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Enhancing the thermoelectric properties of Ca3Co4O9 thin films by Nb ion injection
Journal of Alloys and Compounds 567 (2013) 122–126
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Enhancing the thermoelectric properties of Ca3Co4O9 thin films by Nb ion injection Chunhui Zhu, Heping An, Wenwen Ge, Zhuangzhi Li ⇑, Guide Tang Hebei Advanced Thin Films Laboratory, Department of Physics, Hebei Normal University, Shijiazhuang City 050024, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 28 October 2012 Received in revised form 7 March 2013 Accepted 8 March 2013 Available online 20 March 2013 Keywords: Ion injection Thin films Thermoelectric property Ca3Co4O9
a b s t r a c t High quality Ca3Co4O9 thin films have been grown epitaxially on single crystal Al2O3 substrates with pulsed laser deposition. Nb was implanted into the Ca3Co4O9 films using an ion beam injection technique. The microstructure of the thin films has been investigated by XRD, SEM and AFM. The epitaxial thin films were grown with the c-axis perpendicular to the substrate surface. The effect of Nb doping by ion beam injection was verified using resistivity measurements at room temperature. Resistivity and the Seebeck coefficient were also measured in the temperature range 150–380 K. The results indicate that the power factors of Ca3Co4O9 thin films increase when doped with Nb. When the concentration of doped Nb was 3.65 1019 atoms/cm3, the power factor of the thin films reached 0.10 mW/m K2 at room temperature, and it approached a maximum of 0.17 mW/m K2 at 380 K. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Environmentally-friendly thermoelectric (TE) materials have attracted much attention in recent years [1]. They can be used in electricity generation from waste heat and in electric refrigeration without any environment pollution [2] due to the unique property of converting heat directly into electricity. The conversion efficiency of TE materials is represented by a dimensionless figure of merit, ZT = S2rT/j, where S, r, T and j are the Seebeck coefficient, electrical conductivity, absolute temperature, and thermal conductivity, respectively. For practical applications, the ZT value must be greater than 1. Shikano et al. [3] have reported that the ZT value of Ca3Co4O9 (CCO) exceeds 0.8 at 1000 K. This value is very near the level required for practical application. Moreover, CCO exhibits the advantages of being antioxidant, stable and non-toxic at high temperature. Consequently, it is considered as a potential thermoelectric material for high temperature applications. The Ca3Co4O9 crystal consists of CdI2-type hexagonal CoO2 layers and rock salt type trigonometric Ca2CoO3 layers alternately stacked along the c-axis. The two components have common lattice parameters a = 4.8376 Å and c = 10.833 Å, but different b lattice parameters. For the Ca2CoO3 layers b1 = 4.5565 Å while for the CoO2 layers b2 = 2.8189 Å [4]. This misfit in structure induces a strong anisotropy. Kenfaui et al. [5] have reported that the ZT value of polycrystalline bulk CCO prepared by hot pressing has a noticeably anisotropic ratio ZTab/ZTc = 4.6. However, the direction of growth cannot be strictly controlled in polycrystalline bulk form. ⇑ Corresponding author. Tel.: +86 0311 80787330. E-mail address: [email protected] (Z. Li). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.069
On the other hand, CCO thin films grown on an Al2O3 substrate by pulsed laser deposition (PLD) will grow with the c-axis perpendicular to the surface of the substrate due to the lattice mismatch of the CCO and the Al2O3 substrate [6]. High quality CCO epitaxial thin films with a c-axis orientation yielding highly anisotropic physical properties have been reported [7–10]. Doping is one of the most important methods for improving the performance of semiconductors through modifying the carrier concentration. As a p-type TE material, the carrier concentration of CCO can also be changed by doping with metallic elements. In particular, Wang et al. have reported that Ag, Fe, Mn and Cu doping can improve the thermoelectric properties of CCO [2,11]. Considering that the coordination number of all Co atoms in CCO is six and that Nb and Co cations have similar radii, Nb ions may substitute for Co ions in CCO. In addition, since Nb has a larger atomic mass than does Co, phonon scattering is expected to increase while the thermal conductivity decrease, with the result that the thermoelectric properties may be further improved. In this work, the influences of Nb-doping on the thermoelectric properties of CCO epitaxial thin films are investigated. In traditional approaches, normally there are two steps for preparing CCO thin films doped with metallic elements. The doped precursor compound is first prepared using solid-state reactions, the sol–gel method etc., and secondly, thin films are fabricated using various techniques such as chemical deposition, spin-coating, and PLD. These approaches, however, require significant time to carry out, and it is difficult to achieve high doping concentrations because the thermodynamics tends to be unfavorable. The ion beam injection technique has several advantages over the traditional approaches: (1) The distribution of the dopant ion in the
C. Zhu et al. / Journal of Alloys and Compounds 567 (2013) 122–126
123
probe method at room temperature before injection, after injection and after annealing. The resistivity and Seebeck coefficient of the CCO thin films were measured using the standard four-probe method and the steady state method with the Thermal Transport Option (TTO) module of a Physical Properties Measurement System (PPMS) over the temperature range 150–380 K. The main experimental instruments used in this work are listed below. The pulsed laser deposition and ion beam cleaning system (PLD-450a) was made by the Scientific Instruments Research Center of the Chinese Academy of Sciences in Shenyang. Excimer lasers (Compex Pro 205, KrF k = 248 nm) were produced by Coherent in the USA. The X-ray diffractometer (XRD) (X’Pert Pro MPD, Cu Ka) was made by PANalytical B.V. in the Netherlands. The S-4800 Field Emission Scanning Electron Microscope (SEM) was produced by Hitachi in Japan. The Scanning Probe Microscope (SPM) (NanoScope IV) was produced by Digital Instruments in the USA. The MEVVA ion source was produced by the Tongchuang Applied Plasma Technology Center in Chengdu, China. The physical properties measurement system (PPMS-9T) was made by Quantum Design in the USA.
