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Cite this: DOI: 10.1039/c4cp02957g Received 6th July 2014, Accepted 22nd August 2014
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Direct imaging of layered O3- and P2-NaxFe1/2Mn1/2O2 structures at the atomic scale† Xia Lu,ab Yuesheng Wang,a Pin Liu,a Lin Gu,*a Yong-Sheng Hu,*a Hong Li,a George P. Demopoulosb and Liquan Chena
DOI: 10.1039/c4cp02957g www.rsc.org/pccp
Using aberration-corrected scanning transmission electron microscopy (STEM) with high-angle annular-dark-field (HAADF) and annular-brightfield (ABF) techniques, the atomic-scale structures of the O3 and P2 phases of NaxFe1/2Mn1/2O2 are investigated systematically. The Na, transition metal M (Fe and Mn) and O columns are well revealed and precisely assigned to the O3 and P2 phase layered structures. The O3 phase sample demonstrates larger atomic site fluctuations along [001] direction but with less structural imperfections (e.g. interlayer structure and stacking defaults) than the P2 phase sample. Furthermore, a clear surface with a regular structure is observed for the O3-NaFe1/2Mn1/2O2 sample, while a surface with a large amount of Na–M antisites is observed for the P2-Na2/3Fe1/2Mn1/2O2 sample.
1. Introduction A long-life sodium-ion battery, which has been regarded as a promising electrochemical means to store electrical energy in a large-scale, is a hot topic and an important direction in this field. Although many materials have been proposed,1–22 the sodium storage mechanisms and the relationship between the structural evolutions and the electrochemical performance are far from well understood. In order to speed up the exploration of new and better materials, an in-depth understanding of this electrode material relationship is quite critical for sodium-ion batteries. Among the available structures, the layer-structured materials AMO2 (A = Li, Na; M = transition metals) are particularly promising candidates for Li/Na ion cathodes. Comparing Li and Na materials, the O3 phase is the most readily available structure for LiMO2, while for NaxMO2 (2/3 r x r 1.0), the O3 and/or P2 phases are favorable. Crystallographically, the A ions locate between two a
Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing, 100190, China. E-mail: [email protected], [email protected] b Department of Materials Engineering, McGill University, 3160 University Street, Montre´al (Que´bec), Canada H3A 0C5 † Electronic supplementary information (ESI) available. See DOI: 10.1039/ c4cp02957g
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M–O octahedron slabs along the [001] direction, and they reside at the octahedral sites (Oh coordination) in the O3 structure and at the prismatic sites (D3h coordination) in the P2 structure. Based on the atomic arrangement of oxygen layers, the AMO2 structures can be assigned to the O3, O1, P2, O2, and P3 phases, etc.23,24 Fig. 1 illustrates the O3 and P2 atomic structures of the layered AMO2 viewed from the [010] direction. From Fig. 1, the O columns stack in a ABCABC-mode and the M ion follows the abg-stacking mode in O3 phase. At the same time, the O columns display the ABBA-stacking mode, and only one M ion site is in AAA-stacking mode in the P2 structure along the [001] direction. Concerning the A ion occupations, it only takes the crystallographical 3b site as a regular AO6 octahedra with an edge-shared connection with a MO6 octahedra in the O3 phase, while the A ions fractionally occupy the 2b and 2d sites as AO6 prisms in the P2 structure with some evident distortions. In the P2 structure as shown in Fig. 1b, two different A ion sites demonstrate a zig-zag distribution along the [010] direction.
Fig. 1 Schematic illustrations of the O3 and P2 structures in the [010] direction. (a) The oxygen columns demonstrate the ABCABC-stacking model and the M ions display the abg-stacking model in O3 phase; (b) the oxygen columns give the ABBA-stacking model and the M ions exhibit the aaa-stacking model in P2 phase. The Na ions only occupy the 3b site in the O3 phase, while fractionally taking the 2b and 2d sites in the P2 structure. Crystallographically, the Li/Na ions locate between two M–O octahedron slabs along the [001] direction.
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The 2b site NaO6 octahedron exhibits edge-sharing, while the 2d site NaO6 octahedron exhibits face-sharing with the MO6 octahedra; in the meantime, the 2d Na ion resides on a less stable structure site than the 2b site.25 The A–O octahedra/ prisms share the edges and/or faces with MO6 octahedra, and the empty tetrahedral sites distribute uniformly among them, which is important for the evolution of the structure upon A insertion/extraction and also for A ion diffusion through the host structure.26 As to the Na ion storage and transport mechanism (electrochemistry–structure relationship) in these layered materials, research dates back to the 1980s for the sodium ion battery.27–29 In the early work, Delmas et al. reported phase transitions in the electrochemically active O3-NaCoO2, P 0 3-Na0.66CoO1.92, O0 3-Na0.77CoO1.92 and P2-Na0.70CoO1.96 phases during desodiation/ sodiation.28 They revealed by XRD analysis that the P2 structure is maintained over the whole composition range 0.46 r x r 0.83.28 Then, Molenda et al. demonstrated the initial concentration of electronic carriers to be a Na-dependent insertion property in the case of P2-Na0.70CoO2y cathode.30 After that, Viciu et al. reported several complex Na content-dependent structures of NaxCoO2 with x = 0.32, 0.51, 0.60, 0.75, and 0.92 as determined by powder neutron diffraction, and they revealed that the underlying CoO2 lattice experiences some kind of structural modulation at the x = 0.75 composition.31 Shu et al. investigated the Na+ ion diffusion and ordering properties in a P2-NaxCoO2 single crystal using the PITT technique. They demonstrated that upon excluding the significantly slower diffusion behavior near particular Na ordered phases, the average diffusion coefficients DNa are 107 cm2 s1 for x 4 0.5 and 108 cm2 s1 for x o 0.5.32 Subsequently, Berthelot et al. systematically investigated the P2-NaxCoO2 phase diagram and identified at room temperature nine single-phase domains with narrow sodium composition ranges due to peculiar Na+/Na vacancy ordering.25 Very recently, Yabuuchi et al. reported that the P2-Na2/3Fe1/2Mn1/2O2 phase delivers 190 mA h g1 of reversible capacity with the electrochemically active Fe3+/Fe4+ redox couple, which is much higher than the 110 mA h g1 reversible capacity for O3-Na2/3Fe1/2Mn1/2O2 phase, showing promising potential applications.14 Consequently, it can be summarized that the layered NaxMO2 materials are structure-dependent cathodes with complex phase transitions during desodiation/sodiation, and much research needs to be further focused on clarifying the structure of these layered materials for use in sodium-ion batteries. In terms of electrochemical performance, a solid-solution reaction with (a) an initial short-plateau region and (b) a subsequent long slope region33 is always observed in the O3-LiMO2 (M = Co) cathode for a Li-ion battery. However, there are a series of plateaus appearing on the charge–discharge curves of the P2-NaxMO2 (M = Co) cathode for the Na-ion battery.13,14 Moreover, the structural stabilities of AMO2 cathodes, including the structural evolutions, A–M antisites, phase transitions, and charge compensation from the oxygen matrix are still under dispute. Therefore, it is necessary to get detailed structural information, which is beneficial for dealing with such encountered issues. In this work, the layered O3/P2-NaxFe0.5Mn0.5O2 samples are
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synthesized, and the related atomic structures are carefully investigated using spherical aberration-corrected scanning transmission electron microscopy (STEM) with high-angle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques.
