Article pubs.acs.org/JPCC
Cite This: J. Phys. Chem. C XXXX, XXX, XXX−XXX
Lithium Imide (Li2NH) as a Solid-State Electrolyte for Electrochemical Energy Storage Applications Biswajit Paik*,† and Anna Wolczyk‡ †
WPI-Advanced Institute for Materials Research (WPI-AIMR), Tohoku University, Katahira 2-1-1, Sendai 980-8577, Japan Department of Chemistry and NIS, University of Turin, Via P.Giuria 9, I-10125 Torino, Italy
Downloaded via GRAND VALLEY STATE UNIV on January 11, 2019 at 14:10:56 (UTC). See https://pubs.acs.org/sharingguidelines for options on how to legitimately share published articles.
‡
ABSTRACT: We here report the prospect of lithium imide, Li2NH, as a solid-state electrolyte for the rechargeable Li-ion battery. Our study reveals that the ionic conductivity in Li2NH is greatly improved by the cation and anion ordering, which is a gradual process that is initiated by Li-filling in the LiNH2 lattice (forming nonstoichiometric Li1+xNH2−x) and sequentially forms the orthorhombic α-Li2NH (lower structural ordering) followed by the cubic β-Li2NH (higher structural ordering). The process, thus, enables synchronization between the cation and anion dynamics in β-Li2NH making this Li2NH phase as one of the fastest Liion conducting lightweight ionic alkaline hydrides. Our study also indicates high electrochemical stability for the Li/Li2NH interface in βLi2NH contradicting the existing perception. The study may aid the existing research in developing a cost-effective commercial Li superionic conductor with high electrochemical and thermal stabilities and high energy density at room- to moderate-temperature.
1. INTRODUCTION
Table 1. Reported Structures of Li2NH
Lithium imide (Li2NH) is an important member of the complex hydride Li-N-H system that has been studied intensely for more than a decade as a prospective reversible hydrogen storage material1,2 before being re-emerged as a promising solid-state electrolyte in the Li-ion battery owing to high Li-ion conductivity (>10−4 S/cm at room temperature), reported for the first time by Boukamp et al.3 nearly four decades back. Recent interest in Li2NH as a practical, costeffective, high-density energy storage solid-state electrolyte comes from a number of advantages with Li2NH that include, for example, lightweight (molar weight = 28.9), abundant constituent elements (specially, nitrogen), a thermally stable nitrogen sublattice, a very suitable anion dynamics for cation diffusion, comparatively easy synthesis methods, etc. Nevertheless, a large gray area exists in understanding the electrochemical properties of Li2NH, in general, and ionic conduction, in particular, given the possibility of several crystal structures under the name of lithium imide. Table 1 lists some of these Li2NH structures that have been reported by the first-principles calculations and by the experiments since the first report on the structural determination of the cubic Li2NH by Juza et al.4 The calculated Li2NH structures are predominately of low-symmetry.8,9 However, the experimentally determined structures at room temperature can be classified into two: (i) an orthorhombic phase reported by Balogh et al.7 and (ii) the cubic phases that appeared in several reports.4−6 For convenience, we refer the orthorhombic Li2NH as α-Li2NH and the cubic one as β-Li2NH. It is to be noted that a high-temperature (T > 350 K) Li2NH phase was also © XXXX American Chemical Society
references
structure (space group #)
Juza et al.4 Ohoyama et al.5
cubic Fm3̅m (#225) cubic Fm3̅m (#225), F4̅3m (#216) cubic Fm3̅m (#225) orthorhombic Ima2 (#46)
Noritake et al.6 Balogh et al.7
Magyari-Köpp et al.8 Mueller et al.9 Yang et al.10 Miceli et al.11
cubic Fd3̅m (#227) cubic Fm3̅m (#225) orthorhombic Pnma (#62) orthorhombic Pbca (#61) cubic Fm3̅m (#225) triclinic P1 (#1)
