to float up and precipitate as carbon segregation on the surface of liquid, while the acetone molecules are released and dissolved in melted Cu2 Se forming eutectic. When the eutectic is solidified during cooling process, nucleation of acetone bubbles occurs heterogeneously at the discontinuities of the liquid/solid interfaces [21]. The liquid/solid interfaces keep moving with the temperature gradient, driving the bubbles to move accordingly and leaving tubular pores along the moving trace. Because the crystallization of Cu2 Se and the formation of tubular pores proceed simultaneously, the tubular pores tend to appear at the grain boundaries to minimize the formation energy, thus show preferential orientation within Cu2 Se grains but are generally isotropic in the entire Cu2 Se ingot. Fig. S1 and S2 indicate that the XRD patterns of gasar synthesized porous Cu2 Se could be indexed as single phase (JCPDS #46-1129) without any impurities of void foaming agent. A reversable phase transition from monoclinic (C2/c) structure to cubic (Fm3¯ m) structure is detected at ∼410 K, which is consistent with previously reported Cu2 Se samples [23–25]. Scanning electron microscope (SEM) images of the fractography for P (highly dense sample), P + A (porous sample pre-treated with acetone) and P + A + C (porous sample pre-treated with both acetone and carbon nanopaticles) Cu2 Se samples are presented in Fig. 2(a)–(c). Sample P consists of lamella-grains without any visible voids or impurity, thus having high density of 6.641 gcm−3 , which is nearly 98% of theoretical density. On the contrary, samples P + A and P + A + C are in porous structure. In Fig. 2(b), sample P + A contains few tubular pores with the average diameter of ∼100 μm and length of ∼500 μm. While in Fig. 2(c), sample P + A + C contains serried tubular pores with smaller diameter (∼10 μm) and length (∼80 μm). The tubular pores in both sample P + A and P + A + C are much smaller than that in traditional gasar materials where the pores are in centimetres, which should be attributed to the extreme quick solidification by liquid-nitrogen quenching [20–22]. The tubular pores show preferential orientation within a Cu2 Se grain, while in general they are isotopically dis-
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Fig. 2. SEM images of fractography for (a) melting-solidification synthesized pristine Cu2 Se sample (P), (b) gasar process synthesized Cu2 Se sample with precursor of acetone (P + A) and (c) gasar process synthesized Cu2 Se sample with precursor of acetone absorbed carbon nanoparticle (P + A + C). (d) and (e) refer to area 1 and 2 in (c), respectively, showing the pores parallel or as-normal to the observation. (f) refers to area 3 in (d) showing the corrugated inner surface of the pores. (g) Shows the TEM image of P + A + C sample. (h) and (i) refer to area 4 and 5 in (g), respectively.
tributed. Noticeably, the inner surfaces of the tubular pores are corrugated, as is shown in Fig. 2(f), and the “crest distance” is around 130 nm. Transmission electron microscopy (TEM) image of FIB milled lamella from sample P + A + C is shown in Fig. 2(g–i), in which nanoscale tubular pores along and perpendicular to the observation direction are characterized, indicating the tubular pores are in diameter from nanometres to micrometres hierarchically. Fig. 2(h) and (i) present the nano-sized tubular pores parallel and perpendicular to the observation, respectively. The linear diameter of the nano-sized pores is around 10 nm, and the roughness of their corrugated inner surface is smaller accordingly. The temperature dependence of the electrical conductivity (σ ), Seebeck coefficient (S), power factor (S2 σ ), total thermal conductivity (κ ), lattice thermal conductivity (κ L ) and figure-of-merit (zT) for the tubular porous Cu2 Se samples P + A and P + A + C and the pristine dense sample P are plotted in Fig. 3(a)–(c). The uncertainties of measurement and calculation are determined by the progradation law as is shown in Table S2. All the synthesized samples show evident phase transition from low-temperature monoclinic phase to high-temperature cubic phase at ∼410 K, which is consistent with XRD measurement. The σ values of all the samples first increase with temperature in monoclinic phase, then decrease with temperature in cubic phase. By contrast, S values of all the samples keep upward before and after phase transition, and the S values drop slightly during phase transition temperature. The S increases slightly with porosity, which is attributed to the energy filtering effect by the interfaces between pores and matrix Cu2 Se [26]. The maximum Seebeck coefficient is 275 μV/K at 873 K for sample P + A + C 0.09, which is a relatively high value among the reported Cu2 Se-based thermoelectric materials. Accordingly, the power factor shows a positive correlation with sample porosity with an increase of 36% from 8.0 μW/mK2 to 10.9 μW/mK2 at 823 K in sample P + A + C 0.09.
