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Electrochemical properties of Li2MnO3 nanowires with polycrystalline and monocrystalline states
Journal of Alloys and Compounds 686 (2016) 496e502
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Journal of Alloys and Compounds journal homepage: http://www.elsevier.com/locate/jalcom
Electrochemical properties of Li2MnO3 nanowires with polycrystalline and monocrystalline states Meng Cheng a, b, Kongjun Zhu a, *, Lu Yang a, b, Lei Zhu c, Yong Li c, Weiping Tang c, ** a
State Key Laboratory of Mechanics and Control of Mechanical Structures, College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, PR China b College of Materials Science and Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China c Shanghai Institute of Space Power-Sources, Shanghai 200000, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 March 2016 Received in revised form 1 June 2016 Accepted 9 June 2016 Available online 11 June 2016
In this article, monocrystalline bundles and polycrystalline nanowires were synthesized using a-MnO2 and MnOOH as precursors via an in-situ reaction. The electrochemical properties of the resulting nanocrystals were investigated. Results reveal that the polycrystalline nanowires are more activated than the monocrystalline nanowires. The polycrystalline nanowires also deliver a maximum discharge capacity of 226 mA h/g at 10 mA/g. Compared with the linear charging plateau at >4.3 V in the monocrystalline bundles, the charging platues of (100) and (001) domains in the polycrystalline nanowires are 4.3e4.55 and 4.55e4.6 V, respectively. The high capacity and rate performance of the polycrystalline nanowires are attributed to their good electroconductivity and rapid Liþ diffusion, which originate from defects and small domains. © 2016 Elsevier B.V. All rights reserved.
Keywords: Li2MnO3 Nanowires Grwoth patterns Electrochemical property Liþ diffusion
1. Introduction As a promising choice for powering electric vehicles, rechargeable Li-ion batteries have been extensively investigated and commonly used in the modern society [1e4]. A high specific capacity of cathode materials is required to yield a high energy density of Li-ion batteries. However, the specific capacities of commercial oxide cathode materials, such as LiMn2O4 and LiMO2 (M ¼ Ni, Mn, Co) species, are restricted to less than 200 mA h/g, which largely limits the energy density of Li-ion batteries [5e7]. Considering that Li-rich Mn-based materials present high specific capacities (>250 mA h/g), entail low raw material costs, and provide several other advantages, researchers proposed that Li-rich cathode materials may be used as an alternative to Li-ion batteries for electric vehicles and grid applications [8,9]. Nevertheless, several issues, including performances and safety, should be considered in practical applications; thus, Li-rich materials and Li2MnO3 should be further elucidated. Li2MnO3 presents an a-NaFeO2-layered structure, in which 1/3
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (K. Zhu), [email protected] (W. Tang). http://dx.doi.org/10.1016/j.jallcom.2016.06.081 0925-8388/© 2016 Elsevier B.V. All rights reserved.
Li and 2/3 Mn occupy the 3d position of transition metal layers; as a result, a Li-Mn ordered super structure is formed [10e12]. The electrochemical activation of this material may be attributed to oxidation mechanisms involving Li2O removal and structural reorganization, which are influenced by grain morphology, size, and effects on performance [13e16]. Morphological characteristics and growth processes significantly affect the activation mechanism and electrochemical properties of materials [17e23]. Cho et al. [17,18] synthesized Li0.88[Li0.18Co0.33Mn0.49]O2 and Li[Ni0.25Li0.15Mn0.6]O2 nanowires and 0D Li2MnO3 particles via a hydrothermal process; the capabilities of the resulting materials are closely related to the morphology. For instance, Li[Ni0.25Li0.15Mn0.6]O2 nanowires exhibit 311 mA h/g at 0.3C and a rate capability of 95% at 4C; these parameters indicate that these nanowires exhibit properties superior to nanoplate samples [16]. Wei [19] found that (010) nanoplates were characterized by greater reversible capacity, cycle ability, and rate performance than (001) nanoplates. In the literature reports of Mai’ team, SnO2 quantum dots yielded superior capabilities than nanorods [22,23]. Despite these previous findings, Li2MnO3 particles with varied grwoth patterns and electrochemical properties have yet to be developed. As such, we aim to fill this gap. Our previous paper reported grain size control and its effect on
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the electrochemical properties of Li2MnO3 materials [15]. In this current article, Li2MnO3 nanowires with monocrystalline and polycrystalline states are synthesized, and their capabilities are investigated. The effect of the grwoth patterns on dynamic mechanism is elucidated by investigating the corresponding electrochemical properties upon cycling and by determining Liþ diffusion coefficients. 