Na-Ion Storage and Ultrastable Solid

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Exceptionally reversible Li/Na-ion storage and ultrastable solid-electrolyte interphase in layered GeP5 anode Safa Haghighat-Shishavan, Masoud Nazarian-Samani, Mahboobeh Nazarian-Samani, Ha-Kyung Roh, Kyung Yoon Chung, Si-Hyoung Oh, Byung-Won Cho, Seyed Farshid Kashani-Bozorg, and Kwang-Bum Kim ACS Appl. Mater. Interfaces, Just Accepted Manuscript • DOI: 10.1021/acsami.9b05900 • Publication Date (Web): 13 Aug 2019 Downloaded from pubs.acs.org on August 13, 2019

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ACS Applied Materials & Interfaces

Exceptionally reversible Li/Na-ion storage and ultra-stable solid-electrolyte interphase in layered GeP5 anode Safa Haghighat-Shishavan1, Masoud Nazarian-Samani1,2,, Mahboobeh NazarianSamani1, Ha-Kyung Roh1, Kyung-Yoon Chung3, Si-Hyoung Oh3, Byung-Won Cho3, Seyed Farshid Kashani-Bozorg2, Kwang-Bum Kim1,

1Department

of Materials Science and Engineering, Yonsei University, 134 Sinchondong, Seodaemoon-gu, Seoul 120-749, Republic of Korea.

2School

of Metallurgy and Materials Engineering, College of Engineering, University of Tehran, 11155-4563, Tehran, IR Iran.

3Center

for Energy Convergence, Korea Institute of Science and Technology, Hwarangro, Seongbuk-gu, Seoul 02792, Republic of Korea.

ABSTRACT In this study, we synthesized two layered and amorphous structures of germanium phosphide (GeP5) and compared their electrochemical performances to better understand the role of layered, crystalline structures and their ability to control large volume expansions. We compare results obtained with those of previous, conventional viewpoints addressing the effectiveness of amorphous phases in traditional anodes (Si, Ge, and Sn) to hinder electrode pulverization. By means of both comprehensive experimental characterizations and density functional theory (DFT) calculations, we demonstrate that

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layered, crystalline GeP5 in a hybrid structure with multiwalled carbon nanotubes (CNTs) exhibits exceptionally good transport of electrons and electrolyte ions and tolerance to extensive volume changes, and provides abundant reaction sites relative to an amorphous structure, resulting in a superior solid-electrolyte interphase (SEI) layer and unprecedented initial Coulombic efficiencies in both Li-ion (LIBs) and Na-ion batteries (NIBs). Moreover, the hybrid delivers excellent rate-capability (symmetric and asymmetric) performance and remarkable reversible discharge capacities, even at high current rates, realizing ultradurable cycles in both applications. The findings of this investigation are expected to offer insights into the design and application of layered materials in various devices.

Keywords: Germanium phosphide, Layered structure, Volume expansion control, Mass transportation, Solid electrolyte interphase, Energy storage.

1. INTRODUCTION Germanium–phosphorus intermetallics (germanium phosphides) have been recently introduced as excellent materials for batteries, photocatalysts, and photovoltaic devices on account of their layered graphene-like structure, remarkably high electrical conductivity, high carrier mobility, in-plane anisotropy, and adjustable band gap.1,2,3,4,5,6 In terms of energy-storage applications of LIBs and NIBs, GeP5 delivers the highest theoretical capacities of 2289 and 1888 mAh g−1, respectively, among various germanium phosphides, via combined conversion and alloying mechanisms;2,4 these values are much higher than the theoretical capacities of conventional transition-metal phosphides (e.g., 1280, 1350, 542, and 894 mAh g−1 for CuP2, FeP2, Ni2P, and CoP, respectively, for


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LIBs)7,8,9 because Ge, in contrast to transition metals, contributes to electrochemical charge-storage reactions during charging and discharging. In addition, the configuration of the layered structure of GeP5 is expected to offer several benefits: (1) additional Li+ or Na+ diffusion pathways, which greatly expedite the diffusion of guest ions; (2) outstanding in-plane electrical conductivity, which supports rapid electron transport, thus providing low internal impedance and insignificant polarization; (3) superb mechanical support for the crystalline GeP5 because well-preserved, robust in-plane chemical bonds enable superb mechanical flexibility, while weak interlayer van der Waals forces enable the layers to be freely expanded to accommodate volume expansions; (4) acceptable stability in air (i.e., exposure to O2 and H2O), in contrast to the susceptibility of layered black phosphorus to air exposure, which strengthens its status as a promising candidate for various energy-related applications. Recently, various GeP5 hybrid anodes including GeP5–graphite,2,4,10,11 GeP5–super P,12 GeP5/acetylene black/reduced graphene oxide,13 and GeP5–carbon nanofibers14 have been experimentally investigated for LIB and NIB applications. Zhou and co-workers were the first to introduce GeP5 and report the LIB and NIB performance of a GeP5–C composite.2,4 In their work, the GeP5–C electrode delivered a capacity of around 1500 mAh g−1 at 1.00 A g−1 after 150 cycles in LIB.2 In terms of NIB half-cells, the GeP5–C composite delivered constant discharge capacities of < 1400 mAh g−1 at a current density of 0.05 A g−1 for 120 cycles.4 Furthermore, Wei et al. developed carbon-encapsulated GeP5 nanofibers through combined ball milling and electrospinning approaches, exhibiting reversible discharge capacities of around 1000 mAh g−1 at 0.10 A g−1 after 200 cycles in LIB.14 In addition, Ning et al. suggested a dual-carbon network to improve the electrical conductivity and control the volume changes of GeP5 during Na+ insertion/de-