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Fig. 3. Temperature dependence of (a) electrical conductivity (σ ), (b) Seebeck coefficient (S), (c) power factor (S2 σ ), (d) total thermal conductivity (κ ), (e) lattice thermal conductivity (κ L ) and (f) figure-of-merit (zT).
Fig. 3(d) shows the total thermal conductivity of Cu2 Se samples with different porosity. Although the ĸ values for all the samples display a similar trend with increasing temperature, the ĸ values of P + A and P + A + C samples are remarkably lower than that of pristine Cu2 Se sample. To understand the phonon transport in porous Cu2 Se, lattice thermal conductivity (κ L ) were calculated by subtracting charge carrier thermal conductivity (κ C ) from the total thermal conductivity as shown in Fig. 3(f). The κ C is calculated according to the Wiedemann-Franz law, κC = Lσ T , where L is Lorentz constant and is approximated by empirical function L = 1.5 + exp[−|S|/116] [27]. The calculated κ L values for all the samples are slightly downward in high-temperature phase, implying that both the dense and porous Cu2 Se are thermal stable up to 875 K. With the porosity of P + A + C samples increasing from 3% to 12%, the κ L decreases by 50%, which is consistent with reported values by Zhao et al. [24]. In Fig. S5, the κ L of the tubular porous Cu2 Se calculated by effective medium theory is much higher than the experimental observed κ L , which is attributed to the effect of porosity induced interfacial thermal resistance (ITR), consisting of thermal boundary resistance (TBR) and thermal contact resistance (TCR). The TBR, also known as Kapitza resistance, is determined by phonon transmission probability (α ) in constituent materials. The α value of a phonon i with frequency of ω propagating from phase (A) to phase (B) is:
αi (ωi ) = j
GB (ωi , T )ci,B
j GA (ωi , T )ci,A
where G(ωi , T) is phonon density of states (DOS) at energy ω and temperature T. In the case of porous material, A refers to matrix Cu2 Se and B refers to pores which are considered as vacuum or gas-filled. The giant mismatch of phonon DOS between A and B would greatly decrease α and increase Kapitza resistance [28]. TCR, which is usually caused by the weak mechanical and chemical bonding between constituent phases, would impede the low-frequency phonons by normal boundary scattering and highfrequency phonons by surface bond order imperfections [29]. The boundary scattering arises from the termination of lattice periodicity in the surface with normal direction of matrix material, which can shorten and strengthen the bonds nearby and provide perturbation to the local potential, thus impact the phonon transport dynamics [30]. Moreover, the interfacial roughness can also lead to the reduce of thermal conductivity by forming the surface bond order imperfections to increase the surface-to-volume ratio [31–33], and enhance the phonon scattering. In all, the synergistic effect of TBR and TCR would significantly reduce the lattice thermal conductivity, thus enhance the figure-of-merit. The maximum zT of ∼2.1 is obtained in Cu2 Se bulk with porosity of 9% at 873 K, which is higher than most previous reported values of bulk Cu2 Se materials, indicating gasar process a feasible methodology to synthesize porous TE materials with low thermal conductivity and high figure-of-merit. Mechanical properties are essential for the practical application of TE materials, for they not only determine the cycle durability of TE modulus but also reflect the TE performance to some extent.
J. Liu, M. Li and S. Yang et al. / Scripta Materialia 176 (2020) 104–107
107
Table 1 Vicker’s hardness of P + A + C samples compared with reported results and calculated fracture toughness (KIC ). No.
HV this work (GPa)
Reported HV (GPa) [24,34]
Calculated HV (GPa) [35,36]
Calculated KIC (MPa m1/2 ) [37]
0.43 0.41 0.38 0.29 0.21
0.38∗ 0.34 0.30∗ 0.26∗
0.380 0.336 0.295 0.258
2.25 2.12 1.95 1.81 1.55
Unfortunately, the existence of pores in matrix would inevitably deteriorate the mechanical properties due to the discontinuity of matrix materials. Herein we measure the hardness and fracture toughness of the tubular Cu2 Se samples indirectly based on following formulas [37]:
References
P P+A+C P+A+C P+A+C P+A+C ∗
0.03 0.06 0.09 0.12
Fitted.