2. Experimental section The NaxFe0.5Mn0.5O2 samples were synthesized via a solid-state reaction according to the references.14,25 In a typical synthetic process, the stoichiometric amounts of sodium peroxide (Na2O2), iron oxide (Fe2O3) and manganese(III) oxide (Mn2O3) powders were mixed and ground adequately in a glove box. Then the mixed powder was pelleted into slices. The slices were heated at 700 1C for 36 h to obtain the O3 phase and at 900 1C for 12 h to obtain the P2 phase at a heating rate of 5 1C min1 in the muffle furnace. After the temperature was cooled naturally to room temperature, the O3-/P2-NaxFe0.5Mn0.5O2 samples were obtained. X-ray powder diffraction (XRD) patterns were collected in the range of 10–801 using a Philips X’pert diffractometer with Bragg–Brentano geometry. The measurement was performed in a continuous scan mode, using a Cu Ka X-ray source and a step of 0.01671 s1. A scanning electron microscope (Hitachi S-4800) was used to study the morphology. Aberration-corrected scanning transmission electron microscopy was performed using a JEOL 2100F (JEOL, Tokyo, Japan) transmission electron microscope equipped with a CEOS (CEOS, Heidelberg, Germany) probe aberration corrector. The attainable spatial resolution of the microscope is 90 picometers at an incident angle of 40 mrad.
3. Results and discussion The P2 and O3 type materials are the two main groups for sodium-based layered electrodes. Fig. 2 shows the XRD patterns and the SEM morphologies of the as-prepared P2 and O3 layered samples. From the Fig. 2a, the XRD patterns can be well indexed to the P2 type and O3 type NaxFe0.5Mn0.5O2 samples (hereafter denoted as P2-NaFMO and O3-NaFMO) with minor impurities. The impurities are always indicative of the formation of other types of layer-structured materials and/or
Fig. 2 XRD patterns and SEM images for (a) & (b) layered P2-phase Na2/3Fe1/2Mn1/2O2 and (a) & (c) O3-phase NaFe1/2Mn1/2O2 samples. Most peaks can be precisely indexed to the P2 and O3 NaxFe1/2Mn1/2O2 phases. In addition, the red arrows indicate the minor impurities in the as-prepared samples, which are difficult to eliminate during the synthetic process.
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metal oxides during synthesis.14,25 From the SEM micrographs, the P2-NaFMO samples have a larger particle size than that of the O3-NaFMO samples on average, and the primary particles agglomerate together to form big secondary particles. Concerning the electrochemical performance, Yabuuchi et al. reported that the P2-NaFMO electrode delivers about 190 mA h g1 in the voltage range of 1.5–4.2 V, which is higher than the 110 mA h g1 delivered by the O3-NaFMO sample as a cathode in half cells.14 Furthermore, they found the P2–OP4 phase transition upon desodiation in a P2-NaFMO electrode and O3–P3–OP2 phase transition upon desodiation in a O3-NaFMO electrode using the synchrotron X-ray diffraction method.14 These phase transitions are caused probably by the interaction of the Na+ ion vacancies with the transition metal octahedra, which are the common phenomena in layered materials for rechargeable batteries.33–36 From the viewpoint of structure stability, the layered battery materials are apt to adjust their energetically favorable configuration when x changes in the AxMO2 structure, and after that, the materials try to maintain the corresponding states. Therefore, for a perfect layered crystal, the structure varies with the x value, and phase transitions (or structure evolutions) happen accordingly. In order to elucidate these phase variations, structure characterizations are executed on NaxFe0.5Mn0.5O2 samples using aberrationcorrected STEM with high-angle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques. STEM techniques have proven very successful in rechargeable battery research as they allow for the precise atomic positions (such as light element Li, O), structural defects, phase evolution and phase/grain boundaries to be clearly acquired at an atomic scale,33,37–43 hence, advancing a deeper comprehension of ion storage and transport. It is worth of mentioning that the contrast of the HAADF image follows a Z1.7 dependence as compared with Z1/3 for the ABF image with respect to the atomic number (Z).41,44,45 Fig. 3 and Fig. S1 (ESI†) demonstrate the filtered STEM HAADF and ABF images of P2-NaFMO sample. In Fig. 3a, the HAADF image reveals the ordered arrangement of the transition metal M (Fe and Mn) columns at the [010] zone axis. Along either the [100] or [001] direction, the M columns align straightforwardly, which is totally different from that in the O3-NaFMO samples as discussed later. The distance of the adjacent layer dc is ca. 0.60 nm, which is a little larger than the reported 0.56 nm value14 and is probably caused by nanosizing effects and/or the different synthesis conditions employed. Moreover, from the ABF image in Fig. 3b, the light element Na and O columns can be distinguishably identified besides the M columns, which are also shown in Fig. S1d (ESI†). The Na, O and M columns have also different contrasts in the ABF image as shown in Fig. S1d (ESI†). For the M columns, an aaa-stacking mode is shown, the same as that in Fig. 3a; for the O columns, a distinct ABBA-stacking mode is shown in the acquired ABF picture. Each O column follows the same stacking mode as one of the adjacent O columns but with a site shift (0.5da) to the other adjacent O column along the [001] direction, displaying a mirroring symmetry, when referred to the Na ion layer at the [010] zone axis; for Na columns, it demonstrates a regular ABAB-stacking mode along the [001] direction. Note that
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Fig. 3 STEM (a) HAADF and (b) ABF images for the as-prepared P2-phase Na2/3Fe1/2Mn1/2O2 materials at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original images are presented in the Fig. S1 (ESI†).