insertion processes, resulting in a 400 mAh g−1 capacity at a current density of 0.5 A g−1 after 50 cycles.13 The data obtained in the present study show noticeable improvement of battery performance over previous studies. More importantly, our research is the first study that comprehensively examines structural changes during the synthesis of GeP5 by ball milling and identifies the effects of the layered structure and chemical bonds on battery performance. We discovered, for the first time, that mechanically driven nanocrystallization occurs during prolonged milling and transforms the intermediate amorphous GeP5 to the fully crystalline and layered counterpart. In addition, the most salient novelty of the present study is that we innovatively introduce a completely new viewpoint regarding the utilization of layered structures as excellent active materials to attain ultra-stable battery performance, in comparison with the conventional viewpoint that states that amorphous components could fully control volume variations in anodes subject to very large volume changes during repeated charge/discharge processes.15,16,17 GeP5 is a suitable material with such a structural transition that allows us to design a direct comparison between layered and amorphous structures to reasonably compare two viewpoints under almost similar synthetic conditions. We developed a GeP5–MWCNT composite with robust Ge–O–C and P–O–C chemical bonds and greatly improved electrochemical performance, using a simple, scalable ball-milling synthesis method to fully exploit its outstanding advantages for both LIB and NIB applications. CNTs provide efficient electronic conduction pathways to GeP5 in the composites, mainly due to their one-dimensional (1D) morphology and high electrical conductivity. Network structure of CNTs effectively tolerates severe volume changes of GeP5 during repeated charge/discharge processes, thus providing structural


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stability of the GeP5/CNT composites. Furthermore, CNTs facilitate the formation of chemical bonds with GeP5 at the interfaces of GeP5 and CNTs, principally due to the high specific surface area, which maintains electronic conduction pathways to GeP5 even under large volume changes during repeated charge/discharge processes with Li+ or Na+ ions. We rationally developed two different structures—an amorphous structure of GeP5 and a fully crystalline, layered one—to clearly distinguish the changes in terms of various electrochemical properties. The layered structure was found to lead to unprecedentedly high initial Coulombic efficiency (CE) (94.75 and 94.13 % in LIB and NIB, respectively) and excellent structural reversibility, primarily because of its high in-plane electronic conductivity and abundant structural accessibility to guest Li+ or Na+ ions residing in the interlayer space. Finally, we used various characterization techniques to monitor the chemistry of the SEI layer in the layered GeP5 in the cycled NIB cells. Results revealed that these cells contained thinner and more stable SEI films than cells with amorphous GeP5 electrodes. The feasibility and scalability of the new synthesis method and the unrivaled configuration of the layered GeP5 presented here will facilitate the development of other layered materials for high-performance energy-related applications.

2. RESULTS AND DISCUSSION 2.1. Structural evolution during synthesis First, Raman spectra (Figure 1a) were acquired to trace structural changes during ball milling. Three strong modes of elemental red phosphorus (RP) could be clearly observed in the sample milled for 2 h at 355, 459, and 495 cm-1, indicating that elemental RP remains in this sample. In addition, two modes (at 298 and 414 cm-1) were associated with B3g and A2g of GeP5. After milling for 10 h, the three modes of amorphous RP had



completely disappeared and two extra-small A1g and B2g modes of GeP5 (at 192 and 238 cm-1, respectively) were detected, indicating that the initial elements were completely transformed to the final GeP5 phase. Importantly, further milling for up to 50 h resulted in some structural changes: (1) the intensities of the GeP5 modes were enhanced and the peaks became sharper and more intense, (2) additional GeP5 Raman modes were detected in the range of 300–400 cm-1, and (3) all four major modes of GeP5 were gradually subjected to blue shift, such that they moved toward higher wavelengths, as indicated by the dotted lines. It can thus be inferred that longer milling times alter the GeP5 structure, with a higher degree of crystallinity. In other words, the broad Raman modes and the obvious red shift in the sample milled for 10 h point to its amorphous structure, while powders milled for 50 h became crystalline.

Figure 1. (a) Raman, (b) DSC, and (c) Negative-ion ToF-SIMS characterizations of Ge–P powders that had been ball milled for 2, 10, and 50 h. (d) Schematic explanation of mechanically induced crystallization that occurs during the milling of Ge and RP elemental powders.