H = P/α0 a2 K = P/β0 a3/2 where H is hardness, K is fracture toughness, α 0 and β 0 are numerical constants, P is the peak load, a and c are characteristic dimensions of the “plastic” impression and the radial/median crack, respectively. Compared with traditional porous materials, owing to the existence of corrugated inner surfaces, the gasar process synthesized porous Cu2 Se bulks in this work exhibit favourable mechanical properties. The tubular porous Cu2 Se samples show enhanced microhardness that is even comparable with those dense samples when the porosity is low. The measured thermal expansion coefficient and bending impedance plotted in Fig. S6-7 also demonstrate that the tubular porous structure reinforced with corrugated inner surface can keep a majority of mechanical properties of matrix sample, which is important for thermoelectric application. In this work, Cu2 Se bulks with tubular porous structures were synthesized through a gas reinforced eutectic transformation method. The obtained bulks contained tubular pores scoping from nanometres to micrometres hierarchically, inside which existed the corrugated surface. This unique structure was proposed to slightly modulate carrier transport by interfacial filtering effect, but drastically suppress the phonon progradation due to boundary resistance, surface state localization and thermal radiation. Thus a maximum figure-of-merit of ∼2.1 at 873 K was achieved. Moreover, the Cu2 Se bulks with tubular porous structure also possessed favourable mechanical properties, such as Vickers hardness, fracture toughness and bending impedance compared with traditional porous materials. Our work provided a promising strategy to fabricate porous TE materials with both high figure-of-merit and good mechanical properties Table 1. Acknowledgements This research was supported by funding from the Light of West China (No. XAB2018AW15) and the International Cooperative Project in Shaanxi Province (No. 2018kw-055). Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.09. 009.
[1] J. He, T.M. Tritt, Science 357 (6358) (2017). [2] G.J. Snyder, E.S. Toberer, Nat. Mater. 7 (2) (2008) 105. [3] A.F. Ioffe, Semiconductor Thermoelements and Thermoelectric Cooling, Infosearch Limited, London, 1957. [4] H.J. Goldsmid, Introduction to Thermoelectricity, Springer, Berlin, Germany, 2009. [5] D.M. Rowe, CRC Handbook of Thermoelectrics, CRC Press, Boca Raton, Florida, United States, 1995. [6] T. Zhu, Y. Liu, C. Fu, J.P. Heremans, J.G. Snyder, X. Zhao, Adv. Mater. 29 (14) (2017) 1605884. [7] H. Zhu, C. Xiao, Y. Xie, Adv. Mater. 0 (0) (2018) 18020 0 0. [8] I.I. Ravich, Semiconducting Lead Chalcogenides, Springer Science & Business Media, 2013. [9] J.P. Heremans, M.S. Dresselhaus, L.E. Bell, D.T. Morelli, Nat. Nanotechnol. 8 (2013) 471. [10] L. Hicks, M. Dresselhaus, Phys. Rev. B 47 (24) (1993) 16631. [11] Y. Pei, X. Shi, A. LaLonde, H. Wang, L. Chen, G.J. Snyder, Nature 473 (2011) 66. [12] S.I. Kim, K.H. Lee, H.A. Mun, H.S. Kim, S.W. Hwang, J.W. Roh, D.J. Yang, W.H. Shin, X.S. Li, Y.H. Lee, G.J. Snyder, S.W. Kim, Science 348 (6230) (2015) 109–114. [13] J.P. Heremans, Nat. Phys. 11 (12) (2015) 990. [14] M. Hong, Y. Wang, T. Feng, Q. Sun, S. Xu, S. Matsumura, S.T. Pantelides, J. Zou, Z.-.G. Chen, J. Am. Chem. Soc. 141 (4) (2019) 1742–1748. [15] R. Liu, H. Chen, K. Zhao, Y. Qin, B. Jiang, T. Zhang, G. Sha, X. Shi, C. Uher, W. Zhang, L. Chen, Adv. Mater. 29 (38) (2017) 1702712. [16] P. Giulia, Nat. Rev.Mater. 2 (2017) 17006. [17] X. Shi, A. Wu, W. Liu, R. Moshwan, Y. Wang, Z.-.G. Chen, J. Zou, ACS Nano (2018). [18] M. Hong, Y. Wang, S. Xu, X. Shi, L. Chen, J. Zou, Z.-.G. Chen, Nano Energy 60 (2019) 1–7. [19] B. Xu, T. Feng, M.T. Agne, L. Zhou, X. Ruan, G.J. Snyder, Y. Wu, Angew. Chem. Int. Edit. 56 (13) (2017) 3546–3551. [20] L. Drenchev, J. Sobczak, S. Malinov, W. Sha, Mat. Sci. Technol. 22 (10) (2006) 1135–1147. [21] V. Shapovalov, L. Boyko, Adv. Eng. Mater. 6 (6) (2004) 407–410. [22] J. Apprill, D.R. Poirier, M. Maguire, T. Gutsch, in: MRS Proceedings, 521, 1998. [23] H. Liu, X. Shi, F. Xu, L. Zhang, W. Zhang, L. Chen, Q. Li, C. Uher, T. Day, G.J. Snyder, Nat. Mater. 11 (2012) 422. [24] K. Zhao, H. Duan, N. Raghavendra, P. Qiu, Y. Zeng, W. Zhang, J. Yang, X. Shi, L. Chen, Adv. Mater. 29 (42) (2017) 1701148. [25] H. Liu, X. Yuan, P. Lu, X. Shi, F. Xu, Y. He, Y. Tang, S. Bai, W. Zhang, L. Chen, Y. Lin, L. Shi, H. Lin, X. Gao, X. Zhang, H. Chi, C. Uher, Adv. Mater. 