lattice parameters (Å) a = 5.047 a = 5.077 a = 5.074 a = 7.133, b = 10.087, c = 7.133 a = 10.129 a = 5.090 a = 7.733, b = 3.600, c = 4.872 a = 5.12, b = 10.51, c = 5.27 a = 5.09 a = 21.399, b = 10.087, c = 7.133
reported by Balogh et al.,7 which is structurally similar to the cubic β-Li2NH phase. As mentioned above, the possibility of more than one structure, with very limited studies of the ionic conduction in these structures, makes it difficult to understand the electrochemical properties (including the ion transport) in Li2NH. Further to these difficulties, a number of contradicting reports have been noticed on the structural transition between the orthorhombic and cubic Li2NH.7,12 The first-principles calculations adopted the orthorhombic α-Li2NH as the lowReceived: October 29, 2018 Revised: December 23, 2018
A
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C
2.5. Ionic Conductivity. The ionic conductivity was measured by an AC impedance spectrometer (Hioki 353280) by sandwiching the pressed Li2NH pellet (typical diameter 5−8 mm and thickness 1−3 mm, density >85% of the calculated value) between two lithium plates to collect the impedance of the pellet at ac frequencies ranging from 10−106 Hz at every 5−10 K temperature interval over a temperature window of 300−380 K. The Nyquist plot observed in the symmetrical cell Li/(Li2NH)/Li was fitted into a single R−C circuit to estimate the value of DC resistance R (in ohm), giving the ionic conductivity σ = t/(R·A), where t and A are, respectively, the thickness (in cm) and the circular flat-surface area (in cm2) of the pressed pallet. 2.6. Cyclic Voltammetry/Electrochemical Stability. We obtained the cyclic voltammogram (CV) of the β-Li2NH phase at 343 K using the lithium (the reference electrode) and molybdenum plate (the working electrode) to sandwich the pressed powder. Range of the voltage used was (−) 0.5 V to (+) 5.0 V. The voltage scan rate was 5 mV·s−1. We obtained CV for up to 5 cycles.
temperature Li2NH phase; and, the cubic Li2NH as the hightemperature Li2NH phase;8,9,11 whereas, experimentally, the cubic Li2NH has been well documented at room temperature.4−6 Ionic conductivity data for Li2NH have been obtained only with the cubic phase.3,13 No experimental study on the electrochemical stability has been carried out for any of the Li2NH structures; although a low electrochemical stability has been predicted for the cubic Li2NH by Boukamp et al.3 A detailed study to obtain information on several such aspects of structural and ion transport properties, however, are critical before Li2NH can be considered as a practical solidstate electrolyte for the Li-ion battery applications. In our study, we have been able to clarify the relationship between the cubic and the orthorhombic phases of Li2NH, especially, in terms of the structural evolution to show that the formation of the orthorhombic and the cubic phases is part of a continuous structural ordering. The study also unambiguously identifies that it is the orthorhombic Li2NH phase which undergoes the order−disorder phase transformation; although no abrupt change in ionic conductivity has resulted from this phase transformation as confirmed by the ionic conductivity estimated for these two structures over the temperature range of 300−400 K. The ionic conductivity data also help to compare the possible cation diffusion mechanism based on the available models. Ionic conductivity in cubic Li2NH is higher than that of the orthorhombic phase. The electrochemical stability study reveals that, unlike previous prediction,3 the cubic Li2NH has a wider electrochemical stability window for the Li/Li2NH interface. The study, thus, evaluates the electrochemical properties of Li2NH as a prospective solidstate electrolyte in the rechargeable Li-ion battery.
