Highly dispersed Pt nanoclusters supported on zeolite-templated carbon for the oxygen reduction reaction

The formation of highly dispersed Pt nanoclusters supported on zeolite-templated carbon (PtNC/ZTC) by a facile electrochemical method as an electrocatalyst for the oxygen reduction reaction (ORR) is reported. The uniform micropores of ZTC serve as nanocages to stabilize the PtNCs with a sharp size distribution of 0.8–1.5 nm. The resultant PtNC/ZTC exhibits excellent catalytic activity for the ORR due to the small size of the Pt clusters and high accessibility of the active sites through the abundant micropores in ZTC.

Electrochemically synthesized highly dispersed Pt nanoclusters (PtNCs) stabilized by the nanocages of zeolite-templated carbon (ZTC) exhibit excellent electrocatalytic performance toward the oxygen reduction reaction.

Platinum (Pt) is currently considered one of the best electrocatalysts for the oxygen reduction reaction (ORR), which occurs at the cathode of a fuel cell and is the key process determining the overall performance.^1–5 However, the high cost and scarcity of Pt limit its wide commercialization in this field. According to the US Department of Energy, the total Pt loading is required to be below 0.125 mg cm⁻², in contrast to a presently used Pt loading of 0.4 mg cm⁻² or more for fuel cell application.⁴ Therefore, reducing the Pt loading without loss or with an improvement of the cathode performance has received significant interest in electrocatalytic research for fuel cell systems.^6–10 In this regard, reducing the size of Pt particles to a nanocluster scale (size < 2 nm) and maximizing the Pt dispersion may offer an efficient way to achieve maximum utilization of the Pt electrocatalyst with appropriate consumption.^4,11–15The size of nanomaterials generally plays a critical role in controlling the physical and chemical properties for catalytic applications.^16–20 With a decrease in the particle size to the nanoscale, quantum size effects are induced, which alter the surface energy of the material due to unsaturated coordination and change in the energy level of the d orbital of metal atoms, leading to spatial localization of the electrons.^17–20 This size-induced effect on the electronic structures at the active sites modifies the capability of binding the reactant molecules in catalytic reactions, thereby altering the activity of the nanocatalyst.²⁰ When the particle contains a few to several dozens of atoms with sizes, ranging from sub-nanometer to 2 nm often termed as nanocluster that bridges nanoparticle and a single atom.²¹ However, the Pt single atom is not an appropriate electrocatalyst for the ORR in a fuel cell system as the fast four-electron (4e⁻) pathway for the reduction of O₂ to H₂O requires at least two neighboring Pt atoms.^22,23 Anderson''s group demonstrated that the ratio between the production of H₂O (product of 4e⁻ process) and H₂O₂ (2e⁻) in the ORR strongly depends on the number of atoms in the Pt cluster. Typically, it requires more than 14 atoms in a Pt cluster to produce H₂O efficiently through the 4e⁻ pathway of the ORR.²⁴ Therefore, Pt nanoclusters having more than a dozen atoms have proven to be highly efficient ORR electrocatalysts for fuel cell systems.^13–15 Upon decreasing the size of the nanoparticles to a nanocluster, the electronic state and structure are known to be changed, leading to an increase of the catalytic activity in the ORR. Therefore, it is highly desirable to synthesize a Pt nanocluster-based material as an ORR electrocatalyst with high catalytic performance. To date, several synthesis strategies, such as wet-chemical, atomic-layer deposition, and photochemical methods, have been applied for the preparation of well-dispersed Pt nanoclusters on different types of support, such as dendrimer, metal oxide, and carbon materials.^{13–15,25–31}An alternative approach to synthesize Pt nanocluster (PtNC) is the encapsulation of the cluster within nanosized pores, for example, by utilizing microporous (diameters less than 2 nm) carbon materials.³² Among the microporous carbons, zeolite-templated carbon (ZTC) has been attractive for supporting Pt clusters due to its ordered microporous structure.