25 (45) (2013) 6607–6612. [26] D. Vashaee, A. Shakouri, Phys. Rev. Lett. 92 (10) (2004) 106103. [27] H.-.S. Kim, Z.M. Gibbs, Y. Tang, H. Wang, G.J. Snyder, APL Mater. 3 (4) (2015) 041506. [28] E.T. Swartz, R.O. Pohl, Rev. Mod. Phys. 61 (3) (1989) 605–668. [29] Z. Tešanovic´ , M.V. Jaric´ , S. Maekawa, Phys. Rev. Lett. 57 (21) (1986) 2760–2763. [30] G. Xie, Z. Ju, K. Zhou, X. Wei, Z. Guo, Y. Cai, G. Zhang, NPJ Comput. Mater. 4 (1) (2018) 21. [31] C. Blanc, A. Rajabpour, S. Volz, T. Fournier, O. Bourgeois, Appl. Phys. Lett. 103 (4) (2013) 043109. [32] W. Gao, X. Yi, B. Sun, X. Meng, W. Cai, L. Zhao, Acta Mater. 132 (2017) 405–415. [33] P. Martin, Z. Aksamija, E. Pop, U. Ravaioli, Phys. Rev. Lett. 102 (12) (2009) 125503. [34] K. Tyagi, B. Gahtori, S. Bathula, M. Jayasimhadri, S. Sharma, N.K. Singh, D. Haranath, A.K. Srivastava, A. Dhar, Solid State Commun. 207 (2015) 21–25. [35] H. Nakajima, Mat. Sci. Appl. 9 (2018) 258–267. [36] J. Luo, R. Stevens, Ceram. Int. 25 (3) (1999) 281–286. [37] G. Anstis, P. Chantikul, B.R. Lawn, D. Marshall, J. Am. Ceram. Soc. 64 (9) (1981) 533–538.
Contents lists available at ScienceDirect
Scripta Materialia journal homepage: www.elsevier.com/locate/scriptamat
Enhanced thermoelectric and mechanical properties in hierarchical tubular porous cuprous selenide Jixing Liu a,b,1, Meng Li c,1,∗, Shenghui Yang d,∗, Shengnan Zhang a, Jianqing Feng a, Chengshan Li a,b, Pingxiang Zhang a,b, Lian Zhou a,b a
Superconducting Materials Research Center, Northwest Institute for Nonferrous Metal Research, Xi’an, 710016, China School of Materials Science and Engineering, Northeastern University, Shenyang, 110819, China School of Mechanical and Mining, University of Queensland, Brisbane, QLD 4067, Australia d School of Materials Science and Engineering, Xi’an University of Technology, Xi’an, 710000, China b c
a r t i c l e
i n f o
Article history: Received 18 June 2019 Revised 27 August 2019 Accepted 6 September 2019 Available online 11 October 2019 Keywords: Thermoelectric materials Cu2 Se Structure engineering
a b s t r a c t Superionic thermoelectric Cu2 Se bulks are successfully acquired via an identical gas reinforced eutectic transformation (gasar) process based on its melting-congruent nature. The gasar synthesis process introduces tubular pores to Cu2 Se matrix, which have hierarchical size and corrugated surface, and are distributed homogeneously and isotopically. The Cu2 Se bulk with such a unique porous structure shows enhanced thermoelectric figure-of-merit up to ∼2.1 at 873 K owing to the conspicuously reduced thermal conductivity, which can be attributed to the increased phonon scattering. Moreover, the porous Cu2 Se bulk also possesses higher mechanical robustness compared with traditional porous-structured material because of the corrugated surface, which is promising for making thermoelectric modules. © 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
Thermoelectric (TE) technology provides a promising solution to the current global energy crisis by converting waste heat to electricity or vice versa [1,2]. However, the inferior intrinsic energy conversion efficiency of TE materials, primarily represented by the dimensionless figure-of-merit (zT), restricts TE technology from wide application. The zT of a given material is defined as zT = S2 σ T /κ , where S, σ and κ are Seebeck coefficient, electrical and thermal conductivity, respectively [3–5]. Due to the competing correlations of charge/phonon transport and the complicated entanglements between real/reciprocal space in TE materials, it is experimentally difficult to decouple these parameters and improve zT [6,7]. In recent decades, owing to the development in semiconductor theory [8], the breakthrough on nano-engineering [9] and the application of quantum concept [10], methods to enhance zT have flourishingly emerged, such as band engineering [11], defects scattering [12], bond anharmonicity [13], phonon spectra modifying, nano structuring [14], phase hybridization [15], etc. Recent studies reported high zT in porous TE materials because that the existence of voids in matrix can effectively suppress the phonon mean free path (MFP) to reduce thermal conductivity without sacrificing the power factor (P F = S2 σ ) [16–19]. Besides,
∗
1
Corresponding authors. E-mail addresses: [email protected] (M. Li), [email protected] (S. Yang). These authors contribute equally to this work.