3. RESULTS AND DISCUSSION 3.1. Structural Characterization of α- and β-Li2NH. Figure 1 shows XRD patterns and Raman spectra of α- and βLi2NH. The XRD pattern of α-Li2NH matches closely with the reported structure that may be described with similar accuracy as a cubic fcc structure (Fd3̅m, #227, a = 10.129 Å) as well as an orthorhombic structure (Ima2, #46, a = c = 7.133 Å, b = 10.087 Å).7 On the other hand, the β-Li2NH reflections have close match with the fcc structure (Fm3̅m, #225, a = 5.074 Å) reported by Noritake et al.6 Except for the reflection from dhkl = 5.813 Å representing the (110) plane of the Ima2 structure of α-Li2NH, which stands out distinctly at 2θ ∼ 15° (see Figure 1a(i)), the XRD peaks obtained from the two Li2NH structures are closely spaced. Furthermore, the strongest reflection of α-Li2NH from the plane dhkl = 2.924 Å (the (022) plane in the Ima2 structure) is very close to the strongest reflection of the (111) plane (d111 = 2.929 Å) of βLi2NH as well as the (112) plane (d112 = 2.923 Å) of tetragonal LiNH2 (I4̅, #82). The Raman spectra of these two Li2NH (Figure 1c) show that α-Li2NH has two Raman active vibrations for the [NH]2− anion: a relatively sharper one at 3170 cm−1 and a broader one around 3230 cm−1. In contrast, β-Li2NH shows one [NH]2− vibration at 3165 cm−1. The appearance of the two vibrational modes, instead of one, in α-Li2NH synthesized for our present study supports the first-principles calculation by Herbst et al.,15 who, however, predicted the two peaks to be at 3199−3245 and 3272 cm−1. The previous experimental observations on the vibrational spectrum of Li2NH by different groups show singlepeak16 as well as twin peaks.17,18 We notice that the crystal structures (phase) of Li2NH, in these reports, have been mentioned to be cubic which, presumably, is the one reported by Juza et al.4 and is close to β-Li2NH in the present study. However, our study strongly indicates that the formation of αLi2NH gives rise to the twin peaks in the vibrations of the [NH]2− anion, suggesting two different chemical environments for the [NH]2− anion. In contrast, the appearance of the singlepeak in the [NH]2− vibration (typically, around 3165 cm−1) should be taken as a firm indication of β-Li2NH. 3.2. Evolution of the Li2NH Structures. From the above sections, one may appreciate the differences between α-Li2NH and β-Li2NH to distinguish these two Li2NH structures by
2. EXPERIMENTAL DETAILS 2.1. Sample Preparation. We prepared α-Li2NH by heat treating the ball-milled (5 h, Ar gas, 500 rpm) 1:1 molar mixture of Li3N (Aldrich) and LiNH2 (Aldrich) at 623 K for 1 h following the method reported by Hu et al. following the solid-state reaction LiNH2 + Li3N = 2Li2NH.14 LiNH2 was heat treated at 823 K for 48 h under dynamic vacuum to prepare β-Li2NH according to the synthesis reaction, 2LiNH2 = Li2NH + NH3.6 2.2. X-ray Powder Diffraction (XRD). The crystal structures of the Li2NH phases were determined by the Xray diffraction (PANalytical X’PERT with Cu Kα radiation) by collecting the XRD profiles at room temperature with transmission geometry over the 2θ range of 10−80°. 2.3. Raman Spectroscopy. Raman spectroscopy (Nicolet Almega-HD, Nd:YVO4 laser with a wavelength of 532 nm) was employed to collect the vibrational energies of the [NH]2− anions present in the Li2NH structures. The samples, taken in an Al pan, were loaded in a specially designed container in the glove box to prevent the samples from any exposure to the air/ oxygen during the experiment. 2.4. Differential Scanning Calorimetry (DSC). Differential scanning calorimetry (DSC) measurements were performed using the DSC 204 (Netzsch) running inside a glove box. For each measurement, about 10 mg of sample was sealed in an aluminum crucible under an inert atmosphere. A ramp rate of 1 K/min was used for heating up the α-Li2NH and β-Li2NH phases, respectively, up to 423 and 473 K, kept at the maximum temperature for 10 min and, thereafter, cooled to room temperature continuously under an Ar pressure. B
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C
structure (Ima2, #46) shown separately as a unit cell in Figure 1a. By filling all the vacant tetrahedral Li sites in α-Li2NH, the structure is transformed into β-Li2NH (Figure 2d). In fact, one in every eight tetrahedral Li sites (which is vacated by the Li cations on their migration to the neighboring octahedral positions in α-Li2NH) needs to be filled to transform α-Li2NH into β-Li2NH.7 Such part-filling and ordering of Li vacancy can develop different chemical environments for the [NH]2− anion in the Li2NH structure by perturbing the electron density on N atoms,21,22 supporting multimodal N-H vibrations in the Raman spectrum for α-Li2NH. Two N-H vibrational modes, as shown in Figure 1c, substantiate the above concept. A single Raman vibrational mode in Figure 1c for β-Li 2 NH, consequently, confirms identical and single chemical environment/coordination for the N atoms in β-Li2NH. 3.3. Order−Disorder Phase Transition in Li2NH. In relation to the structural transition of Li 2NH, a few