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We conducted additional characterizations to resolve uncertainties regarding the above-mentioned phase transitions, especially ambiguities concerning the primary amorphous RP phase and the intermediate amorphous GeP5 phase. First, differential scanning calorimetry (DSC) was conducted on the milled samples. Figure 1b shows that there is a broad exothermic peak in the range of 180–220 C for the two samples that were ball milled for 2 and 10 h, while this was not observed for powders milled for 50 h. This peak could be attributed to the crystallization transition of amorphous GeP5 in the two samples.18,19 Given that powders that had been milled for 2 h were mainly composed of the initial RP and Ge elements, the exothermic crystallization peak in the sample milled for 2 h was negligible. Further confirmation was attained from a comparison of the melting transition in the three samples. As shown in Figure 1b, samples milled for 10 and 50 h exhibited a similar sharp endothermic peak at 500 C, while the melting temperature of the sample milled for 2 h was completely different.18,20 The ball milled powders were also examined by time-of-flight secondary ion mass spectrometry (ToFSIMS), with results shown in Figure 1c. The spectra of samples that were ball milled for 2, 10, and 50 h exhibited a GeP ― peak located at m/z = 104.88, confirming the formation of a GeP5 phase. While the peak in the spectrum of the sample that was milled for 2 h is broad and weak, peak intensity in the spectrum of the sample obtained after 10 h of milling is much higher, while there is no noticeable change in the spectrum of the sample prepared with 50 h of milling. These results clearly show that elemental Ge and RP are initially transformed to an amorphous GeP5 phase, with further intensive mechanical deformation during high-energy ball milling leading to mechanically driven nanocrystallization that results in the nucleation and growth of a crystalline GeP5



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phase.18,19,20,21 As a result, the following phase transition was recognized as occurring during ball milling of Ge and RP elemental powders (Figure 1d): Amorphization

Ge + RP

Amorphous GeP5

Crystallization

Crystalline GeP5

(1)

Figure 2 shows bright-field transmission electron microscope (TEM) images and selected-area electron diffraction (SAED) patterns of Ge and RP components after milling for different periods of time. Figures 2a and 2b present typical TEM and high-resolution TEM (HRTEM) micrographs, respectively, of the sample that was ball milled for 2 h. Residual Ge nanoparticles were clearly detected in the structure, and their existence was confirmed through SAED patterns (Figure 2c). The patterns mostly consisted of halos and broad rings, which could be attributed to primary amorphous RP or amorphous GeP5. Barely visible weak rings associated with rhombohedral, layered GeP5 were also detected in this pattern. However, the Ge crystals completely vanished from the sample that was milled for 10 h (Figures 2d-f), and the only lattice fringes were those of GeP5 nanocrystals, which were fully embedded in the amorphous GeP5 matrix. The SAED pattern in Figure 2f shows a few weak crystalline rings and strong halos. The current results, therefore, suggest that it is almost impossible to produce the pure amorphous GeP5 phase, primarily because it necessitates phase evolution at the interfaces and grain boundaries of initial elements. Prolonged milling for 50 h changed all the amorphous regions to nanocrystals, and nanocrystallization resulted in a fully layered, polycrystalline sample, as shown in the HRTEM image in Figure 2h and the SAED pattern in Figure 2i. In addition, TEM elemental mapping analysis confirmed the uniform distribution of Ge and P throughout the samples, as shown in the mapping images in Figure S1. We selected two single-phase milled powders—amorphous (plus minor crystalline GeP5) and fully crystalline, layered GeP5 samples obtained after 10 and 50 h, respectively,


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of milling—for the rest of the experiments. Hereafter, these two samples are referred to as GPA and GPB (Figure S2). These two samples were mixed with MWCNTs in a weight ratio of 7:3 and ball milled for an additional 50 h, as explained in detail in the Experimental Section and Table S1. The two resulting composites, GPA–MWCNT and GPB–MWCNT, were labeled GPCA and GPCB, respectively. We recently demonstrated that prolonged milling is crucial to the formation of stable chemical bonds,22 in that milling results in the fracturing and shortening of CNTs. Ball milling also reduces the number of strong sp2-type C=C bonds and generates open-ended tips and various types of sidewall imperfections, such as vacancies, adatoms, edges, and interstitials (Figures S3-S6).23 These structural defects facilitate in-situ interfacial reactions and increase the number of available active sites, giving rise to firm cross-linking between CNTs through inter-tube dehydration, and between CNTs and GeP5 through two robust chemical bonds, namely, P–O–C and Ge–O–C. It is important to note that the amorphous nature of GPA and the crystalline characteristics of GPB remained almost intact during the second milling step (Figures S6 and S7). We believe that CNTs functioned as a buffer to absorb the high mechanical energy from the milling vial and balls and greatly hindered the transition from amorphous to crystalline GeP5.23,24



Figure 2. TEM, HRTEM, and SAED patterns of Ge–P samples that were ball milled for (a–c) 2 h, (d–f) 10 h, and (g–i) 50 h.

ToF-SIMS analysis was also utilized to probe the presence of chemical bonds between GeP5 and the CNTs in the composites (Figures S8a,b). Two types of chemical bonds, P–O–C and Ge–O–C, were detected at m/z = 58.97 and 101.92, respectively, in GPCA and GPCB, indicating that a robust bond existed between GeP5 and the CNTs, which was greatly beneficial for stable electrochemical properties. Moreover, X-ray 10 ACS Paragon Plus Environment

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photoelectron spectroscopy (XPS) confirmed the interaction between GeP5 and the CNTs. Figures S8c,d show high-resolution P 2p and Ge 3d spectra, respectively, of GPA, GPB, GPCA, and GPCB. When compared with P 2p and Ge 3d spectra of GPA and GPB, it can be seen that there are some peak shifts at binding energies of around 130 and 30 eV in the P 2p and Ge 3d spectra of GPCA and GPCB (highlighted by the dotted lines), signifying that surface interactions between GeP5 and the CNTs. Moreover, the spectra of GPA and GPB exhibit two small peaks at around 134 and 33 eV, which are likely attributed to surface oxidized P (POx) and Ge (GeO 3d), respectively. However, these peaks are much larger in the spectra of GPCA and GPCB, which means that they do not solely correspond to superficial oxides but could be associated with P–O–C and Ge–O–C bonds.