the Na1 and Na2 columns are too close to be distinguishable, and their contrasts overlap with each other as one column to a great extent when projected along the [100] direction as shown in Fig. 3b and Fig. S1e (ESI†). One NaO6 prism is employed as a medium to connect two MO6 octahedra and minimize the coulombic repulsion interaction, hence stabilizing the P2 layered structure. However, in P2-layered structures, it is difficult to get samples with Na ions fully occupying the octahedra/prisms in the (003) face, which means it is hard to obtain the stoichiometric formula NaMO2. If the two Na ions approach closer like the Li–Li distance in the O3-LiCoO2 structure, the Na ion will repel the adjacent one to create Na vacancies due to the larger Na ion radius. Therefore, here in Fig. 3b the Na ion columns are fractionally filled to leave a certain amount of Na vacancies in the as-prepared samples. Furthermore, the Na vacancies show no evidence of clustering or ordered arrangements based on the small contrast changes of each Na column along the [100] direction (refer to Fig. S1e, ESI†). In addition, a certain amount of Na vacancies in the structure are beneficial for: (1) structural stability and (2) Na+ ion transport. The O3-NaFMO sample reveals a different physical picture from that of the P2-NaFMO sample as shown in Fig. 4. From the HAADF image in Fig. 4a, the M columns align in line along the [100] direction, the same as in the P2-NaFMO HAADF image. However, along the [001] direction, each M column shows a site glide of the 2/3da, when compared with the adjacent layers. Therefore, every third layer, we have the same M column occupations, and the M column duplicates the ABCABC-stacking mode to form the O3 structure along the [001] direction. Likewise, the distance of adjacent layer dc is ca. 0.61 nm, which is still a little larger than the reported value of 0.55 nm.14 The results in the ABF image in Fig. 4b and from Fig. S2c (ESI†) demonstrate the unambiguous contrast of Na, M and O columns. All of the three different columns take the same ABCABC-stacking mode in the atomic arrangements. However, there is a little difference for the O columns. As shown in Fig. 4b, the layer distance of each O column is much different from the head-to-head stacking (ABBA) mode of the P2 phase in Fig. 3b along the [001] direction.
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Fig. 4 STEM (a) HAADF and (b) ABF images for the as-prepared O3-phase NaFe1/2Mn1/2O2 sample at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original pictures are presented in the Fig. S2 (ESI†).
Apparently, upon ion extraction/insertion, the O matrix is easier to glide in the P2 structure than in the O3 structure, because the O3 structure is the thermodynamically equilibrated configuration based on theoretical calculations.33 However, structural changes still happen irreversibly, when the alkali metal A vacancies accumulate to some extent in both structures.14 The structure relaxes into new configurations from which it is hard to recover back into the initial structure due to eventual introduction of vacancies upon alkali metal A extraction. In the small regions of the HAADF and ABF images (see Fig. 3 and 4), the P2 and O3 layered structures are shown with nearly ideal Na, O and M atomic arrangements. However, structural distortions and defects are always present in these ordered structures as demonstrated in Fig. 5 and Fig. S3 (ESI†). Actually, these structural imperfections apparently play an important role in ion storage and transport. The HAADF images in Fig. 5a and Fig. S3a (ESI†) reveal the remarkable structural distortion in the O3-NaFMO sample. From the [100] direction, the M site fluctuations cause the variations of the (003) planar distance. The dc is estimated to be 0.61 0.10 nm. Due to negligible Na+ ion vacancies, the Na–Na distance decreases in the (003) plane, which in turn causes strong Na+–Na+ repulsion, introducing larger distortion in the NaO6 octahedron that is ultimately
Fig. 5 STEM HAADF images for (a) O3-phase NaFe1/2Mn1/2O2 sample and (b) P2-phase Na2/3Fe1/2Mn1/2O2 sample at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original pictures are presented in the Fig. S3 (ESI†).
