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High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti
Journal Pre-proof High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti Satoko Abe, Masaya Takagi, Shoko Iwasaki, Yoshinori Satou, Satoko Tezuka, Shigeki Komine, Fumio Munakata PII:
S0022-4596(20)30007-4
DOI:
https://doi.org/10.1016/j.jssc.2020.121177
Reference:
YJSSC 121177
To appear in:
Journal of Solid State Chemistry
Received Date: 27 October 2019 Revised Date:
4 January 2020
Accepted Date: 4 January 2020
Please cite this article as: S. Abe, M. Takagi, S. Iwasaki, Y. Satou, S. Tezuka, S. Komine, F. Munakata, High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti, Journal of Solid State Chemistry (2020), doi: https://doi.org/10.1016/j.jssc.2020.121177. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier Inc.
CRediT author statement
Satoko ABE: Investigation, Writing - Original Draft, Writing - Review & Editing Masaya TAKAGI: Investigation Shoko IWASAKI: Investigation Yoshinori SATOU: Investigation Satoko TEZUKA: Investigation Shigeki KOMINE: Conceptualization, Resources Fumio MUNAKATA: Conceptualization, Supervision
11
50 Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
10 9
0 -50
8
-100
7
-150
6
-200
5
-250
4
-300
3 1.2
1.4
1.6
1.8
2.0 -1
1000/T (K )
2.2
2.4
Li1.05Mn2O4-δ Li1.05Mn2-xNixO4-δ (x=0.2) Li1.05Mn2-xNixO4-δ (x=0.3) Li1.05Mn2-xNixO4-δ (x=0.5) Li1.05Mn2-xTixO4-δ (x=0.1) Li1.05Mn2-xTixO4-δ (x=0.2) Li1.05Mn2-xTixO4-δ (x=0.3)
-350 1.2
1.4
1.6
1.8
2.0 -1
1000/T (K )
2.2
2.4
High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti Satoko ABE1, Masaya TAKAGI1, Shoko IWASAKI1, Yoshinori SATOU2, *, Satoko TEZUKA3, Shigeki KOMINE2, **, and Fumio MUNAKATA1 1
Department of Chemistry and Energy Engineering, Tokyo City University, 1-28-1 Tamazutsumi, Setagaya-ku, Tokyo 158-8557, Japan
2
DENSO CORPORATION, 1-1 Showa-cho, Kariya, Aichi 448-8661, Japan
3
Faculty of Risk and Crisis management, Chiba Institute of Science, 15-8 Shiomi-cho, Choshi, Chiba 288-0025, Japan
*
Present address: Panasonic Corporation, 1006, Oaza Kadoma, Kadoma-shi, Osaka 571-8501, Japan
**
Present address: LG Japan Lab Inc., Glass Cube Shinagawa 3F, 4-13-14, Hagashi Shinagawa, Shinagawa-ku, Tokyo 140-0002, Japan
Abstract Spinel Li1.05Mn2O4−δ substituted with Ni and Ti were fabricated via a solid-state reaction technique. The Rietveld refinement of synchrotron X-ray powder diffraction data confirmed an Fd3m space group of spinel structure for all synthesized materials. Variation in the lattice parameter and M-O distances (M: transition metal at 16d site) due to the dopant content indicated that Ni substitution enhanced the overlap in Mn 3d wave functions and Ti substitution induced structural distortion. Electrical conductivity measurements based on a four-probe method were conducted at high temperatures, demonstrating that both the Ni- and Ti-substituted samples exhibited the same non-adiabatic small-polaron hopping mechanism as that of the non-substituted LiMn2O4. Thermopower measurements were performed at high temperatures via a steady-state technique. All samples with and without substitution were characterized as n-type semiconductors and their Seebeck coefficients were thermally activated. The electrical conductivities and the Seebeck coefficients decreased with increasing Ni content, whereas they were less affected by Ti substitution. The activation energies of electrical conductivity Eσ and thermopower ES were derived from the experimental results and were used to calculate the hopping energies, WH. Ni substitution initially reduced WH but substantially increased WH when Ni = 0.5. In contrast, Ti substitution had a major effect on only ES. Variation in carrier density and structural distortion following the substitution for Mn exerted a large influence on the electrical conductivity and Seebeck coefficients of the doped materials.
Keywords Lithium manganese spinels, Ni and Ti substitution, Small polarons, Electrical conductivities, Thermopower, Synchrotron X-ray powder diffraction
1. Introduction Lithium–transition-metal oxides with spinel structures have been widely investigated as potential cathode materials for rechargeable lithium-ion batteries [1-5]. Spinel LiMn2O4 exhibits properties that are advantageous for commercial battery use, such as high density, large rechargeable capacities, low cost, and non-toxicity. Capacity fading during its charge–discharge cycles, however, still limits the practical application of LiMn2O4-based cathodes [6]. Many reports have suggested that Jahn–Teller distortion [7], oxygen defects 1
[8,9], and Mn dissolution [10] are partly responsible for their poor cyclability. The crystal structure of LiMn2O4 around room temperature belongs to the cubic Fd3m space group [11]. The oxygen atoms at 32e sites form a cubic-close-packed array in which the Li and Mn ions reside at tetrahedral 8a and octahedral 16d sites, respectively. In the stoichiometric spinel structure, the same amounts of Jahn–Teller-active Mn3+ and Mn4+ ions occupy octahedral sites [12]. Jahn–Teller distortion is essentially caused by the cooperative interaction of local distortions around Mn3+ in these octahedral sites. Stoichiometric LiMn2O4 undergoes a reversible structural transition from cubic Fd3m to orthorhombic Fddd around room temperature [13-15]. This transition is attributed to Jahn–Teller distortion [16,17] derived from the ordering of Mn3+ and Mn4+ ions in the octahedral 16d sites of the spinel lattice. LiMn2O4 is an n-type semiconductor with small polarons, which are eg electrons, on Mn3+ trapped in local lattice relaxation sites, and it follows a small-polaron-hopping conduction mechanism [18-22]. Although some reports [20,21] mentioned that this hopping conduction process may be due to adiabatic small polarons, Sugiyama et al. [22] found a linear Arrhenius relationship between σT3/2 vs. 1/T, providing evidence for the presence of non-adiabatic small polarons [23,24]. According to Iguchi et al. [25,26], impedance analyses based on the Debye model [27,28] and dielectric measurements further demonstrated that a non-adiabatic small-polaron-hopping process was vital to the electrical transport in LiMn2O4. More recently, the partial substitution of LiMn2O4 with other transition metals has received considerable attention as a useful strategy to improve electrochemical performance, including poor cyclability [29-35]. The origin of this poor cyclability, however, has not been fully understood. A small-polaron-hopping conduction mechanism exhibits complicated behavior due to the presence of more than two electroactive transition metals [36-38]. In our previous study [39], we focused on Ni and Ti as substitution metals and discussed the effects of Ni substitution and Ni/Ti co-substitution on electrical conductivity at high temperatures. To comprehend the essential roles of Ni and Ti as substitution metals in the spinel lattice, more transport measurements are required. The goal of the present work is to investigate the electrical conductivities and Seebeck coefficients of Ni- and Ti-doped Li1.05Mn2O4−δ at high temperatures. 2. Experimental Li1.05Mn2O4−δ, Li1.05Mn2−xNixO4−δ (x = 0.1, 0.2, 0.3, 0.4, 0.5), and Li1.05Mn2−xTixO4−δ (x = 0.1, 0.2, 0.3) were prepared via a conventional solid-state reaction technique using Li2CO3, Mn2O3, NiO, and TiO2 powders. A Li-rich solid was prepared to effectively evaporate Li2O. The mixed powders were pelletized and calcined in air at 843 K for 5 h and then heated at 1143 K for 10 h. After the heated pellets were ground into powders, they were pressed into pellets again and finally sintered in air at 1143 K for 12 h. The contents of the samples after the second heat treatment were characterized using inductively coupled plasma with optical emission spectroscopy (ICP-OES). The oxidation state of Mn was determined by an iodometric titration method described in detail in Ref. [40,41]. Powder X-ray diffraction was performed at room temperature using synchrotron radiation on the beamline BL5S2 of the Aichi Synchrotron Radiation Center. A wavelength of 1 Å was selected, and the actual wavelength of each measurement was calibrated using a CeO2 standard sample (SRM674b, fabricated by 2
NIST) and refined as 0.999611 Å. DC conductivity was measured using a conventional four-probe method, and thermopower was measured via a steady-state technique. Electrical conductivities and Seebeck coefficients were measured as a function of ) of 2.0 × 104 Pa.