2.2. Electrochemical characterization in LIB Figure S9 presents cyclic voltammetry (CV) curves for the first 10 cycles of the GPCA and GPCB electrodes, respectively, at a scanning rate of 0.10 mV s−1 in the voltage range of 0.01–2.00 V versus Li/Li+. The CV profiles of the two electrodes exhibit almost similar trends, but the GPCB electrode exhibited better reversibility than the GPCA electrode, as also shown in Figure S10. Figure S11 illustrates typical charge–discharge voltage profiles at various cycles for both of the considered electrodes at the current density of 0.50 A g−1. The first charge/discharge capacities (based on the mass of GeP5, Figure S12) were 1927/2191 and 2239/2363 for the GPCA and GPCB electrodes, respectively, corresponding to initial CEs of 87.95% and 94.75%, respectively. Such high CE values in the GPCB electrode are crucial for practical application of this anode in LIBs for commercial purposes. Irreversible capacity is usually due to the formation of an



SEI film from electrolyte decomposition,25 and the SEI film is assumed to be negligible in this sample.

Figure 3. (a) Cycling performance and CE values of GPCA and GPCB electrodes over 350 cycles at 0.50 A g−1. Rate performance comparison of GPCA and GPCB electrodes at (b) symmetric charge–discharge current densities and (c) asymmetric tests. (d) Comparison of high-rate cyclic stability of GPCA and GPCB electrodes at 1.00 A g−1 for 100 cycles, followed by 900 cycles at 2.00 A g−1. The first ten cycles were carried out at 0.20 A g−1 for proper activation.

Figure 3a compares the cycling performance and CE values of the two composites at a current density of 0.50 A g−1. The discharge capacity of the GPCA electrode continuously decreased from the initial value of 2191 mAh g−1 to the final value of 1205 mAh g−1 after over 350 cycles, corresponding to 56.10% capacity retention from the second cycle. On the other hand, the discharge capacity was steadily stable for the GPCB electrode after an initial capacity fading with second discharge capacity of 2118 mAh g−1, which reached 1790 mAh g−1 in the 350th cycle, corresponding to 84.51% capacity retention (from the second cycle). The CE values of the two electrodes were higher than 99.50% after the third cycle and remained completely stable during the cycling tests. 12 ACS Paragon Plus Environment

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Figure 3b shows the discharge capacities of the two electrodes at different current densities in the range of 0.10–10.00 A g−1; it can be seen that when compared to the GPCA electrode, the GPCB electrode exhibited greatly enhanced rate capability. The GPCB electrode successfully delivered high capacities of 2685, 2266, 2143, 2049, 1933, 1836, 1677, 1480, 1222, and 821 mAh g−1 at 0.10, 0.20, 0.50, 1.00, 2.00, 3.00, 4.00, 5.00, 7.00, and 10.00 A g−1, respectively, and successfully bounced back with a slight capacity decrease to 2022 mAh g−1 at 0.50 A g−1. However, the GPCA electrode delivered a discharge capacity of 2488 mAh g−1 at 0.10 A g−1, decaying to 1122, 697, and 375 mAh g−1 at high rates of 5.00, 7.00, and 10.00 A g−1. In addition, the capacity did not return to the previous value at 0.50 A g−1. In addition, we tested the two electrodes in terms of their asymmetric rate capability, i.e., discharging at a constant current rate of 0.20 A g−1 and charging at increasing current rates in the range of 0.20 to 10.00 A g−1. The results are shown in Figure 3c, with a conspicuous discrepancy observed between the two samples. The GPCB electrode delivered charge/discharge capacities of 2437/2560, 2221/2248, and 2196/2219 mAh g−1 at charging current rates of 0.20, 0.50, and 1.00 A g−1, and remained completely stable up to 10.00 A g−1; these values are equal to 2010/2015, 1979/1987, and 1949/1958 mAh g−1 at charging current rates of 8.00, 9.00, and 10.00 A g−1, respectively. However, the GPCA electrode was only stable up to a charging current rate of 6.00 A g−1, and higher current rates resulted in unreliable capacity values. It can therefore be inferred that even though both of the composite materials, GPCA and GPCB, contained plenty of strong chemical bonds (Figure S8), the structural dissimilarity between them caused great differences in the cycling and rate capabilities of the composite electrodes. This difference can be even more clearly seen in the results of



high-rate cycling tests, as illustrated in Figure 3d. The GPCA electrode showed diminishing performance when cycled at 1.00 and 2.00 A g−1 for 100 and 300 cycles, respectively, while the GPCB electrode presented very stable reversible discharge capacities over 1000 cycles, with capacity retention of 84.77% after 100 cycles at 1.00 A g−1 (1755 mAh g−1 at the 100th cycle) and 72.94% after 900 cycles at 2.00 A g−1 (1237 mAh g−1 at the 1000th cycle).