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responsible for the observed layer distance fluctuations in the O3 structure.33 Crystallographically, the layer distance is far larger than the diameter of a Na+ ion (B0.20 nm); therefore, it should be beneficial for Na storage and diffusion. However, the distortion induced by the Na+ ion vacancies is believed to set several changes in motion: (1) macroscopically, the Na, M and O arrays are displaced to find a new thermodynamically equilibrated configuration that as consequence reaches a new stable phase with a new chemical potential m modifying the cell voltage accordingly; (2) microscopically, the Na ion Oh symmetry is broken, causing lattice field interruption and the introduction of coulombic interaction between the two MO6 slabs. That is, the created Na+ ion vacancies contribute a lot to structural instability. Furthermore, by observing Fig. S3a (ESI†), some local microstructure with another M stacking mode can be seen, such as O2-like microstructures (M in ab-stacking and O in ABACstacking, see Fig. S5, ESI†), which are directly related with this local layer structural change/distortion. For the O3-NaFMO sample, the most observed phenomenon is the remarkable layer distortion, while for the P2-NaFMO sample, more information can be acquired from the HAADF images shown in Fig. 5b and Fig. S3b–e (ESI†). In Fig. 5b, the layer distance was measured to be 0.60 0.04 nm with the layer stripe clearly shown along the [100] direction. The fluctuation of layer distance seems smaller than that of the O3-NaFeMO sample along the [001] direction. However, the structural imperfections increase significantly. First, in the yellow arrow region of Fig. S3b–e (ESI†), it can be observed that there are some unexpectedly larger layer distances, which are nearly twice the regular dc value. This kind of layer expansion seems plausibly beneficial for Na+ ion diffusion. Furthermore Fig. S3b (ESI†) also demonstrates the layer stacking default, which gives birth to a ca. 51 angle as shown in the green and yellow underlined parts. This can be treated as the interface of two particles, which will merge together to form a new larger particle. In Fig. S3c (ESI†), a few intercalated layers are witnessed in the bulk of one particle, which are probably caused by Na–M antisites formed during annealing. Here, the adjacent coordination of Na+ ion changes from only the O ions to the O and M ions, which should play an important role in the evolution of structure and even battery failure upon sodiation/de-sodiation. The intercalated structures are also found in the surface region of the P2-NaFMO sample, which extends from the outmost surface to the bulk, about 20 nm far away from the surface as shown in Fig. S3d (ESI†). The O3-NaFMO particles display larger Na, M and O site fluctuations along the [001] direction than do the P2-NaFMO particles. Therefore, concerning surfacal structures as shown in Fig. 6, the O3-NaFMO sample demonstrates a remarkably different structure from that of the P2-NaFMO sample. From Fig. 6a and Fig. S4a and b (ESI†), the HAADF image of the O3-NaFMO sample is seen to exhibit ordered M columns aligned straightforwardly from the bulk to the surface, and thus so do the Na and O columns along the [100] direction (ABF image in Fig. S4c, ESI†). This gives a clear O3-NaFMO surface structure. However, a local region with negligible Na–M antisites was
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For the O3 phase, the O columns follow the ABCABC-stacking mode in the ABF image. The O3 structure demonstrates obvious site fluctuation with the layer distance dc B 0.61 0.10 nm along the [001] direction in the HAADF image. Furthermore, a clear and ordered atomic arrangement was observed with negligible Na–M antisites in several outermost surface layers. For the P2 phase, it was unambiguously observed the regular Na and O in ABBA-stacking, and M in aaa-stacking modes along the [001] direction. The P2 phase also demonstrates a regular layer distance with fewer fluctuations. Moreover, some obvious structural imperfections are detected accordingly from the HAADF and ABF images, including the stacking faults, intercalated structure and remarkable layer distortion. Compared with the O3 sample, the P2 phase Na2/3Fe1/2Mn1/2O2 sample demonstrates a surface structure with a large amount of Na–M antisites. Fig. 6 STEM HAADF images of the surface structures without (a) and with (b) a few Na–M antisites for the as-prepared O3-phase NaFe1/2Mn1/2O2; surface structures without (c) and with (d) obvious Na–M antisites for the as-prepared P2-phase Na2/3Fe1/2Mn1/2O2 at the [010] zone axis. The original pictures are presented in Fig. S4 (ESI†).
observed in Fig. 6b and Fig. S4d (ESI†), which is likely introduced by surface relaxation to excite the Na ions interchanging with M ions. Therefore, the M columns do not strictly follow the abg-stacking fault (Fig. S5a and S6a, ESI†) but have slightly disordered arrangements in several surface layers along the [001] direction. In contrast, this disordered stacking mode seems to facilitate the formation of Na–M antisites in the surface region after cycling.33 In Fig. 6c and d, the HAADF images reveal the surface structure without/with Na–M antisites in a P2-NaFMO sample. A typical difference is that the amount of Na–M antisites increases to a great extent and can be easily found in the available surface region when compared with that on the O3-NaFMO surface. In the enlarged HAADF image of Fig. S4f (ESI†), a large area with fully Na–M antisites was found to cover the surface of the P2-NaFMO sample. This can also be confirmed by the enhanced contrast of the Na columns in the ABF image of Fig. S4g (ESI†). Therefore, it shows that a clear surface is easily formed in the case of the O3-NaFMO sample, while in the case of the P2-NaFMO sample, a surface with significant Na–M antisites forms instead.
4. Conclusions O3 and P2 phase NaxFe1/2Mn1/2O2 samples were synthesized by a solid state method and investigated using aberration-corrected scanning transmission electron microscopy (STEM) with highangle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques. The XRD patterns of the as-prepared samples can be well indexed to the O3 and P2 phases, although some minor impurities were present. The HAADF and ABF images were acquired with clear Na, M (Fe and Mn) and O occupations at the atomic scale. Some new insights concerning the structure and phase transitions during sodiation/de-sodiation are achieved:
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Acknowledgements This work was supported by funding from the ‘‘973’’ Projects (2010CB833102, 2012CB932900), NSFC (51222210, 11234013, 11174334), ‘‘Strategic Priority Research Program’’ of the Chinese Academy of Sciences (Grant No. XDA01020304) and One Hundred Talent Project of the Chinese Academy of Sciences.