contradictory observations have been reported. Balogh et al.7 have observed a reversible phase transition at ∼358 K where the α-Li2NH (Ima2) structure (also termed as LT-Li2NH) is found to transform into a cubic Fm3̅m structure (HT-Li2NH). The structure of HT-Li2NH, determined by Balogh et al.,7 is nearly identical to β-Li2NH. In contrast to this observation, Forman12 reported an order−disorder type phase transition at 356 ± 3 K in the Li2NH structure, which he claimed to match the XRD profile of the fcc Fm3̅m structure of Li2NH reported by Juza et al.4 To address this contradiction, i.e., to clarify which of the two Li2NH phases (α- or β-) show structural transition at around 350−360 K, we run DSC for the pristine α- and β-Li2NH. Figure 3 clearly shows an endo/exothermic peak at 353 K on heating/cooling for α-Li2NH by the heating rate of 1 K/min. No such peak was observed for β-Li2NH within the experimental temperature range of 300−473 K. Our DSC study strongly supports the observation by Balogh et al.7 that the endothermic peak at ∼350−360 K takes place in α-Li2NH (Ima2) when this phase is transformed into the cubic Fm3̅m. Apparently, the Li2NH prepared by Forman12 might have contained α-Li2NH which possibly was mistaken as the fcc Fm3̅m phase. Given the fact that α-Li2NH, which has (except the reflection from the interplanar spacing dhkl = 5.813 Å), otherwise, XRD profiles similar to that of β-Li2NH, was unknown at the time of Forman’s study. 3.4. Ionic Conductivity. The order−disorder type phase transitions in solids are known to alter physical properties including the ionic conduction. One such notable example in complex hydrides is LiBH4 which shows more than three-order of improvement in the ionic conductivity at 390 K undergoing an order−disorder phase transition (orthorhombic → hexagonal).23 In Li2NH, the order−disorder phase transition can be viewed as a result of the dissolution of the Li sublattice into a superionic phase, 24 expecting to improve the conductivity of the Li+ cation. From this specific point of motivation in the fundamental understanding of the Li2NH structures, we study the ionic conduction properties in αLi2NH and compare with that in β-Li2NH, as a part of their general ion transport characterization. Figure 4 shows the ionic conductivity measured for the two Li2NH structures over the temperature range of 300−400 K. The temperature dependence of ionic conductivity is related to the activation energy (Ea) by the relation σ = (A0/T) exp[−Ea/ (kB·T)], where σ is the ionic conductivity at temperature T, kB
Figure 1. (a) Crystal structure of α-Li2NH (Ima2, #46, a = 7.133 Å, b = 10.087 Å, c = 7.133 Å)7 and β-Li2NH (Fm3̅m, a = 5.074 Å);6 (b) XRD profiles and (c) Raman spectra of α-Li2NH and β-Li2NH structures shown in (a).
different characterization methods. However, it is also important to recognize how these two Li2NH structures are related in terms of the structural evolution, which may be viewed as a mechanism that transforms tetragonal LiNH2 into α- and β-Li2NH via the formation of nonstoichiometric Li-NH phases Li1+xNH2−x (0 < x < 1). The mechanism is based on the dehydrogenation model for LiNH2 proposed by David et al.19 According to this model, the filling up of the Li vacant sites in the LiNH2 lattice drives the structural transition from the Li-poor phase of tetragonal LiNH2 into the Li-rich phase of cubic Li2NH and, thus, transforming a poor Li-conductor (LiNH2) into a fast Li-ion conductor (Li2NH). Figure 2 shows a simple scheme of the mechanism,20 on the template of a superstructure of LiNH2 (Figure 2a), depicting the sequence that forms the nonstoichiometric cubic (Fm3̅m) structure of Li1+xNH2−x (Figure 2b) on the way to form first the α-Li2NH (Figure 2c) and eventually to form β-Li2NH (Figure 2d). This gradual structural change takes place by partially filling and ordering the remaining Li vacancies in a large unit cell of α-Li2NH that may be equivalently described either as a partly occupied cubic structure (Fd3̅m, #227, a = 10.129 Å) shown as the large cubic unit cell in Figure 2c for a convenient presentation; or, as a fully occupied orthorhombic C
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C
Figure 2. A simplified scheme to depict the structural change by filling up the vacant Li sites in the tetragonal LiNH2: (a) a superstructure of LiNH2 showing the Li-occupied (red sphere) and Li vacant (blue sphere) sites; (b) newly occupied Li sites (shown by the vertical arrows) locally form the fcc (Fm3̅m) unit cell (a × a × a) of the nonstoichiometric Li1+xNH2−x, the gray atoms/sites indicate that they are identical to those in (a); (c) the nonstoichiometric (a × a × a) unit cell co-exists with the (2a × 2a × 2a) unit cell of the α-Li2NH structure (Fd3̅m) formed by further filling up the Li vacant sites when the remaining vacant sites are orderly arranged (the occupied Li sites in this figure approximately represents only the tetrahedrally coordinated Li atoms of α-Li2NH);7 (d) Li sites are fully occupied to form the β-Li2NH. Hydrogen atoms are omitted for clarity.
