Pomegranate-like C@TiO₂ mesoporous honeycomb spheres for high performance lithium ion batteries

The rapid development of portable electronic devices and electronic vehicles has driven an ever-increasing demand for high-performance lithium ion batteries (LIBs) [1–4]. Due to the low theoretical capacity and slow rate capability of commercial graphite, it is urgent to develop advanced anode materials to replace the traditional graphite. Transition metal oxides (TMOs) have attracted great attention as potential alternative anode materials due to their high theoretical capacity [5–9]. Among them, TiO₂ has been considered as one of the most promising anode materials for its merits, such as low cost, natural abundance, non-toxicity and certain stability [10–13]. Particularly, the high working potential of 1.7 V vs Li/Li⁺ can avoid the decomposition of electrolyte and the formation of SEI film, improving the safety of LIBs. Nevertheless, several intrinsic drawbacks, e.g. low electronic conductivity, sluggish Li⁺ diffusion and poor rate capability, severely limit its practical application in LIBs [14–16].

To improve lithium storage performance of TiO₂, many strategies have been proposed. Designing nanostructured TiO₂ materials, for example, nanotubes [17], nanoparticles [18], nanowires [19], etc, can shorten the diffusion path of Li⁺ ions and the change kinetics of Li⁺ ions insertion into TiO₂ more favorably, resulting in the increase in reversible capacity of TiO₂ [20, 21]. However, pure nanomaterials readily aggregate, which severely weakens the aforesaid nanometer effect. Hybridizing nano TiO₂ and various high-conductive materials, such as carbonaceous materials, can solve the self-aggregation issue of TiO₂ nanomaterials, meanwhile effectively increase electronic conductivity of TiO₂ [22–24]. Accordingly, this hybridization strategy has become very popular. For instance, Zhang et al composited TiO₂ nanowires on reduced graphene oxide host with long cycle stability up to 5000 cycles and high rate performance of 221 mAh g⁻¹ at a current density of 5 A g⁻¹ [25]. We synthesized ultrafine TiO₂ nanocrystalline on N-doped amorphous hollow carbon spheres with a discharge capacity of 242 mAh g⁻¹ at 1 C after 200 cycles [26]. It is found that decreasing the particle size of TiO₂ to a sufficient small range and uniformly dispersing them on high-conductive materials can remarkably increase reversible capacity and rate capability of TiO₂, but very low mass loading of TiO₂ is an underlying defect [27–29]. Constructing TiO₂ nanosized building units to hierarchical porous microsized assemblies is a novel and effective strategy. Not only does the mass loading of TiO₂ increase, but also the various positive effects induced by nanostructures can be completely released and utilized. Cai et al synthesized walnut-like porous core/shell TiO₂ with super lithium storage properties [30]. Our group developed a PVP-boiling process to assemble TiO₂ nanoparticles (P25) into a foam-like, 3D mesoporous structure with an average discharge capacity of 227 mAh g⁻¹ over 200 cycles at 1 C [31]. So far, the novel hierarchical porous microsized assemblies of TiO₂ are highly desirable, but there are still few materials of this type reported. This kind of assembling structure exhibits outstanding lithium storage performance, and deep investigation on them is very necessary [2–6, 13, 23].

Herein, we present a facile method to construct a unique pomegranate-like C@TiO₂ mesoporous honeycomb sphere (denoted as C@TiO₂ MHS) by chemical bath deposition and hydrothermal reaction followed by the coating of polypyrrole (PPy) and carbonization. The composite spheres of TiO₂ and hexadecylamine (HDA) first form by chemical bath deposition and then are further developed to mesoporous honeycomb spheres as pomegranate seeds by hydrothermal reaction. PPy-derived amorphous C shell as the pericarp of the pomegranate encapsulate TiO₂ mesoporous honeycomb sphere to improve conductivity and stability of TiO₂. When used as anode material for LIBs, as expected, the pomegranate-like C@TiO₂ MHSs exhibited outstanding cycling stability and good rate capability. At a current density of 1 C (1 C = 170 mA g⁻¹), a reversible capacity of 184 mAh g⁻¹ was still delivered after 500 cycles. Even at a high current rate of 20 C, a reversible capacity of 60 mAh g⁻¹ was still obtained. These sufficiently demonstrate the structural superiority and great application potential of pomegranate-like C@TiO₂ MHSs as an anode material for LIBs.