temperature from 473 K to 723 K with an oxygen partial pressure (
3. Results The nominal compositions of the samples and those calculated from the ICP-OES analyses were in good agreement with each other. The oxidation state of Mn indicated a Mn valency of approximately 3.5 in Li1.05Mn2O4−δ. The substitution of Ni, a lower valent metal than Mn, in LiMn2O4 increased the oxidation state of Mn to nearly 4 (Table 1); however, the substitution of Ti, a metal with the same valency as Mn4+, in LiMn2O4 made the oxidation state of Mn decrease. Rietveld structural refinements of synchrotron X-ray powder diffraction profiles for all samples were conducted using the RIETAN-FP system [42]. The experimental and calculated patterns of Li1.05Mn2O4−δ, Li1.05Mn1.7Ni0.3O4−δ, and Li1.05Mn1.8Ti0.2O4−δ are shown in Figure 1 and summarized in Table 1. All reflections of samples with and without substitution were indexed to the cubic-spinel-type structure with the Fd3m space group, respectively. These results suggested that substitution by Ni and Ti did not induce structural transformation, and 16d sites in the Fd3m space group were probably replaced with Ni and Ti. The lattice parameter of Li1.05Mn2O4−δ was determined to be 8.24538(5) Å, which is consistent with previous work with stoichiometric samples. Variation in the lattice parameter for doped Li1.05Mn2O4−δ according to the dopant contents, are shown in Figure 2. The Ni-substituted samples showed a substantial decrease in the lattice parameter until the Ni content reached 0.3 equivalents, and the lattice parameter then slowly reached saturation. In contrast, the lattice parameter continuously increased with increasing Ti content for the Ti-doped samples. As shown in Table 1, the M-O distances of the Ni-substituted samples became smaller than Li1.05Mn2O4−δ. In contrast, the Ti-substituted samples maintained M-O distances that were similar to non-doped Li1.05Mn2O4−δ. The temperature dependence of electrical conductivities for Li1.05Mn2O4−δ and the Ni- and Ti-substituted = 2.0 × 104 Pa. The conductivity
samples were measured under a constant partial pressure of oxygen,
due to the hopping conduction of non-adiabatic small polarons is represented by the equation =
⁄
−
=
⁄
−
+
, (1)
where A is the pre-exponential factor, Eσ is the activation energy of electrical conductivity, kB is the Boltzmann constant, ES is the activation energy of thermopower, and WH is the polaron hopping energy of non-adiabatic small polarons [23,24]. Figure 3 illustrates the Arrhenius relations of σT3/2 = Aexp(-Eσ/kBT) for six samples. All samples behaved as n-type semiconductors, characterized by an increase in conductivity with increasing temperature. Six separated lines in Figure 3 show the results of fitting the linear approximation to the conductivity plots of each sample. These plots held linear relationships and followed Arrhenius’ law within the evaluated temperature range of 473 K to 723 K. The high-temperature electrical conduction within all samples, therefore, suggests that the primary mechanism is likely a hopping process of non-adiabatic small 3
polarons. These conductivity results are consistent with our previous report on Li1.05Mn1.5Ni0.5O4−δ [39], which also followed a non-adiabatic small-polaron-hopping mechanism. In addition, the present work demonstrates that Ti substitution does not change the conduction mechanism of Li1.05Mn2O4−δ. We performed thermopower measurements as a function of temperature under
= 2.0 × 104 Pa for the
same samples whose electrical conductivities were measured. The Seebeck coefficient S for a thermally activated case is expressed as !=
+ !# , (2)
"
where q is the charge of a carrier, ES is the activation energy of thermopower, and S0 is the thermopower in the high-temperature limit [27,43,44]. Figure 4 illustrates the Seebeck coefficients vs. 1000/T and the results of fitting a linear approximation by the above equation. The Seebeck coefficients of all six samples were negative and regularly increased with temperature. These negative values indicate that the charge carriers are electrons, which are eg electrons belonging to Mn3+ in this case. 4. Discussion 4.1 Carrier density variation by Ni and Ti substitution Li1.05Mn2O4−δ exhibited the highest electrical conductivity at each temperature for all samples tested (Figure 3). Ni substitution was shown to reduce the conductivities, whereas the conductivities of Ti-substituted samples retained almost the same value as those of Li1.05Mn2O4−δ until the Ti content reached 0.2 equivalents and then slightly decreased. Li1.05Mn1.5Ni0.5O4−δ showed the lowest conductivity of all samples. As shown in Figure 4, Li1.05Mn2O4−δ also had the largest Seebeck coefficient among six samples. The Seebeck coefficients decreased with increasing Ni content, whereas Ti substitution barely affected the Seebeck coefficient values. The trends for electrical conductivity behaviors and Seebeck coefficients with the Ni- and Ti-substituted samples were quite similar. Since the electrical conductivity and the Seebeck coefficient depend on carrier density [45,46], they are not entirely independent of each other. For both systems, the electrical conductivity decreases and the absolute value of the Seebeck coefficient increases logarithmically with decreasing carrier density. The non-adiabatic small-polaron theory described the pre-exponential factor A as =
%
2ℏ
& ' (⁄ ) ⁄
(⁄
, (3)
where n is the carrier density, a is the hopping distance of small polarons, and J is the electron transfer integral between neighboring transition metals [24,27]. This equation indicates that A is directly proportional to the carrier density n and inversely proportional to WH1/2; thus, n contributes more to A than WH. The values of A for all samples, which were derived from the intercepts in the ln (σT3/2) vs. 1000/T Arrhenius plots, are summarized in Table 2. As the Ni content in the samples increased, A significantly decreased in value. In contrast, the A values in the Ti-substituted samples were almost independent of Ti content and were similar to that of Li1.05Mn2O4−δ. Given that Li1.05Mn2−xNixO4−δ is electrically neutral, the loss of an eg electron via the reduction of Mn3+ is required. For the Ti-doped sample, a reduction of Mn4+ is needed such that 4
Li1.05Mn2−xTixO4−δ will hold the same number of eg electrons as Li1.05Mn2O4−δ. There is no discrepancy between the variations of the number of eg electrons and the A values in Table 2, i.e., the carrier density n. Hence, the electrical conductivities and Seebeck coefficients associated with the Ni- and Ti-doped materials were primarily derived from variations in the carrier density n via Ni and Ti substitution.