https://doi.org/10.1016/j.scriptamat.2019.09.009 1359-6462/© 2019 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved.
the economical cost on pristine elements and the improved portability make porous TE materials more promising for making TE modules. In contrast to the aforementioned advantages, however, several issues of porous TE materials need to be settled before being considered for further application. (1) Lack of fast and clean (namely no residual impurity of void foaming medium) synthesis of porous TE materials. (2) Lack of precise control of the size, shape and distribution of pores. (3) Deterioration of chemical stability and mechanical robustness. In all, it should be imperative to develop porous TE materials with both high performance and good stability, while facile and cheap to synthesize. Herein, gas reinforced eutectic transformation (gasar) process was firstly utilized to form porous structure in superionic thermoelectric Cu2 Se due to its melting-congruent nature and comprehensively understood porosity properties [20–24]. The formed pores were in tubular shape scoping from nanometric to micrometric hierarchically, in which the corrugated inner surface was observed. The tubular porous Cu2 Se showed slight fluctuation of Seebeck coefficient and electrical conductivity, but conspicuously reduced thermal conductivity, leading to an enhanced zT up to ∼2.1 at 873 K. It also possessed high mechanical robustness, namely hardness, thermal expansion and fracture toughness, compared with reported porous Cu2 Se. The facile fabrication, optimized TE performance as well as the ameliorated mechanical properties indicate that the gasar process is promising for synthesizing porous TE materials.
J. Liu, M. Li and S. Yang et al. / Scripta Materialia 176 (2020) 104–107
Fig. 1. Schematic diagram illustrating the (a) pristine Cu2 Se with void foaming agent, (b) formation of tubular porous Cu2 Se via gasar process and (c) suppressed phonon transport by interfacial scattering and heat radiation.
Porous Cu2 Se samples are synthesized via a typical gasar process demonstrated in Fig. 1(a) and (b). Carbon nanoparticles (CNPs) pre-treated with acetone and the melting-congruent Cu2 Se powder are hand ground and sealed in an evacuated silica tube, then melted at 1200 °C. In the melted liquid, the CNPs tend to float up and precipitate as carbon segregation on the surface of liquid, while the acetone molecules are released and dissolved in melted Cu2 Se forming eutectic. When the eutectic is solidified during cooling process, nucleation of acetone bubbles occurs heterogeneously at the discontinuities of the liquid/solid interfaces [21]. The liquid/solid interfaces keep moving with the temperature gradient, driving the bubbles to move accordingly and leaving tubular pores along the moving trace. Because the crystallization of Cu2 Se and the formation of tubular pores proceed simultaneously, the tubular pores tend to appear at the grain boundaries to minimize the formation energy, thus show preferential orientation within Cu2 Se grains but are generally isotropic in the entire Cu2 Se ingot. Fig. S1 and S2 indicate that the XRD patterns of gasar synthesized porous Cu2 Se could be indexed as single phase (JCPDS #46-1129) without any impurities of void foaming agent. A reversable phase transition from monoclinic (C2/c) structure to cubic (Fm3¯ m) structure is detected at ∼410 K, which is consistent with previously reported Cu2 Se samples [23–25]. Scanning electron microscope (SEM) images of the fractography for P (highly dense sample), P + A (porous sample pre-treated with acetone) and P + A + C (porous sample pre-treated with both acetone and carbon nanopaticles) Cu2 Se samples are presented in Fig. 2(a)–(c). Sample P consists of lamella-grains without any visible voids or impurity, thus having high density of 6.641 gcm−3 , which is nearly 98% of theoretical density. On the contrary, samples P + A and P + A + C are in porous structure. In Fig. 2(b), sample P + A contains few tubular pores with the average diameter of ∼100 μm and length of ∼500 μm. While in Fig. 2(c), sample P + A + C contains serried tubular pores with smaller diameter (∼10 μm) and length (∼80 μm). The tubular pores in both sample P + A and P + A + C are much smaller than that in traditional gasar materials where the pores are in centimetres, which should be attributed to the extreme quick solidification by liquid-nitrogen quenching [20–22]. The tubular pores show preferential orientation within a Cu2 Se grain, while in general they are isotopically dis-
105
Fig. 2. SEM images of fractography for (a) melting-solidification synthesized pristine Cu2 Se sample (P), (b) gasar process synthesized Cu2 Se sample with precursor of acetone (P + A) and (c) gasar process synthesized Cu2 Se sample with precursor of acetone absorbed carbon nanoparticle (P + A + C). (d) and (e) refer to area 1 and 2 in (c), respectively, showing the pores parallel or as-normal to the observation. (f) refers to area 3 in (d) showing the corrugated inner surface of the pores. (g) Shows the TEM image of P + A + C sample. (h) and (i) refer to area 4 and 5 in (g), respectively.