2. Experimental 2.1. Preparation of Li2MnO3 nanocrystals S1: Li2MnO3 bundles were prepared via a molten-salt growth method. The precursor P1 was prepared via a hydrothermal process [24]. MnSO4 (3 mmol), (NH4)2SO4 (12 mmol), and (NH4)2S2O8 (3 mmol) were dissolved in 30 mL of deionized water in a 50 mL Teflon liner to obtain a mixed solution. The Teflon liner was transferred to an autoclave and maintained at 140 C for 24 h. P1 powders were filtered, washed, and dried at 80 C for 24 h. The obtained precursor (1 g) was dispersed in ethanol and 10 g of LiNO3 was added after 30 min of stirring. The mixture was stirred at 60 C until ethanol evaporated. The obtained dry powders were ground, transferred to a crucible, and calcined at 450 C for 5 h and 550 C for 2 h in air. After the powders were washed and dried, sample S1 was obtained. S2: The molten-salt growth method was applied to prepare S2 powders. P2 was also prepared via a hydrothermal process [25]. KMnO4 (3mmol) and glycol (2mL) were dissolved in 160 mL of deionized water and stirred for 30 min to yield a mixture in a 200 mL Teflon liner. The Teflon liner was transferred to an autoclave and maintained at 120 C for 48 h. P2 powders were filtered, washed, and dried at 80 C for 24 h. The obtained precursor (1 g) was dispersed in ethanol and 10 g of LiNO3 was added after 30 min of stirring. The mixture was stirred at 60 C until ethanol evaporated. The obtained dry powders were ground, transferred to a crucible, and calcined at 400 C for 5 h in air. After the powders were washed, dried, and heated at 450 C for 5 h and at 550 C for 2 h in air, sample S2 was obtained. 2.2. Characterization of the Li2MnO3 products XRD measurements were conducted using a Rigaku D/max2600PC with Cu Ka radiation (l ¼ 0.15406 mm). Particle morphologies were probed by SEM (S-4800) and TEM (JEM-2100F). Raman spectroscopy (InVia-Reflex) was performed to analyze the microstructure of the Li2MnO3 samples. Specific surface areas were determined using a specific surface area and pore size analyzer (ASAP 2020 M). 2.3. Electrochemical measurement The cathode was prepared by thoroughly mixing the active material, super P carbon black, and a polyvinylidene fluoride binder at a weight ratio of 8:1:1. Coin-type battery cells containing a cathode, a Li metal anode, and a microporous polyethylene separator were prepared in a nitrogen-filled glove-box. Electrochemical measurements were conducted with button-type cells by using 1.2 M LiPF6 with ethyl methyl carbonate/ethylene carbonate (7:3, vol. %) as electrolyte. The coin-type cells were evaluated at a constant current density of 20 mA/g by using a Land battery test system. CV curves were performed by Princeton 273A (0.1 mV/s, 2e4.6 V). Electrochemical impedance spectroscopy (EIS) measurements were performed in the frequency range 102-106Hz by VMP3.
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3. Results and discussion 3.1. Structural characterization of the Li2MnO3 nanowires The phase compositions of P1, P2, S1, and S2 were analyzed through XRD in Fig. 1(a). The diffraction lines of P1 and P2 were in accordance with PDF#44-0141 and PDF#88-0649, which were characterized as a-MnO2 and MnOOH, respectively. The XRD patterns of S1 and S2 showed single Li2MnO3 phase without secondary Bragg peaks. The diffraction lines of Li2MnO3 were indexed to a monoclinic unit cell with the space group symmetry of C2/m in S1 and S2. Compared with the peaks in S2, the observed broadened peaks with lower intensity in S1 might be relevant to the particle size, as revealed by TEM characterization. XRD further confirmed that pure Li2MnO3 phase could be produced by using a-MnO2 and MnOOH as precursors via the molten-salt method. Fig. 1(b) illustrated the Raman spectra of S1 and S2. The Raman spectra of both samples exhibited similar bands located at 248, 325, 370, 415, 438, 495, 570, and 613 cm1, which were characteristics of the C2/m phase. These bands belonged to the Mn-O bonding of Li2MnO3. These bands were also consistent with those described in previous reports. The bands were then assigned to phonon vibrations, and the peak at 613 cm1 was attributed to the A1g mode of Mn-O bonding. The samples exhibited minor changes of the Raman peaks locations, full width-at-half maximum (W), and the ratio of the main peaks absolute intensities (I613/I495). Obviously, there was not notably distinction of local crystal structure in S1 and S2 samples [16]. The weaker peak density in S1 sample referred to small particle size, which was coincided with the result of XRD patterns. The combination of XRD and Raman spectra revealed that S1 and S2 were purely dominated by Li2MnO3. 3.2. Morphology of the Li2MnO3 nanowires The morphological characteristics of the Li2MnO3 products were examined through SEM and TEM (Fig. 2). P1 and S1 [Fig. 2(a) and 2(b)] exhibited a bundle morphology, assembled by rods and short wires with 30 nm thickness. In Fig. 2(c), the spacing of the crystal face of approximate 0.47 nm was observed in S1; this value corresponded to the (001) plane. The monocrystalline morphology was evidenced by the SAED pattern. P2 and S2 exhibited similar nanowire morphology with a length of 6 mm and a thickness of 80 nm [Fig. 2(d) and 2(e)]. The specific surface areas of S1 and S2 were 58 and 13 m2/g, respectively. These findings indicated that the particle size of S2 was larger than that of S1. In the high-resolution TEM images, the S2 crystalline composed of small crystalline domains (be separated by dash lines), which exhibited polycrystalline morphology evidenced by the SAED pattern. In the left, the two crystal faces yield spacing of approximate 0.47 and 0.83 nm, corresponding to the (001) plane and cell parameter b of Li2MnO3, respectively. The left region was indexed to the [001] axis, which was referred to as (100) domain. The middle region contained two evident planes, with 0.41 and 0.43 nm spacing corresponding to (110) and (020) planes. The middle region was indexed to the [001] axis, which was referred to as the (100) domain. The right region was similar to the (001) domain. In this regard, two domains piled along the diameter direction and formed nanowire morphology. In summary, HR-TEM and SAED patterns revealed that monocrystalline Li2MnO3 particles assembled in S1; small crystalline domains aggregated and formed polycrystalline nanowire in S2. XRD patterns and SEM images indicated that S1 and S2 were generated through an in situ reaction with a-MnO2 and MnOOH as precursors. Despite the distinct structure and valence values of Mn, a pure Li2MnO3 phase could be obtained via the molten-salt method.
