1. Research background
2.1. Principle and examples of in-situ TEM analysis of diverse materials
There are various holders and in-situ observations to investigate real-time reactions of gas- and liquid- phase reactions and material synthesis like graphene. Among them, I used an open-cell device consisted of two electrodes for in-situ TEM analysis (Figure 19) to observe the lithiation process of battery materials in high-resolution. The nanobattery in the holder is composed of lithium loaded tungsten (W) tip as a counter/reference electrode and target sample loaded copper (Cu) tip as a working electrode.
The W tip side can be moved in 3D directions and thereby possible to contact to the sample side by a connected piezo controller. After attaching and holding two different electrodes, the holder can apply a relative electrical bias between them which is similar charge/discharge reactions in the battery system.
The generated thin Li2O film from Li metal by exposure to air for about 5 s can act as a solid electrolyte inside the TEM. When a relative bias is applied, Li ions and electrons can be transferred from counter electrode to working electrode or the opposite direction through the thin solid electrolyte layer.
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Figure 19. Image of in-situ TEM holder to investigate real-time reactions between battery materials and lithium (Dual-Probe STM-TEM in situ sample holder, Nanofactory Instruments).
i) Silicon nanotube (SiNT)
Among various strategies of structural engineering for alloying anode materials, applying inner pores to make enough space to expand inside is a distinct approach for suppressing volume expansion problem.
Halloysite which is an ideal raw material for that approach has rolled-sheets structure as alumino- silicate layers with inner void space (Figure 20). A typical silicon nanotube was obtained by calcination, etching the residual aluminum oxides, and reduction from SiO2 to Si. Furthermore, a thin carbon layer coated on the synthesized silicon nanotube to overcome its intrinsic chemical reactivity with Li ions and electrons (Figure 21).
Figure 20. A schematic of structural description of halloysite nanotube clay.
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Figure 21. A schematic of the process of synthesizing carbon coated silicon nanotube.
To identify the volume expansion ratio of the synthesized silicon nanotubes, comparable in-situ TEM analysis was carried out to measure the volume changes during the lithiation process. In silicon nanoparticle (SiNP) analysis, one of particles showed a crack at the interface even after the first charge process. The average volume expansion of nanoparticles was about 400% which is slightly larger than the literatures. (Figure 22a). However, a single representative silicon nanotube which has a 100 nm x 410 nm (diameter x length) showed only 210% volume expansion without any cracks on the interface after the first charge process. This different stability was impressive because, in general, larger particles (SiNTs in this case) can be broken easily during the lithiation process. The void space of silicon nanotube was effective in suppress the outer expansion and maintaining its original structural integrity.
(Figure 22b)
Figure 22. Comparable in-situ TEM analysis of volume expansion of (a) silicon nanoparticle (SiNP) and (b) silicon nanotube (SiNT).
42 ii) Red phosphorus (RP)
Red phosphorus is one of the amorphous allotrope of phosphorus which has distinct an advantage that can react with diverse alkali ions. It can make alloys and store high theoretical capacity by reacting with Li ions (2,500 mAh g-1), Na ions (2,500 mAh g-1), and K ions (850 mAh g-1). This alloying material also shows crack and pulverization problem during the charge process with different alkali ions.
Although a lot of strategies such as particle-size control and recrystallization of RP were applied to decrease the problem, the relationship between the electro-, chemo-, and mechanical properties of RP during the reaction with diverse ions is not clear. The diffusivity of Li, Na, K ions during the reaction can be inferred by the measurement of volume expansion of RP using in-situ TEM analysis and the simulation model of stress evolution at the interface (mechanical property change) can be established by the measured diffusivity value (Figure 23).
Figure 23. An example of the correlation establishment process.
To establish the different stress evolution into RP particles, in-situ TEM analysis of charging process of small spherical (near 100 nm) RP with Li, Na, and K ions was conducted. In all small spherical particles, a similar expansion of both axes was observed during alloying reactions with three alkali ions which confirmed that the expansion of small spherical RP was isotropic and no clear cracking process (Figure 24).
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Figure 24. In-situ TEM analysis of small spherical RP during the charging process with different alkali ions (a) Lithium (b) Sodium (c) Potassium.