tributed. Noticeably, the inner surfaces of the tubular pores are corrugated, as is shown in Fig. 2(f), and the “crest distance” is around 130 nm. Transmission electron microscopy (TEM) image of FIB milled lamella from sample P + A + C is shown in Fig. 2(g–i), in which nanoscale tubular pores along and perpendicular to the observation direction are characterized, indicating the tubular pores are in diameter from nanometres to micrometres hierarchically. Fig. 2(h) and (i) present the nano-sized tubular pores parallel and perpendicular to the observation, respectively. The linear diameter of the nano-sized pores is around 10 nm, and the roughness of their corrugated inner surface is smaller accordingly. The temperature dependence of the electrical conductivity (σ ), Seebeck coefficient (S), power factor (S2 σ ), total thermal conductivity (κ ), lattice thermal conductivity (κ L ) and figure-of-merit (zT) for the tubular porous Cu2 Se samples P + A and P + A + C and the pristine dense sample P are plotted in Fig. 3(a)–(c). The uncertainties of measurement and calculation are determined by the progradation law as is shown in Table S2. All the synthesized samples show evident phase transition from low-temperature monoclinic phase to high-temperature cubic phase at ∼410 K, which is consistent with XRD measurement. The σ values of all the samples first increase with temperature in monoclinic phase, then decrease with temperature in cubic phase. By contrast, S values of all the samples keep upward before and after phase transition, and the S values drop slightly during phase transition temperature. The S increases slightly with porosity, which is attributed to the energy filtering effect by the interfaces between pores and matrix Cu2 Se [26]. The maximum Seebeck coefficient is 275 μV/K at 873 K for sample P + A + C 0.09, which is a relatively high value among the reported Cu2 Se-based thermoelectric materials. Accordingly, the power factor shows a positive correlation with sample porosity with an increase of 36% from 8.0 μW/mK2 to 10.9 μW/mK2 at 823 K in sample P + A + C 0.09.
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Fig. 3. Temperature dependence of (a) electrical conductivity (σ ), (b) Seebeck coefficient (S), (c) power factor (S2 σ ), (d) total thermal conductivity (κ ), (e) lattice thermal conductivity (κ L ) and (f) figure-of-merit (zT).
Fig. 3(d) shows the total thermal conductivity of Cu2 Se samples with different porosity. Although the ĸ values for all the samples display a similar trend with increasing temperature, the ĸ values of P + A and P + A + C samples are remarkably lower than that of pristine Cu2 Se sample. To understand the phonon transport in porous Cu2 Se, lattice thermal conductivity (κ L ) were calculated by subtracting charge carrier thermal conductivity (κ C ) from the total thermal conductivity as shown in Fig. 3(f). The κ C is calculated according to the Wiedemann-Franz law, κC = Lσ T , where L is Lorentz constant and is approximated by empirical function L = 1.5 + exp[−|S|/116] [27]. The calculated κ L values for all the samples are slightly downward in high-temperature phase, implying that both the dense and porous Cu2 Se are thermal stable up to 875 K. With the porosity of P + A + C samples increasing from 3% to 12%, the κ L decreases by 50%, which is consistent with reported values by Zhao et al. [24]. In Fig. S5, the κ L of the tubular porous Cu2 Se calculated by effective medium theory is much higher than the experimental observed κ L , which is attributed to the effect of porosity induced interfacial thermal resistance (ITR), consisting of thermal boundary resistance (TBR) and thermal contact resistance (TCR). The TBR, also known as Kapitza resistance, is determined by phonon transmission probability (α ) in constituent materials. The α value of a phonon i with frequency of ω propagating from phase (A) to phase (B) is:
αi (ωi ) = j
GB (ωi , T )ci,B
j GA (ωi , T )ci,A
where G(ωi , T) is phonon density of states (DOS) at energy ω and temperature T. In the case of porous material, A refers to matrix Cu2 Se and B refers to pores which are considered as vacuum or gas-filled. The giant mismatch of phonon DOS between A and B would greatly decrease α and increase Kapitza resistance [28]. TCR, which is usually caused by the weak mechanical and chemical bonding between constituent phases, would impede the low-frequency phonons by normal boundary scattering and highfrequency phonons by surface bond order imperfections [29]. The boundary scattering arises from the termination of lattice periodicity in the surface with normal direction of matrix material, which can shorten and strengthen the bonds nearby and provide perturbation to the local potential, thus impact the phonon transport dynamics [30]. Moreover, the interfacial roughness can also lead to the reduce of thermal conductivity by forming the surface bond order imperfections to increase the surface-to-volume ratio [31–33], and enhance the phonon scattering. In all, the synergistic effect of TBR and TCR would significantly reduce the lattice thermal conductivity, thus enhance the figure-of-merit. The maximum zT of ∼2.1 is obtained in Cu2 Se bulk with porosity of 9% at 873 K, which is higher than most previous reported values of bulk Cu2 Se materials, indicating gasar process a feasible methodology to synthesize porous TE materials with low thermal conductivity and high figure-of-merit. Mechanical properties are essential for the practical application of TE materials, for they not only determine the cycle durability of TE modulus but also reflect the TE performance to some extent.
