Microstructure and thermoelectric properties of n-type Bi2Te2.85Se0.15 prepared by mechanical alloying and plasma activated sintering

Journal of Alloys and Compounds 420 (2006) 256–259

Microstructure and thermoelectric properties of n-type Bi2Te2.85Se0.15 prepared by mechanical alloying and plasma activated sintering X.A. Fan a , J.Y. Yang a,∗ , W. Zhu a , H.S. Yun b , R.G. Chen a , S.Q. Bao a , X.K. Duan a a

State Key Laboratory of Plastic Forming Simulation and Dies Technology, Huazhong University of Science and Technology, 1037 Luoyu Road, Wuhan 430074, PR China b Materials Research and Test Center, Wuhan University of Science and Technology, 122 Luoshi Road, Wuhan 430070, PR China Received 25 September 2005; received in revised form 9 October 2005; accepted 18 October 2005 Available online 23 November 2005

Abstract Starting from elemental bismuth, tellurium and selenium powders, n-type Bi2 Te2.85 Se0.15 solid solution with fine microstructure was prepared by mechanical alloying (MA) and plasma activated sintering (PAS) in the present work. The effect of PAS process on microstructure and thermoelectric properties of the sintered samples was investigated. The sintering temperature of PAS process (683 K) was 80–100 K lower than that of conventional hot pressing and the whole PAS process was also greatly shortened to about 30 min. A preferentially orientated microstructure with the basal planes (0 0 l) perpendicular to pressing direction was formed in the PASed sample and the maximum figure of merit (Z) at room temperature was 1.80 × 10−3 K−1 . © 2005 Elsevier B.V. All rights reserved. Keywords: Thermoelectric materials; Bi2 Te3 –Bi2 Se3 ; Grain orientation; Plasma activated sintering; Mechanical alloying

1. Introduction As one of the most excellent thermoelectric materials near room temperature [1], bismuth telluride based compounds have a remarkable anisotropy and the thermoelectric properties of these alloys along a-axis is superior to that in c-axis, which originates from the rhombohedral structure composed of quintuple atomic layer series in the order of Te(1) –Bi–Te(2) –Bi–Te(1) along the caxis [2]. Currently, unidirectional crystal growth methods such as zone melting or Bridgman technology [3,4] are widely used for preparing these alloys. Although, the resulting single crystal materials present excellent thermoelectric properties, they have poor mechanical property due to weak Van der Waals bonding between Te(1) –Te(1) layers and coarse grain size [5,6]. With the requirement of miniaturization and complication of thermoelectric module, isotropic mechanically strong materials with good thermoelectric properties are more preferable. Therefore, powder metallurgical methods, which produce randomly oriented polycrystalline and fine microstructure and thus good mechanical properties, such as hot pressing, hot extrusion and etc. [5–9],

∗

Corresponding author. Tel.: +86 27 87540944; fax: +86 27 87543776. E-mail address: [email protected] (J.Y. Yang).

0925-8388/$ – see front matter © 2005 Elsevier B.V. All rights reserved. doi:10.1016/j.jallcom.2005.10.025

are extensively studied. However, conventional powder metallurgical methods need to melt the green components to get an alloy ingot and then powderize the ingot and sinter it, some drawbacks, such as high energy cost, impurity contamination and complicated procedures should be overcome. In addition, the cross-sectional size limitation and materials waste at both ends of the products by hot extrusion, also offset the abovementioned advantages for hot extrusion to some extent [10–12]. Plasma activated sintering (PAS) is a novel consolidation technology, which enables rapid synthesis various materials by applying pulse current and an axial mechanical load. Plasma generates due to discharge between powders, the surface of powders is activated and purified, a self-heating phenomenon is achieved between the particles, heat- and mass-transfer can be completed instantaneously [13]. Therefore, it can obtain materials with very fine microstructure and high density at a relatively lower sintering temperature in a very short time [14–16]. The n-type Bi2 Te2.85 Se0.15 is studied widely in Bi2 Te3 related materials by conventional methods, which shows very good thermoelectric properties [9,17]. In our previous work [18], the figures of merit of Bi2 Te3 based materials prepared by MA and hot pressing (HP) are comparatively low compared with conventional single crystal materials, especially for n-type Bi2 Te3 -based materials. To improve the thermoelectric