^33–37 ZTC is a potentially promising material as catalyst support as it offers the advantages of extremely large surface area and high electrical conductivity of graphene-like carbon frameworks constituting a three-dimensional (3D) interconnected pore structure.³⁶ Moreover, the micropores of ZTC can serve as nanocages for stabilization of the Pt nanoclusters. Coker et al. used Pt²⁺ ion-exchanged zeolite as a carbon template to synthesize Pt nanoparticles in ZTC with size in a range of 1.3 to 2.0 nm.³³ Recently, atomically dispersed Pt ionic species was synthesized via a simple wet-impregnation method on ZTC containing a large amount of sulfur (17 wt%).²³ Itoi et al. synthesized PtNC consisting of 4–5 atoms and a single Pt atom in ZTC using the organoplatinum complex.³⁷ Although these methods produced Pt nanoclusters with narrow size distribution and atomic dispersion, they required multi-step processes and/or high-temperature treatment (>300 °C). High-temperature treatment often induces the sintering of nanoclusters to aggregated clusters. Therefore, it is highly desirable to develop a simple and low-cost method for the preparation of PtNC supported on ZTC (PtNC/ZTC) for use as an efficient ORR electrocatalyst. The electrochemical reduction approach offers an alternate and efficient route for the synthesis of PtNC in the micropores of ZTC. The electrochemical method is one of the popular ways to prepare electrocatalysts because it is a simple single-step procedure and ensures electrical contact between the nanoparticles and the support.^38,39Herein, we report a facile electrochemical method for the formation of PtNC with a narrow size range of 0.8–1.5 nm supported on ZTC. The resultant PtNC/ZTC shows higher electrocatalytic activities towards ORR compared to that of commercial Pt/C. Here, ZTC plays two important roles: (i) it provides nanocages to stabilize the PtNC and (ii) it accelerates the ORR activity by enhancing the accessibility of active sites through its abundant micropores. Fig. 1a shows a schematic representation of the typical electrochemical synthesis of PtNC/ZTC. In the first step, ZTC was impregnated with a Pt-precursor dissolved in a water–ethanol mixture. As ZTC possess ordered micropores (Fig. S1a^†) with high Brunauer–Emmett–Teller (BET) surface area of 3400 m² g⁻¹ (vide infra), the uniform adsorption and anchorage of PtCl₆²⁻ ions into the micropores of ZTC was favored. After impregnating and drying, the resultant ZTC–PtCl₆²⁻ was mixed with water–ethanol and Nafion to make the ink for the preparation of the electrode. Using the prepared electrode, a potential of 0.77 V vs. reversible hydrogen electrode (RHE) (Fig. 1b) was applied followed by potential cycling between 1.12 to −0.02 V vs. RHE until the cyclic voltammogram was stabilized. The Pt content of PtNC/ZTC was determined to be ∼10 wt% (Fig. S2^†) by thermogravimetric analysis (TGA). The obtained PtNC/ZTC was electrochemically characterized by cyclic voltammetry and electrochemical impedance spectroscopy. The cyclic voltammogram (Fig. 1c) after potential cycling in fresh KOH electrolyte shows the characteristic Pt peaks corresponding to hydrogen adsorption and desorption. The Nyquist plots (Fig. 1d) demonstrate that PtNC/ZTC has lower electrolyte resistance (42 Ω) than that of ZTC (70 Ω), implying an improvement in the conductivity of ZTC by the presence of PtNC. Due to the increase in the conductivity, PtNC/ZTC could facilitate the electron transfer more effectively than ZTC, enhancing its electrocatalytic activity.Open in a separate windowFig. 