4.2 Activation energies and hopping energies The activation energies of electrical conductivity Eσ and thermopower ES were derived from the slopes of the ln (σT3/2) vs. 1000/T Arrhenius (Figure 3) and S vs. 1000/T plots (Figure 4), respectively. Eσ and ES values are summarized in Table 2. Eσ contains ES and WH, as seen in equation (1). In addition to activation energies, we calculated the hopping energies WH using the relation WH = Eσ − ES (Table 2). The hopping energy WH corresponds to the minimum energy to produce the equivalency of the two neighboring sites involved in the polaron hopping [47-49]. A functionality of WH = 1/2WP – t, where WP is the polaron binding energy and 2t is the 3d-narrow bandwidth, is maintained for polaronic conduction [24]. As t is negligible in the non-adiabatic case [24], it is unlikely that the substitution of Ni and Ti for Mn expands the bandwidth 2t. The WH behavior of the Ni- and Ti- doped material, therefore, is predominantly governed by WP. To maintain electrical neutrality, Ni2+ displaced Mn3+ in the Ni-substituted samples. When the dopant content of Ni2+ increased, Mn3+ decreased and Mn4+ increased, affording Li1.05Mn1.5Ni0.5O4−δ with the minimum amount of Mn3+ of all Ni-substituted samples. According to Shannon [50], although Ni2+ has a larger ionic radius than Mn3+, Mn4+ has a smaller ionic radius than Mn3+. The average octahedral site cation radius decreases as increasing Ni content. This explained the decrease in lattice parameter and M-O distance accompanied by the substitution of Ni. The overlap in Mn 3d wave functions via O 2p wave functions can be enhanced due to the contraction of M-O distance. The carrier localization, hence, weakens and Wp decreases. Consequently, the WH values of Li1.05Mn2−xNixO4−δ (x = 0.2 and 0.3) became smaller than that of Li1.05Mn2O4−δ. For Li1.05Mn1.5Ni0.5O4−δ, the value of WH dramatically increased and showed the highest Eσ for all samples. When the Ni content is greater than 0.3, the rate of decrease in lattice parameter undergoes a drastic change and becomes smaller. We speculated that the state of Li1.05Mn1.5Ni0.5O4−δ might differ from that of Li1.05Mn2−xNixO4−δ (x = 0.2 and 0.3). Lee et al. [51] found a small degree of ordering in 16d sites in Li1.05Mn1.5Ni0.5O4−δ using fourier transform infrared spectroscopy, although the Rietveld refinement of the synchrotron X-ray diffraction pattern proved their Li1.05Mn1.5Ni0.5O4−δ to be a disordered spinel structure (Fd3m) without any long-range ordering. Infrared spectroscopy is generally recognized as a useful tool to qualitatively prove the cation ordering in 16d sites even if it is a small degree of ordering [52]. This localized inhomogeneity could explain why our Li1.05Mn1.5Ni0.5O4−δ exhibited the highest WH of all Ni-substituted samples. The Ti-substituted samples, in contrast, maintained nearly equivalent WH values to those of Li1.05Mn2O4−δ, and the electrical neutrality of the system replaced Mn4+ with Ti4+ in Li1.05Mn2−xTixO4−δ. Ti4+ has a larger ionic radius than Mn4+ and the lattice parameters of Ti-doped samples become larger. Ti substitution, however, keeps the M-O distances almost the same as those of Li1.05Mn2O4−δ, and it does not essentially affect the M-O bonds. Consequently, the WH values of Ti-doped samples did not differ from those of non-doped Li1.05Mn2O4−δ. 5
The activation energy of thermopower ES is the energy difference between identical lattice distortions with and without the carrier, i.e., the energy required to generate a free hopping polaron [47-49]. The ES values of Ni- and Ti-substituted samples slowly increased with dopant content. Because both substitution metals have a larger ionic radius than Mn described above, substituting these metals for Mn should increase a level of structural distortion into the lattice. This condition was supposed to exercise an influence on the increase in E S. Given the small observed effects of Ni and Ti substitution on Eσ, ES, and WH values, the contents of these dopants do not appear to cause a significant change to the electronic structure.
