Strategic approach for NR/HMS conversion into MCS nanocomposites
The distribution of organic carbon precursors in mesostructured silica is a crucial factor to prepare MCS materials with good structural and textural properties via the carbonization process. The encapsulation of organic substance in mesoporous silica channels via wet impregnation1,2 and chemical functionalization of mesostructured silica surface with organosilanes or organic compounds4,13 resulted in the poor distribution of carbon phase in obtained nanocomposites. However, mesostructured organic–inorganic nanocomposites prepared using the evaporation-induced co-assembly method in the presence of phenolic resin-based precursors15 provided MCS nanocomposites with high carbon content and well-dispersed carbon moieties15 due to the high dispersion of the polymeric carbon precursor in the mesostructured silicate framework.
In this study, the incorporation and distribution of rubber phase in the mesostructured silicate framework of NR/HMS were confirmed by small-angle X-ray scattering (SAXS) and high-resolution transmission electron microscopy (HRTEM), respectively. Figure 1 compares the SAXS patterns of organic–inorganic mesophases in the mixtures during the synthesis of HMS and NR/HMS. For NR/HMS (Fig. 1a), the peak at q = 1.46 nm−1, which corresponds to the mesophase structure with a hexagonal unit cell of 0.68 nm, was observed when the rubber gel was mixed with tetraethyl orthosilicate (TEOS), dodecylamine (DDA), and water. It was larger than the mesophase formed during the synthesis of pure silica HMS (unit cell = 0.58 nm) with the peak at q = 1.71 nm−1 (Fig. 1b). The obtained result indicates that rubber molecules were incorporated into the organic–inorganic mesophase of NR/HMS. Figure 2 indicates the proposed mechanistic model for the formation of NR/HMS nanocomposites synthesized via the in situ sol–gel process. First, TEOS and DDA were homogeneously dissolved in the NR solution using tetrahydrofuran as the synthesis media. Upon the addition of H2O, TEOS was partially hydrolyzed to silicate species simultaneously with the rearrangement of DDA molecules into hexagonal rod-like micelles, induced by H-bonding between the amine groups of DDA and hydroxyl groups of silicate species. The ethoxy groups adjacent to rubber molecules remained nonhydrolyzed and acted as linkers between the rubber chains and silicate oligomers, which resulted in mesostructured entrapped NR/silica composite framework. The HRTEM image indicates the high dispersion of rubber lamellae in the NR/HMS nanocomposite obtained by this approach (Fig. 1c).
Our strategy for the conversion of the NR/HMS precursor into MCS composites via carbonization is shown in Fig. 2. In the first step, the as-synthesized NR/HMS nanocomposite was pretreated with an H2SO4/ethanol solution. H2SO4 adsorbed on the NR/HMS precursor was supposed to work as a catalyst to convert rubber phase into carbonaceous residue during the subsequent carbonization process. The thermal degradation of NR, a polymer of cis 1,4-isoprene, is a radical reaction, which is initiated by the random-chain scission of the β bond with respect to the double bonds of the polymeric backbone30. As shown in Fig. 3, two different allylic radicals are generated at temperatures below 380 °C, which results in lower molecular weight polyisoprene31. At higher temperatures (410–430 °C), both radicals undergo depropagation or unzipping toward isoprene monomers, simultaneously with intramolecular cyclization, followed by scission, to yield dipentene and other cycloalkene derivatives30,32,33. A further increase in the temperature to > 450 °C enhances fragmentation and aromatization of formed products, affording isoprene, terpenes, and aromatic compounds such as p-cymene34. In the presence of acid catalysts, the yield of aromatics was increased35, and the crude products from NR degradation could be repolymerized31. Therefore, the content and chemical nature of carbon moieties formed in the resulting MCS materials should be determined by carbonization temperature, concentration of H2SO4 solution used in the pretreatment step, and initial NR content of the NR/HMS precursor.