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37 L. Gu, C. B. Zhu, H. Li, Y. Yu, C. L. Li, S. Tsukimoto, J. Maier and Y. Ikuhara, Direct Observation of Lithium Staging in Partially Delithiated LiFePO4 at Atomic Resolution, J. Am. Chem. Soc., 2011, 133(13), 4661–4663. 38 X. Lu, L. Zhao, X. Q. He, R. J. Xiao, L. Gu, Y. S. Hu, H. Li, Z. X. Wang, X. F. Duan, L. Q. Chen, J. Maier and Y. Ikuhara, Lithium Storage in Li4Ti5O12 Spinel: The Full Static Picture from Electron Microscopy, Adv. Mater., 2012, 24(24), 3233–3238. 39 X. Lu, Z. L. Jian, Z. Fang, L. Gu, Y. S. Hu, W. Chen, Z. X. Wang and L. Q. Chen, Atomic-scale investigation on lithium storage mechanism in TiNb2O7, Energy Environ. Sci., 2011, 4(8), 2638–2644. 40 Y. Sun, L. Zhao, H. L. Pan, X. Lu, L. Gu, Y. S. Hu, H. Li, M. Armand, Y. Ikuhara, L. Q. Chen and X. J. Huang, Direct atomic-scale confirmation of three-phase storage mechanism in Li4Ti5O12 anodes for room-temperature sodium-ion batteries, Nat. Commun., 2013, 4, 1870. 41 X. Q. He, L. Gu, C. B. Zhu, Y. Yu, C. L. Li, Y. S. Hu, H. Li, S. Tsukimoto, J. Maier, Y. Ikuhara and X. F. Duan, Direct Imaging of Lithium Ions Using Aberration-Corrected
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Cite this: DOI: 10.1039/c4cp02957g Received 6th July 2014, Accepted 22nd August 2014
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Direct imaging of layered O3- and P2-NaxFe1/2Mn1/2O2 structures at the atomic scale† Xia Lu,ab Yuesheng Wang,a Pin Liu,a Lin Gu,*a Yong-Sheng Hu,*a Hong Li,a George P. Demopoulosb and Liquan Chena
DOI: 10.1039/c4cp02957g www.rsc.org/pccp
Using aberration-corrected scanning transmission electron microscopy (STEM) with high-angle annular-dark-field (HAADF) and annular-brightfield (ABF) techniques, the atomic-scale structures of the O3 and P2 phases of NaxFe1/2Mn1/2O2 are investigated systematically. The Na, transition metal M (Fe and Mn) and O columns are well revealed and precisely assigned to the O3 and P2 phase layered structures. The O3 phase sample demonstrates larger atomic site fluctuations along [001] direction but with less structural imperfections (e.g. interlayer structure and stacking defaults) than the P2 phase sample. Furthermore, a clear surface with a regular structure is observed for the O3-NaFe1/2Mn1/2O2 sample, while a surface with a large amount of Na–M antisites is observed for the P2-Na2/3Fe1/2Mn1/2O2 sample.
1. Introduction A long-life sodium-ion battery, which has been regarded as a promising electrochemical means to store electrical energy in a large-scale, is a hot topic and an important direction in this field. Although many materials have been proposed,1–22 the sodium storage mechanisms and the relationship between the structural evolutions and the electrochemical performance are far from well understood. In order to speed up the exploration of new and better materials, an in-depth understanding of this electrode material relationship is quite critical for sodium-ion batteries. Among the available structures, the layer-structured materials AMO2 (A = Li, Na; M = transition metals) are particularly promising candidates for Li/Na ion cathodes. Comparing Li and Na materials, the O3 phase is the most readily available structure for LiMO2, while for NaxMO2 (2/3 r x r 1.0), the O3 and/or P2 phases are favorable. Crystallographically, the A ions locate between two a
Beijing National Laboratory for Condensed Matter Physics, Institute of Physics, Chinese Academy of Sciences, Beijing, 100190, China. E-mail: [email protected], [email protected] b Department of Materials Engineering, McGill University, 3160 University Street, Montre´al (Que´bec), Canada H3A 0C5 † Electronic supplementary information (ESI) available. See DOI: 10.1039/ c4cp02957g
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M–O octahedron slabs along the [001] direction, and they reside at the octahedral sites (Oh coordination) in the O3 structure and at the prismatic sites (D3h coordination) in the P2 structure. Based on the atomic arrangement of oxygen layers, the AMO2 structures can be assigned to the O3, O1, P2, O2, and P3 phases, etc.23,24 Fig. 1 illustrates the O3 and P2 atomic structures of the layered AMO2 viewed from the [010] direction. From Fig. 1, the O columns stack in a ABCABC-mode and the M ion follows the abg-stacking mode in O3 phase. At the same time, the O columns display the ABBA-stacking mode, and only one M ion site is in AAA-stacking mode in the P2 structure along the [001] direction. Concerning the A ion occupations, it only takes the crystallographical 3b site as a regular AO6 octahedra with an edge-shared connection with a MO6 octahedra in the O3 phase, while the A ions fractionally occupy the 2b and 2d sites as AO6 prisms in the P2 structure with some evident distortions. In the P2 structure as shown in Fig. 1b, two different A ion sites demonstrate a zig-zag distribution along the [010] direction.
Fig. 1 Schematic illustrations of the O3 and P2 structures in the [010] direction. (a) The oxygen columns demonstrate the ABCABC-stacking model and the M ions display the abg-stacking model in O3 phase; (b) the oxygen columns give the ABBA-stacking model and the M ions exhibit the aaa-stacking model in P2 phase. The Na ions only occupy the 3b site in the O3 phase, while fractionally taking the 2b and 2d sites in the P2 structure. Crystallographically, the Li/Na ions locate between two M–O octahedron slabs along the [001] direction.