2.1. Material synthesis

2.2. Materials characterization

2.3. Electrochemical measurements

The synthesis process of pomegranate-like C@TiO₂ MHSs is schematically illustrated in figure 1. HDA is a crucial material for the formation of TiO₂ MHSs. On the one hand, HDA and hydrous titania (hydrolysis product of TIP) together form an organic-inorganic composite sphere through a hydrogen-bonding interaction. On the other hand, HDA acts as a pore-forming additive. It can be completely removed through the hydrothermal reaction. Meanwhile hydrous titania is also crystallized. Thus, TiO₂ MHSs are obtained. TiO₂ MHSs are further coated by a layer of PPy and then carbonized, forming a novel pomegranate-like structure. It retains high specific surface area of TiO₂ MHSs; meanwhile significantly improves electronic conductivity and structure stability of TiO₂ MHSs.

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Figure 2(a) shows the SEM image of the organic-inorganic composite spheres of HDA and hydrous titania. They are of uniform shape and size, smooth surface as well as obvious monodisperse features. The TEM image (figure 2(b)) clearly exhibits their fine structure, presenting a homogeneous solid sphere structure. It is difficult to distinguish HDA and hydrous titania inside these spheres. After the hydrothermal reaction, the solid spheres transform to MHSs with a very rough surface (figure 2(c)). In the surface of these small spheres exist the abundant pores. The hydrothermal crystallization and the removal of HDA do not induce the collapse of these MHSs. The TEM image (figure 2(d)) and high-magnification TEM image (figure 2(e)) distinctly reveal that these MHSs are in fact composed of numerous uniform nanocrystalline TiO₂ with the diameter of about 10 nm. Among these nanocrystalline TiO₂, abundant mesopores exist. Clearly, these mesopores are throughout the interior and surface of these small spheres. The HRTEM image (figure 2(f)) displays a distinct crystal lattice with an interplanar spacing of 0.352 nm, well coinciding with (101) crystal planes of the anatase TiO₂, confirming good crystallinity of TiO₂. The phase analysis on the products after the hydrothermal treatment was conducted by XRD (figure 2(g)). The XRD pattern shows well-defined diffraction peaks. All of them can be well assigned to the anatase phase TiO₂ (JCPDS NO 04-0477). The short and widened diffraction peaks suggest the small size of TiO₂ crystalline. The calculation based on the Scherrer equation indicates the crystal size of TiO₂ is around 10 nm, which is in accordance with the TEM result. No other diffraction peaks corresponding to HDA and no other materials were detected, demonstrating the hydrothermal reaction totally removed the pore-forming agent of HDA. The microstructure of TiO₂ MHSs was further examined by N₂ adsorption/desorption isotherms measurement (figure 2(h)). The isotherms can be classified as a typical type-IV curve with an obvious hysteresis loop in the high-relative pressure region (0.68–0.89), demonstrating the presence of the mesopores. The BET specific surface area reaches 153 m² g⁻¹. The high specific surface area can offer more active sites for electrochemical reaction, improving reaction kinetics of TiO₂. The pore size distribution curve calculated by the BJH method further confirms the existence of the mesopores. The main pore size ranges from 3 to 24 nm, centering at 13.6 nm. The total pore volume reaches 0.49 cm³ g⁻¹. The abundant mesopores greatly facilitate penetration of electrolyte, as well as buffer expansion and shrinkage of TiO₂ during the (de)lithiation process.