Figure 3. DSC profile for α-Li2NH showing within the temperature window of 330−380 K. The Endothermic peak is at ∼352 K during heating/cooling (heating rate 1 K/min). Inset: the DSC profile for βLi2NH obtained for the temperature range of 300−473 K. Figure 4. Ionic conductivities of α-Li2NH and β-Li2NH estimated by impedance spectroscopy in a symmetric cell Li/(Li2NH)/Li. Typical Nyquist plot for α-Li2NH showing real and imaginary components of the complex impedances indicated, respectively, by Re Z and Im Z at temperatures 310, 315, and 330 K.
is the Boltzmann constant and A0 is the pre-exponential factor. The following general features are noticed in Figure 4: (i) the ionic conductivity of α-Li2NH is lower than that in β-Li2NH within the experimental temperature, (ii) estimated activation energy (Ea) in α- is larger than that in β-Li2NH (∼60 kJ/mol in α-Li2NH vs 52 kJ/mol in β-Li2NH), (iii) no jump in the
ionic conductivity is observed at the order−disorder phase transition temperature of ∼352 K (see Figure 3). D
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C Ionic conductivity in β-Li2NH agrees well with the previous reports by Boukamp et al.3 and Li et al.13 included in Figure 4 for comparison. The ionic conductivity data we obtained from this study suggest that Li2NH synthesized by Li et al.13 was the cubic (β) Li2NH phase (although, the conduction mechanism described in their study was for the orthorhombic phase). In our study, the room temperature ionic conductivity in βLi2NH (∼5 × 10−4 S/cm) is several times higher than that in α-Li2NH (∼10−4 S/cm). To achieve an ionic conductivity ≥10−2 S/cm the temperature in α- and β-Li2NH needs to be, respectively, above 360 and 400 K. 3.5. Li Diffusion in Orthorhombic (α)-Li2NH. Considering the activation energy to be the sum of the Frankel defect formation energy (Ed) and the energy for the cation migration (Em), the estimated activation energy 60 kJ/mol for α-Li2NH may be compared with the microscopic view of the ionic conduction in α-Li2NH. According to the model of Li ion conduction in orthorhombic Li2NH reported by Li et al.,13 the preferred diffusion mechanism of the Li cation, viz., the interstitialcy mechanism, has Em ∼ 50 kJ/mol. On the other hand, the Frankel defects may be formed via a number of mechanisms having the formation energy (Ed) ranging from as little as ∼1 kJ/mol to a high value >100 kJ/mol.13 Given the estimated Ea of 60 kJ/mol in our study (which may include ∼50 kJ/mol Em for the interstitialcy mechanism), we believe that in the αLi2NH compound the Frankel defect formation possibly involves a number of defect formation pathways that may need the formation energy of ∼10 kJ/mol. According to the model of Li et al.,13 the formation energy of ∼10 kJ/mol may prefer a specific defect formation mechanism that moves the Li+ cation from one specific tetrahedral site (termed as T4, by Li et al.,13 out of five symmetrically nonequivalent tetrahedral Li positions in α-Li2NH) to either another neighboring tetrahedral site (Ed ∼ 1 kJ/mol) or only to the octahedral site (Ed < 20 kJ/mol) in α-Li2NH. In practice, the defect formation may be a combination of these two mechanisms with the average defect formation energy, presumably, around 10 kJ/mol that may match an estimated Ea of 60 kJ/mol when added to the Li+ cation migration energy of ∼50 kJ/mol. Conversely, the estimated value of Ea in this study excludes the possibility of the other defect formation mechanisms (with high formation energies) as described by Li et al.13 3.6. Li Diffusion in Cubic (β-) Li2NH. The model to describe ionic conduction in the cubic Fm3̅m structure is available for the temperature above 400 K.24,25 Following the structural transition (LT-Li2NH → HT-Li2NH) at ∼350−360 K, where the LT-Li2NH phase is α-Li2NH and the HT-Li2NH phase has been identified as the cubic structure (identical to the