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The 2b site NaO6 octahedron exhibits edge-sharing, while the 2d site NaO6 octahedron exhibits face-sharing with the MO6 octahedra; in the meantime, the 2d Na ion resides on a less stable structure site than the 2b site.25 The A–O octahedra/ prisms share the edges and/or faces with MO6 octahedra, and the empty tetrahedral sites distribute uniformly among them, which is important for the evolution of the structure upon A insertion/extraction and also for A ion diffusion through the host structure.26 As to the Na ion storage and transport mechanism (electrochemistry–structure relationship) in these layered materials, research dates back to the 1980s for the sodium ion battery.27–29 In the early work, Delmas et al. reported phase transitions in the electrochemically active O3-NaCoO2, P 0 3-Na0.66CoO1.92, O0 3-Na0.77CoO1.92 and P2-Na0.70CoO1.96 phases during desodiation/ sodiation.28 They revealed by XRD analysis that the P2 structure is maintained over the whole composition range 0.46 r x r 0.83.28 Then, Molenda et al. demonstrated the initial concentration of electronic carriers to be a Na-dependent insertion property in the case of P2-Na0.70CoO2y cathode.30 After that, Viciu et al. reported several complex Na content-dependent structures of NaxCoO2 with x = 0.32, 0.51, 0.60, 0.75, and 0.92 as determined by powder neutron diffraction, and they revealed that the underlying CoO2 lattice experiences some kind of structural modulation at the x = 0.75 composition.31 Shu et al. investigated the Na+ ion diffusion and ordering properties in a P2-NaxCoO2 single crystal using the PITT technique. They demonstrated that upon excluding the significantly slower diffusion behavior near particular Na ordered phases, the average diffusion coefficients DNa are 107 cm2 s1 for x 4 0.5 and 108 cm2 s1 for x o 0.5.32 Subsequently, Berthelot et al. systematically investigated the P2-NaxCoO2 phase diagram and identified at room temperature nine single-phase domains with narrow sodium composition ranges due to peculiar Na+/Na vacancy ordering.25 Very recently, Yabuuchi et al. reported that the P2-Na2/3Fe1/2Mn1/2O2 phase delivers 190 mA h g1 of reversible capacity with the electrochemically active Fe3+/Fe4+ redox couple, which is much higher than the 110 mA h g1 reversible capacity for O3-Na2/3Fe1/2Mn1/2O2 phase, showing promising potential applications.14 Consequently, it can be summarized that the layered NaxMO2 materials are structure-dependent cathodes with complex phase transitions during desodiation/sodiation, and much research needs to be further focused on clarifying the structure of these layered materials for use in sodium-ion batteries. In terms of electrochemical performance, a solid-solution reaction with (a) an initial short-plateau region and (b) a subsequent long slope region33 is always observed in the O3-LiMO2 (M = Co) cathode for a Li-ion battery. However, there are a series of plateaus appearing on the charge–discharge curves of the P2-NaxMO2 (M = Co) cathode for the Na-ion battery.13,14 Moreover, the structural stabilities of AMO2 cathodes, including the structural evolutions, A–M antisites, phase transitions, and charge compensation from the oxygen matrix are still under dispute. Therefore, it is necessary to get detailed structural information, which is beneficial for dealing with such encountered issues. In this work, the layered O3/P2-NaxFe0.5Mn0.5O2 samples are
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synthesized, and the related atomic structures are carefully investigated using spherical aberration-corrected scanning transmission electron microscopy (STEM) with high-angle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques.
2. Experimental section The NaxFe0.5Mn0.5O2 samples were synthesized via a solid-state reaction according to the references.14,25 In a typical synthetic process, the stoichiometric amounts of sodium peroxide (Na2O2), iron oxide (Fe2O3) and manganese(III) oxide (Mn2O3) powders were mixed and ground adequately in a glove box. Then the mixed powder was pelleted into slices. The slices were heated at 700 1C for 36 h to obtain the O3 phase and at 900 1C for 12 h to obtain the P2 phase at a heating rate of 5 1C min1 in the muffle furnace. After the temperature was cooled naturally to room temperature, the O3-/P2-NaxFe0.5Mn0.5O2 samples were obtained. X-ray powder diffraction (XRD) patterns were collected in the range of 10–801 using a Philips X’pert diffractometer with Bragg–Brentano geometry. The measurement was performed in a continuous scan mode, using a Cu Ka X-ray source and a step of 0.01671 s1. A scanning electron microscope (Hitachi S-4800) was used to study the morphology. Aberration-corrected scanning transmission electron microscopy was performed using a JEOL 2100F (JEOL, Tokyo, Japan) transmission electron microscope equipped with a CEOS (CEOS, Heidelberg, Germany) probe aberration corrector. The attainable spatial resolution of the microscope is 90 picometers at an incident angle of 40 mrad.
3. Results and discussion The P2 and O3 type materials are the two main groups for sodium-based layered electrodes. Fig. 2 shows the XRD patterns and the SEM morphologies of the as-prepared P2 and O3 layered samples. From the Fig. 2a, the XRD patterns can be well indexed to the P2 type and O3 type NaxFe0.5Mn0.5O2 samples (hereafter denoted as P2-NaFMO and O3-NaFMO) with minor impurities. The impurities are always indicative of the formation of other types of layer-structured materials and/or
Fig. 2 XRD patterns and SEM images for (a) & (b) layered P2-phase Na2/3Fe1/2Mn1/2O2 and (a) & (c) O3-phase NaFe1/2Mn1/2O2 samples. Most peaks can be precisely indexed to the P2 and O3 NaxFe1/2Mn1/2O2 phases. In addition, the red arrows indicate the minor impurities in the as-prepared samples, which are difficult to eliminate during the synthetic process.