Effect of carbonization temperatures on MCS formation
The weight loss and differential thermal analysis (DTA) curves of HMS, NR/HMS, and MCS nanocomposites obtained at different carbonization temperatures are compared in Fig. 4a. The NR/HMS showed the major weight loss of ~ 17 wt% at 150–420 °C owing to the decomposition of NR incorporated into the mesostructured silicate framework. The small weight loss (~ 7 wt%) in the range of 420–650 °C was related to the rubber-derived carbon residue16 and dehydroxylation of silicate network36. Some rubber fractions still remained in the resulting nanocomposites obtained at 350 °C. This decomposition step disappeared at the carbonization temperature of 450 °C. As shown in Table 1, the carbon content of the MCS nanocomposites prepared at 700 °C and 800 °C was insignificantly different (3.4–3.5 wt%) and not affected by increased carbonization temperature. This result was similar to that of Kim et al.37, who observed an insignificant change in the carbon content of carbon materials prepared from cellulose and carbonized at 500–800 °C. In general, the majority of weight loss for polymers up to 500 °C was due to the release of volatile matter, after which the structure of carbon residue was rearranged from amorphous carbon toward a more crystalline phase at higher temperatures38.
The conversion of the NR phase into carbon residue during the carbonization process was also confirmed by FT-IR (Fig. 4b) and solid-state 13C NMR (Fig. 4c). The presence of rubber in the NR/HMS precursor was confirmed by FT-IR bands related to C–H stretching (2,800–3,000 cm−1) and C–H deformation vibrations (1,370 cm−1 and 1,430 cm−1)16, the 13C-NMR spectrum showed chemical shifts at 23, 27, 33, 127, and 136 ppm, which were attributed to –CH3, –CH2–, –CH2–, =CH–, and > C=, respectively39,40,41, of poly(cis-1,4-isoprene) in the NR structure. Compared to pure silica HMS, the presence of rubber in NR/HMS decreased the free silanol band (3,750 cm−1) because some remnant ethoxy groups remained in the silicate framework16,17. The carbonization at 350–450 °C restored silanol groups, while the characteristic bands of NR gradually decreased with an increase in temperature. Above 700 °C, the band at 1,580–1615 cm−1, which corresponded to C=C–C stretching42, revealed the presence of aromatic carbon residue in MCS nanocomposites. This observation agreed with a new 13C-NMR signal at 128 ppm, which was attributed to aromatic carbon species7, observed for MCS-0.5G-1.00M-700. These results indicate the decomposition and transformation of rubber to aromatic carbon during carbonization at 700 °C. Of note, the temperature, at which aromatic carbon species formed in this study, was lower than that used for the preparation of sucrose-derived carbon materials via carbonization at 800 °C43.
Structural and textural properties of MCS materials
The small-angle X-ray diffraction (XRD) patterns of pure silica HMS and representative nanocomposites exhibited an intense reflection at 2θ in the range of 1.5°–3.0° (Fig. 5a), which corresponded to the (100) plane of hexagonal mesostructure with a wormhole-like silicate framework16. Basically, silica-based materials consisted of Si(OSi)2(OH)2, Si(OSi)3(OH), and Si(OSi)4 species, which corresponded to Q2 (− 92 ppm), Q3 (− 101 ppm), and Q4 (− 110 ppm) resonances, respectively, as revealed in the 29Si magic angle spinning (MAS) NMR spectra (Fig. 5b). The ratio of Q2:Q3:Q4 species of HMS, NR/HMS, and MCS were 1:7.5:33, 1:11:21, and 1:6.4:36, respectively. The presence of rubber incorporated into the HMS structure enhanced the wall thickness (Wt) of NR/HMS (Table 1) but hampered the hexagonal ordering of mesoporous structure16,17,18. Moreover, the higher content of isolated silanol sites (Q3) in NR/HMS than in HMS indicated that rubber molecules present in the synthesis mixture lowered the degree of silicate condensation17.
An increase in the carbonization temperature up to 450 °C increased the (001) reflection and led to reduced Wt and contraction of hexagonal unit cell (a0) of MCS nanocomposites. The comparison of intensities of Qn groups in composites revealed that the obtained MCS had lower Q2 and Q3 but higher four siloxane-bonded site (Q4) than the NR/HMS precursor. The overall result indicated that the carbonization process enhanced the condensation of incompletely hydrolyzed silicate species to form siloxane bonds44. Although MCS-0.5G-0.05M-700 exhibited the loss of mesostructure ordering owing to the dehydroxylation of silicate networks, which was enhanced at > 500 °C45,46, its a0 and Wt values increased compared to those of HMS calcined at 700 °C (Table 1). This result suggests that the formed carbon layer acted as a rigid support in the wall and probably limited structural shrinkage. This result was similar to that reported by Liu et al.15, who observed an interpenetrating network of carbon–silica nanocomposites prepared from a soluble resol polymer as a carbon source via evaporation-induced triconstituent co-assembly. The carbon layer acted as a rigid support and decreased framework shrinkage. Unfortunately, the mesostructured framework of MCS-0.5G-0.05M-800 collapsed due to severe dehydroxylation.