β-Li2NH), these reports attempted to describe the dynamics of the ionic conduction in the cubic Fm3̅m phase existing above this transition temperature. This leaves βLi2NH, which exists at room temperature, without any model to describe the ionic transport. Nevertheless, combining the following observations that (i) β-Li2NH and the cubic phase above 400 K are structurally identical, (ii) the activation energy and, thereby, the ionic transport mechanism in β-Li2NH is likely to be unaltered over the temperature range of 300−400 K, as observed by us and by others,3,13 we may argue that the high ionic conduction in β-Li2NH is governed by the Li sublattice in coordination with the [NH]2− dynamics described in the HT-Li2NH structure.24 According to this model, the Li sublattice may be dissolved into a superionic phase in the
antifluorite (Fm3̅m) cubic structure in which the Li+ cations diffuse through the channels formed by the [NH]2− sublattice. However, our estimated activation energy (60 kJ/mol) for the β-Li2NH is higher than the calculated activation energy (5 V), which is in contrast to previously postulated weak electrochemical stability for cubic Li2NH. The study, thus, confirms the high prospect of Li2NH as an electrochemically stable fast Li-ion conductor suitable for use in the all-solid-state Li-ion battery at around room temperature.
5.0 V using Li as reference and molybdenum (Mo) as working electrodes (Figure 5). The anodic and the cathodic currents, at
■
AUTHOR INFORMATION
Corresponding Author
*E-mail:
[email protected]. ORCID
Biswajit Paik: 0000-0002-1631-7913 Figure 5. Cyclic voltammogram of β-Li2NH at 343 K in an asymmetric cell of Li/(β-Li2NH)/Mo.
Notes
The authors declare no competing financial interest.
around 0 V, are primarily due to the deposition (dissociation) of Li+ cation on (from) the Mo electrode. Although, there is no significant decomposition of this complex hydride as evidenced by the absence of additional anodic/cathodic current, we noticed an increase in the DC resistance during the charging/ discharging (slope in the anodic/cathodic current around 0 V). Compared with the dc resistance (several tens of ohm) in βLi2NH estimated by impedance spectroscopy, the estimated dc resistance (>500 ohm) during charging/discharging in CV is significant. The source of this additional resistance is likely to come from the electrolyte−electrode interfaces. We believe that the charge-transfer resistance at Li/Li2NH may not be negligible since we have noticed polarization at the lowfrequency range of impedance spectroscopy collected for βLi2NH. According to the predictions mentioned above,3,13 the Li/ Li2NH interface cannot survive a potential difference of >0.7 V between the two electrodes. Although, this value of electrochemical potential window is higher than that estimated for Li3N,29 our experiment gives hints to an electrochemical stability of β-Li2NH that is even better than the predicted 0.7 eV limit.3,13 In fact, the CV of β-Li2NH in Figure 5 shows that β-Li2NH may have a comparable electrochemical stability window demonstrated by some newly developed complex compounds30 and is larger than those of some well-known nitrides31 and oxides.32 Further investigation, especially, on the electrolyte−electrode interface will be critical to evaluate the electrochemical stability of Li2NH to propose a suitable set of electrodes for battery design.
■
ACKNOWLEDGMENTS
■
REFERENCES
Authors acknowledge the fund by the Japan Society for the Promotion of Science (JSPS) KAKENHI under Grant-in-aid Nos. 2522091, 26820311, and 26820313. A.W. acknowledges European Marie Curie Actions under ECOSTORE Grant Agreement No. 607040 for supporting this work. The authors also gratefully acknowledge the support from S. Orimo (Tohoku University, Japan) and M. Baricco (University of Turin, Italy) for allowing the laboratory facilities for the study.