Volume expansion ratio of RP after reacting with different alkali ions was calculated by measuring particle size before and after reaction (Figure 25). In the lithiation process, both axes were increased about 27 ± 4% evenly, resulting in an average volume expansion of about 105 ± 20%. In the sodiation process, the expansion value was about 47 ± 9% in axes and 217 ± 60% in a volume expansion. In the potassiation process, the expansion value was about 43 ± 10% in axes and 192 ± 60% in a volume expansion. The large deviation value in sodiation and potassiation was caused by poor image contrast due to the overlapped particles after volume expansion. The calculated volume expansion value of RP was different with the theoretically calculated expansion ratio of reacting with each alkali ion (Table 1).
These results could be caused by the native surface oxide layer, but can suggest that small and spherical properties can suppress large volume expansion and crack formation.
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Figure 25. Measurement of size changes of small spherical RP using in-situ TEM analysis during the charging process with different alkali ions (a) Lithium (b) Sodium (c) Potassium.
Table 1. Theoretical volume expansion ratio of RP during the charge process with different alkali ions (a) Lithium (b) Sodium (c) Potassium.
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In-situ TEM analysis of ball-milled large commercial RP particles was also conducted to compare with the small spherical RP. Representative morphological properties of the particles were not spherical, larger (over 200 nm), and rough interface. The calculated volume expansion ratio of commercial RP during the charge process with different alkali ions was quite different with the spherical particles. The measured volume expansion ratio was 72%, 179%, and 217% for lithium, sodium, and potassium, respectively (Figure 26). The different volume expansion values can occur due to the shape of particles which has different lengths in the 3D directions or presence of a thick oxide layer at the interface. Above all, the apparent cracking process was observed during the lithiation and potassiation process (white arrows in images, Figure 26a and 26c). In particular, the crack formed during the potassiation process grew to near 50 nm after fully reacting with K ions (Figure 26c).
Figure 26. In-situ TEM analysis of large commercial RP during the charging process with different alkali ions (a) Lithium (b) Sodium (c) Potassium.
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Based on the in-situ TEM analysis, three different model of alloying reactions in RP can be established (Figure 27). In this system, alkali ions can be transferred from the point that two electrodes of holder attached and it can make positional gradient of ions (ion flux distribution) inside the particles. In small spherical RP case, the effect of ion flux distribution inside the particles can be relieved easily because of the isotropic shape as well as the case of ion transport through a short-edge part of the particles (Figure 27a and 27b). However, high tensile stress (red color) can be occurred when ions entered through the narrow interface that leads to crack and fracture of particles (Figure 27c) as shown in the in-situ TEM analysis.
Figure 27. Three different alloying model of RP derived by in-situ TEM analysis.
iii) iron fluoride (FeF2)
The application of conversion cathode materials as high-energy density materials (~570 mAh g-1) for lithium-ion batteries has been hindered by less understanding of their reaction mechanisms during the cycling. Single crystalline iron (II) fluoride (FeF2) nanorods with an ionic liquid (1m LiFSI/Pyr1,3FSI) suggested that a proper environment for reaction mechanism studies of reversible conversion reactions (Figure 28). A stable solid electrolyte interphase formed by the ionic liquid electrolyte made the state more reversible through the interfacial ability to prevent the fusing of FeF2 particles. In-situ TEM analysis was conducted to compare the difference of reversibility between the conventional electrolyte system and the ionic liquid electrolyte system focusing on the interfacial difference.
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Figure 28. HRTEM images of synthesized FeF2 nanorod with different zone axis. Scale bars (a) 5nm, (c) 5nm.
FeF2 conversion materials show the phase separation process during reacting with Li ions into metallic Fe and LiF matrix (FeF2 + 2Li++ 2e- → Fe + 2LiF). Once the phase separation occurs, it is hard to come back to the original structural state due to the high activation energy for breaking formed LiF matrix.
FeF2 nanorods after the first discharge process in the conventional carbonate electrolyte system (1m LiPF6 in 1:1 ethylene carbonate (EC): dimethyl carbonate (DMC), LP30) showed particle aggregation that can decrease the possibility of sufficiently reconversion to the inside (Figure 29a). Observation of the discharge process using in-situ TEM analysis also revealed that two nanorods reacted as one particle across interparticle boundaries to form a large particle (Figure 29b).