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Fig. 1. XRD patterns (a) and Raman spectra (b) of S1 and S2.
Fig. 2. SEM and TEM images of P1 (a), P2 (d), S1 (b, c), and S2 (e, f) samples.
3.3. Electrochemical properties of Li2MnO3 nanowires The first charge/discharge voltage profiles of S1 and S2 were determined at a constant current density of 10 mA/g in Fig. 3(a). The initial discharge capacities were 178 and 226 mA h/g for S1 and S2, respectively. The first-cycle coulombic efficiencies of S1 and S2 were 68.4% and 67%, respectively. Thus, approximately 78% Liþ returned to the original position. S1 exhibited a linear charge plateau above 4.3 V. By contrast, the charge plateau of S2 contained two sections at 4.3e4.55 V and 4.55e4.6 V. The lower plateau might contribute to the (100) domain, where Liþ was extracted perpendicular to the [001] direction, because the plateau voltage corresponded to the energy of Liþ extracted from the lattice. Fig. 3(b) illustrated the cycle performance of S1 and S2 at a rate of 10 mA/g (1st-50th). S1 and S2 electrodes showed poor cycle performances. The discharge capacity retentions of S1 and S2 were 70.6% and 91.7% in the 6th cycle and 29% and 27.1% in the 50th cycle, respectively. A logarithmic degradation was observed in the S1 electrode before the 30th cycle; the S2 capacity demonstrated good
retention before the 6th cycle and then rapidly faded in the 6th30th cycles. In general, S2 was more active and might exhibit a distinct mechanism of capacity degradation than S1. The rate performance was determined at 10 mA/g, which was increased to 200 mA/g stepwise. S2 showed a stable capacity at a high current rate. A drastic decrease in the specific capacity was observed in S1 cycled at a high current rate. The varied cycle and rate performance might be partly attributed to the activation of the particles. S1 and S2 were subjected to CV measurements [Fig. 4(a) and (b)]. The redox peaks exhibited evident distinctions in terms of location and strength. For S1, four distinct peaks were observed in oxidation and reduction reactions during the cycles. The oxidation peaks near 4.6 V represented Li2O extraction from the Li2MnO3 structure [14,15]. The 4.2 V peak after the first cycle might arise from oxygen. The reduction peaks near 3.3 and 2.95 V were attributed to Mn4þ/3þ and Mn3þ/2þ reduction in the layered structure, respectively. The increasing intensity of the Mn4þ/3þ peak illustrated the activation of the Li2MnO3 structure upon cycling. The strength of the Mn3þ/2þ
M. Cheng et al. / Journal of Alloys and Compounds 686 (2016) 496e502
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Fig. 3. Charge/discharge curves (a), cycle performance (b), and rate performance (c) of the S1 and S2 samples.
Fig. 4. CV curves of the S1 (a) and S2 (b) samples.
peak also demonstrated an uptrend when cycling was extended. In comparison to the locations of the peaks in S1, the locations of the Mn3þ-Mn2þ redox peaks shifted in S2 sample. The characteristic spinel peak at 2.8 and 3.0 V indicated the formation of a spinel phase upon cycling. More spinel phases might be generated in the powders with a larger particle size [15]. The sharper Mn3þ-Mn2þ redox peaks implied that more Mn2þ participated in electrochemical reaction, which might contribute to capacity degradation. Consistent with the cycle performance, phase transformation and Mn2þ participation might degrade the capacity of S2 upon cycling. With the aim to analyze the structural evolution of the samples, XRD patterns, TEM images and SAED patterns were carried out. In Fig. S1, XRD patterns of the electrodes represented that minor spinel phase formed after the initial cycle, with well-maintained layered structure. The peak area ratios of spinel (400) peak and
layered (-202) peak implied more spinel phase in S2 electrode after the initial cycle. Some distortion structure was observed on the edge of the crystalline in both samples. Spinel (400) plane was detected in the SAED pattern of the S2 electrode. XRD patterns and TEM images indicated that spinel phase, distortion structure and Jahn Teller effect might cause the capacity degradation [28]. More spinel phase and Mn2þ participation might have an influence on the mechanism of capacity degradation, which needed more experiments and analysis. Electrochemical impedance spectroscopy (EIS) was performed to identify the changes in the impedance after the initial discharge (V vs. Liþ/Li) and clarify electrochemical degradation upon cycling. In Fig. 5, the semicircle in the high-middle frequency region was attributed to the lithium-ion migration through the SEI film (Rf and Cf) and charge transfer reaction (Rct and Cdl), and the slope in the
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3.4. Liþ chemical diffusion coefficients Fig. 6 illustrated the open-circuit voltage (OCV) curves of S1 and S2. In the initial charge stage of S1 and S2, the evident over potential, especially above 4.3 V, suggested that the reaction dynamics of Li2MnO3 activation was more sluggish than that of Li extraction on the surface below 4.4 V. The over potential of 4.3e4.55 V and 4.55e4.6 V plateau indicated the distinct activation kinetics of the varied regions in S2. DLiþ could be calculated by the following simplified equation when the relationship of the cell voltage versus S1/2 was approximately linear [27]:
DLiþ ¼
4
mVM pt MS
Fig. 5. Nyquist plots of S1 and S2.