5. Conclusion In the present work, the spinel Li1.05Mn2O4−δ, Li1.05Mn2−xNixO4−δ, and Li1.05Mn2−xTixO4−δ of space group Fd3m were synthesized and their transport properties were analyzed. The ionic radii of these transition metals and the observed variation in the samples’ lattice parameters suggest that Ni substitution increases the overlap in Mn 3d wave functions, and Ti substitution introduces structural distortion. Electrical conductivities measured in the high-temperature range of 473 K to 723 K suggest that the primary mechanism for electrical conduction in all samples is a non-adiabatic hopping process by small polarons. The thermopower measurements demonstrated that all samples with and without substitution are n-type semiconductors and their Seebeck coefficients are thermally activated. The electrical conductivity and Seebeck coefficient values decreased with added Ni content, while the Ti-substituted samples maintained the same values as Li1.05Mn2O4−δ. The electrical neutrality in the Ni-substituted samples necessitated a loss of eg electrons from Mn3+, while Ti-substituted samples maintained the same eg electron density as that of Li1.05Mn2O4−δ. The electrical conductivity and Seebeck coefficient behaviors of the Ni- and Ti-substituted compounds were considered to be caused by a variation in the carrier density. Furthermore, the hopping energy WH of the Ni-substituted samples decreased because the overlap in Mn 3d wave functions increased and the carriers were less localized. An exception to this rule was observed for the sample with the lowest Mn3+ content, Li1.05Mn1.5Ni0.5O4−δ, which experienced a sudden increase in WH. Ti substitution induced structural distortion into the lattice, and it therefore had almost no effect on WH but a small effect on the activation energy of the thermopower ES. References [1] M. M. Thackeray, W. I. F. David, P. G. Bruce, and J. B. Goodenough, Mater. Res. Bull. 18 (1983) 461-472. [2] M. M. Thackeray, P. J. Johnson, L. A. de Picciotto, P. G. Bruce, and J. B. Goodenough, Mater. Res. Bull. 19 (1984) 179-187. [3] J. M. Tarascon and D. Guyomard, J. Electrochem. Soc. 138 (1991) 2864-2868. [4] D. Guyomard and J. M. Tarascon, J. Electrochem. Soc. 139 (1992) 937-948. [5] M. M. Thackeray, A. de Kock, M. H. Rossouw, D. Liles, R. Bittihn and D. Hoge, J. Electrochem. Soc. 139 (1992) 363-366. [6] A. Yamada, J. Solid State Chem. 122 (1996) 160-165. 6
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8
Table 1. Mn valency, space groups, lattice parameters, M–O distances, and Rietveld factors (Rwp and RB) for Li1.05Mn2O4−δ, Li1.05Mn1.7Ni0.3O4−δ, and Li1.05Mn1.8Ti0.2O4−δ. Sample
Li1.05Mn2O4−δ
Li1.05Mn1.7Ni0.3O4−δ
Li1.05Mn1.8Ti0.2O4−δ
Mn valency
3.52
3.70
3.35
Space group
Fd3m
Fd3m
Fd3m
Lattice parameter (Å)
8.24538(5)
8.19287(2)
8.26125(5)
M-O distance (Å)
1.9571(8)
1.9496(4)
1.9582(5)
Rwp
5.52
3.57
4.04
RB
2.16
3.45
1.74
Rietveld factors (%)
10
15
20
25
30
35
40
45
50
55
60
65
70
75
2θ (deg)
Fig. 1. Observed and calculated synchrotron X-ray diffraction patterns at room temperature: (a) Li1.05Mn2O4−δ, (b) Li1.05Mn1.7Ni0.3O4−δ, and (c) Li1.05Mn1.8Ti0.2O4−δ.
9
8.28
Lattice parameter (Å)
8.26 8.24 8.22 8.20 Li1.05Mn2O4-δ Li1.05Mn2-xNixO4-δ Li1.05Mn2-xTixO4-δ
8.18 8.16
0.0
0.1
0.2
0.3
0.4
0.5
x in Li1.05Mn2-xMxO4-δ
Fig. 2. Lattice parameters as a function of dopant content for Li1.05Mn2O4, Ni-substituted sample Li1.05Mn2−xNixO4−δ, and Ti-substituted sample Li1.05Mn2−xTixO4−δ.
11 Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
10 9 8 7 6 5 4 3 1.2
1.4
1.6
1.8
2.0
2.2
2.4
-1
1000/T (K )
Fig. 3. Arrhenius plots of ln(σT3/2) vs. 1000/T for Li1.05Mn2 O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3).
10
50
Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
0 -50 -100 -150 -200 -250 -300 -350 1.2
1.4
1.6
1.8
2.0
2.2
2.4
-1
1000/T (K )
Fig. 4. Seebeck coefficients vs. 1000/T for Li1.05Mn2 O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3).
Table 2. Pre-exponential factors and activation energies Eσ of electrical conductivity, activation energies ES of thermopower, and hopping energies WH for Li1.05Mn2O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3). Sample
Composition
Li1.05Mn2.0-xTixO4−δ
Thermoelectric power
Hopping energy
Pre-exponential factor
Eσ
ES
WH
(Ω-1cm-1K3/2)
(eV)
(eV)
(eV)
5.1×106
0.41
0.05
0.36
x=0.2
2.6×106
0.41
0.07
0.33
x=0.3
1.7×106
0.41
0.08
0.34
x=0.5
1.6×106
0.45
0.08
0.37
x=0.1
5.8×106
0.42
0.06
0.36
x=0.2
5.3×106
0.42
0.06
0.36
x=0.3
4.3×106
0.43
0.07
0.36
Li1.05Mn2O4−δ Li1.05Mn2.0-xNixO4−δ
Electrical conductivity
11
Graphical abstract legend (TOC figure)
The figure on the left is Arrhenius plots of ln(σT3/2) vs. 1000/T, and the figure on the right is Seebeck coefficients vs. 1000/T.
Highlights · Ni- and Ti- doped lithium manganese spinels were synthesized. · The doping effect on structure and transport properties was studied systematically. · Ni- and Ti-doped samples followed a non-adiabatic small-polaron-hopping mechanism. · Ni dopant decreased the hopping energy and Ti dopant induced structural distortion.
S0022-4596(20)30007-4
DOI:
https://doi.org/10.1016/j.jssc.2020.121177
Reference:
YJSSC 121177
To appear in:
Journal of Solid State Chemistry
Received Date: 27 October 2019 Revised Date:
4 January 2020
Accepted Date: 4 January 2020
Please cite this article as: S. Abe, M. Takagi, S. Iwasaki, Y. Satou, S. Tezuka, S. Komine, F. Munakata, High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti, Journal of Solid State Chemistry (2020), doi: https://doi.org/10.1016/j.jssc.2020.121177. This is a PDF file of an article that has undergone enhancements after acceptance, such as the addition of a cover page and metadata, and formatting for readability, but it is not yet the definitive version of record. This version will undergo additional copyediting, typesetting and review before it is published in its final form, but we are providing this version to give early visibility of the article. Please note that, during the production process, errors may be discovered which could affect the content, and all legal disclaimers that apply to the journal pertain. © 2020 Published by Elsevier Inc.