Figure 29. Ex-situ and in-situ TEM analysis of the discharge process of FeF2 nanorods. Scale bars (a) 200 nm, (a inset) 10 nm, (b) 50 nm.
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FeF2 nanorod after the first discharge process in the ionic liquid electrolyte system (1m LiFSI in Pyr1,3FSI) showed distinct SEI layer derived by the ionic liquid electrolyte (Figure 30a). Unlike the carbonate electrolyte system, FeF2 nanorods can maintain their single-particle level and electrostatic interaction among particles through the strongly bound interfacial layer during cycling. Furthermore, reversibility of the conversion material was significantly improved due to its high ionic conductive and robust properties (Figure 30b and 30c). An intact interfacial layer formed by the ionic liquid electrolyte during the discharge process (Figure 30a and 4 in 30c) was observed after starting the charge process (5 and 6 in Figure 30c) which can increase the reversibility of the conversion materials. Although in- situ TEM measurements could not directly show the formation of SEI layer and improvement of reversibility from the ionic liquid electrolyte system, the suggested reaction pathways along the grain boundaries among nanorods and particle aggregation could provide the reason of irreversibility of diverse conversion materials.
Figure 30. (a) Ex-situ analysis of FeF2 nanorod after the first discharge process. (b) The galvanostatic and corresponding dQ/dV profile for the first discharge process. (c) Ex-situ analysis of FeF2 nanorod
of each state in (b). Scale bars (a) 200 nm, (a inset) 10 nm, (c) 20 nm.
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2.2. New insight into how inter-particle interaction can impact the battery performance by analysis of the different reaction mechanism
Alloying materials have gathered a lot of interest as a promising anode material for the next-generation of high-energy density battery system. However, they show huge volume expansion during the lithiation process, resulting in cracks and pulverization of the particles. During the repeatable cycles, loss of electrical contact among particles and harmful side reactions can occur because of the formed cracks and pulverized particles. There are numerous efforts to solve the issues above such as size control, applying hollow or porous substrate, and adopting coating layer. Applying coating layers on alloying anode materials provide the most effective improvement to tolerate stress and mitigate the side reactions with liquid electrolyte inside the battery system. The surface layer could mechanically suppress the volume expansion during the cycles and enhance the electrical properties for transportation of electrons and Li ions. [16 – 24]
In general, carbonaceous materials have been widely used for improving electrochemical properties of alloying anode materials which usually have low electrical conductivity. Other coating materials like conducting polymers or metal oxides also were mainly focused on the enhancement of their electrical properties. Although those materials contributed to some extent in suppressing volume expansion of alloying materials, most of studies were focused on the improvement on single particle level. However, the interparticle interaction is as important as the enhancement of electrochemical performance on single particle level because a lot of particles aggregate to form an electrode inside the battery system.
In fact, the role of interparticle interaction of alloying materials with diverse coating layers remains less understood.
Diverse coating layers may show different interparticle interaction because of their intrinsic physical properties such as flexibility and hardness. I conducted in-situ TEM analysis of tin oxide (SnO2) with polymeric coating layer (polypyrrole, ppy) and inorganic (manganese oxide, MnO2) layer focusing on the interparticle interaction. Real-time lithiation process of SnO2 cluster with two typical flexible (ppy) and rigid (MnO2) coating layers was observed using in-situ TEM technique focusing on their contribution to stable cycling as well as interparticle interaction in a cluster (Figure 31).
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Figure 31. A schematic of in-situ TEM analysis of SnO2 alloying material with different coating layers.
51 2.2.1. Experimental section
i) Synthesis of SnO2@MnO2, and SnO2@ppy nanoparticles
200 mg of SnO2 powders were dispersed in 50 mL of deionized water and sonicated them for 20 min.