1(a) Illustration for the formation of PtNC/ZTC:Pt-precursor was impregnated into ZTC micropores, and then a potential (0.77 V vs. RHE) was exerted on the ZTC–PtCl₆²⁻ composite in a 0.1 M KOH solution to form PtNC/ZTC (b) Chronoamperometric response of ZTC–PtCl₆²⁻ at a constant potential of 0.77 V (vs. RHE) in 0.1 M KOH electrolyte. (c) Cyclic voltammogram of PtNC/ZTC in a fresh 0.1 M KOH at a scan rate of 20 mV s⁻¹. (d) Nyquist plots of ZTC and PtNC/ZTC in 0.1 M KOH. Fig. 2a and b show images from aberration-corrected scanning transmission electron microscope (STEM) with high-angle annular dark-field (HAADF). The HAADF-STEM images exhibit the typical morphology of the final product (PtNC/ZTC) after electrochemical reduction. As shown in Fig. 2a, it is very clear that isolated PtNCs are uniformly dispersed in ZTC. These PtNCs have a homogeneous distribution with a narrow size range (0.8–1.5 nm, Fig. 2b). On further magnification, the STEM image shows a cluster-like structure of Pt (Fig. 2c). The STEM image of selected PtNC (Fig. 2d) reveals that it consists of ∼20 atoms. The number of atom content in PtNC was further determined by matrix-assisted laser-desorption-ionization time-of-flight (MALDI-TOF) mass spectrometry using trans-2-[3-(4-test-butylphenyl)-2-methyl-2-propenylidene]malononitrile as the matrix.^40,41 As shown in Fig. S3,^† MALDI-TOF measurement produces a mass spectra with a predominant peak centered at ∼3700 Da corresponding to the Pt₁₉ cluster. The TEM image (Fig. S4 a and b^†) validates the formation of PtNC with an average size of 0.9 nm. In addition, the energy dispersive X-ray spectrometer (EDS) mapping images clearly shows the uniform dispersion of Pt nanocluster in ZTC (Fig. S4c^†). The X-ray powder diffraction (XRD) pattern (Fig. 2e) of PtNC/ZTC showed three broad peaks associated with small size metallic Pt corresponds to (111), (200), and (311) planes (Fig. 2e, inset), along with peaks of ZTC at 2θ = 7.8° and 14.9° corresponding to the ordered microporous structure. Along with the structural analysis, the porous texture of PtNC/ZTC was examined by Ar adsorption (Fig. 2f). PtNC/ZTC had a high BET surface area of 2360 m² g_ZTC⁻¹, which is 1.4 times lower than that of pristine ZTC (3400 m² g_ZTC⁻¹). The decrease in Ar adsorption capacity after the formation of PtNC in ZTC is interpreted as a result of the filling of ZTC micropores by PtNC. This micropore filling was confirmed in the pore size distributions of the pristine ZTC and the metal-loaded carbon (inset of Fig. 2f). The X-ray photoelectron spectroscopy (XPS) results reveal the signature of Pt in ZTC (Fig. S5^†). The elemental survey (Fig. S5a^†) shows the signature of C 1s, O 1s, F 1s (Nafion), and Pt 4f. The chemical nature of Pt in PtNC/ZTC was inspected by a detailed Pt 4f XPS analysis. The deconvoluted Pt 4f XPS spectra (Fig. S5b^†) reveals the presence of both metallic and ionic Pt species. The peaks observed at 71.0 (4f_7/2) and 74.2 (4f_5/2) eV correspond to metallic Pt whereas the other peaks positioned at 72.6 (4f_7/2) and 76.0 (4f_5/2) are attributed to Pt²⁺ and the peaks at 74.9 (4f_7/2) and 77.8 (4f_5/2) eV are attributed to Pt⁴⁺ originating from the surface oxidation of metallic Pt.⁴²Open in a separate windowFig. 2(a–d) Representative spherical aberration-corrected HAADF-STEM images of PtNC/ZTC at various magnifications. (e) XRD pattern of PtNC/ZTC and (f) Ar adsorption–desorption isotherms of ZTC and PtNC/ZTC. Inset in (e) shows a 30 times magnified high-angle region of XRD of PtNC/ZTC. Inset in (f) shows the pore size distributions of the ZTC and PtNC/ZTC.The formation of narrow sized PtNC by the electrochemical method can be ascribed to the stabilization of PtNC in the ZTC micropores, which serve as cages to impose a spatial limitation on the size of the Pt clusters. For comparison, Pt supported on ZTC was also prepared by the conventional incipient wetness impregnation and subsequent H₂-reduction at high temperature (300 °C). The Pt obtained by this incipient wetness impregnation method shows the formation of Pt nanoparticles on the exterior surface of ZTC (PtNP/ZTC) (Fig. S6^†). The formation of larger Pt nanoparticles is due to the sintering at high temperature, showing that even ZTC micropores could not prevent the aggregation of PtNCs at high temperatures. Fig. 3 shows the electrochemical ORR activity of PtNC/ZTC using linear sweep voltammetry (LSV) technique on a rotating disc electrode (RDE) in a 0.1 M KOH solution saturated with O₂ at a scan rate of 5 mV s⁻¹. The ORR activity of ZTC (without PtNC) was measured for comparison as well. As shown in Fig. 3a, PtNC/ZTC exhibited higher diffusion limiting current density and higher positive onset and