2.3. Electrochemical characterization in NIB Figure S13 compares the CV curves for the first 10 cycles of the GPCA and GPCB electrodes, respectively, in the electrolyte containing Na+ ions. The curves were obtained at a scanning rate of 0.10 mV s−1 in the voltage range of 0.01–2.00 V versus Na/Na+. The CV curves of the GPCB electrode completely overlapped after the third cycle, indicating excellent electrochemically reversible charge and discharge reactions (Figure S14). Figure S15 presents the typical charge–discharge voltage profiles of the GPCA and GPCB samples after 1, 2, 5, 10, 50, 100, and 200 cycles at 0.50 A g−1. The GPCA electrode delivered a capacity of 1309 and 1440 mAh g−1 for the first charge and discharge, respectively, corresponding to an initial CE of 90.90%. Higher capacities of 1635 and 1737 mAh g−1 were obtained for the first charge and discharge cycles, respectively, of the GPCB electrode, corresponding to an initial CE of 94.13% (Figure 4a). As expected, the charge–discharge voltage curves of the GPCB electrode overlap each other and exhibit scarcely any decline, demonstrating the excellent cyclic stability of the electrode over 200 cycles, with a discharge capacity of 1430 mAh g−1 at the 200th cycle, equivalent to 91.43% capacity retention from the second cycle. Conversely, the GPCA electrode delivered a capacity of 1339 mAh g−1 at the second cycle; the capacity remained constant for only 50


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cycles and then gradually decayed to reach a value of 993 mAh g−1 after 200 cycles. As seen in Figure 4a, the CE values of both electrodes increased to 99.50% after only a few cycles and remained constant at around 100.00% in later cycles.

Figure 4. Comparison of cycling stability of GPCA and GPCB electrodes in the voltage range of 0.01– 2.00 V at a current density of (a) 0.50 A g−1 for 200 cycles and (b) 1.00 A g−1 for 100 cycles, followed by another 300 cycles at 2.00 A g−1. The first ten cycles in (b) were tested at 0.20 A g−1 for the activation process. Comparison of rate performance of GPCA and GPCB electrodes (c) at symmetric charge– discharge current densities and (d) in asymmetric tests.

We further evaluated the charge–discharge cycles of the two composites at high current densities of 1.00 and 2.00 A g−1 for 100 and 300 cycles, respectively (Figure 4b). Similar to its performance in the high-rate tests in the Li+ ions containing electrolyte, the GPCB electrode remained completely stable as it delivered an initial capacity of 1500 mAh g−1 at 1.00 A g−1, and the capacity remained almost constant after 100 cycles (1361 mAh g−1, equal to 90.73% capacity retention). A further 300 cycles at 2.00 A g−1 resulted in a high final capacity of 982 mAh g−1, implying excellent capacity retention of 75.28%



at this current density. However, the GPCA electrode delivered an initial capacity of 1210 mAh g−1 at 1.00 A g−1, which decreased to 825 mAh g−1 after 100 cycles. Further cycling at 2.00 A g−1 led to initial and final capacities of 791 and 319 mAh g−1, corresponding to capacity retentions of only 68.18% and 40.33%, respectively. Furthermore, the electrodes were evaluated at various current densities ranging from 0.10 to 5.00 A g−1, as illustrated in Figure 4c. The results confirm that the GPCB electrode delivered higher rate capability than the GPCA electrode. The GPCB electrode successfully delivered discharge capacities of 1798, 1681, 1571, 1461, 1381, 1251, 1175, and 1091 mAh g−1 at current densities of 0.10, 0.20, 0.50, 1.00, 2.00, 3.00, 4.00, and 5.00 A g−1, respectively, and rebounded to 1540 mAh g−1 when the current density switched back to 0.50 A g−1 after 48 cycles. In contrast, the GPCA electrode delivered 1660, 1506, 1360, 1171, 1040, 904, 758, and 650 mAh g−1 at the same current rates, and returned to 1350 mAh g−1 at 0.50 A g−1 after a similar number of cycles. In addition, the asymmetric rate capability of the two composites was evaluated at a constant discharge current density of 0.20 A g−1 and increasing charge current densities in the range of 0.20 to 10.00 A g−1, as shown in Figure 4d. It can be seen that the GPCB electrode showed completely stable performance up to 10.00 A g−1, reaching charge/discharge capacities of 1436/1456 mAh g−1 at this rate and delivered 1639/1659 mAh g−1 when the charging current rate switched back to 0.20 A g−1. The charge/discharge capacities of the GPCA electrode, however, exhibited a diminishing trend from 1480/1595 mAh g−1 at 0.20 A g−1 to 1046/1067 mAh g−1 at 10.00 A g−1, and increased to 1318/1338 mAh g−1 at 0.20 A g−1 after 72 cycles. It can thus be deduced that the crystalline, layered structure of GPCB (compared to the amorphous structure of GPCA) also made a large positive contribution to overall NIB electrochemical performance. Such ultradurable cyclic performance and rate capability at


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high charge–discharge rates of the GPCB electrode are among the best ever reported battery performances for various phosphides in both LIB and NIB applications.7,8,9 In addition, the electrochemical performance of all Ge-P intermetallics as anodes in LIB and NIB is compared with that of the GPCB electrode (Tables S2 and S3), which confirms the excellent LIB and NIB performance of the proposed material compared to that of previously reported intermetallics.