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metal oxides during synthesis.14,25 From the SEM micrographs, the P2-NaFMO samples have a larger particle size than that of the O3-NaFMO samples on average, and the primary particles agglomerate together to form big secondary particles. Concerning the electrochemical performance, Yabuuchi et al. reported that the P2-NaFMO electrode delivers about 190 mA h g1 in the voltage range of 1.5–4.2 V, which is higher than the 110 mA h g1 delivered by the O3-NaFMO sample as a cathode in half cells.14 Furthermore, they found the P2–OP4 phase transition upon desodiation in a P2-NaFMO electrode and O3–P3–OP2 phase transition upon desodiation in a O3-NaFMO electrode using the synchrotron X-ray diffraction method.14 These phase transitions are caused probably by the interaction of the Na+ ion vacancies with the transition metal octahedra, which are the common phenomena in layered materials for rechargeable batteries.33–36 From the viewpoint of structure stability, the layered battery materials are apt to adjust their energetically favorable configuration when x changes in the AxMO2 structure, and after that, the materials try to maintain the corresponding states. Therefore, for a perfect layered crystal, the structure varies with the x value, and phase transitions (or structure evolutions) happen accordingly. In order to elucidate these phase variations, structure characterizations are executed on NaxFe0.5Mn0.5O2 samples using aberrationcorrected STEM with high-angle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques. STEM techniques have proven very successful in rechargeable battery research as they allow for the precise atomic positions (such as light element Li, O), structural defects, phase evolution and phase/grain boundaries to be clearly acquired at an atomic scale,33,37–43 hence, advancing a deeper comprehension of ion storage and transport. It is worth of mentioning that the contrast of the HAADF image follows a Z1.7 dependence as compared with Z1/3 for the ABF image with respect to the atomic number (Z).41,44,45 Fig. 3 and Fig. S1 (ESI†) demonstrate the filtered STEM HAADF and ABF images of P2-NaFMO sample. In Fig. 3a, the HAADF image reveals the ordered arrangement of the transition metal M (Fe and Mn) columns at the [010] zone axis. Along either the [100] or [001] direction, the M columns align straightforwardly, which is totally different from that in the O3-NaFMO samples as discussed later. The distance of the adjacent layer dc is ca. 0.60 nm, which is a little larger than the reported 0.56 nm value14 and is probably caused by nanosizing effects and/or the different synthesis conditions employed. Moreover, from the ABF image in Fig. 3b, the light element Na and O columns can be distinguishably identified besides the M columns, which are also shown in Fig. S1d (ESI†). The Na, O and M columns have also different contrasts in the ABF image as shown in Fig. S1d (ESI†). For the M columns, an aaa-stacking mode is shown, the same as that in Fig. 3a; for the O columns, a distinct ABBA-stacking mode is shown in the acquired ABF picture. Each O column follows the same stacking mode as one of the adjacent O columns but with a site shift (0.5da) to the other adjacent O column along the [001] direction, displaying a mirroring symmetry, when referred to the Na ion layer at the [010] zone axis; for Na columns, it demonstrates a regular ABAB-stacking mode along the [001] direction. Note that
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Fig. 3 STEM (a) HAADF and (b) ABF images for the as-prepared P2-phase Na2/3Fe1/2Mn1/2O2 materials at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original images are presented in the Fig. S1 (ESI†).
the Na1 and Na2 columns are too close to be distinguishable, and their contrasts overlap with each other as one column to a great extent when projected along the [100] direction as shown in Fig. 3b and Fig. S1e (ESI†). One NaO6 prism is employed as a medium to connect two MO6 octahedra and minimize the coulombic repulsion interaction, hence stabilizing the P2 layered structure. However, in P2-layered structures, it is difficult to get samples with Na ions fully occupying the octahedra/prisms in the (003) face, which means it is hard to obtain the stoichiometric formula NaMO2. If the two Na ions approach closer like the Li–Li distance in the O3-LiCoO2 structure, the Na ion will repel the adjacent one to create Na vacancies due to the larger Na ion radius. Therefore, here in Fig. 3b the Na ion columns are fractionally filled to leave a certain amount of Na vacancies in the as-prepared samples. Furthermore, the Na vacancies show no evidence of clustering or ordered arrangements based on the small contrast changes of each Na column along the [100] direction (refer to Fig. S1e, ESI†). In addition, a certain amount of Na vacancies in the structure are beneficial for: (1) structural stability and (2) Na+ ion transport. The O3-NaFMO sample reveals a different physical picture from that of the P2-NaFMO sample as shown in Fig. 4. From the HAADF image in Fig. 4a, the M columns align in line along the [100] direction, the same as in the P2-NaFMO HAADF image. However, along the [001] direction, each M column shows a site glide of the 2/3da, when compared with the adjacent layers. Therefore, every third layer, we have the same M column occupations, and the M column duplicates the ABCABC-stacking mode to form the O3 structure along the [001] direction. Likewise, the distance of adjacent layer dc is ca. 0.61 nm, which is still a little larger than the reported value of 0.55 nm.14 The results in the ABF image in Fig. 4b and from Fig. S2c (ESI†) demonstrate the unambiguous contrast of Na, M and O columns. All of the three different columns take the same ABCABC-stacking mode in the atomic arrangements. However, there is a little difference for the O columns. As shown in Fig. 4b, the layer distance of each O column is much different from the head-to-head stacking (ABBA) mode of the P2 phase in Fig. 3b along the [001] direction.
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Fig. 4 STEM (a) HAADF and (b) ABF images for the as-prepared O3-phase NaFe1/2Mn1/2O2 sample at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original pictures are presented in the Fig. S2 (ESI†).
Apparently, upon ion extraction/insertion, the O matrix is easier to glide in the P2 structure than in the O3 structure, because the O3 structure is the thermodynamically equilibrated configuration based on theoretical calculations.33 However, structural changes still happen irreversibly, when the alkali metal A vacancies accumulate to some extent in both structures.14 The structure relaxes into new configurations from which it is hard to recover back into the initial structure due to eventual introduction of vacancies upon alkali metal A extraction. In the small regions of the HAADF and ABF images (see Fig. 3 and 4), the P2 and O3 layered structures are shown with nearly ideal Na, O and M atomic arrangements. However, structural distortions and defects are always present in these ordered structures as demonstrated in Fig. 5 and Fig. S3 (ESI†). Actually, these structural imperfections apparently play an important role in ion storage and transport. The HAADF images in Fig. 5a and Fig. S3a (ESI†) reveal the remarkable structural distortion in the O3-NaFMO sample. From the [100] direction, the M site fluctuations cause the variations of the (003) planar distance. The dc is estimated to be 0.61 0.10 nm. Due to negligible Na+ ion vacancies, the Na–Na distance decreases in the (003) plane, which in turn causes strong Na+–Na+ repulsion, introducing larger distortion in the NaO6 octahedron that is ultimately
Fig. 5 STEM HAADF images for (a) O3-phase NaFe1/2Mn1/2O2 sample and (b) P2-phase Na2/3Fe1/2Mn1/2O2 sample at the [010] zone axis. The dc and da represent the distance of the (003) plane and the distance of two M columns along the [010] direction. The original pictures are presented in the Fig. S3 (ESI†).