3. Results and discussion Fig. 1. XRD patterns for the CCO thin films. The inset shows an expanded view for the angular range 16.3–17°.
Fig. 2. Cross section microstructure of CCO thin film on (0 0 0 1) Al2O3 substrate with SEM.
sample can be controlled by the injection energy; (2) The doping concentration can be controlled through the injection doses, and high concentrations are easy to achieve; (3) Less time is required. In this paper, the ion beam injection technique has been used for doping CCO thin films with high concentrations of Nb ions, and the thermoelectric properties of CCO thin films with different injection doses have been measured.
Fig. 1 shows the XRD pattern of CCO thin films on (0 0 0 l) single crystal Al2O3 substrates after annealing. The X’Pert HighScore Plus software was used to search for peaks and carry out data processing. The (0 0 2), (0 0 3), (0 0 4), (0 0 5), (0 0 6) and (0 0 8) crystal plane diffraction peaks appear at 16.63°, 25.10°, 33.41°, 41.99°, 51.21° and 69.79°, respectively. This indicates perfect c-axis oriented growth of CCO thin films due to the fact that diffraction peaks from (00l) planes are observed in the patterns and no other plane peaks are found. Two points should be noted. First, the positions of the diffraction peaks show a clear drift as the Nb injection dose rises as shown in the inset of Fig. 1. This means that the Nb ions enter into the crystal lattice and demonstrates that Nb doping has been achieved by ion beam injection. In addition, no peak for cobalt oxide is found in the patterns possibly due to the cobalt oxide appearing in the films as non-crystalline nanoclusters. The SEM image in Fig. 2 shows the cross section microstructure of a CCO thin film grown on a (0 0 0 l) single crystal Al2O3 substrate. The dense CCO thin film and the growth interface can be clearly seen. The AFM (5 5 lm2) image in Fig. 3 shows the surface morphology of the CCO thin film. The image indicates that the film is smooth and the grain size is uniform. Both pictures show that the film has good crystallinity and has grown with its c-axis perpendicular to the Al2O3 substrate surface. In our experience, there are large differences in the crystalline quality of CCO thin films prepared by PLD. Therefore, we prepared
2. Experimental details The CCO precursor powder was prepared using the sol–gel method and then pressed into high density solid targets to be used for PLD. The CCO thin films were grown on single crystal Al2O3 substrates by PLD with a substrate temperature of about 1023 K, a laser energy density of 1.5 J/cm2, a deposition frequency of 3 Hz and under an oxygen pressure of 20 Pa. The commercial Al2O3 substrates (10 3 mm2) were cleaned with pure alcohol, acetone and deionized water for 5 min each in an ultrasonic bath and blown dry with N2. The phase composition and structure of the CCO thin films were characterized by X-ray Diffraction (XRD). The cross section microstructure and surface morphology of the films were observed using a Scanning Electron Microscope (SEM) and an Atomic Force Microscope (AFM). Nb ion injection into the CCO thin films was implemented using a Metal Vapor Vacuum Arc (MEVVA) ion source. The implantation depth and distribution were simulated with the SRIM software [12]. The injection voltage was 34 kV and the injection dip angle was 7°. In this work, the average Nb implantation concentrations in CCO thin films were 3.65 1019 atoms/cm3 (referred to below as CCONb4), 1.46 1020 atoms/cm3 (CCONb16), 7.33 1020 atoms/cm3 (CCONb80) and 3.65 1021 atoms/cm3 (CCONb400). The samples were annealed at 923 K under 1 atm oxygen pressure for 30 min after injection. The phase composition and surface morphology were observed using XRD, SEM and AFM. In order to quantify the effects of doping, the resistivity of the samples was measured using the four-
Fig. 3. Surface morphology of CCO thin film with AFM.
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Table 1 Resistivity of CCO thin films with different Nb injection doses at room temperature. Rb, Ri and Ra are the resistivities for the films before injection, after injection and after annealing, respectively. Sample
Rb (mX cm)
Ri (mX cm)
Ra (mX cm)
(Rb–Ra)/Rb (%)
CCONb4 CCONb16 CCONb80 CCONb400
10.51 10.53 10.67 10.64
34.79 44.09 66.47 643.53
7.78 9.59 12.43 21.34
25.97 8.92 16.49 100.56
Fig. 4. Temperature dependence of resistivity for different Nb injection doses.