40 mg of cetyltrimethyl ammonium bromide (CTAB) and 1 mL of ammonia solution (28%) were added into the sonicated batch and stirred for 30 min. Then, 100 mg of resorcinol and 140 μL of formaldehyde solution (37wt%) were added and stirred for 5 h to form SnO2@C nanoparticles. A final carbon coated SnO2 nanoparticles were obtained after calcination under an Ar atmosphere at 700℃ for 2 h. The obtained carbon coated SnO2 nanoparticles were redispersed in 50 mL of deionized water and 80 mg of KMnO4 was added. The solution was transferred to a 100 mL of Teflon-lined reactor after 5 min stirring.
Kept the heating temperature at 180℃ for 8h and cooled down to room temperature. The SnO2@MnO2
nanoparticles were obtained by centrifuge after washing with deionized water 3 times and overnight drying at 70℃.
SnO2@ppy nanoparticles were synthesized through an in-situ chemical-polymerization method. 200 mg of SnO2 powders with 40 mg of dodecyl sulfonate were dispersed in 50 mL of deionized water and sonicated the solution for 20 min. After continuous stirring for 3 h, 15 μL of pyrrole monomer was added and stirred for 30 min more. Then, 1 mL of solution containing 150 mg of ammonium persulfate was added by dropwise. After continuous stirring for 4 h, SnO2@ppy nanoparticles were obtained by centrifuge after washing with deionized water 3 times.
ii) Characterizations
Transmission electron microscopy (TEM, tecnai G2 F20 X-TWIN, FEI) was used to observe the morphology and structure of bare SnO2 and synthesized products. X-ray diffraction (XRD) patterns were obtained using Kα radiation (40 kV, 40 mA) on an X-ray diffractometer (Ultima III, Rigaku, Japan).
X-ray photoelectron spectroscopy (XPS, K-Alpha, Thermo Scientific, U.K.) was used to observe the elemental composition of the samples.
iii) Battery performance test and electrochemical measurements
A slurry for electrodes of bare SnO2, SnO2@MnO2, and SnO2@ppy were composed of carbon black and sodium carboxymethyl cellulose (CMC) binder at a mass loading ratio of 8:1:1 with NMP solvent.
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The slurry was cast onto copper current collector with mass loading of 1.0 – 1.5 mg cm-2, and the cast slurry was dried overnight. Coin-type (2032) half cells were fabricated in glove box under Ar atmosphere using Li metal as counter/reference electrode and 1.0 M LiPF6 in 1:1 w/w ethylene carbonate (EC)/diethyl carbonate (DEC) was used as an electrolyte. Galvanostatic cycling tests were conducted between 0.005 – 3.0 V using Neware battery cycler (CT-4008T-5V10mA-164, Shenzehn, China). Cyclic voltammetry (CV) and electrochemical impedance spectroscopy (EIS) (from 100 kHz to 10 mHz) tests were conducted using Bio-logic VMP3.
iv) In-situ TEM analysis
In-situ TEM analysis was carried out at an acceleration voltage of 200 kV in same TEM (tecnai G2 F20 X-TWIN, FEI) as used for the morphological analysis. Dual-Probe STM-TEM in situ sample holder (Nanofactory Intruments) was used to apply bias between nanobattery inside the TEM. The nanobattery was composed of Li metal counter/reference electrode on tungsten (W) tip and SnO2, SnO2@MnO2, and SnO2@ppy nanoparticles on copper (Cu) tip. A formed thin (~20 nm) Li2O layer on Li metal electrode during exposure to air for 5 s was performed a role of conductive solid electrolyte. A relative electrical bias of 1.0 – 1.5 V was applied between the two electrodes to transfer Li ions and electrons through the solid electrolyte layer. It is possible to observe the real-time changes during the lithiation/delithiation process with this mimicked nanobattery on the in-situ TEM holder. For in-situ voltage scanning and flexible test experiments, Pt/Ir STM tip and Au tip which have lower electrical resistance were used instead of W and Cu tip to measure more exact resistance values of SnO2, SnO2@MnO2, and SnO2@ppy nanoparticles/cluster.