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responsible for the observed layer distance fluctuations in the O3 structure.33 Crystallographically, the layer distance is far larger than the diameter of a Na+ ion (B0.20 nm); therefore, it should be beneficial for Na storage and diffusion. However, the distortion induced by the Na+ ion vacancies is believed to set several changes in motion: (1) macroscopically, the Na, M and O arrays are displaced to find a new thermodynamically equilibrated configuration that as consequence reaches a new stable phase with a new chemical potential m modifying the cell voltage accordingly; (2) microscopically, the Na ion Oh symmetry is broken, causing lattice field interruption and the introduction of coulombic interaction between the two MO6 slabs. That is, the created Na+ ion vacancies contribute a lot to structural instability. Furthermore, by observing Fig. S3a (ESI†), some local microstructure with another M stacking mode can be seen, such as O2-like microstructures (M in ab-stacking and O in ABACstacking, see Fig. S5, ESI†), which are directly related with this local layer structural change/distortion. For the O3-NaFMO sample, the most observed phenomenon is the remarkable layer distortion, while for the P2-NaFMO sample, more information can be acquired from the HAADF images shown in Fig. 5b and Fig. S3b–e (ESI†). In Fig. 5b, the layer distance was measured to be 0.60 0.04 nm with the layer stripe clearly shown along the [100] direction. The fluctuation of layer distance seems smaller than that of the O3-NaFeMO sample along the [001] direction. However, the structural imperfections increase significantly. First, in the yellow arrow region of Fig. S3b–e (ESI†), it can be observed that there are some unexpectedly larger layer distances, which are nearly twice the regular dc value. This kind of layer expansion seems plausibly beneficial for Na+ ion diffusion. Furthermore Fig. S3b (ESI†) also demonstrates the layer stacking default, which gives birth to a ca. 51 angle as shown in the green and yellow underlined parts. This can be treated as the interface of two particles, which will merge together to form a new larger particle. In Fig. S3c (ESI†), a few intercalated layers are witnessed in the bulk of one particle, which are probably caused by Na–M antisites formed during annealing. Here, the adjacent coordination of Na+ ion changes from only the O ions to the O and M ions, which should play an important role in the evolution of structure and even battery failure upon sodiation/de-sodiation. The intercalated structures are also found in the surface region of the P2-NaFMO sample, which extends from the outmost surface to the bulk, about 20 nm far away from the surface as shown in Fig. S3d (ESI†). The O3-NaFMO particles display larger Na, M and O site fluctuations along the [001] direction than do the P2-NaFMO particles. Therefore, concerning surfacal structures as shown in Fig. 6, the O3-NaFMO sample demonstrates a remarkably different structure from that of the P2-NaFMO sample. From Fig. 6a and Fig. S4a and b (ESI†), the HAADF image of the O3-NaFMO sample is seen to exhibit ordered M columns aligned straightforwardly from the bulk to the surface, and thus so do the Na and O columns along the [100] direction (ABF image in Fig. S4c, ESI†). This gives a clear O3-NaFMO surface structure. However, a local region with negligible Na–M antisites was
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For the O3 phase, the O columns follow the ABCABC-stacking mode in the ABF image. The O3 structure demonstrates obvious site fluctuation with the layer distance dc B 0.61 0.10 nm along the [001] direction in the HAADF image. Furthermore, a clear and ordered atomic arrangement was observed with negligible Na–M antisites in several outermost surface layers. For the P2 phase, it was unambiguously observed the regular Na and O in ABBA-stacking, and M in aaa-stacking modes along the [001] direction. The P2 phase also demonstrates a regular layer distance with fewer fluctuations. Moreover, some obvious structural imperfections are detected accordingly from the HAADF and ABF images, including the stacking faults, intercalated structure and remarkable layer distortion. Compared with the O3 sample, the P2 phase Na2/3Fe1/2Mn1/2O2 sample demonstrates a surface structure with a large amount of Na–M antisites. Fig. 6 STEM HAADF images of the surface structures without (a) and with (b) a few Na–M antisites for the as-prepared O3-phase NaFe1/2Mn1/2O2; surface structures without (c) and with (d) obvious Na–M antisites for the as-prepared P2-phase Na2/3Fe1/2Mn1/2O2 at the [010] zone axis. The original pictures are presented in Fig. S4 (ESI†).
observed in Fig. 6b and Fig. S4d (ESI†), which is likely introduced by surface relaxation to excite the Na ions interchanging with M ions. Therefore, the M columns do not strictly follow the abg-stacking fault (Fig. S5a and S6a, ESI†) but have slightly disordered arrangements in several surface layers along the [001] direction. In contrast, this disordered stacking mode seems to facilitate the formation of Na–M antisites in the surface region after cycling.33 In Fig. 6c and d, the HAADF images reveal the surface structure without/with Na–M antisites in a P2-NaFMO sample. A typical difference is that the amount of Na–M antisites increases to a great extent and can be easily found in the available surface region when compared with that on the O3-NaFMO surface. In the enlarged HAADF image of Fig. S4f (ESI†), a large area with fully Na–M antisites was found to cover the surface of the P2-NaFMO sample. This can also be confirmed by the enhanced contrast of the Na columns in the ABF image of Fig. S4g (ESI†). Therefore, it shows that a clear surface is easily formed in the case of the O3-NaFMO sample, while in the case of the P2-NaFMO sample, a surface with significant Na–M antisites forms instead.
4. Conclusions O3 and P2 phase NaxFe1/2Mn1/2O2 samples were synthesized by a solid state method and investigated using aberration-corrected scanning transmission electron microscopy (STEM) with highangle annular-dark-field (HAADF) and annular-bright-field (ABF) techniques. The XRD patterns of the as-prepared samples can be well indexed to the O3 and P2 phases, although some minor impurities were present. The HAADF and ABF images were acquired with clear Na, M (Fe and Mn) and O occupations at the atomic scale. Some new insights concerning the structure and phase transitions during sodiation/de-sodiation are achieved:
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Acknowledgements This work was supported by funding from the ‘‘973’’ Projects (2010CB833102, 2012CB932900), NSFC (51222210, 11234013, 11174334), ‘‘Strategic Priority Research Program’’ of the Chinese Academy of Sciences (Grant No. XDA01020304) and One Hundred Talent Project of the Chinese Academy of Sciences.
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