Table 2 Carrier concentration n of CCO thin films with different Nb injection doses at 300 K. Sample
CCO
3
n (cm
)
3.5 10
CCONb4 21
22
1.2 10
CCONb16 8.5 10
21
CCONb80 21
5.5 10
CCONb400 1.4 1021
a number of CCO thin films, and the resistivity of the films was measured at room temperature by the four-probe method to pick out five similar samples. The resistivity of these five samples was around 10.6 mX cm. After Nb was injected into these samples using the MEVVA ion source, the resistivity was measured again. The resistivity was measured a third time after annealing to show the full effect of Nb doping. As shown in Table 1, the resistivity after annealing of the CCONb4, CCONb16, CCONb80 and CCONb400 samples was changed by 25.97%, 8.92%, 16.49% and 100.56% with respect to the as-prepared samples. Conductivity can be written as
r ¼ nel
Fig. 4 shows the resistivity characteristics of CCO thin films with different doses of injected Nb in the temperature range 150–380 K. As can be seen, the resistivity of the samples first decreases and then increases as the Nb injection dose rises. At 300 K, CCONb4 had the smallest resistivity, 7.67 mX cm, of all the samples. For CCONb0, CCONb16, CCONb80 and CCONb400, the resistivities were 10.49 mX cm, 9.49 mX cm, 12.31 mX cm and 21.13 mX cm, respectively. There are small differences between these values and the values obtained using the four-probe method as shown in Table 1 due to experimental error. The Ca3Co4O9 molecular structure can be denoted as (Ca2CoO3)(CoO2)1.62 [6,13]. The CoO2 layer which has an octahedral structure is a conductive layer while the Ca2CoO3 layer which has a rock salt type trigonometric structure is a carrier reservoir layer. For Ca cations, the coordination number is 8, the valence state is +2, and the ionic radius is 1.12 Å. In CCO thin films Nb will appear as Nb5+. When the coordination number is 6, the ionic radius of Nb5+ is 0.64 Å, and when the coordination number is 8, the ionic radius is 0.74 Å. The substitution of Nb for Ca is therefore be improbable due to the large ionic radius difference between Nb5+ and Ca2+ with the same coordination number. For all Co cations in CCO, the coordination number is 6. X-ray absorption (XAS) and photoemission studies of CCO show there is no Co2+ in the Ca2CoO3 crystal structure and indicate that the hole doped Co–O triangular lattice has low spin Co4+ species in a low spin Co3+ background [14,15]. Ultraviolet photoemission spectroscopy (USP) results suggest that phase separation plays an important role in CCO [14]. Since the radius of Nb5+ (0.64 Å) is close to that of Co3+/Co4+ (0.545 Å/0.53 Å, low spin states), Nb may randomly substitute for Co cations. The spin-state transition plays a leading role in the hopping transport process for CCO, so hopping transportation should become more difficult due to the differences in the electronic structures of Nb and Co cations, which will weaken the superexchange interactions. On the other hand, it is obvious that with the substitution of Nb for Co, high-valence cations replace low-valence cations and donate electrons so that the hole concentration should decrease. Both processes are expected to lead to a resistivity increase with an increase in the Nb concentration. However, experimentally, the opposite results have been observed. It is found that the resistivity decreases to a minimum for CCONb4 and then increases, eventually reaching a maximum at the highest Nb concentration, as shown in Fig. 4. In addition, Hall measurements indicate that the hole concentration first increases and then decreases with Nb doping concentration as listed in Table 2. One possible reason is that with Nb doping, the oxygen deficit [16] disappears and the oxygen content increases in CCO films and the in-
ð1Þ
where n is the carrier concentration, e is the electron charge and l is the carrier mobility. It is clear that the changes in the conductivity may be due to changes in the carrier concentration, the mobility or both. The Seebeck coefficient can be expressed as
SðTÞ ¼
2 c p2 kB T @ ln lðeÞ þ n 3e @e e¼EF
ð2Þ
where c, n, kB, and l(e) are the specific heat, carrier concentration, Boltzmann constant and carrier mobility, respectively [2]. Eq. (2) shows that S is determined by n and l(e) for a specific temperature. It is therefore expected that the Seebeck coefficient S will also be modified by Nb doping as it depends on the same two parameters as does the conductivity. Consequently, it is possible that the value of ZT for CCO thin films could be improved by Nb doping.
Fig. 5. LnrT vs. 1000/T for CCO thin films with different Nb injection doses.
C. Zhu et al. / Journal of Alloys and Compounds 567 (2013) 122–126
Fig. 6. Temperature dependence of Seebeck coefficient for CCO thin films with different Nb injection doses.
Fig. 7. Temperature dependence of power factor for CCO thin films with different Nb injection doses.
creased number of holes provided by the increased oxygen content is larger than the number of electrons injected. However, the number of oxygen deficiencies is limited, so the hole concentration eventually stops increasing and begin to decrease as the Nb injection dose rises. Mikami et al. [17] have reported a similar result in their study of Bi doping of single crystal CCO. It is found that Bi3+ replaces Ca2+ and can effectively decrease the oxygen deficit. Consequently the oxygen content increases in the CCO structure. In addition, it is suggested that the replaced Co atoms could react with oxygen and produce CoO non-crystalline nanoclusters. The CoO nanoclusters formed are expected to enhance the scattering of holes in CCO films, and thereby lead to a further increase in the resistivity. Hopping of a hole between Co3+ and Co4+ is the elementary charge transport mechanism in CCO. According to small polaron theory [18,19], the electrical conductivity of small polarons can be described by the expression r = C/T exp(Ea/kBT), where C is a pre-exponential term, Ea is the activation energy for hopping, kB is the Boltzmann constant, and T is the absolute temperature. As shown in Fig. 5, the experimental data show a good linear relationship between lnrT and 1000/T, which indicates a small polaron hopping conduction mechanism in the CCO films. The fitted values of the activation energies for CCONb0, CCONb4, CCONb16, CCONb80 and CCONb400 are 31 meV, 40 meV, 45 meV, 46 meV and 40 meV, respectively. The
125