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After the coating of PPy, the independent TiO₂ MHSs are assembled into clusters of grapes due to the adhesion effect of PPy (figure 3(a)). The small spheres again become smooth without burr phenomenon, and the electronic conductivity is obviously improved. This indicates PPy has successfully coated TiO₂ MHSs. The TEM image (figure 3(b)) discloses that each TiO₂ MHSs is coated by a layer of PPy, presenting a unique pomegranate-like internal structure. The thickness of the PPy layer is around 50 nm. After carbonization, clusters of grapes are well preserved but the electronic conductivity is further increased because the display effect of SEM is remarkably improved (figure 3(c)). The TEM image (figure 3(d)) demonstrates TiO₂ nanoparticles do not grow up during the carbonization process, and TiO₂ MHSs are also well retained. Nevertheless, the thickness of the carbon shell seems to decrease to 30–40 nm due to the shrinkage caused by the carbonization. The compressive pre-stress induced by this shrinkage makes the carbon shell better resist the volume expansion of TiO₂ MHSs during the lithiation process, which is favorable for increasing the structure stability of TiO₂ MHSs. A high-magnification TEM image (figure 3(e)) reveals the local fine structure of C@TiO₂ MHSs. TiO₂ nanoparticles tightly contact the carbon shell without any interspace, which can effectively improve electronic conductivity of TiO₂ MHSs. The internal mesoporous honeycomb structure can be still distinguished clearly, implying the mesopores inside the TiO₂ MHSs should not be filled by PPy and carbon. It is worth noting that in the carbon shell abundant micropores exist, which can allow the infiltration of the electrolyte.

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The phase of the final product was examined by XRD (figure 4(a)). The XRD pattern only displays diffraction peaks of anatase TiO₂ without any impurity. However, at 10–35°, the base line of the XRD pattern obviously rises, presenting a small hump. This means the presence of carbon materials. To further determine the structures of TiO₂ and carbon in the final product, the Raman spectra was recorded (figure 4(b)). Five typical Raman peaks appear at 152, 198, 392, 501 and 628 cm⁻¹, which well matches the E_g, E_g(2), B_1g, A_1g and E_g(3) characteristic peaks of anatase TiO₂. At 1351 and 1577 cm⁻¹ appear two obvious humps that correspond to the D-band and G-band of carbonaceous material. This indicates the presence of the carbon. Compared with the Raman peaks of TiO₂, the intensity of the D-band and G-band are weaker, inflecting the lower content of carbon. The I_D/I_G ratio may be used to evaluate the graphitization degree of carbonaceous material. The high ratio value of 1.03 confirms the amorphous disordered structure of PPy-derived carbon. To obtain the content of the carbon, TGA was performed in air (figure 4(c)). C@TiO₂ MHSs exhibit two distinct weight loss. The first weight loss is relatively slight, and occurs from room temperature to 200 °C, which is ascribed to the evaporation of the adsorbed moisture. It accounts for 7.5% of the total weight. The second weight loss occurs from 300 °C to 700 °C, which is attributed to the combustion of the carbon in air. After water is excluded, the content of the carbon is 56.17% and the rest is TiO₂. Furthermore, EDS was also performed to further verify the constituent elements and their content (figure 4(d)). There exist four elements, C, N, Ti and O. This means the amorphous carbon should be doped by N [17]. The doping of N can greatly increase electronic conductivity of the amorphous carbon. The EDS result indicates the weight percentages of C, N and TiO₂ are 53.96%, 2.70% and 43.34%, respectively, which agrees with the TGA result.