2
DEs DEt
2
Fig. 6. GITT curves and Liþ chemical diffusion coefficients in the initial cycle.
Table 1 The morphologies, BET, capacities, Rct, and DLiþ of S1 and S2. Samples S1 S2
Morphology Monocrystalline bundle Polycrystalline nanowire
BET 2
58 m /g 13 m2/g
low frequency region was attributed to the lithium-ion diffusion in the bulk electrodes (w) [26]. Corresponding Rct of S1 and S2 were 581 and 534U after the initial discharge, which were 139 and 187U in fresh cells. After 3 cycles, Rcts were increased to 2093 and 1557U for S1 and S2 electrodes. The lower Rct of S2 demonstrated better electroconductivity presumably as a consequence of crystal boundary and stacking defects, which were necessary to obtain high capacity and rate performance.
Capacity (10 mA/g)
Rct (after 3st cycle)
Dþ Li (after 1st cycle)
178 mA h/g 226 mA h/g
2093 U 1557 U
1016 cm2/s 1015 cm2/s
Where VM was the molar volume of active materials; M and m are the molecular weight and the mass of the active material, respectively; S was the surface area of the interface between an electrode and an electrolyte; and DEs and DEt were the physical interpretations in Fig. S3. The calculated DLiþ as a function of OCV in the initial cycle was shown in Fig. 6. In the initial charge of the S1 electrode, DLiþ
M. Cheng et al. / Journal of Alloys and Compounds 686 (2016) 496e502
decreased remarkably as its capacity increased and reached a minimal value of 1018 cm2/s at approximately 20% capacity (approximately 4.4 V). The DLiþ of Li2MnO3 at the activation plateau was as low as 1018 cm2/s. This finding suggested that the sluggish kinetics was an activation bottleneck of Li2MnO3. By comparison, S2 exhibited a DLiþ of 1016 cm2/s from 10% to 45% capacity and approximately 1017 cm2/s beyond 45% capacity; these findings corresponded to the two sections of 4.3e4.55 and 4.55e4.6 V plateau. During discharging, DLiþ demonstrated a downtrend as OCV decreased because of the increased concentration polarization accompanied by the insertion reaction of Li-ions in S1 and S2. Compared with the value obtained via the charging process, the DLiþ of S1 and S2 increased to 1016 and 1015 cm2/s, respectively. These values indicated that the structure was activated. The large DLiþ of S2 might be attributed to the grwoth patterns. Crystal boundaries and defects could introduce active reaction points and small domains with varied orientation, which were beneficial for Li extraction and accommodates the volume change in the crystal cell [15]. Liþ atoms extracted along the direction perpendicular to the c axis of the layered structure cathode materials. In the monocrystalline bundles, Liþ ions were directly delivered in [001] direction. While in polycrystalline nanowires, Liþ must reach the surface or the grain boundary along the direction perpendicular to the c axis, and then migrated on the surface or the grain boundary. Thus, the larger particle size and polycrystalline morphology endowed the S2 sample with longer Liþ-extraction path. Furthermore the low energy barrier of the crystal boundaries and stacking defects provided more active points for electrochemical reaction, which facilitated Liþ-extraction. In summary, the crystal boundaries and stacking defects in polycrystalline nanowires enormously promoted Liþ chemical diffusion coefficients. The morphologies, BET, capacities, Rct, and DLiþ of S1 and S2 were listed in Table 1. Despite the larger particle size observed in S2, the crystal boundaries, stacking defects, and small domains with varied orientation could enhance the electroconductivity and Liþ diffusion rate; thus, discharge capability, especially at high rates, could be improved.
4. Conclusion Monocrystalline bundles and polycrystalline nanowires with a pure layered structure were synthesized via an in-situ reaction. The polycrystalline nanowires were more activated than the monocrystalline nanowires. The former also delivered higher capacity and rate performance than the latter. Monocrystalline bundles exhibited a linear charging plateau at 4.3e4.6 V. By comparison, the charging plateau of the polycrystalline nanowires contained two sections at 4.3e4.55 and 4.55e4.6 V, which corresponded to (100) and (001) domains. The stacking defects and grain boundary introduced more active points for the electrochemical reaction. Higher electroconductivity and Liþ diffusion coefficient endowed polycrystalline nanowires with enhanced capability.
Acknowledgements This work was supported by the National Nature Science Foundation of China (NSFC No. 51372114), the Research Fund of State Key Laboratory of Mechanics and Control of Mechanical Structures (Nanjing University of Aeronautics and astronautics) (Grant No. 0514Y01), the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD).