result indicates that the activation energy of the carriers increases slightly when Nb is injected. This may be due to the fact that Nb substitution for Co cations makes hole transportation more difficult. Fig. 6 shows the temperature dependence of the Seebeck coefficient for CCO thin films with different Nb injection doses in the temperature range 150–380 K. Compared with the pure CCO sample, the Seebeck coefficient shows little change for CCONb4, CCONb16, and CCONb80. At room temperature, CCO and CCONb4 have a similar Seebeck coefficient of about 90 lV/K. CCONb16 and CCONb80 have a Seebeck coefficient of about 96 lV/K which is slightly higher than that for CCO. The Seebeck coefficient for CCONb400 is only about 72 lV/K which is obviously lower than that for CCO. Koshibae et al. [20] have studied the thermopower in strongly correlated electron systems with orbital degeneracy and suggested that the triangular lattice of CCO is advantageous to the thermopower that is induced by the degeneracy. Nb substitution for Co may increase the difficulty of the superexchange interactions. Substitution of small amount of Nb has a weak effect on the thermopower as may be concluded from the observation that CCO, CCO4, CCONb16 and CCONb80 have similar thermopower values. Replacement of Co by Nb to a greater extent may well reduce the superexchange between Co3+ and Co4+ and hence decrease the Seebeck coefficient as may be seen in the case of CCONb400. The temperature dependence of the power factor, S2r, for CCO thin films with different Nb injection doses in the temperature range 150–380 K is shown in Fig. 7. At room temperature, CCONb4 shows the highest power factor of about 0.10 mW/m K2 due to the facts that Seebeck coefficient changes little and the resistivity decreases dramatically compared to the other samples. The power factor of CCONb16 also improved over that of CCO because it has a slightly higher Seebeck coefficient and lower resistivity than does CCO. The reduced power factor of CCONb80 is due to the fact that the resistivity increases more than does the Seebeck coefficient. For CCONb400, the decrease of the Seebeck coefficient and the increase in the resistivity both lead to a dramatic reduction in the power factor. Comparing the power factors of pure CCO and the Nb injected CCO thin films, CCONb4 has the highest power factor at room temperature. The maximum power factor observed in the present work was 0.17 mW/m K2 for CCONb4 at 380 K. Thermal conductivity depends on both the conductivity of the lattice and of the charge carries, and may be expressed as j = jL + je, where jL is the lattice thermal conductivity and je is the carrier thermal conductivity. je can be calculated from the Widemann– Franz law as je = LrT = LT/q, where L is the Lorenz constant, L = 2.45 108 V2/K2. Consequently, je shows the same trend as does r. Therefore je is increased for CCONb4 and CCONb16, with respect to CCO, but is still only about 0.1 W/m K, as calculated from the resistivity. Shikano et al. [3] reported that the thermal conductivity in the a–b plane of a single crystal of CCO is 3 W/m K, which is much larger than 0.1 W/m K. Therefore, it is probable that the thermal conductivity of CCO thin films comes mainly from lattice vibrations. Lattice thermal conductivity can be described by the equation jL = (1/3)CDvLlL, where CD, vL and lL are the lattice specific heat, phonon group velocity and mean free path of phonons, respectively. Because the mean free path of the phonon is expected to be significantly decreased due to increased scattering [21] by CoO nanoclusters, the lattice thermal conductivity is also decreased, and the thermoelectric figure of merit, ZT, can be improved, if CD and vL are assumed to remain unchanged.
4. Conclusions An ion beam injection technique has been applied to dope CCO thin films with the metallic element, Nb, and the experimen-
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tal results demonstrate that the method is effective in increasing the power factor of CCO films. Different doses of Nb ions were injected into CCO epitaxial thin films grown on single crystal Al2O3 substrates using a MEVVA ion source, and the samples were annealed in an oxygen atmosphere for 30 min in order to remove implantation damage and activate implanted ions. In this way, a series of samples with different Nb doping concentrations was obtained. The resistivity and Seebeck coefficients of the samples were measured using a PPMS in the temperature range 150– 380 K. Measurement results indicate that the resistivity decreases for small doses of injected Nb, but the resistivity increases for large doses. At 300 K, CCONb4 has a resistivity of 7.67 mX cm which was the smallest observed across all the samples studied. The Seebeck coefficient changes in an irregular manner for CCO, CCONb4, CCONb16 and CCONb80, but decreases dramatically for CCONb400. This demonstrates that small doses of injected Nb do not significantly modify the Seebeck coefficient, while larger doses of Nb can cause a large decrease. At room temperature, CCONb4 had the highest power factor, 0.10 mW/m K2, while the maximum power factor at any temperature studied was 0.17 mW/m K2 at 380 K.
Acknowledgments This work is supported by the Key item Science Foundation of Hebei Province (Grant No. 10965125D), the National Natural Science Foundation of China, under the Contract No. NSF-11174069.
References [1] M. Mikami, N. Ando, R. Funahashi, J. Solid State Chem. 178 (2005) 2186–2190. [2] Y. Wang, Y. Sui, J.J. Cheng, X.J. Wang, W.H. Su, J. Alloy Comp. 477 (2009) 817– 821. [3] M. Shikano, R. Funahashi, Appl. Phys. Lett. 82 (2003) 1851. [4] A.C. Masset, C. Michel, A. Maignan, M. Hervieu, O. Toulemonde, F. Studer, B. Raveau, J. Hejtmanek, Phys. Rev. B 62 (2000) 166. [5] D. Kenfaui, B. Lenoir, D. Chateigner, B. Ouladdiaf, M. Gomina, J.G. Noudem, J. Eur. Ceram. Soc. 32 (10) (2012) 2405–2414. [6] H.W. Eng, W. Prellier, S. Hebert, D. Grebille, J. Appl. Phys. 97 (2005) 013706. [7] T. Sun, H.H. Hng, Q.Y. Yan, J. Ma, J. Alloy Comp. 512 (2012) 133–138. [8] T. Sun, H.H. Hng, Q.Y. Yan, J. Ma, J. Electron. Mater. 39 (9) (2010) 1611–1615. [9] Q. Qiao, A. Gulee, T. Paulauskas, S. Kolesnik, B. Dabrowski, M Ozedemir, C. Boyraz, D. Mazumdar, A. Gupta, R.F. Klie, J. Phys.:Condens. Matter. 23 (2011) 305005. [10] R. Moubah, S. Colis, C. Ulhaq-Bouillet, G. Schmerber, N. Viart, M. Drillon, A. Dinia, D. Muller, J.J. Grob, Eur. Phys. J. B 66 (2008) 315–319. [11] Y. Wang, Y. Sui, P. Ren, L. Wang, X.J. Wang, W.H. Su, H.J. Fan, Chem. Mater. 22 (2010) 1155–1163. [12] J.F. Ziegler, Software for the stopping and range of ions in matter, homepage:. [13] G. Yang, Q. Ramasse, R.F. Klie, Phys. Rev. B 78 (2008) 153109. [14] Y. Wakisawa, S. Hirata, T. Mizokawa, Y. Suzuki, Y. Miyazaki, T. Kajitani, Phys. Rev. B 78 (2008) 235107. [15] T. Mizokawa, L.H. Tjeng, H.J. Lin, C.T. Chen, R. Kitawaki, I. Terasaki, S. Lambert, C. Michel, Phys. Rev. B 71 (2005) 193107. [16] J. Shioyama, S. Horii, K. Otzshi, M. Sano, K. Kishino, Jpn. J. Appl. Phys. 42 (2003) L194. [17] M. Mikami, K.J. Chong, Y. Miyazaki, T. Kajitani, T. Inoue, S. Sodeoka, R. Funahashi, Jpn. J. Appl. Phys. 45 (5A) (2006) 4131–4136. [18] A.J. Bosman, H.J. Van Daal, Adv. Phys. 19 (1970) 1–117. [19] T. Holstein, Ann. Phys-New York 281 (2000) 725–773. [20] W. Koshibae, S. Maekawa, Phys. Rev. Lett. 87 (2001) 23. [21] G.S. Nolas, J. Sharp, H.J. Goldsmid, Thermoelectrics basic principles and new materials developments, Springer, New York, 2001.