CRediT author statement
Satoko ABE: Investigation, Writing - Original Draft, Writing - Review & Editing Masaya TAKAGI: Investigation Shoko IWASAKI: Investigation Yoshinori SATOU: Investigation Satoko TEZUKA: Investigation Shigeki KOMINE: Conceptualization, Resources Fumio MUNAKATA: Conceptualization, Supervision
11
50 Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
10 9
0 -50
8
-100
7
-150
6
-200
5
-250
4
-300
3 1.2
1.4
1.6
1.8
2.0 -1
1000/T (K )
2.2
2.4
Li1.05Mn2O4-δ Li1.05Mn2-xNixO4-δ (x=0.2) Li1.05Mn2-xNixO4-δ (x=0.3) Li1.05Mn2-xNixO4-δ (x=0.5) Li1.05Mn2-xTixO4-δ (x=0.1) Li1.05Mn2-xTixO4-δ (x=0.2) Li1.05Mn2-xTixO4-δ (x=0.3)
-350 1.2
1.4
1.6
1.8
2.0 -1
1000/T (K )
2.2
2.4
High-temperature transport properties of lithium manganese spinels substituted with Ni and Ti Satoko ABE1, Masaya TAKAGI1, Shoko IWASAKI1, Yoshinori SATOU2, *, Satoko TEZUKA3, Shigeki KOMINE2, **, and Fumio MUNAKATA1 1
Department of Chemistry and Energy Engineering, Tokyo City University, 1-28-1 Tamazutsumi, Setagaya-ku, Tokyo 158-8557, Japan
2
DENSO CORPORATION, 1-1 Showa-cho, Kariya, Aichi 448-8661, Japan
3
Faculty of Risk and Crisis management, Chiba Institute of Science, 15-8 Shiomi-cho, Choshi, Chiba 288-0025, Japan
*
Present address: Panasonic Corporation, 1006, Oaza Kadoma, Kadoma-shi, Osaka 571-8501, Japan
**
Present address: LG Japan Lab Inc., Glass Cube Shinagawa 3F, 4-13-14, Hagashi Shinagawa, Shinagawa-ku, Tokyo 140-0002, Japan
Abstract Spinel Li1.05Mn2O4−δ substituted with Ni and Ti were fabricated via a solid-state reaction technique. The Rietveld refinement of synchrotron X-ray powder diffraction data confirmed an Fd3m space group of spinel structure for all synthesized materials. Variation in the lattice parameter and M-O distances (M: transition metal at 16d site) due to the dopant content indicated that Ni substitution enhanced the overlap in Mn 3d wave functions and Ti substitution induced structural distortion. Electrical conductivity measurements based on a four-probe method were conducted at high temperatures, demonstrating that both the Ni- and Ti-substituted samples exhibited the same non-adiabatic small-polaron hopping mechanism as that of the non-substituted LiMn2O4. Thermopower measurements were performed at high temperatures via a steady-state technique. All samples with and without substitution were characterized as n-type semiconductors and their Seebeck coefficients were thermally activated. The electrical conductivities and the Seebeck coefficients decreased with increasing Ni content, whereas they were less affected by Ti substitution. The activation energies of electrical conductivity Eσ and thermopower ES were derived from the experimental results and were used to calculate the hopping energies, WH. Ni substitution initially reduced WH but substantially increased WH when Ni = 0.5. In contrast, Ti substitution had a major effect on only ES. Variation in carrier density and structural distortion following the substitution for Mn exerted a large influence on the electrical conductivity and Seebeck coefficients of the doped materials.
Keywords Lithium manganese spinels, Ni and Ti substitution, Small polarons, Electrical conductivities, Thermopower, Synchrotron X-ray powder diffraction
1. Introduction Lithium–transition-metal oxides with spinel structures have been widely investigated as potential cathode materials for rechargeable lithium-ion batteries [1-5]. Spinel LiMn2O4 exhibits properties that are advantageous for commercial battery use, such as high density, large rechargeable capacities, low cost, and non-toxicity. Capacity fading during its charge–discharge cycles, however, still limits the practical application of LiMn2O4-based cathodes [6]. Many reports have suggested that Jahn–Teller distortion [7], oxygen defects 1
[8,9], and Mn dissolution [10] are partly responsible for their poor cyclability. The crystal structure of LiMn2O4 around room temperature belongs to the cubic Fd3m space group [11]. The oxygen atoms at 32e sites form a cubic-close-packed array in which the Li and Mn ions reside at tetrahedral 8a and octahedral 16d sites, respectively. In the stoichiometric spinel structure, the same amounts of Jahn–Teller-active Mn3+ and Mn4+ ions occupy octahedral sites [12]. Jahn–Teller distortion is essentially caused by the cooperative interaction of local distortions around Mn3+ in these octahedral sites. Stoichiometric LiMn2O4 undergoes a reversible structural transition from cubic Fd3m to orthorhombic Fddd around room temperature [13-15]. This transition is attributed to Jahn–Teller distortion [16,17] derived from the ordering of Mn3+ and Mn4+ ions in the octahedral 16d sites of the spinel lattice. LiMn2O4 is an n-type semiconductor with small polarons, which are eg electrons, on Mn3+ trapped in local lattice relaxation sites, and it follows a small-polaron-hopping conduction mechanism [18-22]. Although some reports [20,21] mentioned that this hopping conduction process may be due to adiabatic small polarons, Sugiyama et al. [22] found a linear Arrhenius relationship between σT3/2 vs. 1/T, providing evidence for the presence of non-adiabatic small polarons [23,24]. According to Iguchi et al. [25,26], impedance analyses based on the Debye model [27,28] and dielectric measurements further demonstrated that a non-adiabatic small-polaron-hopping process was vital to the electrical transport in LiMn2O4. More recently, the partial substitution of LiMn2O4 with other transition metals has received considerable attention as a useful strategy to improve electrochemical performance, including poor cyclability [29-35]. The origin of this poor cyclability, however, has not been fully understood. A small-polaron-hopping conduction mechanism exhibits complicated behavior due to the presence of more than two electroactive transition metals [36-38]. In our previous study [39], we focused on Ni and Ti as substitution metals and discussed the effects of Ni substitution and Ni/Ti co-substitution on electrical conductivity at high temperatures. To comprehend the essential roles of Ni and Ti as substitution metals in the spinel lattice, more transport measurements are required. The goal of the present work is to investigate the electrical conductivities and Seebeck coefficients of Ni- and Ti-doped Li1.05Mn2O4−δ at high temperatures. 2. Experimental Li1.05Mn2O4−δ, Li1.05Mn2−xNixO4−δ (x = 0.1, 0.2, 0.3, 0.4, 0.5), and Li1.05Mn2−xTixO4−δ (x = 0.1, 0.2, 0.3) were prepared via a conventional solid-state reaction technique using Li2CO3, Mn2O3, NiO, and TiO2 powders. A Li-rich solid was prepared to effectively evaporate Li2O. The mixed powders were pelletized and calcined in air at 843 K for 5 h and then heated at 1143 K for 10 h. After the heated pellets were ground into powders, they were pressed into pellets again and finally sintered in air at 1143 K for 12 h. The contents of the samples after the second heat treatment were characterized using inductively coupled plasma with optical emission spectroscopy (ICP-OES). The oxidation state of Mn was determined by an iodometric titration method described in detail in Ref. [40,41]. Powder X-ray diffraction was performed at room temperature using synchrotron radiation on the beamline BL5S2 of the Aichi Synchrotron Radiation Center. A wavelength of 1 Å was selected, and the actual wavelength of each measurement was calibrated using a CeO2 standard sample (SRM674b, fabricated by 2
NIST) and refined as 0.999611 Å. DC conductivity was measured using a conventional four-probe method, and thermopower was measured via a steady-state technique. Electrical conductivities and Seebeck coefficients were measured as a function of ) of 2.0 × 104 Pa.