half-wave potential compared to ZTC alone, indicating that PtNC is the active center for the ORR. To investigate the effect of the Pt loading amount on the ORR activity, PtNC/ZTC with various Pt loadings, 2–20 wt%, was used for the measurement of LSV at 1600 rpm. With an increase in Pt content, both the onset and half-wave potential shifted towards more positive potential up to 10 wt% loading of Pt (Fig. 3a and S7^†). Upon further increase of loading of Pt on ZTC to 20 wt%, both the onset and half-wave potential of PtNC/ZTC shifted towards less positive potential along with a slight decrease in the diffusion limiting current density (Fig. 3a). The decrease in the ORR activity of PtNC/ZTC at high loading of Pt (20 wt%) was attributed to the decrease in the electrochemically active surface area (Fig. S8^†) and decrease in the specific surface area (Fig. S9^†). The STEM image clearly shows that the aggregated Pt clusters were formed on the exterior surface of ZTC at 20 wt% loading of Pt (Fig. S10c^†), blocking the accessibility of active sites. Therefore, PtNC/ZTC with the optimum loading of 10 wt% of Pt leads to superior ORR activity with a high positive onset potential of 0.99 V, which is similar to commercial Tanaka Pt/C (Pt/C-TKK) (Fig. 3b), and a half-wave potential of 0.87 V, which is ∼10 mV more positive than that of commercial Pt/C-TKK (0.86 V) (Fig. 3b). Compared to the case of PtNC/ZTC, both the onset and half-wave potential of PtNP/ZTC prepared by the conventional incipient wetness impregnation and subsequent H₂-reduction with the same loading of Pt exhibited a less positive value (Fig. S11^†). The poorer activity of PtNP/ZTC is due to the blockage of active sites by larger PtNPs formed on the exterior surface of ZTC (Fig. S6^†).Open in a separate windowFig. 3(a) RDE ORR polarization curves of PtNC/ZTC with different mass loading of Pt. (b) Comparison of PtNC/ZTC (PtNC_10%/ZTC) with commercial Pt/C-TKK at the same loading of 40 μg_Pt cm⁻². (c) RDE ORR polarization curves of PtNC/ZTC at different rotation speeds. Inset in (c) shows the corresponding K–L plots at different potentials. (d) Represents the kinetic current density values of Pt/C-TKK and PtNC/ZTC at the potential of 0.8 V vs. RHE.To investigate the kinetics of the ORR activity of PtNC/ZTC, LSV measurements were performed with RDE at different rotating rates (Fig. 3c), and the kinetics was analyzed using a Koutecký–Levich (K–L) plot (Fig. 3c, inset). From Fig. 3c, it was observed that the current density increases with the increasing speed of rotation of the electrode, which is characteristic of a diffusion-controlled reaction. The corresponding linear K–L plots (Fig. 3c, inset) with a similar slope at different potentials reveal that the number of transferred electrons was ∼4, indicating that O₂ is directly reduced to OH⁻ and the ORR is dominated by the H₂O₂-free 4e⁻ pathway. To estimate the amount of produced peroxide ion, rotating ring-disc electrode (RRDE) measurement was performed and the produce peroxide ion calculated from RRDE curve was < 4% (Fig. S12^†). The kinetic current density (J_k) obtained from K–L plot at the potential of 0.8 V (Fig. 3d) for PtNC/ZTC (J_k = 50 mA cm⁻²) is 2.2 times higher than that of commercial Pt/C-TKK (J_k = 22 mA cm⁻²).As Pt-based electrocatalysts are known to be highly active in an acidic medium, the ORR activity of PtNC/ZTC in O₂-saturated 0.1 M HClO₄ was also evaluated by comparing it with that of commercial Pt/C-TKK with the same loading of Pt on the electrode surface using RDE at a scan rate of 5 mV s⁻¹. The PtNC/ZTC-based electrode exhibited ORR activity with an onset potential of 0.96 V (Fig. 4), which is close to that of Pt/C-TKK (0.98 V), and half-wave potentials of 0.84 V, which is 20 mV more positive than that of Pt/C-TKK (0.82 V). PtNC/ZTC showed a slightly higher diffusion-limiting current density of ∼5.9 mA cm⁻² (0.4–0.7 V) compared with that of the Pt/C-TKK catalyst (∼5.6 mA cm⁻²). The kinetics of the ORR in an acidic medium was further analyzed using RDE at different rotation rates (Fig. S13^†) and it was observed that the current density increases with the increasing speed of rotation of the electrode, as in the case of the alkaline medium. The number of electron involved and the amount of produced H₂O₂ estimated by RRDE measurement were ∼4 and < 5%, respectively (Fig. S14^†). The mass activity of PtNC/ZTC obtained using the mass transport corrected kinetic current at 0.8 V is 0.15 A mg⁻¹, which is 3.2 times higher than that of Pt/C-TKK (0.046 A mg⁻¹).Open in a separate windowFig. 