2.4. Benefits of layered structure of GeP5 There are several reasons for the exceptional electrochemical properties of the GPCB electrode. First, the robust P–O–C and Ge–O–C chemical bonds between GeP5 and functionalized CNTs afford a stable and flexible network to accommodate the intensive volumetric variations during the charge and discharge processes. The presence of CNTs is also crucial to guaranteeing the reversibility and recombination of the reactions during repeated charge and discharge processes.22,26,27,28 Second, compared to the amorphous structure in GPCA, the intrinsic layered structure of GeP5 in GPCB was electrochemically beneficial for kinetically easier Li+ insertion and extraction reactions and provided greatly enhanced pathways for ion transportation during the cycling tests, especially at high current rates. Third, it appears that the layered structure could effectively endure volume changes during cycling, and there was less pulverization of the electrode components in the layered crystalline structure than the amorphous counterpart. Fourth, the electrical conductivity of layered GeP5 is 10000 times that of layered black phosphorus and 10 times that of graphite.2 In addition, it is completely reasonable for the amorphous GeP5 structure to have lower electrical conductivity than the fully crystalline GeP5 structure.



The layered GeP5 efficiently supplied electron pathways, thereby significantly enhancing redox reactions during the charge and discharge processes in GPCB. To better comprehend the underlying reasons for the greatly enhanced electrochemical properties of GPCB compared to GPCA, especially for steady high-rate cycles, electrochemical impedance spectroscopy (EIS) analysis was conducted after various charge–discharge cycles; the typical Nyquist plots of the two composite electrodes are compared in Figures S16a,b. The EIS curves were analyzed using the complex nonlinear least square (CNLS) fitting method for the equivalent circuit shown in Figure S16c. Results revealed that the Re value (which includes the electrolyte, contact, and electrode resistances) of the GPCA electrode was approximately 5–6 times that of the GPCB electrode (22.00 vs. 4.00 ), and resistance increased with further cycling. Moreover, cycling up to 50 cycles yielded no obvious difference in the overall shape of the Nyquist plots of the GPCB electrode, while two noticeable semicircles appear in the high- and middle-frequency regions of the Nyquist plots of the GPCA electrode. The smaller semicircle in the high-frequency area characterizes the impedance of the SEI layer (RSEI), and the larger semicircle in the middle-frequency region describes the chargetransfer resistance at the electrode–electrolyte interfaces (Rct). In addition, the straight line in the low-frequency region, known as the Warburg impedance (Zw), is related to the diffusion of Na+ ions in the electrode. The RSEI and Rct values of the two composites are compared in Figure 5a. As mentioned earlier, the GPCB electrode displayed lower values for both RSEI and Rct than those displayed by the GPCA electrode: Rct of the GPCB electrode only slightly increased from 13.00  in the 1st cycle to 18.00  in the 50th cycle. However, Rct of the GPCA electrode in the 1st cycle was as high as 20.00 , increasing to 22.00  in the 2nd cycle. Extra cycling to the 5th cycle again increased Rct to 28.00 ,


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and continuous increases to 34.00 and 55.50  were observed at the 10th and 50th cycles, respectively.

Figure 5. (a) Comparison of Rct and RSEI of GPCA and GPCB at various cycles, derived from Nyquist plots. (b) EPR spectra of GPCA and GPCB in the charged state in the 50th cycle. (c) Schematic representation of mass transportation in amorphous and layered GeP5 during battery tests of the GPCA and GPCB electrodes.

Similarly, RSEI was 3.00  in the 1st cycle of the GPCB electrode and slightly increased in subsequent cycles (4.00, 5.00, 5.50, and 6.00  in the 2nd, 5th, 10th, and 50th cycles, respectively). In contrast, RSEI of the GPCA electrode was 6.00  in the 1st cycle, and increased to 7.00  in the 2nd cycle and further to 8.00, 11.00, and 22.00  in the 5th, 19 ACS Paragon Plus Environment


10th, and 50th cycles, respectively. This different change in RSEI mostly originated from the formation and stability of the SEI film on the surface of GPCB, which effectively enhanced the cycling performance over that of the amorphous GPCA electrode. In fact, it can be inferred that the layered structure of GeP5 could properly tolerate the extensive volume changes that occurred during insertion or extraction of Na+ ions. On the contrary, the amorphous GeP5 structure had relatively poor structural integrity to endure volume variations. This hypothesis will be further investigated in greater detail. To demonstrate the better charge transfer in the GPCB electrode than in the GPCA electrode during cycling, electron paramagnetic resonance (EPR) analysis was conducted at 4 K on the charged samples after the 50th cycle (Figure 5b), because a strong EPR peak is an indication of the conduction of electrons.29 Both of the mass-normalized samples exhibited a single Lorentzian line in a magnetic field, the strength of which varied from 3420 to 3450 G. The nearly symmetric shapes and almost similar widths of the spectra of both composites indicate that the electrons were fairly localized in the materials, with uniform distributions. Moreover, the results revealed stronger EPR spin intensity for GPCB, showing that there was a higher concentration of unpaired electrons or free radicals in this composite than in GPCA. Another difference was the shift of the GPCB spectrum to a higher g-value (2.0026 vs. 2.0038, respectively, Equation S1), suggesting that the electrons in GPCB were in a more electropositive chemical environment, thereby being able to occupy higher energy levels and thus more easily paired. This could result in enhanced cyclability and capacity of GPCB,22,30 in complete agreement with the EIS analysis and battery performance. Additional confirmation of the positive effect of the layered structure on eventual battery properties was obtained from a comparison of the Li+ and Na+ ion diffusion coefficients (D), determined from the CV curves at various