53 2.2.2. Results and discussions
i) Synthesized SnO2@MnO2 and SnO2@ppy nanoparticles
Figure 32a – 32c were schematics and low magnification TEM images of bare SnO2, SnO2@MnO2, and SnO2@ppy nanoparticles. Bare SnO2 nanoparticles in Figure 32a showed a smooth and clean surface with the average size of 50 – 100 nm. In Figure 32b, fine 5 – 10 nm of MnO2 coating layers derived by KMnO4 precursor were uniformly deposited on bare SnO2 nanoparticles. A relative thick (~20 nm) polypyrrole layers from in-situ chemical polymerization were also well coated on each nanoparticle. The average particles size did not change a lot from 50 – 100 nm even after adopting two different coating layers as shown in Figure 32b and 32c.
Figure 32. A schematic and low magnification TEM images of (a) bare SnO2, (b) SnO2@MnO2, and (c) SnO2@ppy nanoparticles.
The crystal structure of bare and synthesized products was analyzed using an X-ray diffractometer (XRD) as shown in Figure 33a – 33c. There is no different distinct diffraction pattern with SnO2
reference (JCPDS no. 77-0450) in every XRD pattern except the SnO2@MnO2 nanoparticles. Only small amount of additional diffraction peaks derived from β-MnO2 (JCPDS no. 44-0141) were observed in the inorganic coated SnO2 (Figure 33b). Almost same diffraction pattern property of SnO2@ppy with a bare SnO2 sample indicated that polymeric coating layers are amorphous shell (Figure 33c).
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Figure 33. XRD patterns of (a) bare SnO2, (b) SnO2@MnO2, and (c) SnO2@ppy nanoparticles.
In X-ray photoelectron spectroscopy (XPS) was conducted to verify the elements’ chemical composition at the surface of synthesized samples (Figure 34). Strong peaks related to Mn2p at near 642 eV with Sn3d indicated MnO2 coating layers were successfully deposited on the surface of SnO2
nanoparticles (Figure 34a). In Figure 34b, much stronger peaks of C1s at near 284 eV with N1s at near 400 eV which are derived from polypyrrole were observed in SnO2@ppy nanoparticles. The reason that a typical Sn3d peak did not show in Figure 34b maybe because the formed ppy shell was too thick.
Figure 34. XPS results of (a) SnO2@MnO2 and (b) SnO2@ppy nanoparticles.
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ii) Electrochemical performance of bare SnO2, SnO2@MnO2, and SnO2@ppy nanoparticles Electrochemical half-cell tests of bare SnO2 and SnO2 with different coating layers were conducted using coin-type (2032) cell. The battery cycling test was tested under 0.005 – 3.0 V range which is within the reaction region of SnO2. Figure 35 showed the 1st and 10th charge/discharge profiles from the half-cell tests of (black) bare SnO2, (blue) SnO2@MnO2, and (red) SnO2@ppy electrode, respectively. All specific capacities of electrodes (x-axis in Figure 35a) were calculated based on the total mass of SnO2 of each sample. Although three samples showed similar specific capacity at the 1st cycle (780 mAh g-1), the delithiation capacity of bare SnO2 decreased rapidly to near 270 mAh g-1 at the 10th cycle maybe due to the cracking and electrical loss of active materials on the electrode.
However, the delithiation capacity of SnO2 with coating layers after 10th cycles maintained well near 680 mAh g-1 in SnO2@MnO2 and almost similar capacity in SnO2@ppy electrode. This capacity comparison result after 10th cycles indicates that polypyrrole coating layer was more efficient to maintain the original capacity of SnO2. Furthermore, cyclic voltammetry (CV) results suggested that coating layers did not show any electrochemical activity to increase the specific capacity in the same voltage range with the galvanostatic cycle Figure 35b).
Figure 35. (a) The 1st and 10th charge/discharge profiles of (black) bare SnO2, (blue) SnO2@MnO2, and (red) SnO2@ppy electrode (b) Cyclic voltammetry results of each electrode.
The SnO2@ppy electrode exhibited much superior rate capability than bare SnO2 and MnO2 coated SnO2 electrode. At the first (0.1 C), the delithiation capacity was 780 mAh g-1 and slightly decreased with the C rate increased. However, relatively higher capacities than bare SnO2 and SnO2@MnO2
maintained well as 676, 614, and 530 mAh g-1 at 0.2, 0.5, and 1.0 C rate (Figure 36a). The specific capacity showed near 440 mAh g-1 at 2.0 C rate which is even higher than the theoretical capacity of