X.A. Fan et al. / Journal of Alloys and Compounds 420 (2006) 256–259

properties of Bi2 Te3 based compounds, obtain materials with very fine microstructure and high density at a relatively lower sintering temperature in a shorter time, n-type Bi2 Te2.85 Se0.15 alloy was firstly pre-treated via mechanical alloying (MA) and subsequently consolidated by PAS method and the effect of processing parameters on microstructure and thermoelectric properties was reported in the present work. 2. Experiments Element Bi (99.9 wt.%, 300 mesh), Te (99.99 wt.%, 100 mesh) and Se (99.9 wt.%, 100 mesh) were subjected to MA according to the nominal composition of Bi2 Te2.85 Se0.15 in a QM-4F planetary ball mill under purified argon atmosphere. Stainless steel vessels and balls were used, the milling time was 6 h and the ball to powder weight ratio was 10:1. To minimize oxygen contamination, all powders weighing, loading and unloading were operated in a glove-box under purified argon. Subsequently, the as-MAed powder was consolidated by plasma activated sintering under a pressure of 30 MPa. The heating rate was 40 K/min and the sintering process was held at different temperatures (593, 623, 653, 683 and 713 K) for 15 min. The sintered compacts were column with dimensions φ21 mm × 10mm. Phase identification and crystallographic orientation were analyzed with XRD in a Philips X’Pert PRO diffractometer by using ˚ The fractographs were observed in a FEI Sirion Cu K␣ radiation (λ = 1.5406 A). 200 FE-SEM. The thermoelectric properties were measured at room temperature (300 K). A 10 K temperature difference was applied to between two ends of a 3 mm × 3 mm × 15 mm bar specimen to evaluate Seebeck coefficients (α). Electrical resistivity (ρ) was measured by a standard two-probe method. A TC-7000 Laser Flash apparatus was served for thermal conductivity (κ) measurement. The figure of merit (Z) was determined by the equation of Z = α2 /ρκ.

257

tion, indicating that the basal planes are preferentially orientated perpendicular to pressing direction. Fig. 2 shows SEM fractographs of the sample sintered at 683 K from the fractures parallel and perpendicular to pressing direction. It is in good consistent with the XRD result shown in Fig. 1. The parallel to pressing section presents a lightly orientated microstructure (Fig. 2a); in contrast, a very fine and equiaxed microstructure is observed in the perpendicular to pressing section (Fig. 2b). The preferential orientation formation is mainly attributed to the lamellar crystallographic structure and the Van der Walls bonding between Te(1) –Te(1) . The particles with basal planes bonded by Van der Waals force would rotate perpendicular to pressing direction. Furthermore, the pulse current induced a high electric field in samples and might promote the preferential orientation [15]. The orientation degree of the (0 0 l) planes can be determined by the orientation factor F, which can be calculated using Lotgering method [19]: F=

(P − P0 ) (1 − P0 )

(1)

3. Results and discussion Fig. 1 shows the XRD patterns of the as-PASed samples from sections parallel and perpendicular to pressing direction at 683 K. Both the as-milled powder and the as-PASed sample all show a single-phase structure. Furthermore, XRD peaks of the as-PASed n-type Bi2 Te2.85 Se0.15 solid solution become sharper and stronger after plasma activated sintering, indicating the improvement of the crystallinity of the sintered samples. The relative intensities of (0 0 l) planes including (0 0 6), (0 0 15) and (0 0 18) from section perpendicular to pressing are lightly stronger than those from the section parallel to pressing direc-

Fig. 1. XRD patterns of the as-PASed samples from sections parallel and perpendicular to pressing direction at sintering temperature 683 K.