temperature from 473 K to 723 K with an oxygen partial pressure (
3. Results The nominal compositions of the samples and those calculated from the ICP-OES analyses were in good agreement with each other. The oxidation state of Mn indicated a Mn valency of approximately 3.5 in Li1.05Mn2O4−δ. The substitution of Ni, a lower valent metal than Mn, in LiMn2O4 increased the oxidation state of Mn to nearly 4 (Table 1); however, the substitution of Ti, a metal with the same valency as Mn4+, in LiMn2O4 made the oxidation state of Mn decrease. Rietveld structural refinements of synchrotron X-ray powder diffraction profiles for all samples were conducted using the RIETAN-FP system [42]. The experimental and calculated patterns of Li1.05Mn2O4−δ, Li1.05Mn1.7Ni0.3O4−δ, and Li1.05Mn1.8Ti0.2O4−δ are shown in Figure 1 and summarized in Table 1. All reflections of samples with and without substitution were indexed to the cubic-spinel-type structure with the Fd3m space group, respectively. These results suggested that substitution by Ni and Ti did not induce structural transformation, and 16d sites in the Fd3m space group were probably replaced with Ni and Ti. The lattice parameter of Li1.05Mn2O4−δ was determined to be 8.24538(5) Å, which is consistent with previous work with stoichiometric samples. Variation in the lattice parameter for doped Li1.05Mn2O4−δ according to the dopant contents, are shown in Figure 2. The Ni-substituted samples showed a substantial decrease in the lattice parameter until the Ni content reached 0.3 equivalents, and the lattice parameter then slowly reached saturation. In contrast, the lattice parameter continuously increased with increasing Ti content for the Ti-doped samples. As shown in Table 1, the M-O distances of the Ni-substituted samples became smaller than Li1.05Mn2O4−δ. In contrast, the Ti-substituted samples maintained M-O distances that were similar to non-doped Li1.05Mn2O4−δ. The temperature dependence of electrical conductivities for Li1.05Mn2O4−δ and the Ni- and Ti-substituted = 2.0 × 104 Pa. The conductivity
samples were measured under a constant partial pressure of oxygen,
due to the hopping conduction of non-adiabatic small polarons is represented by the equation =
⁄
−
=
⁄
−
+
, (1)
where A is the pre-exponential factor, Eσ is the activation energy of electrical conductivity, kB is the Boltzmann constant, ES is the activation energy of thermopower, and WH is the polaron hopping energy of non-adiabatic small polarons [23,24]. Figure 3 illustrates the Arrhenius relations of σT3/2 = Aexp(-Eσ/kBT) for six samples. All samples behaved as n-type semiconductors, characterized by an increase in conductivity with increasing temperature. Six separated lines in Figure 3 show the results of fitting the linear approximation to the conductivity plots of each sample. These plots held linear relationships and followed Arrhenius’ law within the evaluated temperature range of 473 K to 723 K. The high-temperature electrical conduction within all samples, therefore, suggests that the primary mechanism is likely a hopping process of non-adiabatic small 3
polarons. These conductivity results are consistent with our previous report on Li1.05Mn1.5Ni0.5O4−δ [39], which also followed a non-adiabatic small-polaron-hopping mechanism. In addition, the present work demonstrates that Ti substitution does not change the conduction mechanism of Li1.05Mn2O4−δ. We performed thermopower measurements as a function of temperature under
= 2.0 × 104 Pa for the
same samples whose electrical conductivities were measured. The Seebeck coefficient S for a thermally activated case is expressed as !=
+ !# , (2)
"
where q is the charge of a carrier, ES is the activation energy of thermopower, and S0 is the thermopower in the high-temperature limit [27,43,44]. Figure 4 illustrates the Seebeck coefficients vs. 1000/T and the results of fitting a linear approximation by the above equation. The Seebeck coefficients of all six samples were negative and regularly increased with temperature. These negative values indicate that the charge carriers are electrons, which are eg electrons belonging to Mn3+ in this case. 4. Discussion 4.1 Carrier density variation by Ni and Ti substitution Li1.05Mn2O4−δ exhibited the highest electrical conductivity at each temperature for all samples tested (Figure 3). Ni substitution was shown to reduce the conductivities, whereas the conductivities of Ti-substituted samples retained almost the same value as those of Li1.05Mn2O4−δ until the Ti content reached 0.2 equivalents and then slightly decreased. Li1.05Mn1.5Ni0.5O4−δ showed the lowest conductivity of all samples. As shown in Figure 4, Li1.05Mn2O4−δ also had the largest Seebeck coefficient among six samples. The Seebeck coefficients decreased with increasing Ni content, whereas Ti substitution barely affected the Seebeck coefficient values. The trends for electrical conductivity behaviors and Seebeck coefficients with the Ni- and Ti-substituted samples were quite similar. Since the electrical conductivity and the Seebeck coefficient depend on carrier density [45,46], they are not entirely independent of each other. For both systems, the electrical conductivity decreases and the absolute value of the Seebeck coefficient increases logarithmically with decreasing carrier density. The non-adiabatic small-polaron theory described the pre-exponential factor A as =
%
2ℏ
& ' (⁄ ) ⁄
(⁄
, (3)
where n is the carrier density, a is the hopping distance of small polarons, and J is the electron transfer integral between neighboring transition metals [24,27]. This equation indicates that A is directly proportional to the carrier density n and inversely proportional to WH1/2; thus, n contributes more to A than WH. The values of A for all samples, which were derived from the intercepts in the ln (σT3/2) vs. 1000/T Arrhenius plots, are summarized in Table 2. As the Ni content in the samples increased, A significantly decreased in value. In contrast, the A values in the Ti-substituted samples were almost independent of Ti content and were similar to that of Li1.05Mn2O4−δ. Given that Li1.05Mn2−xNixO4−δ is electrically neutral, the loss of an eg electron via the reduction of Mn3+ is required. For the Ti-doped sample, a reduction of Mn4+ is needed such that 4
Li1.05Mn2−xTixO4−δ will hold the same number of eg electrons as Li1.05Mn2O4−δ. There is no discrepancy between the variations of the number of eg electrons and the A values in Table 2, i.e., the carrier density n. Hence, the electrical conductivities and Seebeck coefficients associated with the Ni- and Ti-doped materials were primarily derived from variations in the carrier density n via Ni and Ti substitution.