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scanning rates and the well-known Randles–Sevcik equation (Equation S2), as described in detail in Supporting Information (Figures S17-S19). The above-mentioned differences are schematically compared in Figure 5c, in terms of the inconsistencies in the transport of electrons and Na+ (or Li+) ions. In addition to the highly beneficial effect of the layered structure in providing rapid and stress-free transport of electrons and Na+ (or Li+) ions, it also effectively prevented irretrievable confinement of the Na+ (or Li+) ions and side reactions, eventually resulting in excellent and stable electrochemical properties. Since the sizes of Li+ and Na+ ions (1.52 and 2.04 Å, respectively) are sufficiently smaller than the interlayer spacing in GeP5 (3.09 Å), this large, inherent pathway was greatly advantageous to overcoming obstacles for Li+ and Na+ insertion, leading to completely reversible reactions with enhanced kinetics and insignificant volume variations during charge and discharge processes. This is in contrast to conventional electrode materials, in which the amorphous structure delivers a better battery performance than the crystalline phase.15,16

2.5. Exploration of SEI film chemistry and stability We conducted further TEM observations to visually explore the recharging products and the SEI film in the two composite electrodes after several cycles. First, two cells were compared in terms of the reversibility of the formation of GeP5 after recharging. HRTEM images, as well as SAED patterns after 50 cycles (in the charged state), revealed that the GPCA electrode predominantly contained amorphous phases (Figure 6a), and its SAED pattern presented broad, halo-like diffuse rings (Figure 6b). However, clear lattice fringes and strong spots and rings associated with the crystalline GeP5 phase were observed in the GPCB electrode, as shown in Figures 6d and 6e. Another important point was the



presence of CNTs in the images, which showed that they could retain their structure and not change to amorphous or disordered carbons. For additional confirmation, highresolution X-ray diffraction (HRXRD) analysis was conducted on the same electrodes (Figure S20), which showed obvious crystalline (003), (104), (015), (113), and (202) diffraction peaks of GeP5 in the GPCB, while only two (003) and (202) peaks were detected in the GPCA electrode with some peak shifts (structural distortions). Importantly, an amorphous halo in a 2 range of 28–36  appeared for both electrodes, suggesting that the GPCB electrode probably also has minor amorphous recharging products. It can thus be inferred that the initial layered structure of GeP5 results in more reversible reactions during durable cycling, guaranteeing excellent battery performance.


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Figure 6. HRTEM micrographs and SAED patterns of (a, b) GPCA and (d, e) GPCB after the 50th cycle in the charged state. TEM micrographs of (c) GPCA and (f) GPCB after the 50th cycle in the discharged state. The SEI layer in the two samples is outlined by dotted lines for comparison. Low- and highmagnification SEM images of cycled (g,h) GPCA and (i,j) GPCB electrodes in NIB after 50 cycles.

As illustrated in Figures 6c,f, the GPCA electrode consisted of an uneven amorphous SEI film after 50 cycles (discharged state), while the GPCB electrode contained a homogeneous, thin SEI layer around the active material, in complete agreement with EIS data. In fact, a thin passivating SEI layer appropriately protected the electrolyte from direct contact with the active material in the GPCB electrode and prevented extra electrolyte decomposition. On the other hand, inevitable severe volume changes and pulverization of GeP5 in the GPCA electrode resulted in continued growth and formation of the unstable SEI layer on newly exposed electrode surfaces, causing unremitting consumption of cations on the electrode–electrolyte interface owing to the incessant passivation reaction on the electrode’s surface. In addition, excessive growth of the SEI layer consumed a considerable amount of Na+ ions in the electrolyte and further obstructed electron conduction pathways, resulting in continuous increases in the impedance of the GPCA cell during cycling, as also demonstrated by XPS analysis (Figure S21). The presence of the stable SEI film was also demonstrated by SEM images of the cycled electrodes. After 50 cycles, the GPCA electrode had very large, continuous cracks (Figures 6g,h), while the surface of the cycled GPCB electrode was smooth and showed only tiny cracks (Figures 6i,j). In addition, positive-ion surface mode ToF-SIMS analysis was performed on the samples in discharge state after the 50th cycle to explicitly determine the chemistry of the SEI layer in the samples. As expected, the SEI layer predominantly consisted of inorganic



species such as NaOH, Na2O, NaCl, NaF, and Na2CO3. Figure 7a shows a few highintensity fragments in the range of m/z = 50–140: m/z = 61.97 (Na2O + ), 62.98 (Na2OH + ), 64.97 (Na2F + ), 77.97 (Na2O2+ ), 80.95 (Na2Cl + ), 84.96 (Na3O + ), 102.97 (Na2F3+ ), 104.97 (Na3FOH + ), 106.96 (Na3F2+ ), 108.95 (Na2PO2+ ), 112.96 (Na3CO2+ ), 124.94 (Na2PO3+ ), and 128.96 (Na3CO3+ ). Minor indications of organic species were also observed, such as fragments at m/z = 50.97 (C4H3+ ), 54.98 (C4H7+ ), 66.97 (C5H7+ ), 68.97 + + (C5H9+ ), 70.99 (C5H11 ), 71.98 (C4H8O + ), 76.97 (C6H5+ ), 82.95 (C6H11 ), 86.98 (C5H11

O + ), 90.98 (C7H7+ ), 116.97 (C4H2O2Cl + ), and 126.97 (C2H2O2Na3+ ). For the first time, it can therefore be inferred that the SEI film in the NIB cells of GeP5–CNT composites consists of both organic and inorganic components. It is worth mentioning that additional features were observed in Figure 7a: first, the intensity of organic components in GPCB is slightly higher than in GPCA, which might have partially affected enhanced NIB performance; second, consistent with XPS and EIS analyses and TEM observations, the intensities of inorganic fragments of GPCA are vividly higher than the intensities of those in GPCB, suggesting again that the SEI film in GPCA was thicker and less monotonic.