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Electrochemical properties of Li2MnO3 nanowires with polycrystalline and monocrystalline states Meng Cheng a, b, Kongjun Zhu a, *, Lu Yang a, b, Lei Zhu c, Yong Li c, Weiping Tang c, ** a
State Key Laboratory of Mechanics and Control of Mechanical Structures, College of Aerospace Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, PR China b College of Materials Science and Engineering, Nanjing University of Aeronautics and Astronautics, Nanjing 210016, China c Shanghai Institute of Space Power-Sources, Shanghai 200000, China
a r t i c l e i n f o
a b s t r a c t
Article history: Received 15 March 2016 Received in revised form 1 June 2016 Accepted 9 June 2016 Available online 11 June 2016
In this article, monocrystalline bundles and polycrystalline nanowires were synthesized using a-MnO2 and MnOOH as precursors via an in-situ reaction. The electrochemical properties of the resulting nanocrystals were investigated. Results reveal that the polycrystalline nanowires are more activated than the monocrystalline nanowires. The polycrystalline nanowires also deliver a maximum discharge capacity of 226 mA h/g at 10 mA/g. Compared with the linear charging plateau at >4.3 V in the monocrystalline bundles, the charging platues of (100) and (001) domains in the polycrystalline nanowires are 4.3e4.55 and 4.55e4.6 V, respectively. The high capacity and rate performance of the polycrystalline nanowires are attributed to their good electroconductivity and rapid Liþ diffusion, which originate from defects and small domains. © 2016 Elsevier B.V. All rights reserved.
Keywords: Li2MnO3 Nanowires Grwoth patterns Electrochemical property Liþ diffusion
1. Introduction As a promising choice for powering electric vehicles, rechargeable Li-ion batteries have been extensively investigated and commonly used in the modern society [1e4]. A high specific capacity of cathode materials is required to yield a high energy density of Li-ion batteries. However, the specific capacities of commercial oxide cathode materials, such as LiMn2O4 and LiMO2 (M ¼ Ni, Mn, Co) species, are restricted to less than 200 mA h/g, which largely limits the energy density of Li-ion batteries [5e7]. Considering that Li-rich Mn-based materials present high specific capacities (>250 mA h/g), entail low raw material costs, and provide several other advantages, researchers proposed that Li-rich cathode materials may be used as an alternative to Li-ion batteries for electric vehicles and grid applications [8,9]. Nevertheless, several issues, including performances and safety, should be considered in practical applications; thus, Li-rich materials and Li2MnO3 should be further elucidated. Li2MnO3 presents an a-NaFeO2-layered structure, in which 1/3
* Corresponding author. ** Corresponding author. E-mail addresses: [email protected] (K. Zhu), [email protected] (W. Tang). http://dx.doi.org/10.1016/j.jallcom.2016.06.081 0925-8388/© 2016 Elsevier B.V. All rights reserved.
Li and 2/3 Mn occupy the 3d position of transition metal layers; as a result, a Li-Mn ordered super structure is formed [10e12]. The electrochemical activation of this material may be attributed to oxidation mechanisms involving Li2O removal and structural reorganization, which are influenced by grain morphology, size, and effects on performance [13e16]. Morphological characteristics and growth processes significantly affect the activation mechanism and electrochemical properties of materials [17e23]. Cho et al. [17,18] synthesized Li0.88[Li0.18Co0.33Mn0.49]O2 and Li[Ni0.25Li0.15Mn0.6]O2 nanowires and 0D Li2MnO3 particles via a hydrothermal process; the capabilities of the resulting materials are closely related to the morphology. For instance, Li[Ni0.25Li0.15Mn0.6]O2 nanowires exhibit 311 mA h/g at 0.3C and a rate capability of 95% at 4C; these parameters indicate that these nanowires exhibit properties superior to nanoplate samples [16]. Wei [19] found that (010) nanoplates were characterized by greater reversible capacity, cycle ability, and rate performance than (001) nanoplates. In the literature reports of Mai’ team, SnO2 quantum dots yielded superior capabilities than nanorods [22,23]. Despite these previous findings, Li2MnO3 particles with varied grwoth patterns and electrochemical properties have yet to be developed. As such, we aim to fill this gap. Our previous paper reported grain size control and its effect on
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the electrochemical properties of Li2MnO3 materials [15]. In this current article, Li2MnO3 nanowires with monocrystalline and polycrystalline states are synthesized, and their capabilities are investigated. The effect of the grwoth patterns on dynamic mechanism is elucidated by investigating the corresponding electrochemical properties upon cycling and by determining Liþ diffusion coefficients. 2. Experimental 2.1. Preparation of Li2MnO3 nanocrystals S1: Li2MnO3 bundles were prepared via a molten-salt growth method. The precursor P1 was prepared via a hydrothermal process [24]. MnSO4 (3 mmol), (NH4)2SO4 (12 mmol), and (NH4)2S2O8 (3 mmol) were dissolved in 30 mL of deionized water in a 50 mL Teflon liner to obtain a mixed solution. The Teflon liner was transferred to an autoclave and maintained at 140 C for 24 h. P1 powders were filtered, washed, and dried at 80 C for 24 h. The obtained precursor (1 g) was dispersed in ethanol and 10 g of LiNO3 was added after 30 min of stirring. The mixture was stirred at 60 C until ethanol evaporated. The obtained dry powders were ground, transferred to a crucible, and calcined at 450 C for 5 h and 550 C for 2 h in air. After the powders were washed and dried, sample S1 was obtained. S2: The molten-salt growth method was applied to prepare S2 powders. P2 was also prepared via a hydrothermal process [25]. KMnO4 (3mmol) and glycol (2mL) were dissolved in 160 mL of deionized water and stirred for 30 min to yield a mixture in a 200 mL Teflon liner. The Teflon liner was transferred to an autoclave and maintained at 120 C for 48 h. P2 powders were filtered, washed, and dried at 80 C for 24 h. The obtained precursor (1 g) was dispersed in ethanol and 10 g of LiNO3 was added after 30 min of stirring. The mixture was stirred at 60 C until ethanol evaporated. The obtained dry powders were ground, transferred to a crucible, and calcined at 400 C for 5 h in air. After the powders were washed, dried, and heated at 450 C for 5 h and at 550 C for 2 h in air, sample S2 was obtained. 