(1) Chen, P.; Xiong, Z. T.; Luo, J. Z.; Lin, J. Y.; Tan, K. L. Interaction of hydrogen with metal nitrides and imides. Nature 2002, 420, 302−304. (2) Orimo, S.; Nakamori, Y.; Eliseo, J. R.; Züttel, A.; Jensen, C. M. Complex hydrides for hydrogen storage. Chem. Rev. 2007, 107, 4111− 4132. (3) Boukamp, B. A.; Huggins, R. A. Ionic conductivity in lithium imide. Phys. Lett. A 1979, 72, 464−466. (4) Juza, R.; Opp, K. Metallamide und metallnitride, 25. Mitteilung. zur kenntnis des lithiumimides. Z. Anorg. Allg. Chem. 1951, 266, 325− 330. (5) Ohoyama, K.; Nakamori, Y.; Orimo, S.; Yamada, K. Revised crystal structure model of Li2NH by neutron powder diffraction. J. Phys. Soc. Jpn. 2005, 74, 483−487. (6) Noritake, T.; Nozaki, H.; Aoki, M.; Towata, S.; Kitahara, G.; Nakamori, Y.; Orimo, S. Crystal structure and charge density analysis of Li2NH by synchrotron X-ray diffraction. J. Alloys Compd. 2005, 393, 264−268. (7) Balogh, M. P.; Jones, C. Y.; Herbst, J. F.; Hector, L. G., Jr.; Kundrat, M. Crystal structures and phase transformation of deuterated lithium imide, Li2ND. J. Alloys Compd. 2006, 420, 326− 336. (8) Magyari-Köpp, B.; Ozolinš, V.; Wolverton, C. Theoretical prediction of low-energy crystal structures and hydrogen storage energetics in Li2NH. Phys. Rev. B 2006, 73, No. 220101. (9) Mueller, T.; Ceder, G. Effective interactions between the N−H bond orientations in lithium imide and a proposed ground-state structure. Phys. Rev. B 2006, 74, No. 134104. (10) Yang, J.; Lamsal, J.; Cai, Q.; Yelon, W. B.; James, W. J. Study of the crystal structure and phase transition of Li2NH system. MRS Online Proc. Libr. 2008, 1098, HH03−06.
4. CONCLUSIONS To assess the prospect of Li2NH as a solid-state electrolyte for the rechargeable Li-ion battery, we have reported here a systematic study on the orthorhombic (α-) and the cubic (β-) phases of Li2NH. The study has enabled us to understand the structural evolution that leads to α- to β-phase transformation. A schematic model for the structural modification has been proposed to comply with the structural difference between these two phases observed by XRD and Raman spectroscopy. We also report ionic conductivity for these two phases over the F
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
Article
The Journal of Physical Chemistry C (11) Miceli, G.; Cucinotta, C. S.; Bernasconi, M.; Parrinnello, M. First principles study of the LiNH2/Li2NH transformation. J. Phys. Chem. C 2010, 114, 15174−15183. (12) Forman, R. A. Order−disorder transition in lithium imide, Li2NH. J. Chem. Phys. 1971, 55, 1987−1988. (13) Li, W.; Wu, G.; Xiong, Z.; Feng, Y. P.; Chen, P. Li+ ionic conductivities and diffusion mechanisms in Li-based imides and lithium amide. Phys. Chem. Chem. Phys. 2012, 14, 1596−1606. (14) Hu, Y. H.; Ruckenstein, E. Ultrafast reaction between Li3N and LiNH2 to prepare the effective hydrogen storage material Li2NH. Ind. Eng. Chem. Res. 2006, 45, 4993−4998. (15) Herbst, J. F.; Hector, L. G., Jr. Energetics of the Li amide/Li imide hydrogen storage reaction. Phys. Rev. B 2005, 72, No. 125120. (16) Chen, P.; Xiong, Z.; Luo, J.; Lin, J.; K. L. Tan, K. L. Interaction between lithium amide and lithium hydride. J. Phys. Chem. B 2003, 107, 10967−10970. (17) Kojima, Y.; Kawai, Y. IR characterizations of lithium imide and amide. J. Alloys Compd. 