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Table 1 Vicker’s hardness of P + A + C samples compared with reported results and calculated fracture toughness (KIC ). No.
HV this work (GPa)
Reported HV (GPa) [24,34]
Calculated HV (GPa) [35,36]
Calculated KIC (MPa m1/2 ) [37]
0.43 0.41 0.38 0.29 0.21
0.38∗ 0.34 0.30∗ 0.26∗
0.380 0.336 0.295 0.258
2.25 2.12 1.95 1.81 1.55
Unfortunately, the existence of pores in matrix would inevitably deteriorate the mechanical properties due to the discontinuity of matrix materials. Herein we measure the hardness and fracture toughness of the tubular Cu2 Se samples indirectly based on following formulas [37]:
References
P P+A+C P+A+C P+A+C P+A+C ∗
0.03 0.06 0.09 0.12
Fitted.
H = P/α0 a2 K = P/β0 a3/2 where H is hardness, K is fracture toughness, α 0 and β 0 are numerical constants, P is the peak load, a and c are characteristic dimensions of the “plastic” impression and the radial/median crack, respectively. Compared with traditional porous materials, owing to the existence of corrugated inner surfaces, the gasar process synthesized porous Cu2 Se bulks in this work exhibit favourable mechanical properties. The tubular porous Cu2 Se samples show enhanced microhardness that is even comparable with those dense samples when the porosity is low. The measured thermal expansion coefficient and bending impedance plotted in Fig. S6-7 also demonstrate that the tubular porous structure reinforced with corrugated inner surface can keep a majority of mechanical properties of matrix sample, which is important for thermoelectric application. In this work, Cu2 Se bulks with tubular porous structures were synthesized through a gas reinforced eutectic transformation method. The obtained bulks contained tubular pores scoping from nanometres to micrometres hierarchically, inside which existed the corrugated surface. This unique structure was proposed to slightly modulate carrier transport by interfacial filtering effect, but drastically suppress the phonon progradation due to boundary resistance, surface state localization and thermal radiation. Thus a maximum figure-of-merit of ∼2.1 at 873 K was achieved. Moreover, the Cu2 Se bulks with tubular porous structure also possessed favourable mechanical properties, such as Vickers hardness, fracture toughness and bending impedance compared with traditional porous materials. Our work provided a promising strategy to fabricate porous TE materials with both high figure-of-merit and good mechanical properties Table 1. Acknowledgements This research was supported by funding from the Light of West China (No. XAB2018AW15) and the International Cooperative Project in Shaanxi Province (No. 2018kw-055). Supplementary materials Supplementary material associated with this article can be found, in the online version, at doi:10.1016/j.scriptamat.2019.09. 009.