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The electrochemical performances of C@TiO₂ MHSs as anode material are shown in figure 5. Figure 5(a) exhibits the CV curves of C@TiO₂ MHSs at a scan rate of 0.2 mV s⁻¹ in the potential window of 1.0–3.0 V. One pair of well-defined reduction/oxidation peak appears at 1.66 and 2.08 V, corresponding to Li⁺ intercalation and extraction in the TiO₂ lattice [32]. Except for the reduction/oxidation peak, the other region of CV curves is nearly rectangular, demonstrating a strong pseudocapacitive lithium storage behavior. Aside from the first cathodic scanning, the CV curves basically overlap each other, implying good reversibility and stability. The CV curves of C@TiO₂ MHSs clearly differ from those of TiO₂ MHSs (figure 5(b)). (1) The reduction/oxidation peaks of C@TiO₂ MHSs are not prominent but those of TiO₂ MHSs are very sharp and prominent, indicating lithium intercalation storage is not dominant for the former, but dominant for the latter. This could be due to the inhibition effect of the amorphous carbon shell that limits the rapidly entering of the electrolyte, so suppresses sufficient contact between the inner TiO₂ nanocrystalline and the electrolyte. (2) The rectangular shape of C@TiO₂ MHSs is more remarkable, and the corresponding area is larger. These suggest that the amorphous carbon shell can inspire the pseudocapacitive storage effect of TiO₂ MHSs [27, 33–35].

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Figure 5(c) shows the discharge-charge profiles of C@TiO₂ MHSs in the first three cycles at a current density of 1 C (1 C = 170 mA g⁻¹). Generally, the discharge profile of TiO₂ includes three typical regions [36, 37]. The first region is a fast potential decrease from open circuit voltage to about 1.75 V, corresponding to the homogeneous Li insertion into the bulk by a solid-solution Li storage mechanism of TiO₂. The potential plateau at 1.75 V is the second region which reflects Li⁺ insertion and occupation of interstitial octahedral sites of anatase TiO₂. The third region is a potential slope from 1.75 V to 1.0 V. For C@TiO₂ MHSs, their charge and discharge plateaus shrink to be a very short plateau and an inflection point, meaning the lithium intercalation reaction is suppressed. This point obviously differs from the long charge and discharge plateaus of pure TiO₂ MHSs, as shown in figure 5(d), which possibly results from the enclosed environment induced by the carbon shell that limits rapid sufficient contact between TiO₂ MHSs and electrolyte. However, the first and the third regions of the discharge profiles are especially remarkable, and the corresponding discharge capacity reaches 47/47 mAh g⁻¹ and 262/282 mAh g⁻¹ in the second and third discharge process, much higher than 11 and 95 mAh g⁻¹ of TiO₂ MHSs. This indicates the amorphous carbon shell changes the common lithium storage mechanism of TiO₂ MHSs and enhances the solid-solution and pseudocapacitive Li storage effects of TiO₂. These characteristics of the discharge-charge profiles agree well with those of the corresponding CV curves.

Figure 5(e) demonstrates the cycling properties of C@TiO₂ MHSs at a current density of 1 C. The first discharge capacity reaches 554 mAh g⁻¹ but the second one drops to 343 mAh g⁻¹. The initial irreversible capacity loss is mainly due to the intercalation of Li⁺ ions into irreversible sites of TiO₂ and irreversible decomposition of the electrolyte [1, 38]. Afterwards, the discharge capacity slowly declines, but is always much higher than the theoretical specific capacity of TiO₂. At the 500th cycle, the discharge capacity still reaches 184 mAh g⁻¹, exhibiting high electrochemical activity. The coulombic efficiencies in most cycles are higher than 99% except the initial 20 cycles, showing superior reversibility. By contrast, the discharge capacity of TiO₂ MHSs is lower, and the decline rate is faster. At the 500th cycle, TiO₂ MHSs only deliver a discharge capacity of 74 mAh g⁻¹. This comparison demonstrates the amorphous carbon shell has a remarkable effect on improving reversible capacity and cycling stability of TiO₂ MHSs. Figure 5(e) also provides the cycling performance of PPy-derived carbon at 1.0–3.0 V. The negligible reversible capacity (around 13 mAh g⁻¹) of PPy-derived carbon demonstrates the discharge capacity of C@TiO₂ MHSs mainly comes from TiO₂. In comparison with other reported high-performance TiO₂-based composites, the cycling performance of C@TiO₂ MHSs exhibits a prominent advantage (table 1), highlighting the important significance of the pomegranate-like structure. Figure 5(f) displays the rate performance of C@TiO₂ MHSs. At a current density of 0.5 C, C@TiO₂ MHSs can deliver a discharge capacity as high as 500 to 368 mAh g⁻¹, presenting high electrochemical activity. At 1 C, the discharge capacity stabilizes at 298 mAh g⁻¹. At 5 C, 10 C and 20 C, the discharge capacity decreases to 150, 98 and 60 mAh g⁻¹, presenting a certain high-rate capability. When the current returns to 1 C, the discharge capacity may reach about 273 mAh g⁻¹, recovering to 91.6% of the initial level at 1 C. The impressive capacity retention ratio indicates C@TiO₂ MHSs can withstand high-current charge-discharge reactions, meanwhile retaining the stability of the material structure.