4(a) RDE ORR polarization curves at 1600 rpm and (b) mass activity at 0.8 V of PtNC/ZTC and Pt/C-TKK in 0.1 M HClO₄.Furthermore, the methanol tolerance of PtNC/ZTC was assessed by intentionally adding methanol to the oxygen saturated electrolyte solution (both in alkaline and acidic media). The commercial Pt/C-TKK was used for comparison as well. The peak current densities for methanol oxidation with PtNC/ZTC were ∼2.8 and ∼3 times lower than that of Pt/C-TKK in alkaline (Fig. 5a) and acidic (Fig. 5b) media, respectively. These results indicate that PtNC/ZTC has much higher tolerance towards methanol than Pt/C-TKK does. This higher methanol tolerance of PtNC/ZTC can be attributed to the small size of the Pt cluster, which may not be sufficient to catalyze the oxidation of methanol efficiently, as the oxidation of methanol requires Pt ensemble sites.⁴³Open in a separate windowFig. 5ORR polarization curves of PtNC/ZTC and Pt/C-TKK in the absence (solid line) and presence (dotted line) of 0.1 M of CH₃OH at a rotation rate of 1600 rpm in (a) alkaline and (b) acid media.The durability of PtNC/ZTC was also investigated by the amperometric technique. The test was performed at a constant voltage of the half-wave potential in an O₂-saturated alkaline medium and at 0.7 V in an O₂-saturated acidic medium at a rotation rate of 1600 rpm (Fig. S15a and b^†). The durability of the PtNC/ZTC catalyst in the alkaline medium was higher than that of Pt/C-TKK, exhibiting a 30% decrease compared to a 40% decrease of Pt/C-TKK in 5.5 h of ORR operation (Fig. S15a^†). The higher durability of PtNC/ZTC compared to Pt/C-TKK in the alkaline medium may be due to the stabilization of PtNC by pore entrapment. In the acidic medium, however, PtNC/ZTC exhibited a 54% decrease in the initial current after 5.5 h of operation while a 33% decrease was observed in the case of Pt/C-TKK (Fig. S15b^†). The decrease in ORR activity in the acidic medium may be due to the leaching out of tiny Pt nanoclusters in acid electrolyte from the ZTC micropores. To understand the decrease in the ORR durability with time, STEM measurements of PtNC/ZTC after 5.5 h of ORR operation were performed. In the alkaline medium, the STEM image of post-ORR PtNC/ZTC shows a slight change in the size of PtNC (Fig. S15c^†) while the STEM image of PtNC/ZTC after ORR in the acidic medium exhibited sintering of PtNC into large particles with an average size of 30 nm (Fig. S15d^†), resulting in a decrease of the ORR activity. In the alkaline medium, the decrease in ORR activity with time may be due to the oxidation of the ZTC support in KOH.⁴⁴We attributed the excellent ORR activity of PtNC/ZTC to the interplay between the following: (1) the structure of the Pt cluster possessing a high ratio of surface atoms that benefits the surface reactions,^45–47 (2) the microporous 3D graphene-like structure of the ZTC support that enables easy access of O₂ and electrolyte molecules to the active sites,⁴⁸ and (3) the high conductivity and large accessible surface area of ZTC that facilitates the electron transfer.^49–51     

Microwave-assisted synthesis of mutually embedded Rh concave nanocubes with enhanced electrocatalytic activity

Novel mutually embedded Rh concave nanocubes were synthesized by reducing Rh(acac)₃ in tetraethylene glycol in the presence of benzyldimethylhexadecylammonium, KI and polyvinylpyrrolidone under microwave irradiation for 120 s. KI and HDBAC were crucial to the formation of mutually embedded nanostructures. The as-prepared Rh nanocrystals exhibited higher electrocatalytic activity and stability.

Mutually embedded Rh concave nanocubes were synthesized by reducing Rh(acac)₃ with tetraethylene glycol (TEG) as both a solvent and a reducing agent under microwave irradiation for 120 s.