Contents lists available at SciVerse ScienceDirect
Journal of Alloys and Compounds journal homepage: www.elsevier.com/locate/jalcom
Enhancing the thermoelectric properties of Ca3Co4O9 thin films by Nb ion injection Chunhui Zhu, Heping An, Wenwen Ge, Zhuangzhi Li ⇑, Guide Tang Hebei Advanced Thin Films Laboratory, Department of Physics, Hebei Normal University, Shijiazhuang City 050024, People’s Republic of China
a r t i c l e
i n f o
Article history: Received 28 October 2012 Received in revised form 7 March 2013 Accepted 8 March 2013 Available online 20 March 2013 Keywords: Ion injection Thin films Thermoelectric property Ca3Co4O9
a b s t r a c t High quality Ca3Co4O9 thin films have been grown epitaxially on single crystal Al2O3 substrates with pulsed laser deposition. Nb was implanted into the Ca3Co4O9 films using an ion beam injection technique. The microstructure of the thin films has been investigated by XRD, SEM and AFM. The epitaxial thin films were grown with the c-axis perpendicular to the substrate surface. The effect of Nb doping by ion beam injection was verified using resistivity measurements at room temperature. Resistivity and the Seebeck coefficient were also measured in the temperature range 150–380 K. The results indicate that the power factors of Ca3Co4O9 thin films increase when doped with Nb. When the concentration of doped Nb was 3.65 1019 atoms/cm3, the power factor of the thin films reached 0.10 mW/m K2 at room temperature, and it approached a maximum of 0.17 mW/m K2 at 380 K. Ó 2013 Elsevier B.V. All rights reserved.
1. Introduction Environmentally-friendly thermoelectric (TE) materials have attracted much attention in recent years [1]. They can be used in electricity generation from waste heat and in electric refrigeration without any environment pollution [2] due to the unique property of converting heat directly into electricity. The conversion efficiency of TE materials is represented by a dimensionless figure of merit, ZT = S2rT/j, where S, r, T and j are the Seebeck coefficient, electrical conductivity, absolute temperature, and thermal conductivity, respectively. For practical applications, the ZT value must be greater than 1. Shikano et al. [3] have reported that the ZT value of Ca3Co4O9 (CCO) exceeds 0.8 at 1000 K. This value is very near the level required for practical application. Moreover, CCO exhibits the advantages of being antioxidant, stable and non-toxic at high temperature. Consequently, it is considered as a potential thermoelectric material for high temperature applications. The Ca3Co4O9 crystal consists of CdI2-type hexagonal CoO2 layers and rock salt type trigonometric Ca2CoO3 layers alternately stacked along the c-axis. The two components have common lattice parameters a = 4.8376 Å and c = 10.833 Å, but different b lattice parameters. For the Ca2CoO3 layers b1 = 4.5565 Å while for the CoO2 layers b2 = 2.8189 Å [4]. This misfit in structure induces a strong anisotropy. Kenfaui et al. [5] have reported that the ZT value of polycrystalline bulk CCO prepared by hot pressing has a noticeably anisotropic ratio ZTab/ZTc = 4.6. However, the direction of growth cannot be strictly controlled in polycrystalline bulk form. ⇑ Corresponding author. Tel.: +86 0311 80787330. E-mail address: [email protected] (Z. Li). 0925-8388/$ - see front matter Ó 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.jallcom.2013.03.069
On the other hand, CCO thin films grown on an Al2O3 substrate by pulsed laser deposition (PLD) will grow with the c-axis perpendicular to the surface of the substrate due to the lattice mismatch of the CCO and the Al2O3 substrate [6]. High quality CCO epitaxial thin films with a c-axis orientation yielding highly anisotropic physical properties have been reported [7–10]. Doping is one of the most important methods for improving the performance of semiconductors through modifying the carrier concentration. As a p-type TE material, the carrier concentration of CCO can also be changed by doping with metallic elements. In particular, Wang et al. have reported that Ag, Fe, Mn and Cu doping can improve the thermoelectric properties of CCO [2,11]. Considering that the coordination number of all Co atoms in CCO is six and that Nb and Co cations have similar radii, Nb ions may substitute for Co ions in CCO. In addition, since Nb has a larger atomic mass than does Co, phonon scattering is expected to increase while the thermal conductivity decrease, with the result that the thermoelectric properties may be further improved. In this work, the influences of Nb-doping on the thermoelectric properties of CCO epitaxial thin films are investigated. In traditional approaches, normally there are two steps for preparing CCO thin films doped with metallic elements. The doped precursor compound is first prepared using solid-state reactions, the sol–gel method etc., and secondly, thin films are fabricated using various techniques such as chemical deposition, spin-coating, and PLD. These approaches, however, require significant time to carry out, and it is difficult to achieve high doping concentrations because the thermodynamics tends to be unfavorable. The ion beam injection technique has several advantages over the traditional approaches: (1) The distribution of the dopant ion in the