4.2 Activation energies and hopping energies The activation energies of electrical conductivity Eσ and thermopower ES were derived from the slopes of the ln (σT3/2) vs. 1000/T Arrhenius (Figure 3) and S vs. 1000/T plots (Figure 4), respectively. Eσ and ES values are summarized in Table 2. Eσ contains ES and WH, as seen in equation (1). In addition to activation energies, we calculated the hopping energies WH using the relation WH = Eσ − ES (Table 2). The hopping energy WH corresponds to the minimum energy to produce the equivalency of the two neighboring sites involved in the polaron hopping [47-49]. A functionality of WH = 1/2WP – t, where WP is the polaron binding energy and 2t is the 3d-narrow bandwidth, is maintained for polaronic conduction [24]. As t is negligible in the non-adiabatic case [24], it is unlikely that the substitution of Ni and Ti for Mn expands the bandwidth 2t. The WH behavior of the Ni- and Ti- doped material, therefore, is predominantly governed by WP. To maintain electrical neutrality, Ni2+ displaced Mn3+ in the Ni-substituted samples. When the dopant content of Ni2+ increased, Mn3+ decreased and Mn4+ increased, affording Li1.05Mn1.5Ni0.5O4−δ with the minimum amount of Mn3+ of all Ni-substituted samples. According to Shannon [50], although Ni2+ has a larger ionic radius than Mn3+, Mn4+ has a smaller ionic radius than Mn3+. The average octahedral site cation radius decreases as increasing Ni content. This explained the decrease in lattice parameter and M-O distance accompanied by the substitution of Ni. The overlap in Mn 3d wave functions via O 2p wave functions can be enhanced due to the contraction of M-O distance. The carrier localization, hence, weakens and Wp decreases. Consequently, the WH values of Li1.05Mn2−xNixO4−δ (x = 0.2 and 0.3) became smaller than that of Li1.05Mn2O4−δ. For Li1.05Mn1.5Ni0.5O4−δ, the value of WH dramatically increased and showed the highest Eσ for all samples. When the Ni content is greater than 0.3, the rate of decrease in lattice parameter undergoes a drastic change and becomes smaller. We speculated that the state of Li1.05Mn1.5Ni0.5O4−δ might differ from that of Li1.05Mn2−xNixO4−δ (x = 0.2 and 0.3). Lee et al. [51] found a small degree of ordering in 16d sites in Li1.05Mn1.5Ni0.5O4−δ using fourier transform infrared spectroscopy, although the Rietveld refinement of the synchrotron X-ray diffraction pattern proved their Li1.05Mn1.5Ni0.5O4−δ to be a disordered spinel structure (Fd3m) without any long-range ordering. Infrared spectroscopy is generally recognized as a useful tool to qualitatively prove the cation ordering in 16d sites even if it is a small degree of ordering [52]. This localized inhomogeneity could explain why our Li1.05Mn1.5Ni0.5O4−δ exhibited the highest WH of all Ni-substituted samples. The Ti-substituted samples, in contrast, maintained nearly equivalent WH values to those of Li1.05Mn2O4−δ, and the electrical neutrality of the system replaced Mn4+ with Ti4+ in Li1.05Mn2−xTixO4−δ. Ti4+ has a larger ionic radius than Mn4+ and the lattice parameters of Ti-doped samples become larger. Ti substitution, however, keeps the M-O distances almost the same as those of Li1.05Mn2O4−δ, and it does not essentially affect the M-O bonds. Consequently, the WH values of Ti-doped samples did not differ from those of non-doped Li1.05Mn2O4−δ. 5
The activation energy of thermopower ES is the energy difference between identical lattice distortions with and without the carrier, i.e., the energy required to generate a free hopping polaron [47-49]. The ES values of Ni- and Ti-substituted samples slowly increased with dopant content. Because both substitution metals have a larger ionic radius than Mn described above, substituting these metals for Mn should increase a level of structural distortion into the lattice. This condition was supposed to exercise an influence on the increase in E S. Given the small observed effects of Ni and Ti substitution on Eσ, ES, and WH values, the contents of these dopants do not appear to cause a significant change to the electronic structure.
5. Conclusion In the present work, the spinel Li1.05Mn2O4−δ, Li1.05Mn2−xNixO4−δ, and Li1.05Mn2−xTixO4−δ of space group Fd3m were synthesized and their transport properties were analyzed. The ionic radii of these transition metals and the observed variation in the samples’ lattice parameters suggest that Ni substitution increases the overlap in Mn 3d wave functions, and Ti substitution introduces structural distortion. Electrical conductivities measured in the high-temperature range of 473 K to 723 K suggest that the primary mechanism for electrical conduction in all samples is a non-adiabatic hopping process by small polarons. The thermopower measurements demonstrated that all samples with and without substitution are n-type semiconductors and their Seebeck coefficients are thermally activated. The electrical conductivity and Seebeck coefficient values decreased with added Ni content, while the Ti-substituted samples maintained the same values as Li1.05Mn2O4−δ. The electrical neutrality in the Ni-substituted samples necessitated a loss of eg electrons from Mn3+, while Ti-substituted samples maintained the same eg electron density as that of Li1.05Mn2O4−δ. The electrical conductivity and Seebeck coefficient behaviors of the Ni- and Ti-substituted compounds were considered to be caused by a variation in the carrier density. Furthermore, the hopping energy WH of the Ni-substituted samples decreased because the overlap in Mn 3d wave functions increased and the carriers were less localized. An exception to this rule was observed for the sample with the lowest Mn3+ content, Li1.05Mn1.5Ni0.5O4−δ, which experienced a sudden increase in WH. Ti substitution induced structural distortion into the lattice, and it therefore had almost no effect on WH but a small effect on the activation energy of the thermopower ES. References [1] M. M. Thackeray, W. I. F. David, P. G. Bruce, and J. B. Goodenough, Mater. Res. Bull. 18 (1983) 461-472. [2] M. M. Thackeray, P. J. Johnson, L. A. de Picciotto, P. G. Bruce, and J. B. Goodenough, Mater. Res. Bull. 19 (1984) 179-187. [3] J. M. Tarascon and D. Guyomard, J. Electrochem. Soc. 138 (1991) 2864-2868. [4] D. Guyomard and J. M. Tarascon, J. Electrochem. Soc. 139 (1992) 937-948. [5] M. M. Thackeray, A. de Kock, M. H. Rossouw, D. Liles, R. Bittihn and D. Hoge, J. Electrochem. Soc. 139 (1992) 363-366. [6] A. Yamada, J. Solid State Chem. 122 (1996) 160-165. 6