As important catalytic materials, the controlled syntheses of platinum group metal nanocatalysts have attracted wide attention for many years. The catalytic performances of platinum group metal nanomaterials are highly dependent upon their morphologies, compositions and surface structures. Their nanocrystals with controlled shapes have been extensively explored in order to promote their catalytic performance and reduce their cost because of their scarcity and high prices. As an important platinum group metal, rhodium (Rh) is often used as a typical catalyst with high activity and selectivity in hydroreduction,¹ hydroformylation,^1f,2 NO_x reduction,³ CO oxidation,⁴ cross coupling,⁵ and fuel cell^1g,4d,6 and other chemical reactions.⁷ In addition, Rh has a strong resistance to acids and bases as well as a high melting point. However, nanoscaled Rh exhibits high thermodynamic instability owing to its high surface free energy, although it is more stable than many other catalytically active metals. So, the shape-controlled synthesis of Rh nanocrystals is still one of the challenges in this field, though Pt and Pd nanocrystals with many different morphologies have been obtained. In the past decade, great efforts have been devoted to tailoring the sizes, morphologies and surface structures of Rh nanoparticles to improve their catalytic efficiency because of the scarcity and preciousness. Up to now, many Rh nanomaterials with various morphologies such as nanosheets,^1f,4a,8 nanotetropods,⁹ hierarchical dendrites,⁶ hyperbranched nanoplates,^1g,10 ultrathin nanosheet assemblies,^1e,7a and cubic,¹¹ tetrahedral,^1d icosahedral,¹² tetrahexahedral,^4d concave cubic,¹³ concave tetrahedral nanocrystals,¹⁴ as well as cubic nanoframes,^5b,5c truncated octahedral nanoframes,¹⁵ multipods¹⁶ and mesoporous³ Rh nanoparticles have been successfully synthesized. Moreover, all the obtained Rh nanoparticles were monodispersed and displayed enhanced catalytic activities. Although various nanoparticles with concave, frame, branched, or hierarchical structures as well as normal flat or convex surfaces have been created, it is still worthy to develop further novel Rh nanostructures.Herein, we report a facile microwave-assisted strategy for a one-pot synthesis of mutually embedded Rh concave nanocubes, a unique hierarchical nanostructure with several identical concave nanocubes embedded in each other. The co-adsorption of I⁻ ions and benzyldimethylhexadecylammonium (HDBAC) was dominantly responsible for the generation of the hierarchical concave cubic Rh nanocrystals. The as-prepared Rh nanocrystals displayed an enhanced electrochemical activity for formic acid electro-oxidation. Fig. 1a and b as well as Fig. S1 (ESI^†) display TEM images of the typical Rh nanocrystals synthesized under microwave irradiation for 120 s in the presence of PVP and an appropriate amount of KI and HDBAC. Interestingly, as can be seen, the resulting Rh nanocrystals were present in mutually embedded concave nanocubic morphologies under TEM, showing a unique hierarchical nanostructure feature. This hierarchical structure was consisted of at least two concave cubes (Fig. 1b) which were embedded in each other. If considering an individual concave nanocube, the average diameter was about 85 nm. The high-resolution TEM image, as shown in Fig. 1c and d, showed the lattice fringes with an interplanar spacing of 0.192 and 0.135 nm, which can be indexed to the (200) and (220) planes of face-centered cubic (fcc) Rh. The corresponding fast Fourier transform (FFT) pattern for the selected box area in Fig. 1c is shown in the inset in Fig. 1d, indicating a single crystal structure and good crystallinity. In fact, the as-prepared concave cube demonstrated octapod characteristics because of its showing both concave faces and concave edges. The SEM image further confirmed hierarchical structure feature. As shown in Fig. 1e and the inset for partially enlarged picture as well as Fig. S2 (ESI^†), the concave nanocubic structures and their mutually embedded feature are present clearly.Open in a separate windowFig. 1TEM (a and b), HRTEM (c and d) and SEM (e) images of the as-prepared mutually embedded Rh concave nanocubes. (d) shows the selected box area in (c). The insets in (d and e) show the FFT pattern and a partially enlarged SEM picture, respectively.The typical XRD pattern of the as-synthesized mutually embedded Rh concave nanocubes is shown in Fig. 2. The characteristic peaks at 41.28, 48.05, 70.18 and 84.45° are corresponding well with the (111), (200), (220) and (311) lattice planes according to the standard diffraction file (JCPDS 05-0685), respectively. The sharp and strong (111) diffraction peak, which indicated the preferential orientation of (111) planes and the consistency with the HRTEM observation, suggested its high purity and crystallinity of the obtained Rh nanocrystals. In addition, XPS measurement demonstrated the binding energy of Rh 3d_5/2 and Rh 3d_3/2 at 307.16 and 311.91 eV (Fig. 3), respectively, with an interval of 4.75 eV, which was coincident with the reference values (307.0 and 311.75 eV),¹⁷ indicating Rh⁰ with zero oxidation for the as-prepared mutually embedded concave nanocubes.Open in a separate windowFig. 