Figure 7. (a) Positive-ion ToF-SIMS of GPCA and GPCB samples of the 50th discharge cycle. (b) Schematic illustration of chemistry of SEI layer in GeP5–CNT composites using 1 M NaClO4 in EC–DEC electrolyte containing FEC; both organic and inorganic species were present.


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Figure 7b presents a simplified illustration of the SEI layer consisting of the abovementioned organic and inorganic species. The thinner SEI formed on GPCB along with higher content of organic components on its exterior interface could effectively enhance transport of Na+ ions and highly inhibit electrolyte decomposition on the surface of the electrode. In addition, the relatively uniform mixture and distribution of the two organic and inorganic components in the GPCB’s SEI layer offered better transport of Na+ ions. Another likely effect of the thinner, unbroken SEI layer in GPCB is better conservation of surface defects from irreversible reactions, as shown in the EPR spectrum (Figure 5b).

2.6. DFT calculations We eventually performed DFT calculations to better understand the reasons behind the excellent structural integrity of the layered GeP5 structure over ultra-long cycles, compared to the amorphous structure that failed to preserve its integrity (Figures 6g-j). Simulations of both layered and amorphous structures were conducted by adopting the GeP5’s rhombohedral structure as the layered counterpart, as well as its representative amorphous structure, constructed by ab-initio molecular dynamics (MD) simulations. Figures 8a,b illustrate the sodiation of GeP5 at the atomic scale for different NaxP and NayGe (x + y ≤ 4) up to full Na concentrations for Ge and P constituents (i.e. Na3P and NaGe) for both layered and amorphous networks, respectively. We calculated the volume changes of both layered and amorphous structures during sodiation up to full Na insertion (x + y = 4), with an almost linear increase with Na content (Figure 8c). Full sodiation caused a 386% volume expansion in the amorphous structure, but only 182% in the layered one. This is fully consistent with experimental characterizations and clearly demonstrates that pulverization due to volume variations



could be successfully alleviated to a greater degree in layered GeP5 than in its amorphous counterpart.

Figure 8. Structures of (a) layered and (b) amorphous GeP5 before and after Na insertion with different concentrations (considering NaxP and NayGe, and x + y ≤ 4). (c) Comparison of volume change ratios in Na-inserted layered and amorphous structures.

3. CONCLUSION In summary, we successfully designed a composite of layered, crystalline GeP5 and MWCNTs using a simple, large-scale ball-milling method. This composite consisted of robust P–O–C and Ge–O–C chemical bonds, resulting in greatly improved mechanical integrity of the final electrodes as anodes for LIB and NIB applications. The intrinsic layered nature of GeP5 with large interlayer spacing and its extremely high electrical conductivity provided ultra-efficient Li+/Na+ and electron pathways for electrochemical


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reactions. It also provided sufficient internal void space to adequately control large volume variations during charge and discharge processes, as demonstrated experimentally as well as through DFT calculations, causing a thin, monotonic SEI layer to form on the surface of the GPCB electrode. It was interesting to find that the SEI layer consisted of organic and inorganic species, both of which were beneficial for enhanced battery performance. As a result of these synergic characteristics, the anode electrode material delivered excellent LIB and NIB performance in terms of unprecedented initial CE, noticeable capacity, durable cycling, and high rate capability (symmetric and asymmetric evaluations) performance. This innovative fabrication approach is expected to offer general guidance for the design and engineering of other layered materials with highly exposed active sites for high-efficiency energy applications.

ASSOCIATED CONTENT Supporting Information Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsami. Experimental section, additional Raman, EELS, ToF-SIMS, TEM, EDS-TEM, SEM, XRD, and XPS analyses, multi-point BET plots, as well as CV curves, charge–discharge voltage profiles and EIS data in LIB and NIB.

AUTHOR INFORMATION Corresponding Authors [email protected] [email protected]



Author Contributions SHS, MNS, MNS, and KBK conceived the original idea. SHS, MNS, and MNS carried out all experiments, calculations, and characterizations and wrote the manuscript. All authors discussed the results and have approved this final edition. Notes The authors declare no competing financial interests.

ACKNOWLEDGMENTS This work was supported by the Korea Institute of Energy Technology Evaluation and Planning (KETEP) and the Ministry of Trade, Industry & Energy (MOTIE) of the Republic of Korea (No. 20172420108590) and the National Research Foundation of Korea Grant funded by the Korean Government (MSIP) (NRF-2011-0030542).

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Table of Contents Only

Layered germanium phosphide (GeP5): Provision of additional space with excellent properties in both Li-ion and Na-ion batteries.


Na-Ion Storage and Ultrastable Solid

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