2.2. Characterization of the Li2MnO3 products XRD measurements were conducted using a Rigaku D/max2600PC with Cu Ka radiation (l ¼ 0.15406 mm). Particle morphologies were probed by SEM (S-4800) and TEM (JEM-2100F). Raman spectroscopy (InVia-Reflex) was performed to analyze the microstructure of the Li2MnO3 samples. Specific surface areas were determined using a specific surface area and pore size analyzer (ASAP 2020 M). 2.3. Electrochemical measurement The cathode was prepared by thoroughly mixing the active material, super P carbon black, and a polyvinylidene fluoride binder at a weight ratio of 8:1:1. Coin-type battery cells containing a cathode, a Li metal anode, and a microporous polyethylene separator were prepared in a nitrogen-filled glove-box. Electrochemical measurements were conducted with button-type cells by using 1.2 M LiPF6 with ethyl methyl carbonate/ethylene carbonate (7:3, vol. %) as electrolyte. The coin-type cells were evaluated at a constant current density of 20 mA/g by using a Land battery test system. CV curves were performed by Princeton 273A (0.1 mV/s, 2e4.6 V). Electrochemical impedance spectroscopy (EIS) measurements were performed in the frequency range 102-106Hz by VMP3.
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3. Results and discussion 3.1. Structural characterization of the Li2MnO3 nanowires The phase compositions of P1, P2, S1, and S2 were analyzed through XRD in Fig. 1(a). The diffraction lines of P1 and P2 were in accordance with PDF#44-0141 and PDF#88-0649, which were characterized as a-MnO2 and MnOOH, respectively. The XRD patterns of S1 and S2 showed single Li2MnO3 phase without secondary Bragg peaks. The diffraction lines of Li2MnO3 were indexed to a monoclinic unit cell with the space group symmetry of C2/m in S1 and S2. Compared with the peaks in S2, the observed broadened peaks with lower intensity in S1 might be relevant to the particle size, as revealed by TEM characterization. XRD further confirmed that pure Li2MnO3 phase could be produced by using a-MnO2 and MnOOH as precursors via the molten-salt method. Fig. 1(b) illustrated the Raman spectra of S1 and S2. The Raman spectra of both samples exhibited similar bands located at 248, 325, 370, 415, 438, 495, 570, and 613 cm1, which were characteristics of the C2/m phase. These bands belonged to the Mn-O bonding of Li2MnO3. These bands were also consistent with those described in previous reports. The bands were then assigned to phonon vibrations, and the peak at 613 cm1 was attributed to the A1g mode of Mn-O bonding. The samples exhibited minor changes of the Raman peaks locations, full width-at-half maximum (W), and the ratio of the main peaks absolute intensities (I613/I495). Obviously, there was not notably distinction of local crystal structure in S1 and S2 samples [16]. The weaker peak density in S1 sample referred to small particle size, which was coincided with the result of XRD patterns. The combination of XRD and Raman spectra revealed that S1 and S2 were purely dominated by Li2MnO3. 3.2. Morphology of the Li2MnO3 nanowires The morphological characteristics of the Li2MnO3 products were examined through SEM and TEM (Fig. 2). P1 and S1 [Fig. 2(a) and 2(b)] exhibited a bundle morphology, assembled by rods and short wires with 30 nm thickness. In Fig. 2(c), the spacing of the crystal face of approximate 0.47 nm was observed in S1; this value corresponded to the (001) plane. The monocrystalline morphology was evidenced by the SAED pattern. P2 and S2 exhibited similar nanowire morphology with a length of 6 mm and a thickness of 80 nm [Fig. 2(d) and 2(e)]. The specific surface areas of S1 and S2 were 58 and 13 m2/g, respectively. These findings indicated that the particle size of S2 was larger than that of S1. In the high-resolution TEM images, the S2 crystalline composed of small crystalline domains (be separated by dash lines), which exhibited polycrystalline morphology evidenced by the SAED pattern. In the left, the two crystal faces yield spacing of approximate 0.47 and 0.83 nm, corresponding to the (001) plane and cell parameter b of Li2MnO3, respectively. The left region was indexed to the [001] axis, which was referred to as (100) domain. The middle region contained two evident planes, with 0.41 and 0.43 nm spacing corresponding to (110) and (020) planes. The middle region was indexed to the [001] axis, which was referred to as the (100) domain. The right region was similar to the (001) domain. In this regard, two domains piled along the diameter direction and formed nanowire morphology. In summary, HR-TEM and SAED patterns revealed that monocrystalline Li2MnO3 particles assembled in S1; small crystalline domains aggregated and formed polycrystalline nanowire in S2. XRD patterns and SEM images indicated that S1 and S2 were generated through an in situ reaction with a-MnO2 and MnOOH as precursors. Despite the distinct structure and valence values of Mn, a pure Li2MnO3 phase could be obtained via the molten-salt method.