2XRD pattern of the typical mutually embedded Rh concave nanocubes.Open in a separate windowFig. 3XPS spectrogram of the typical mutually embedded Rh concave nanocubes.It was worth noting that the use of KI was much essential for creating the mutually embedded Rh concave nanocubes. As shown in Fig. 4a, except flower ball structures connected with each other, neither concave cube nor hierarchical structure was observed in the absence of KI. When 0.4 mmol of KI was added, the embedded Rh concave nanocubes accompanying with some irregular nanostructures were generated (Fig. 4b). With further increasing the amount of KI from 0.8 to 1.2 mmol, complex inter-embedded nanostructures with obscure polyhedral outlines were formed (Fig. 4c and d). Accordingly, an excessive amount of KI was unfavorable for the generation of the mutually embedded Rh concave nanocubes. According to the previous report,¹⁸ the addition of KI would manipulate the reducing kinetics to generate Rh concave nanostructures under microwave irradiation. In the presence of KI, the precursor was transformed to a more stably coordinated anion [RhI₆]³⁻. As a result, the reducing rate of Rh(iii) to Rh(0) as well as both the nucleation and growth rate of Rh nanoparticles decreased, which would be favorable for the oriented growth of Rh concave cubes. The role of I⁻ ions was elucidated by using an equivalent amount of KBr or KCl in stead of KI, respectively, under the same other conditions. As can be seen (Fig. S3, ESI^†), no single or embedded concave nanocubes but amorphous Rh nanoparticles with agglomeration were observed under these two alternative experiments. These results suggested that different halides would result in different Rh nanostructures and the existence of an appropriate amount of I⁻ ions was beneficial for the formation of the mutually embedded Rh concave nanocubes. According to the literature,¹⁹ halides tend to selectively adsorb to {100} planes. Generally, the six surfaces of a Rh cube are 〈100〉 oriented. So, we suggested that the selective adsorption of I⁻ ions on Rh {100} planes confined a growth along 〈100〉 direction and facilitated the formation of Rh concave structures with growth along {111} facets.Open in a separate windowFig. 4TEM images of the products prepared with different amount of KI while keeping the same other conditions. (a) Without KI; (b) 0.4 mmol KI; (c) 1.2 mmol KI; (d) 1.6 mmol KI.Moreover, the effect of HDBAC on the creation of the embedded Rh concave nanocubes was also investigated. As shown in Fig. 5a, no shaped Rh nanocrystal was produced except agglomerated nanoparticles without using HDBAC. While 0.1 mmol of HDBAC was used relatively to the parameters in the typical experimental procedure, Rh nanostructures with an nonuniform cross-sectional dimension and overly embedded each other were generated (Fig. 5b). With further increasing the amount of HDBAC from 0.4 to 0.6 mmol, the cross-section dimension and the concavity of the concave Rh nanocubes decreased gradually though embedded Rh nanostructures were still generated (Fig. 5c and d). Whereas no concave nanostructure was observed with using an equivalent amount of CTAB or CTAC in stead of HDBAC, respectively (Fig. S4, ESI^†). These results implied that the formation as well as the size and surface structure of the mutually embedded Rh concave nanocubes were dependent upon the confinement effect of HDBAC. On the one hand, the existence of HDBAC would contribute to creation of the concave cubes and their mutually embedded structures, on the other hand, the growth of shaped Rh nanoparticles was confined and the adsorption of I⁻ ions on Rh {100} planes was disturbed, resulting in less concavity and smaller size, due to an excessive amount of HDBAC.Open in a separate windowFig. 5TEM images of the products prepared with different amount of HDBAC while keeping the same other conditions. (a) Without HDBAC; (b) 0.1 mmol HDBAC; (c) 0.4 mmol HDBAC; (d) 0.6 mmol HDBAC.In order to investigate whether its formation was related to oxidation etching of Rh surface by I⁻ ion/O₂,^19,20 nitrogen was filled into the reaction bottle to remove oxygen before reaction and the same results were obtained. So, the oxidative etching can be negligible, which may be ascribed to the extremely short time under microwave irradiation.According to the previous report,²¹ based on the results of the experiments with dependent I⁻ ions and HDBAC, the formation of the mutually embedded Rh concave nanocubes may be ascribed to symmetry breaking due to asymmetric passivation and attachment of Rh