The N2 adsorption–desorption isotherms revealed that both pure silica HMS and nanocomposites exhibited type IV sorption isotherms (Fig. 5c), according to the IUPAC classification, which is characteristic of framework-confined mesoporous materials. NR/HMS exhibited lower textural properties than HMS (Table 1), which suggested that part of rubber molecules might cover the pore mouths or occupy the porous channels of mesostructured silica. This conclusion was supported by the FE-SEM analysis (Fig. 5d), which revealed the presence of particle agglomerates (102–167 nm). Compared to the NR/HMS precursor, the MCS nanocomposites obtained at 350 °C and 450 °C showed enhanced textural properties. Carbonization decomposed rubber covering the precursor surface and enhanced the condensation of silicate framework. The resulting MCS materials possessed reduced particle size and lower aggregation such as MCS-0.5G-0.05M-700 with an average size of 20–32 nm (Fig. 5d). In addition, the representative Transmission electron microscopy (TEM) image of this MCS sample (Fig. 5e) showed conventional uniform wormhole-like mesopores, as observed for HMS materials16,17,18.
Chemical structure of carbon species in MCS materials
Raman spectroscopy and X-ray photoelectron spectroscopy (XPS) were applied to acquire information about the evolution of the chemical nature of carbon species on MCS nanocomposites. As shown in Fig. 6a, the NR/HMS precursor exhibited a broad band in the range of 1,090–1,780 cm−1, which was attributed to the polymeric backbone of rubber incorporated into the HMS structure47. It corresponded to the C1s XPS signals (Fig. 6b) of NR/HMS, which represented the C=C bond (282.9 eV) and C–C/C–H bonds (284.8 eV)17. The C–O/C–O–C bonds (285.8 eV) and C=O bond (287.2 eV) were the surface functional groups formed by the oxidative cleavage of NR chains during storage48. Residual rubber moieties still existed on MCS-0.5G-0.05M-350. A weak band at approximately 1,150 cm−1 was attributed to the O=S=O symmetric stretching vibration of the sulfonic acid group grafted on the carbon surface49,50. An increase in temperature matured the carbon residue structure, as deduced from increased band intensity.
The resulting MCS nanocomposites exhibited two well-developed bands at approximately 1,370 cm−1 (D band) and 1,590 (G band) cm−1, which revealed the disorder or defects in the organization of carbon atoms and the sp2 in-plane vibration of graphitic carbon atoms, respectively51. Moreover, the O=S=O band disappeared at the carbonization temperature of 700 °C, which supported the rearrangement of amorphous carbon residue to graphitic carbon structure6. In the high Raman shift region (Fig. S1 in SI), two bands, at 2,650 cm−1 (D′1 band) and 2,850 cm−1 (D′3 band), were observed and attributed to the overtone of D or D1 and the combination of D and G, respectively52. These bands were related to the second order scattering of imperfect graphite and disordered carbons such as graphene oxide-like carbon species53. The result agreed with the C1s XPS spectrum of representative MCS materials, which showed carbon groups (C=C, CHx, and C–C) at 283.6 eV, hydroxyl groups or ether linkages (C–O, C–O–C) at 285.3 eV, carbonyl groups (C=O) at 287.2 eV, and carboxyl or ester groups (COO) at 289.0 eV. In addition, the O1s XPS spectrum (Fig. S2 in SI) confirmed the presence of carboxyl or ester groups at 530.3 eV, and hydroxyl or ether groups at 531.7 eV54. These results suggest that the resulting MCS nanocomposites exhibited the high dispersion of graphene oxide-like carbonaceous moieties with different types of surface oxygen-containing groups (C–O, C=O, C–O–O, and O–C=O).