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Fig. 1. XRD patterns (a) and Raman spectra (b) of S1 and S2.
Fig. 2. SEM and TEM images of P1 (a), P2 (d), S1 (b, c), and S2 (e, f) samples.
3.3. Electrochemical properties of Li2MnO3 nanowires The first charge/discharge voltage profiles of S1 and S2 were determined at a constant current density of 10 mA/g in Fig. 3(a). The initial discharge capacities were 178 and 226 mA h/g for S1 and S2, respectively. The first-cycle coulombic efficiencies of S1 and S2 were 68.4% and 67%, respectively. Thus, approximately 78% Liþ returned to the original position. S1 exhibited a linear charge plateau above 4.3 V. By contrast, the charge plateau of S2 contained two sections at 4.3e4.55 V and 4.55e4.6 V. The lower plateau might contribute to the (100) domain, where Liþ was extracted perpendicular to the [001] direction, because the plateau voltage corresponded to the energy of Liþ extracted from the lattice. Fig. 3(b) illustrated the cycle performance of S1 and S2 at a rate of 10 mA/g (1st-50th). S1 and S2 electrodes showed poor cycle performances. The discharge capacity retentions of S1 and S2 were 70.6% and 91.7% in the 6th cycle and 29% and 27.1% in the 50th cycle, respectively. A logarithmic degradation was observed in the S1 electrode before the 30th cycle; the S2 capacity demonstrated good
retention before the 6th cycle and then rapidly faded in the 6th30th cycles. In general, S2 was more active and might exhibit a distinct mechanism of capacity degradation than S1. The rate performance was determined at 10 mA/g, which was increased to 200 mA/g stepwise. S2 showed a stable capacity at a high current rate. A drastic decrease in the specific capacity was observed in S1 cycled at a high current rate. The varied cycle and rate performance might be partly attributed to the activation of the particles. S1 and S2 were subjected to CV measurements [Fig. 4(a) and (b)]. The redox peaks exhibited evident distinctions in terms of location and strength. For S1, four distinct peaks were observed in oxidation and reduction reactions during the cycles. The oxidation peaks near 4.6 V represented Li2O extraction from the Li2MnO3 structure [14,15]. The 4.2 V peak after the first cycle might arise from oxygen. The reduction peaks near 3.3 and 2.95 V were attributed to Mn4þ/3þ and Mn3þ/2þ reduction in the layered structure, respectively. The increasing intensity of the Mn4þ/3þ peak illustrated the activation of the Li2MnO3 structure upon cycling. The strength of the Mn3þ/2þ
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Fig. 3. Charge/discharge curves (a), cycle performance (b), and rate performance (c) of the S1 and S2 samples.
Fig. 4. CV curves of the S1 (a) and S2 (b) samples.
peak also demonstrated an uptrend when cycling was extended. In comparison to the locations of the peaks in S1, the locations of the Mn3þ-Mn2þ redox peaks shifted in S2 sample. The characteristic spinel peak at 2.8 and 3.0 V indicated the formation of a spinel phase upon cycling. More spinel phases might be generated in the powders with a larger particle size [15]. The sharper Mn3þ-Mn2þ redox peaks implied that more Mn2þ participated in electrochemical reaction, which might contribute to capacity degradation. Consistent with the cycle performance, phase transformation and Mn2þ participation might degrade the capacity of S2 upon cycling. With the aim to analyze the structural evolution of the samples, XRD patterns, TEM images and SAED patterns were carried out. In Fig. S1, XRD patterns of the electrodes represented that minor spinel phase formed after the initial cycle, with well-maintained layered structure. The peak area ratios of spinel (400) peak and
layered (-202) peak implied more spinel phase in S2 electrode after the initial cycle. Some distortion structure was observed on the edge of the crystalline in both samples. Spinel (400) plane was detected in the SAED pattern of the S2 electrode. XRD patterns and TEM images indicated that spinel phase, distortion structure and Jahn Teller effect might cause the capacity degradation [28]. More spinel phase and Mn2þ participation might have an influence on the mechanism of capacity degradation, which needed more experiments and analysis. Electrochemical impedance spectroscopy (EIS) was performed to identify the changes in the impedance after the initial discharge (V vs. Liþ/Li) and clarify electrochemical degradation upon cycling. In Fig. 5, the semicircle in the high-middle frequency region was attributed to the lithium-ion migration through the SEI film (Rf and Cf) and charge transfer reaction (Rct and Cdl), and the slope in the
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3.4. Liþ chemical diffusion coefficients Fig. 6 illustrated the open-circuit voltage (OCV) curves of S1 and S2. In the initial charge stage of S1 and S2, the evident over potential, especially above 4.3 V, suggested that the reaction dynamics of Li2MnO3 activation was more sluggish than that of Li extraction on the surface below 4.4 V. The over potential of 4.3e4.55 V and 4.55e4.6 V plateau indicated the distinct activation kinetics of the varied regions in S2. DLiþ could be calculated by the following simplified equation when the relationship of the cell voltage versus S1/2 was approximately linear [27]:
DLiþ ¼
4
mVM pt MS
Fig. 5. Nyquist plots of S1 and S2.