C. Zhu et al. / Journal of Alloys and Compounds 567 (2013) 122–126
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probe method at room temperature before injection, after injection and after annealing. The resistivity and Seebeck coefficient of the CCO thin films were measured using the standard four-probe method and the steady state method with the Thermal Transport Option (TTO) module of a Physical Properties Measurement System (PPMS) over the temperature range 150–380 K. The main experimental instruments used in this work are listed below. The pulsed laser deposition and ion beam cleaning system (PLD-450a) was made by the Scientific Instruments Research Center of the Chinese Academy of Sciences in Shenyang. Excimer lasers (Compex Pro 205, KrF k = 248 nm) were produced by Coherent in the USA. The X-ray diffractometer (XRD) (X’Pert Pro MPD, Cu Ka) was made by PANalytical B.V. in the Netherlands. The S-4800 Field Emission Scanning Electron Microscope (SEM) was produced by Hitachi in Japan. The Scanning Probe Microscope (SPM) (NanoScope IV) was produced by Digital Instruments in the USA. The MEVVA ion source was produced by the Tongchuang Applied Plasma Technology Center in Chengdu, China. The physical properties measurement system (PPMS-9T) was made by Quantum Design in the USA.
3. Results and discussion Fig. 1. XRD patterns for the CCO thin films. The inset shows an expanded view for the angular range 16.3–17°.
Fig. 2. Cross section microstructure of CCO thin film on (0 0 0 1) Al2O3 substrate with SEM.
sample can be controlled by the injection energy; (2) The doping concentration can be controlled through the injection doses, and high concentrations are easy to achieve; (3) Less time is required. In this paper, the ion beam injection technique has been used for doping CCO thin films with high concentrations of Nb ions, and the thermoelectric properties of CCO thin films with different injection doses have been measured.
Fig. 1 shows the XRD pattern of CCO thin films on (0 0 0 l) single crystal Al2O3 substrates after annealing. The X’Pert HighScore Plus software was used to search for peaks and carry out data processing. The (0 0 2), (0 0 3), (0 0 4), (0 0 5), (0 0 6) and (0 0 8) crystal plane diffraction peaks appear at 16.63°, 25.10°, 33.41°, 41.99°, 51.21° and 69.79°, respectively. This indicates perfect c-axis oriented growth of CCO thin films due to the fact that diffraction peaks from (00l) planes are observed in the patterns and no other plane peaks are found. Two points should be noted. First, the positions of the diffraction peaks show a clear drift as the Nb injection dose rises as shown in the inset of Fig. 1. This means that the Nb ions enter into the crystal lattice and demonstrates that Nb doping has been achieved by ion beam injection. In addition, no peak for cobalt oxide is found in the patterns possibly due to the cobalt oxide appearing in the films as non-crystalline nanoclusters. The SEM image in Fig. 2 shows the cross section microstructure of a CCO thin film grown on a (0 0 0 l) single crystal Al2O3 substrate. The dense CCO thin film and the growth interface can be clearly seen. The AFM (5 5 lm2) image in Fig. 3 shows the surface morphology of the CCO thin film. The image indicates that the film is smooth and the grain size is uniform. Both pictures show that the film has good crystallinity and has grown with its c-axis perpendicular to the Al2O3 substrate surface. In our experience, there are large differences in the crystalline quality of CCO thin films prepared by PLD. Therefore, we prepared
2. Experimental details The CCO precursor powder was prepared using the sol–gel method and then pressed into high density solid targets to be used for PLD. The CCO thin films were grown on single crystal Al2O3 substrates by PLD with a substrate temperature of about 1023 K, a laser energy density of 1.5 J/cm2, a deposition frequency of 3 Hz and under an oxygen pressure of 20 Pa. The commercial Al2O3 substrates (10 3 mm2) were cleaned with pure alcohol, acetone and deionized water for 5 min each in an ultrasonic bath and blown dry with N2. The phase composition and structure of the CCO thin films were characterized by X-ray Diffraction (XRD). The cross section microstructure and surface morphology of the films were observed using a Scanning Electron Microscope (SEM) and an Atomic Force Microscope (AFM). Nb ion injection into the CCO thin films was implemented using a Metal Vapor Vacuum Arc (MEVVA) ion source. The implantation depth and distribution were simulated with the SRIM software [12]. The injection voltage was 34 kV and the injection dip angle was 7°. In this work, the average Nb implantation concentrations in CCO thin films were 3.65 1019 atoms/cm3 (referred to below as CCONb4), 1.46 1020 atoms/cm3 (CCONb16), 7.33 1020 atoms/cm3 (CCONb80) and 3.65 1021 atoms/cm3 (CCONb400). The samples were annealed at 923 K under 1 atm oxygen pressure for 30 min after injection. The phase composition and surface morphology were observed using XRD, SEM and AFM. In order to quantify the effects of doping, the resistivity of the samples was measured using the four-
Fig. 3. Surface morphology of CCO thin film with AFM.