To gain further insight into the electrochemical reaction kinetics of C@TiO₂ MHSs, EIS measurements were performed (figure 5(g)). All Nyquist plots consist of a depressed semicircle in the high-frequency region and a sloping line in the low-frequency region. They correspond to charge transfer resistance (R_ct) and Li⁺ diffusion in the electrode material. In comparison with pure TiO₂ MHSs, C@TiO₂ MHSs exhibit a smaller semicircle, meaning a smaller R_ct. The fitted R_ct value is 184 Ω versus 197 Ω of TiO₂ MHSs. This decrease in R_ct results from the improvement effect of the amorphous carbon shell on electronic conductivity of TiO₂ MHSs [23, 44]. After 1 cycle, the contact resistance of C@TiO₂ MHSs decreases from 16.88 Ω to 3.64 Ω due to the activation effect. The R_ct value slightly decreases to 172 Ω, indicating good structure stability. Furthermore, the three sloping lines are basically parallel to each other. This suggests the amorphous carbon shell and the charge-discharge reaction should not distinctly influence Li⁺ diffusion in the inner TiO₂, demonstrating good kinetic characteristics of C@TiO₂ MHSs.

To verify the structure stability of C@TiO₂ MHSs during charge-discharge cycling, we disassembled a coin-type cell after 500 cycles at 1 C, and then observed the morphology and microstructure of the electrode material through SEM and TEM. It was found in the SEM image (figure 6(a)) that C@TiO₂ MHSs still exhibit the initial small sphere morphology without any distinct structure decay, presenting good structural integrity. The TEM image (figure 6(b)) discloses the internal structure of these small spheres. The pomegranate-like inner structure is well-preserved. 500 cycles did not damage the mesoporous honeycomb sphere structure of TiO₂. C@TiO₂ MHSs present excellent structural robustness, which is very favorable for long-term cycling stability of TiO₂.

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The outstanding electrochemical performance of C@TiO₂ MHSs mainly benefits from the unique pomegranate-like structure. The mesoporous honeycomb sphere is of large specific surface area and small crystal size, which endows TiO₂ with high electrochemical activity and kinetics performance. The amorphous carbon shell significantly improves the electronic conductivity of TiO₂; well buffering the volume variation of TiO₂ and greatly increasing the structure stability of TiO₂. Although the lithium intercalation storage of TiO₂ is suppressed due to the coating of the amorphous carbon shell, solid-solution and pseudocapacitive Li storage effects are remarkably increased. These factors bring about higher reversible capacity and more stable cycling performance of C@TiO₂ MHSs.

In summary, C@TiO₂ MHSs are successfully synthesized and show the unique pomegranate-like structure that is composed of the internal TiO₂ mesoporous honeycomb sphere and the external amorphous carbon shell. The mesoporous honeycomb sphere structure provides TiO₂ with high electrochemical activity and kinetic performance. The amorphous carbon shell further enhances the electronic conductivity and structure stability of TiO₂, and is very crucial for electrochemical performance improvement of TiO₂ MHSs. These favorable structural features lead to excellent lithium storage performance of C@TiO₂ MHSs, including high reversible capacity, stable cycling performance and good rate capability. This work gains new insight into the structure design of high-performance TiO₂-carbon-based composite materials.

This work was funded by State Key Lab of Silicon Materials, Zhejiang University.