2005, 395, 236−239. (18) Anderson, P. A.; Chater, P. A.; Hewett, D. R.; Slater, P. R. Hydrogen storage and ionic mobility in amide−halide systems. Faraday Discuss. 2011, 151, 271−284. (19) David, W. I. F.; Jones, M. O.; Gregory, D. H.; Jewell, C. M.; Johnson, S. R.; Walton, A.; Edwards, P. P. A mechanism for nonstoichiometry in the lithium amide/lithium imide hydrogen storage reaction. J. Am. Chem. Soc. 2007, 129, 1594−1601. (20) Paik, B.; Matsuo, M.; Sato, T.; Qu, L.; Wolczyk, A. R.; Orimo, S. Effect of the structural evolution on the ionic conductivity of Li-NH system during the dehydrogenation. Appl. Phys. Lett. 2016, 108, No. 213903. (21) Wolczyk, A. Mixed Li-Based Complex Hydrides for Energy Storage. Ph.D. Thesis, Università degli Studi di Torino: Italy, 2017. (22) Kamakura, N.; Takeda, Y.; Saitoh, Y.; Yamagami, H.; Tsubota, M.; Paik, B.; Ichikawa, T.; Kojima, Y.; Muro, T.; Kato, Y.; Kinoshita, T. Electronic structure of lithium amide. Phys. Rev. B 2011, 83, No. 033103. (23) Matsuo, M.; Nakamori, Y.; Orimo, S.; et al. Lithium superionic conduction in lithium borohydride accompanied by structural transition. Appl. Phys. Lett. 2007, 91, No. 224103. (24) Araújo, C. M.; Blomqvist, A.; Scheicher, R. H.; Chen, P.; Ahuja, R. Superionicity in the hydrogen storage material Li2NH: Molecular dynamics simulations. Phys. Rev. B 2009, 79, No. 172101. (25) Blomqvist, A.; Araújo, C. M.; Scheicher, R. H.; Srepusharawoot, P.; Li, W.; Chen, P.; Ahuja, R. Hydrogen as promoter and inhibitor of superionicity: A case study on Li-N-H systems. Phys. Rev. B 2010, 82, No. 024304. (26) Wolczyk, A.; Paik, B.; Sato, T.; Nervi, C.; Brighi, M.; GharibDoust, S. P.; Chierotti, M.; Matsuo, M.; Li, G.; Gobetto, R.; Jensen, T.; Č erný, R.; Orimo, S.; Baricco, M. Li5(BH4)3NH: lithiumrich mixed anion complex hydride. J. Phys. Chem. C 2017, 121, 11069−11075. (27) Paik, B.; Hasegawa, T.; Ishii, I.; Michigoe, A.; Suzuki, T.; Ogita, N.; Udagawa, M.; Ichikawa, T.; Kojima, Y. The anharmonic vibration of Li in LiNH2. Appl. Phys. Lett. 2012, 100, No. 151911. (28) Paik, B. Modifying specific Li-sites of LiNH2 by Na: Study of multication hydride Li3Na(NH2)4 in electrochemical applications. J. Phys. Chem. C 2017, 121, 23906−23910. (29) Yonco, R. M.; Valeckis, E.; Maroni, V. A. Solubility of nitrogen in liquid lithium and thermal decomposition of solid Li3N. J. Nucl. Mater. 1975, 57, 317−324. (30) Kamaya, N.; Homma, K.; Yamakawa, Y.; Hirayama, M.; Kanno, R.; Yonemura, M.; Kamiyama, T.; Kato, Y.; Hama, S.; Kawamoto, K.; Mitsui, A. A lithium superionic conductor. Nat. Mater. 2011, 10, 682−686. (31) Rabenau, A. Lithium nitride and related materials case study of the use of modern solid state research techniques. Solid State Ionics 1982, 6, 277−293. (32) García-Martín, S.; Rojo, J. M.; Tsukamoto, H.; Moran, E.; Alario-Franco, M. A. Lithium-ion conductivity in the novel
La1/3−xLi3xNbO3 solid solution with perovskite-related structure. Solid State Ionics 1999, 116, 11−18.
G
DOI: 10.1021/acs.jpcc.8b10528 J. Phys. Chem. C XXXX, XXX, XXX−XXX