Fig. 2. SEM fractographs of the as-PASed samples sintered at 683 K (a) parallel and (b) perpendicular to pressing direction.

258

X.A. Fan et al. / Journal of Alloys and Compounds 420 (2006) 256–259

Fig. 3. Sintering temperature dependence of orientation factors (F) of the (0 0 l) planes of the as-PASed samples.

I(0 0 l) P=
I(h k l)

(2)

I0 (0 0 l) P0 =
I0 (h k l)

(3)

Where P and P0 are the ratios of the integrated intensities of all (0 0 l) planes to those of all (h k l) planes for the preferentially orientated and randomly orientated samples, respectively. The orientation factors (F) of the (0 0 l) planes of the samples sintered at different temperature are shown in Fig. 3. The orientation factors (F) decreases lightly from 0.13 to 0.10 when increasing sintering temperature, indicating the influence of sintering temperature on the orientation factors is not obvious. Fig. 4 shows the SEM fractographs of the samples sintered at 623 and 713 K from sections perpendicular to pressing direction. It is clear that the grains experience less growth even increasing the PAS temperature from 623 to 713 K. Seebeck coefficient (α) and electrical resistivity (ρ) of the samples sintered at different temperatures are shown in Fig. 5. It can be seen that the electrical resistivity (ρ) and Seebeck coefficient (α) decrease with increasing PAS temperature. As we know, electrical resistivity can be expressed as ρ = 1/(nc eµ) [20], Seebeck coefficient (α) of thermoelectric materials in the extrinsic conduction region can be expressed as [15]:   kB 2(2πm∗ kB T )3/2 α= (4) δ + 2 + ln e n c h3 Where nc is the carrier concentration; µ the carrier mobility; kB the Boltzmann constant; δ the scattering parameter; m* the effective mass and h is the Planck constant. With increasing of the sintering temperature, the concentration of lattice defects generated during MA process decreases and the density of the as-PASed samples also increases. So the carrier mobility (µ) increases or carrier effective mass decreases, therefore the electrical resistivity and the magnitude of Seebeck coefficient (α) decrease with increasing sintering temperature. Fig. 6 displays the dependence of thermal conductivity (κ) of the sintered samples on the PAS temperature. It increases with increasing sintering temperature. As we know, thermal conduc-

Fig. 4. SEM fractographs of the as-PASed samples at sintering temperatures (a) 623 K and (b) 713 K from sections perpendicular to pressing direction.

tivity (κ) of semiconductor can be expressed as κ = κel + κph [7], where κel and κph correspond to carrier and phonon contribution to thermal conductivity, κel is related to ρ according to Wiedemann–Franz law: κel = LT/ρ, where L is Lorentz constant (L = 2.48 × 10−8 W  K−2 ). As shown in Fig. 5, ρ decreases with increasing PAS temperature, therefore κel increases with the PAS temperature in Fig. 6. However, the influence of sintering temperature on κph is not obvious (Fig. 6). The sintering temperature of PAS process is 80–100 K lower than that of conventional methods such as hot pressing and the sintering time

X.A. Fan et al. / Journal of Alloys and Compounds 420 (2006) 256–259

259

increasing PAS temperature from 593 to 683 K and it decreases when further increasing PAS temperature over 683 K. The sample sintered at 683 K shows the maximum figure of merit (Z = 1.80 × 10−3 K−1 ) at room temperature (300 K). The result is encouraging, although it is still lower than the reported best value for single crystal n-type bismuth telluride based materials. Further work on microstructure control, doping and PAS process optimization is under way in our lab. 4. Conclusions

Fig. 5. Sintering temperature dependence of Seebeck coefficient (α) and electrical resistivity (ρ) of the as-PASed samples.