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Table 1 Resistivity of CCO thin films with different Nb injection doses at room temperature. Rb, Ri and Ra are the resistivities for the films before injection, after injection and after annealing, respectively. Sample
Rb (mX cm)
Ri (mX cm)
Ra (mX cm)
(Rb–Ra)/Rb (%)
CCONb4 CCONb16 CCONb80 CCONb400
10.51 10.53 10.67 10.64
34.79 44.09 66.47 643.53
7.78 9.59 12.43 21.34
25.97 8.92 16.49 100.56
Fig. 4. Temperature dependence of resistivity for different Nb injection doses.
Table 2 Carrier concentration n of CCO thin films with different Nb injection doses at 300 K. Sample
CCO
3
n (cm
)
3.5 10
CCONb4 21
22
1.2 10
CCONb16 8.5 10
21
CCONb80 21
5.5 10
CCONb400 1.4 1021
a number of CCO thin films, and the resistivity of the films was measured at room temperature by the four-probe method to pick out five similar samples. The resistivity of these five samples was around 10.6 mX cm. After Nb was injected into these samples using the MEVVA ion source, the resistivity was measured again. The resistivity was measured a third time after annealing to show the full effect of Nb doping. As shown in Table 1, the resistivity after annealing of the CCONb4, CCONb16, CCONb80 and CCONb400 samples was changed by 25.97%, 8.92%, 16.49% and 100.56% with respect to the as-prepared samples. Conductivity can be written as
r ¼ nel
Fig. 4 shows the resistivity characteristics of CCO thin films with different doses of injected Nb in the temperature range 150–380 K. As can be seen, the resistivity of the samples first decreases and then increases as the Nb injection dose rises. At 300 K, CCONb4 had the smallest resistivity, 7.67 mX cm, of all the samples. For CCONb0, CCONb16, CCONb80 and CCONb400, the resistivities were 10.49 mX cm, 9.49 mX cm, 12.31 mX cm and 21.13 mX cm, respectively. There are small differences between these values and the values obtained using the four-probe method as shown in Table 1 due to experimental error. The Ca3Co4O9 molecular structure can be denoted as (Ca2CoO3)(CoO2)1.62 [6,13]. The CoO2 layer which has an octahedral structure is a conductive layer while the Ca2CoO3 layer which has a rock salt type trigonometric structure is a carrier reservoir layer. For Ca cations, the coordination number is 8, the valence state is +2, and the ionic radius is 1.12 Å. In CCO thin films Nb will appear as Nb5+. When the coordination number is 6, the ionic radius of Nb5+ is 0.64 Å, and when the coordination number is 8, the ionic radius is 0.74 Å. The substitution of Nb for Ca is therefore be improbable due to the large ionic radius difference between Nb5+ and Ca2+ with the same coordination number. For all Co cations in CCO, the coordination number is 6. X-ray absorption (XAS) and photoemission studies of CCO show there is no Co2+ in the Ca2CoO3 crystal structure and indicate that the hole doped Co–O triangular lattice has low spin Co4+ species in a low spin Co3+ background [14,15]. Ultraviolet photoemission spectroscopy (USP) results suggest that phase separation plays an important role in CCO [14]. Since the radius of Nb5+ (0.64 Å) is close to that of Co3+/Co4+ (0.545 Å/0.53 Å, low spin states), Nb may randomly substitute for Co cations. The spin-state transition plays a leading role in the hopping transport process for CCO, so hopping transportation should become more difficult due to the differences in the electronic structures of Nb and Co cations, which will weaken the superexchange interactions. On the other hand, it is obvious that with the substitution of Nb for Co, high-valence cations replace low-valence cations and donate electrons so that the hole concentration should decrease. Both processes are expected to lead to a resistivity increase with an increase in the Nb concentration. However, experimentally, the opposite results have been observed. It is found that the resistivity decreases to a minimum for CCONb4 and then increases, eventually reaching a maximum at the highest Nb concentration, as shown in Fig. 4. In addition, Hall measurements indicate that the hole concentration first increases and then decreases with Nb doping concentration as listed in Table 2. One possible reason is that with Nb doping, the oxygen deficit [16] disappears and the oxygen content increases in CCO films and the in-
ð1Þ
where n is the carrier concentration, e is the electron charge and l is the carrier mobility. It is clear that the changes in the conductivity may be due to changes in the carrier concentration, the mobility or both. The Seebeck coefficient can be expressed as
SðTÞ ¼
2 c p2 kB T @ ln lðeÞ þ n 3e @e e¼EF
ð2Þ
where c, n, kB, and l(e) are the specific heat, carrier concentration, Boltzmann constant and carrier mobility, respectively [2]. Eq. (2) shows that S is determined by n and l(e) for a specific temperature. It is therefore expected that the Seebeck coefficient S will also be modified by Nb doping as it depends on the same two parameters as does the conductivity. Consequently, it is possible that the value of ZT for CCO thin films could be improved by Nb doping.
Fig. 5. LnrT vs. 1000/T for CCO thin films with different Nb injection doses.
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Fig. 6. Temperature dependence of Seebeck coefficient for CCO thin films with different Nb injection doses.