[1] J. He, T.M. Tritt, Science 357 (6358) (2017). [2] G.J. Snyder, E.S. Toberer, Nat. Mater. 7 (2) (2008) 105. [3] A.F. Ioffe, Semiconductor Thermoelements and Thermoelectric Cooling, Infosearch Limited, London, 1957. [4] H.J. Goldsmid, Introduction to Thermoelectricity, Springer, Berlin, Germany, 2009. [5] D.M. Rowe, CRC Handbook of Thermoelectrics, CRC Press, Boca Raton, Florida, United States, 1995. [6] T. Zhu, Y. Liu, C. Fu, J.P. Heremans, J.G. Snyder, X. Zhao, Adv. Mater. 29 (14) (2017) 1605884. [7] H. Zhu, C. Xiao, Y. Xie, Adv. Mater. 0 (0) (2018) 18020 0 0. [8] I.I. Ravich, Semiconducting Lead Chalcogenides, Springer Science & Business Media, 2013. [9] J.P. Heremans, M.S. Dresselhaus, L.E. Bell, D.T. Morelli, Nat. Nanotechnol. 8 (2013) 471. [10] L. Hicks, M. Dresselhaus, Phys. Rev. B 47 (24) (1993) 16631. [11] Y. Pei, X. Shi, A. LaLonde, H. Wang, L. Chen, G.J. Snyder, Nature 473 (2011) 66. [12] S.I. Kim, K.H. Lee, H.A. Mun, H.S. Kim, S.W. Hwang, J.W. Roh, D.J. Yang, W.H. Shin, X.S. Li, Y.H. Lee, G.J. Snyder, S.W. Kim, Science 348 (6230) (2015) 109–114. [13] J.P. Heremans, Nat. Phys. 11 (12) (2015) 990. [14] M. Hong, Y. Wang, T. Feng, Q. Sun, S. Xu, S. Matsumura, S.T. Pantelides, J. Zou, Z.-.G. Chen, J. Am. Chem. Soc. 141 (4) (2019) 1742–1748. [15] R. Liu, H. Chen, K. Zhao, Y. Qin, B. Jiang, T. Zhang, G. Sha, X. Shi, C. Uher, W. Zhang, L. Chen, Adv. Mater. 29 (38) (2017) 1702712. [16] P. Giulia, Nat. Rev.Mater. 2 (2017) 17006. [17] X. Shi, A. Wu, W. Liu, R. Moshwan, Y. Wang, Z.-.G. Chen, J. Zou, ACS Nano (2018). [18] M. Hong, Y. Wang, S. Xu, X. Shi, L. Chen, J. Zou, Z.-.G. Chen, Nano Energy 60 (2019) 1–7. [19] B. Xu, T. Feng, M.T. Agne, L. Zhou, X. Ruan, G.J. Snyder, Y. Wu, Angew. Chem. Int. Edit. 56 (13) (2017) 3546–3551. [20] L. Drenchev, J. Sobczak, S. Malinov, W. Sha, Mat. Sci. Technol. 22 (10) (2006) 1135–1147. [21] V. Shapovalov, L. Boyko, Adv. Eng. Mater. 6 (6) (2004) 407–410. [22] J. Apprill, D.R. Poirier, M. Maguire, T. Gutsch, in: MRS Proceedings, 521, 1998. [23] H. Liu, X. Shi, F. Xu, L. Zhang, W. Zhang, L. Chen, Q. Li, C. Uher, T. Day, G.J. Snyder, Nat. Mater. 11 (2012) 422. [24] K. Zhao, H. Duan, N. Raghavendra, P. Qiu, Y. Zeng, W. Zhang, J. Yang, X. Shi, L. Chen, Adv. Mater. 29 (42) (2017) 1701148. [25] H. Liu, X. Yuan, P. Lu, X. Shi, F. Xu, Y. He, Y. Tang, S. Bai, W. Zhang, L. Chen, Y. Lin, L. Shi, H. Lin, X. Gao, X. Zhang, H. Chi, C. Uher, Adv. Mater. 25 (45) (2013) 6607–6612. [26] D. Vashaee, A. Shakouri, Phys. Rev. Lett. 92 (10) (2004) 106103. [27] H.-.S. Kim, Z.M. Gibbs, Y. Tang, H. Wang, G.J. Snyder, APL Mater. 3 (4) (2015) 041506. [28] E.T. Swartz, R.O. Pohl, Rev. Mod. Phys. 61 (3) (1989) 605–668. [29] Z. Tešanovic´ , M.V. Jaric´ , S. Maekawa, Phys. Rev. Lett. 57 (21) (1986) 2760–2763. [30] G. Xie, Z. Ju, K. Zhou, X. Wei, Z. Guo, Y. Cai, G. Zhang, NPJ Comput. Mater. 4 (1) (2018) 21. [31] C. Blanc, A. Rajabpour, S. Volz, T. Fournier, O. Bourgeois, Appl. Phys. Lett. 103 (4) (2013) 043109. [32] W. Gao, X. Yi, B. Sun, X. Meng, W. Cai, L. Zhao, Acta Mater. 132 (2017) 405–415. [33] P. Martin, Z. Aksamija, E. Pop, U. Ravaioli, Phys. Rev. Lett. 102 (12) (2009) 125503. [34] K. Tyagi, B. Gahtori, S. Bathula, M. Jayasimhadri, S. Sharma, N.K. Singh, D. Haranath, A.K. Srivastava, A. Dhar, Solid State Commun. 207 (2015) 21–25. [35] H. Nakajima, Mat. Sci. Appl. 9 (2018) 258–267. [36] J. Luo, R. Stevens, Ceram. Int. 25 (3) (1999) 281–286. [37] G. Anstis, P. Chantikul, B.R. Lawn, D. Marshall, J. Am. Ceram. Soc. 64 (9) (1981) 533–538.