is only several percents of that of conventional methods. Therefore, the grains experience less growth even increasing the PAS temperature from 623 to 713 K, as shown in Fig. 4, thus phonongrain boundary scattering and the resultant κph do not change obviously when increasing sintering temperature (Fig. 6). The figures of merit (Z) of the samples at different sintering temperatures are shown in Fig. 7. It increases with

The n-type Bi2 Te2.85 Se0.15 compounds were obtained by MA and subsequently PAS in this work. A fine and preferentially orientated microstructure that the basal planes (0 0 l) perpendicular to pressing direction was formed in the PASed samples. The maximum orientation factor of the (0 0 l) planes was obtained as 0.13 when sintered at 593 K and the influence of sintering temperature on the orientation factors (F) was not obvious. The maximum figures of merit (Z) at room temperature (300 K) was 1.80 × 10−3 K−1 after plasma activated sintering. Acknowledgements This work is co-financed by the Nation Basic Research Project (2004CCA03200) and Natural Science Foundation of China (50401008). References

Fig. 6. Sintering temperature dependence of thermal conductivity (κ) of the as-PASed samples.

Fig. 7. Sintering temperature dependence of figures of merit (Z) of the as-PASed samples.

[1] D.Y. Chung, T. Hogan, P. Brazis, M.R. Lane, C. Kannewurf, M. Bastea, C. Uher, M.G. Kanatzidis, Science 287 (2000) 1024–1027. [2] H. Kitagawa, T. Nagamori, T. Tatsuta, T. Kitamura, Y. Shinohara, Y. Noda, Scripta Mater. 49 (2003) 309–313. [3] D.B. Hyun, T.S. Oh, J.S. Hwang, J.D. Shim, N.V. Kolomoets, Scripta Mater. 40 (1999) 49–56. [4] O. Yamashita, S. Tomiyoshi, K. Makita, J. Appl. Phys. 93 (2003) 368–374. [5] J. Yang, T. Aizawa, A. Yamamoto, T. Ohta, J. Alloys Compd. 309 (2000) 225–228. [6] S.J. Hong, B.S. Chun, Mater. Sci. Eng. A 356 (2003) 345–351. [7] J. Yang, T. Aizawa, A. Yamamoto, T. Ohta, Mater. Chem. Phys. 70 (2001) 90–94. [8] S. Miura, Y. Satob, K. Fukuda, K. Nishimura, K. Ikeda, Mater. Sci. Eng. A 277 (2000) 244–249. [9] J.Y. Yang, T. Aizawa, A. Yamamoto, T. Ohta, J. Alloys Compd. 312 (2000) 326–330. [10] J.T. Im, K.T. Hartwig, J. Sharp, Acta Mater. 52 (2004) 49–55. [11] J. Seo, K. Park, D. Lee, C. Lee, Mater. Sci. Eng. B 49 (1997) 247–250. [12] T.S. Kim, I.S. Kim, T.K. Kim, S.J. Hong, Mater. Sci. Eng. B 90 (2002) 42–46. [13] K. Yamazaki, S.H. Risbud, H. Aoyama, K. Shoda, J. Mater. Process. Technol. 56 (1996) 955–965. [14] J. Jiang, L.D. Chen, S.Q. Bai, Q. Yao, Q. Wang, Scripta Mater. 52 (2005) 347–351. [15] N. Keawprak, Z.M. Sun, H. Hashimoto, M.W. Barsoum, J. Alloys Compd. 397 (2005) 236–244. [16] L.D. Chen, J. Jiang, X. Shi, Mater. Res. Soc. 793 (2004) 931–939. [17] J. Seo, K. Park, C. Lee, Mater. Res. Bull. 33 (1998) 553–559. [18] J. Yang, R. Chen, X. Fan, S. Bao, W. Zhu, J. Alloys Compd. (in press) (corrected proof available online 18 August 2005). [19] F.K. Lotgering, J. Inorg. Nucl. Chem. 9 (1959) 113. [20] J. Seo, K. Park, D. Lee, C. Lee, Scripta Mater. 38 (1998) 477–484.