Tuning physicochemical properties of MCS materials
The concentration of H2SO4 during the pretreatment step and the initial NR content of the NR/HMS precursor were varied to tune the physicochemical properties of resulting MCS nanocomposites. As summarized in Table 1, the carbon residue in the composite structure (3.4–16.1 wt%) was systematically increased with increasing H2SO4 concentrations from 0.5 to 2.0 M. The XRD patterns of MCS in this series exhibited characteristic peaks at lower 2θ positions compared to those of pure silica HMS (Fig. 7a), which corresponded to enhanced a0 and Wt values (Table 1), concomitantly with an enlarged particle size (Fig. 7c), owing to an increase in the amount of carbon incorporated into the mesostructured silica framework. However, an increase in the initial NR content of the NR/HMS precursor from 0.5 to 1.5 g did not affect the carbon content and a0 value of resulting MCS nanocomposites. This result was ascribed to the limited incorporation of NR molecules into the mesostructure of HMS.
The MCS nanocomposites prepared under different conditions exhibited a typical N2 physisorption isotherm of mesostructured materials (Fig. 7b). The use of higher H2SO4 concentration decreased SBET, Dp, and Vt (Table 1), which correlated well with the trend in the carbon content of this MCS series because some carbon residue may occupy the mesopores of resulting nanocomposites. By combining this result with XRD analysis confirmed that MCS prepared with a high NR content had a higher carbon phase fraction occluding into the mesopores than MCS prepared from the NR/HMS precursor with high NR dispersion.
Valle-Vigón et al.6 have reported that MCS materials prepared using a Pluronic P123 triblock co-polymer template as a carbon source via 800 °C carbonization had the highest carbon content of 13 wt%, SBET of 460 m2 g−1, and Vt of 0.58 cm3 g−1. Furthermore, the sucrose-derived MCS7, which was prepared by the co-assembly method, containing 16 wt% carbon, exhibited SBET and Vt of 316 m2 g−1 and 0.82 cm3 g−1, respectively. This result indicates that MCS nanocomposites prepared from the NR/HMS precursor not only had a high amount of carbon but also superior textural properties compared to previously reported MCS materials with a similar carbon content.
The Raman spectra of MCS materials obtained using different H2SO4 concentrations and initial NR content of the NR/HMS precursor are compared in Fig. 6a. The D and G bands were deconvoluted into five components, which were assigned to polyenes (D4; 1,208 cm−1), graphene edges (D1; 1,352 cm−1), amorphous carbon (D3; 1,462 cm−1), graphitic carbon (G; 1572 cm−1), and graphene sheets (D2; 1,610 cm−1)7,52. The fitting parameters obtained from the Raman analysis are summarized in Table S1 (SI). An increase in the H2SO4 concentration from 0.05 M to 2.00 M increased the fraction of amorphous carbon (D3), whereas the formation of crystalline carbon (G) decreased from 21.9% (MCS-0.5G-0.05M-700) to 18.7% (MCS-0.5G-2.00M-700). The use of high H2SO4 concentration may result in the high degree of sulfonic acid groups being grafted onto carbon residue6, which prevented the rearrangement of amorphous carbon to graphitic carbon. Furthermore, the C1s XPS analysis (Fig. 6b) indicated an increased content of oxygen-containing groups on the carbon surface due to the oxidation of carbonaceous residue by H2SO4. A similar observation was reported in the sulfonation of biochar using a concentrated H2SO4 solution, which not only introduced sulfonic acid groups but also generated carboxyl and hydroxyl groups on the resulting acidic carbon55,56,57. The fraction of oxygen-containing groups on MCS nanocomposites prepared from the NR/HMS precursor (21.3–32.0%) was higher than that on MCS materials prepared using furfuryl alcohol as a carbon precursor (12.0%)58. An increase in the initial NR content of the NR/HMS precursor from 0.5 to 1.5 g resulted in a greater graphitized carbon phase (G) in resulting MCS nanocomposites (Table S1 in SI).
H2SO4 enhanced dehydration reactions and the substitution of sulfonic acid groups onto the resulting amorphous carbon at low temperatures6. With an increase in the carbonization temperature, the oxygen-containing functional groups decomposed to H2O, carbon monoxide, carbon dioxide, and sulfur dioxide6,59. Moreover, H2SO4 was essential in generating carbonaceous residues from NR molecules by promoting the formation of aromatic structures and cross-linking processes6. It is worth noting that MCS nanocomposites prepared from the NR/HMS precursor had the highest carbon yield of 67.4 wt%. Nishihara et al.3 have observed a 22.9–37.3 wt% carbon yield for MCS materials prepared by coating mesoporous silica SBA-15 with 2,3-dihydroxynaphthalene as a carbon source followed by carbonization at 800 °C. Raman analysis showed that the degree of graphitization of carbon phase [ID1/(IG + ID1 + ID2] contained in MCS materials was in the range of 0.53–0.57, which indicated a similarity in their carbon structure. The relatively large fraction of D1 (40.75–48.23%) suggested that carbon residue was present as highly dispersed nanosized graphene in mesostructured nanocomposites.