2
DEs DEt
2
Fig. 6. GITT curves and Liþ chemical diffusion coefficients in the initial cycle.
Table 1 The morphologies, BET, capacities, Rct, and DLiþ of S1 and S2. Samples S1 S2
Morphology Monocrystalline bundle Polycrystalline nanowire
BET 2
58 m /g 13 m2/g
low frequency region was attributed to the lithium-ion diffusion in the bulk electrodes (w) [26]. Corresponding Rct of S1 and S2 were 581 and 534U after the initial discharge, which were 139 and 187U in fresh cells. After 3 cycles, Rcts were increased to 2093 and 1557U for S1 and S2 electrodes. The lower Rct of S2 demonstrated better electroconductivity presumably as a consequence of crystal boundary and stacking defects, which were necessary to obtain high capacity and rate performance.
Capacity (10 mA/g)
Rct (after 3st cycle)
Dþ Li (after 1st cycle)
178 mA h/g 226 mA h/g
2093 U 1557 U
1016 cm2/s 1015 cm2/s
Where VM was the molar volume of active materials; M and m are the molecular weight and the mass of the active material, respectively; S was the surface area of the interface between an electrode and an electrolyte; and DEs and DEt were the physical interpretations in Fig. S3. The calculated DLiþ as a function of OCV in the initial cycle was shown in Fig. 6. In the initial charge of the S1 electrode, DLiþ
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decreased remarkably as its capacity increased and reached a minimal value of 1018 cm2/s at approximately 20% capacity (approximately 4.4 V). The DLiþ of Li2MnO3 at the activation plateau was as low as 1018 cm2/s. This finding suggested that the sluggish kinetics was an activation bottleneck of Li2MnO3. By comparison, S2 exhibited a DLiþ of 1016 cm2/s from 10% to 45% capacity and approximately 1017 cm2/s beyond 45% capacity; these findings corresponded to the two sections of 4.3e4.55 and 4.55e4.6 V plateau. During discharging, DLiþ demonstrated a downtrend as OCV decreased because of the increased concentration polarization accompanied by the insertion reaction of Li-ions in S1 and S2. Compared with the value obtained via the charging process, the DLiþ of S1 and S2 increased to 1016 and 1015 cm2/s, respectively. These values indicated that the structure was activated. The large DLiþ of S2 might be attributed to the grwoth patterns. Crystal boundaries and defects could introduce active reaction points and small domains with varied orientation, which were beneficial for Li extraction and accommodates the volume change in the crystal cell [15]. Liþ atoms extracted along the direction perpendicular to the c axis of the layered structure cathode materials. In the monocrystalline bundles, Liþ ions were directly delivered in [001] direction. While in polycrystalline nanowires, Liþ must reach the surface or the grain boundary along the direction perpendicular to the c axis, and then migrated on the surface or the grain boundary. Thus, the larger particle size and polycrystalline morphology endowed the S2 sample with longer Liþ-extraction path. Furthermore the low energy barrier of the crystal boundaries and stacking defects provided more active points for electrochemical reaction, which facilitated Liþ-extraction. In summary, the crystal boundaries and stacking defects in polycrystalline nanowires enormously promoted Liþ chemical diffusion coefficients. The morphologies, BET, capacities, Rct, and DLiþ of S1 and S2 were listed in Table 1. Despite the larger particle size observed in S2, the crystal boundaries, stacking defects, and small domains with varied orientation could enhance the electroconductivity and Liþ diffusion rate; thus, discharge capability, especially at high rates, could be improved.
4. Conclusion Monocrystalline bundles and polycrystalline nanowires with a pure layered structure were synthesized via an in-situ reaction. The polycrystalline nanowires were more activated than the monocrystalline nanowires. The former also delivered higher capacity and rate performance than the latter. Monocrystalline bundles exhibited a linear charging plateau at 4.3e4.6 V. By comparison, the charging plateau of the polycrystalline nanowires contained two sections at 4.3e4.55 and 4.55e4.6 V, which corresponded to (100) and (001) domains. The stacking defects and grain boundary introduced more active points for the electrochemical reaction. Higher electroconductivity and Liþ diffusion coefficient endowed polycrystalline nanowires with enhanced capability.
Acknowledgements This work was supported by the National Nature Science Foundation of China (NSFC No. 51372114), the Research Fund of State Key Laboratory of Mechanics and Control of Mechanical Structures (Nanjing University of Aeronautics and astronautics) (Grant No. 0514Y01), the Priority Academic Program Development of Jiangsu Higher Education Institutions (PAPD).
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