[7] J.M. Tarascon, D. Guyomard, Electrochim. Acta 38 (1993) 1221-1231. [8] Y. Gao, J.R. Dahn, Solid State Ionics 84 (1996) 33-40. [9] B. Deng, H. Nakamura, and M. Yoshio, J. Power Sources 180 (2008) 864-868. [10] J-M. Tarascon, W. R. McKinnon, F. Coowar, T. N. Bowner, G. Amatucci, and D. Guyomard, J. Electrochem. Soc. 141 (1994) 1421-1431. [11] A. Yamada and M. Tanaka, Mater. Res. Bull. 30 (1995) 715-721. [12] K. Miura, A. Yamada, and M. Tanaka, Electrochimica Acta 41 (1996) 249-256. [13] J. Rodriguez-Carvajal, G. Rousse, C. Masquelier, and M. Hervieu, Phys. Rev. Lett. 81 (1998) 4660-4663. [14] G. Rousse, C. Masquelier, J. Rodriguez-Carvajal, E. Elkaim, J. -P. Lauriat, and J. L. Martinez, Chem. Mater. 11 (1999) 3629-3635. [15] K. Oikawa, T. Kamiyama, F. Izumi, B. C. Chakoumakos, H. Ikuta, M. Wakihara, J. Li, and Y. Matsui, Solid State Ionics 109 (1998) 35-41. [16] A. Yamada, J. Solid State Chem. 122, 160 (1996). [17] A. Yamada and M. Tanaka, Mater. Res. Bull. 30, 715 (1995). [18] J. B. Goodenough, A. Manthiran, P. Wnetrzewski, J. Power Sources 43 (1993) 269-275. [19] L. Schutte, G. Colsmann, and B. Reuter, J. Solid State Chem. 27 (1979) 227-231. [20] V. Massarotti, D. Capsoni, M. Bini, G. Chiodelli, C. B. Azzoni, M. C. Mozzati, A. Paleari, J. Solid State Chem. 131 (1997) 94-100. [21] Y. Shimakawa, T. Numata, and J. Tabuchi, J. Solid State Chem. 131 (1997) 138-143. [22] J. Sugiyama, T. Atsumi, A. Koiwai, T. Sasaki, T. Hioki, S. Noda, N. Kamegashira, J. Phys.: Condens. Matter 9 (1997) 1729-1742. [23] T. Holstein, Ann. Phys. (N.Y.) 8 (1959) 343-389. [24] I. G. Austin and N. F. Mott, Adv. Phys. 18 (1969) 41-102. [25] E. Iguchi, N. Nakamura, H. Nakatsugawa, and A. Aoki, Philos. Mag. B 78 (1998) 65-77. [26] E. Iguchi, Y. Tokuda, H. Nakatsugawa, and F. Munakata, J. Appl. Phys. 91 (2002) 2149-2154. [27] D. Emin, and T. Holstein, Ann. Phys. 53 (1969) 439-520. [28] H. Flölich, Theory of Dielectric (Clarendon, Oxford, 1958), p. 70. [29] T. Ohzuku, S. Takeda, M. Iwanaga, J. Power Sources 81–82 (1999) 90-94. [30] J. Molenda, J. Marzec, K. S´wierczek, W. Ojczyk, M. Ziemnicki, P. Wilk, M. Molenda, M. Drozdek, R. Dziembaj, Solid State Ionics 171 (2004) 215-227. [31] M. Wakihara, L. Guohua, H. Ikuta, T. Uchida, Solid State Ionics 86–88 (1996) 907-909. [32] D. Capsoni, M. Bini, G. Chiodell, V. Massarotti, P. Mustarell, L. Linati, M.C. Mozzati, C.B. Azzoni, Solid State Commun. 126 (2003) 169-174. [33] J.-H. Kim, S.-T. Myung, C. S. Yoon, I.-H. Oh, and Y.-K. Sun, J. Electrochem. Soc. 151 (2004) A1911-1918. [34] J. B. Goodenough and Y. Kim, Chem. Mater. 22 (2010) 587-603. [35] J. Song, D. W. Shin, Y. Lu, C. D. Amos, A. Manthiram, and J. B. Goodenough, Chem. Mater. 24 (2012) 3101-3109. 7
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8
Table 1. Mn valency, space groups, lattice parameters, M–O distances, and Rietveld factors (Rwp and RB) for Li1.05Mn2O4−δ, Li1.05Mn1.7Ni0.3O4−δ, and Li1.05Mn1.8Ti0.2O4−δ. Sample
Li1.05Mn2O4−δ
Li1.05Mn1.7Ni0.3O4−δ
Li1.05Mn1.8Ti0.2O4−δ
Mn valency
3.52
3.70
3.35
Space group
Fd3m
Fd3m
Fd3m
Lattice parameter (Å)
8.24538(5)
8.19287(2)
8.26125(5)
M-O distance (Å)
1.9571(8)
1.9496(4)
1.9582(5)
Rwp
5.52
3.57
4.04
RB
2.16
3.45
1.74
Rietveld factors (%)
10
15
20
25
30
35
40
45
50
55
60
65
70
75
2θ (deg)
Fig. 1. Observed and calculated synchrotron X-ray diffraction patterns at room temperature: (a) Li1.05Mn2O4−δ, (b) Li1.05Mn1.7Ni0.3O4−δ, and (c) Li1.05Mn1.8Ti0.2O4−δ.
9
8.28
Lattice parameter (Å)
8.26 8.24 8.22 8.20 Li1.05Mn2O4-δ Li1.05Mn2-xNixO4-δ Li1.05Mn2-xTixO4-δ
8.18 8.16
0.0
0.1
0.2
0.3
0.4
0.5
x in Li1.05Mn2-xMxO4-δ
Fig. 2. Lattice parameters as a function of dopant content for Li1.05Mn2O4, Ni-substituted sample Li1.05Mn2−xNixO4−δ, and Ti-substituted sample Li1.05Mn2−xTixO4−δ.
11 Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
10 9 8 7 6 5 4 3 1.2
1.4
1.6
1.8
2.0
2.2
2.4
-1
1000/T (K )
Fig. 3. Arrhenius plots of ln(σT3/2) vs. 1000/T for Li1.05Mn2 O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3).
10
50
Li 1.05Mn2O4-δ Li 1.05Mn2-xNixO4-δ (x=0.2) Li 1.05Mn2-xNixO4-δ (x=0.3) Li 1.05Mn2-xNixO4-δ (x=0.5) Li 1.05Mn2-xTixO4-δ (x=0.1) Li 1.05Mn2-xTixO4-δ (x=0.2) Li 1.05Mn2-xTixO4-δ (x=0.3)
0 -50 -100 -150 -200 -250 -300 -350 1.2
1.4
1.6
1.8
2.0
2.2
2.4
-1
1000/T (K )
Fig. 4. Seebeck coefficients vs. 1000/T for Li1.05Mn2 O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3).
Table 2. Pre-exponential factors and activation energies Eσ of electrical conductivity, activation energies ES of thermopower, and hopping energies WH for Li1.05Mn2O4−δ, Li1.05Mn2−xNix O4−δ (x = 0.2, 0.3, and 0.5), and Li1.05Mn2−xTix O4−δ (x = 0.1, 0.2, and 0.3). Sample
Composition
Li1.05Mn2.0-xTixO4−δ
Thermoelectric power
Hopping energy
Pre-exponential factor
Eσ
ES
WH
(Ω-1cm-1K3/2)
(eV)
(eV)
(eV)
5.1×106
0.41
0.05
0.36
x=0.2
2.6×106
0.41
0.07
0.33
x=0.3
1.7×106
0.41
0.08
0.34
x=0.5
1.6×106
0.45
0.08
0.37
x=0.1
5.8×106
0.42
0.06
0.36
x=0.2
5.3×106
0.42
0.06
0.36
x=0.3
4.3×106
0.43
0.07
0.36
Li1.05Mn2O4−δ Li1.05Mn2.0-xNixO4−δ
Electrical conductivity
11
Graphical abstract legend (TOC figure)
The figure on the left is Arrhenius plots of ln(σT3/2) vs. 1000/T, and the figure on the right is Seebeck coefficients vs. 1000/T.
Highlights · Ni- and Ti- doped lithium manganese spinels were synthesized. · The doping effect on structure and transport properties was studied systematically. · Ni- and Ti-doped samples followed a non-adiabatic small-polaron-hopping mechanism. · Ni dopant decreased the hopping energy and Ti dopant induced structural distortion.