Hydrophobicity analysis by the H2O adsorption measurement
The behavior of H2O adsorption at a low relative pressure was used to evaluate the effects of carbonaceous residue on the hydrophobic properties of MCS nanocomposites. Figure 8 showcases the H2O adsorption isotherms of samples at P/P0 of 0–0.6. NR/HMS exhibited a lower amount of adsorbed H2O than HMS owing to not only the depletion of exposed surface silanol groups, as revealed by the FT-IR analysis, but also the hydrophobic environment created by the rubber phase. Compared with pure silica HMS, MCS materials showed a lower adsorbed volume owing to the depletion of silanol groups via dehydroxylation during the carbonization and hydrophobicity of carbon moieties dispersed in the resulting nanocomposites. Of note, MCS-0.5G-1.0M-700 and MCS-1.5G-1.0M-700 were more hydrophobic than their NR/HMS precursor. An increase in the H2SO4 concentration to 2.0 M lowered the hydrophobicity of MCS materials because this high acidic condition enhanced the content of oxygen-functional groups on the nanocomposite surface, as evidenced by XPS results. The analysis of adsorption data at a low relative pressure (Table 1) indicated a decreased monolayer adsorption volume (Vm) in the following order: HMS (82.9 cm3 g−1) > MCS-0.5G-2.00M-700 (45.2 cm3 g−1) > MCS-0.5G-0.05M-700 (41.6 cm3 g−1) > NR/HMS (40.1 cm3 g−1) > MCS-0.5G-1.0M-700 (22.0 cm3 g−1) ≈ MCS-1.5G-1.0M-700 (22.0 cm3 g−1).
Preliminary investigation of HMS and nanocomposites as drug carriers
HMS-D, NR/HMS-D, and MCS-0.5G-1.00M-700-D were characterized by various techniques to assure successful diclofenac loading on these carriers. The presence of diclofenac decreased the characteristic hexagonal mesostructure (Fig. 9a). The FT-IR spectra of the carriers after drug loading revealed new bands at 1,280–1,350 cm−1 and 1,450–1,600 cm−1, which corresponded to C–N and aromatic stretching, respectively (Fig. 9b), of the diclofenac molecule22. Moreover, the N2 physisorption measurement indicated a decrease in SBET and Vt after diclofenac loading (Table 2), which was more pronounced for HMS. Presumably, a larger portion of diclofenac molecules blocked the pore mouth of pure silica carrier, while the hydrophobic properties of NR/HMS and MCS nanocomposites promoted the diffusion and dispersion of diclofenac in mesostructured pores.
Diclofenac dissolved very fast in a simulated intestinal environment (pH = 6.8), its cumulative release profile reached an equilibrium of approximately 80% in the first 60 min (Fig. 10a). Using HMS and nanocomposite carriers strongly controlled the sustained release of diclofenac in the intestinal environment (Fig. 10b), which agreed with the previous report by Brovo et al.60. The rate of diclofenac released from the three carriers was ranked in the following descending order: HMS-D > NR/HMS-D > MCS-0.5G-100M-700-D, which matched the hydrophobicity trend of these materials (Fig. 8). Then, the cumulative release profiles were fitted with different kinetic models, and the calculated parameters and corresponding correlation coefficients (R2) are summarized in Table 3. Diclofenac released from each carrier was best represented by the Higuchi model, which indicated that the release of diclofenac from carriers as a square root of time-dependent process and diffusion control61. From the Korsmeyer–Peppas kinetic model, the diffusional exponent (n) value was used to characterize different release mechanisms for cylindrical-shaped matrices. HMS-D had n higher than 0.89, which indicated that diclofenac release followed the super case II transport61, while the smaller n values of NR/HMS-D and MCS-0.5G-100M-700-D suggested anomalous diffusion or nonFickian diffusion61. These results suggest that the hydrophobicity of NR/HMS and MCS-0.5G-100M-700 slowed down solvent diffusion into their porous structure and the dissolution of loaded diclofenac62.