Fig. 7. Temperature dependence of power factor for CCO thin films with different Nb injection doses.
creased number of holes provided by the increased oxygen content is larger than the number of electrons injected. However, the number of oxygen deficiencies is limited, so the hole concentration eventually stops increasing and begin to decrease as the Nb injection dose rises. Mikami et al. [17] have reported a similar result in their study of Bi doping of single crystal CCO. It is found that Bi3+ replaces Ca2+ and can effectively decrease the oxygen deficit. Consequently the oxygen content increases in the CCO structure. In addition, it is suggested that the replaced Co atoms could react with oxygen and produce CoO non-crystalline nanoclusters. The CoO nanoclusters formed are expected to enhance the scattering of holes in CCO films, and thereby lead to a further increase in the resistivity. Hopping of a hole between Co3+ and Co4+ is the elementary charge transport mechanism in CCO. According to small polaron theory [18,19], the electrical conductivity of small polarons can be described by the expression r = C/T exp(Ea/kBT), where C is a pre-exponential term, Ea is the activation energy for hopping, kB is the Boltzmann constant, and T is the absolute temperature. As shown in Fig. 5, the experimental data show a good linear relationship between lnrT and 1000/T, which indicates a small polaron hopping conduction mechanism in the CCO films. The fitted values of the activation energies for CCONb0, CCONb4, CCONb16, CCONb80 and CCONb400 are 31 meV, 40 meV, 45 meV, 46 meV and 40 meV, respectively. The
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result indicates that the activation energy of the carriers increases slightly when Nb is injected. This may be due to the fact that Nb substitution for Co cations makes hole transportation more difficult. Fig. 6 shows the temperature dependence of the Seebeck coefficient for CCO thin films with different Nb injection doses in the temperature range 150–380 K. Compared with the pure CCO sample, the Seebeck coefficient shows little change for CCONb4, CCONb16, and CCONb80. At room temperature, CCO and CCONb4 have a similar Seebeck coefficient of about 90 lV/K. CCONb16 and CCONb80 have a Seebeck coefficient of about 96 lV/K which is slightly higher than that for CCO. The Seebeck coefficient for CCONb400 is only about 72 lV/K which is obviously lower than that for CCO. Koshibae et al. [20] have studied the thermopower in strongly correlated electron systems with orbital degeneracy and suggested that the triangular lattice of CCO is advantageous to the thermopower that is induced by the degeneracy. Nb substitution for Co may increase the difficulty of the superexchange interactions. Substitution of small amount of Nb has a weak effect on the thermopower as may be concluded from the observation that CCO, CCO4, CCONb16 and CCONb80 have similar thermopower values. Replacement of Co by Nb to a greater extent may well reduce the superexchange between Co3+ and Co4+ and hence decrease the Seebeck coefficient as may be seen in the case of CCONb400. The temperature dependence of the power factor, S2r, for CCO thin films with different Nb injection doses in the temperature range 150–380 K is shown in Fig. 7. At room temperature, CCONb4 shows the highest power factor of about 0.10 mW/m K2 due to the facts that Seebeck coefficient changes little and the resistivity decreases dramatically compared to the other samples. The power factor of CCONb16 also improved over that of CCO because it has a slightly higher Seebeck coefficient and lower resistivity than does CCO. The reduced power factor of CCONb80 is due to the fact that the resistivity increases more than does the Seebeck coefficient. For CCONb400, the decrease of the Seebeck coefficient and the increase in the resistivity both lead to a dramatic reduction in the power factor. Comparing the power factors of pure CCO and the Nb injected CCO thin films, CCONb4 has the highest power factor at room temperature. The maximum power factor observed in the present work was 0.17 mW/m K2 for CCONb4 at 380 K. Thermal conductivity depends on both the conductivity of the lattice and of the charge carries, and may be expressed as j = jL + je, where jL is the lattice thermal conductivity and je is the carrier thermal conductivity. je can be calculated from the Widemann– Franz law as je = LrT = LT/q, where L is the Lorenz constant, L = 2.45 108 V2/K2. Consequently, je shows the same trend as does r. Therefore je is increased for CCONb4 and CCONb16, with respect to CCO, but is still only about 0.1 W/m K, as calculated from the resistivity. Shikano et al. [3] reported that the thermal conductivity in the a–b plane of a single crystal of CCO is 3 W/m K, which is much larger than 0.1 W/m K. Therefore, it is probable that the thermal conductivity of CCO thin films comes mainly from lattice vibrations. Lattice thermal conductivity can be described by the equation jL = (1/3)CDvLlL, where CD, vL and lL are the lattice specific heat, phonon group velocity and mean free path of phonons, respectively. Because the mean free path of the phonon is expected to be significantly decreased due to increased scattering [21] by CoO nanoclusters, the lattice thermal conductivity is also decreased, and the thermoelectric figure of merit, ZT, can be improved, if CD and vL are assumed to remain unchanged.
4. Conclusions An ion beam injection technique has been applied to dope CCO thin films with the metallic element, Nb, and the experimen-
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tal results demonstrate that the method is effective in increasing the power factor of CCO films. Different doses of Nb ions were injected into CCO epitaxial thin films grown on single crystal Al2O3 substrates using a MEVVA ion source, and the samples were annealed in an oxygen atmosphere for 30 min in order to remove implantation damage and activate implanted ions. In this way, a series of samples with different Nb doping concentrations was obtained. The resistivity and Seebeck coefficients of the samples were measured using a PPMS in the temperature range 150– 380 K. Measurement results indicate that the resistivity decreases for small doses of injected Nb, but the resistivity increases for large doses. At 300 K, CCONb4 has a resistivity of 7.67 mX cm which was the smallest observed across all the samples studied. The Seebeck coefficient changes in an irregular manner for CCO, CCONb4, CCONb16 and CCONb80, but decreases dramatically for CCONb400. This demonstrates that small doses of injected Nb do not significantly modify the Seebeck coefficient, while larger doses of Nb can cause a large decrease. At room temperature, CCONb4 had the highest power factor, 0.10 mW/m K2, while the maximum power factor at any temperature studied was 0.17 mW/m K2 at 380 K.
Acknowledgments This work is supported by the Key item Science Foundation of Hebei Province (Grant No. 10965125D), the National Natural Science Foundation of China, under the Contract No. NSF-11174069.
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