nuclei. I⁻ ions were responsible for retarding the growth of {100} and {110} facets of cubic nuclei and promoting the preferential overgrowth on {111} planes, resulting in the formation of the concave structure with concave faces and edges. However, the existence of an appropriate amount of HDBAC may retard the deposition of Rh atoms on one or two corners with (111) facets due to the confinement, resulting in symmetry breaking. Meanwhile, with the confinement of HDBAC, the inevitable collision of nuclei leads to attachment of nuclei one another along the confined corners due to surface defects or dislocations. As a result, the mutually embedded Rh concave nanocubes would be generated with the growth of nuclei.The electrochemical performances of the as-synthesized Rh nanocrystals were examined by electrocatalytic oxidation of formic acid. The specific current density was normalized to the electrochemical surface area (ECSA). According to the cyclic voltammetry (CV) curves (Fig. S5^†), the ECSAs were calculated as 66.5 and 52.3 m² g⁻¹ for the mutually embedded Rh concave nanocubes and the commercial Rh black, respectively. Fig. 6a shows the CV curves for the electro-oxidation of formic acid in HClO₄ by the as-prepared mutually embedded Rh concave nanocubes and the commercial Rh black. The peak current density for the mutually embedded Rh concave nanocubes was measured to be 2.988 mA cm⁻² at 0.667 V, while it was 1.379 mA cm⁻² at 0.674 V for Rh black. The electrocatalytic activity of the hierarchical Rh nanostructures, though with a larger size, was about 2.2 times that of Rh black. Obviously, the mutually embedded Rh concave nanocubes exhibited an enhanced electrocatalytic activity for formic acid comparing with the commercial Rh black, which should be ascribed to their special surface structure with more edges, corners and terraces. Fig. 6b shows the CA curves of the electrocatalytic oxidation of formic acid for these catalysts. Compared with Rh black, a slower current attenuation as well as a higher retention of current after 800 s was observed for the as-prepared embedded Rh concave nanocubes, revealing a good electrochemical stability.Open in a separate windowFig. 6The CV (a) and CA (b) curves of the mutually embedded Rh concave nanocubes and Rh black in 0.1 M HClO₄ + 0.5 M HCOOH solutions with the cyclic potential between −0.2 and 1.0 V at a sweep rate of 50 mV s⁻¹.In summary, a novel hierarchical Rh nanostructure with several concave nanocubes embedded mutually could be rapidly prepared by reducing Rh(acac)₃ in TEG under microwave irradiation for 120 s in the presence of PVP and an appropriate amount of KI and HDBAC. In the preparing process, TEG was used as both a solvent and a reducing agent. The existence of KI and HDBAC was critical to the formation of the mutually embedded Rh concave nanocubes. The as-prepared mutually embedded Rh concave nanocubes demonstrated higher electrocatalytic activity and stability than commercial Rh black in the electro-oxidation of formic acid.     

Photoswitchable probe with distinctive characteristics for selective fluorescence imaging and long-term tracing

A photo-switchable and high-contrast bio-imaging indicator 4,4′-(1E,1′E)-(4,4′-(cyclopentene-1,2-diyl)bis(5-methylthiophene-4,2-diyl))bis(methan-1-yl-1-ylidene)bis(azan-1-yl-1-ylidene)bis(2-(benzo[d]thiazol-2-yl)phenol) (BMBT) has been demonstrated, by integrating photochromophore with excited-state intramolecular proton transfer (ESIPT) moiety. The ability of reversible emission switching enables arbitrarily selective labeling or concealing of cells simply by controlling light irradiation. Besides, when the emission was switched on, BMBT is demonstrated to exhibit unique characteristics of aggregation induced emission (AIE), providing a high on–off ratio for favorable bio-imaging. Thus, the non-labeling and easily-controlled selective imaging, as well as good biocompatibility indicates BMBT to be a favorable cell probe with great potentials for functional bio-imaging fluorophore.

In this work, a photoswitchable probe was synthesized by integrating a photochromophore with an excited-state intramolecular proton transfer (ESIPT) moiety. It was explored to be a favorable fluorophore for selective fluorescence imaging and long-term tracing.

Although there are many conventional fluorescent probes used for fluorescent imaging in the past few years,^1–5 such as rhodamine,⁶ cyanine dye,⁷ quantum dots,^8,9 and lanthanide probes,¹⁰ these fluorescent probes can only respond irreversibly to one event.^11,12 In comparison, photochromophores,¹³ which can reversibly response with UV and visible light,¹⁴ are more valuable fluorescent probes for regional optical marking of interested cells.^15,16 Because of their favourable characteristics, such as excellent thermal stability, good fatigue resistance and fast response time, diarylethenes derivatives have drawn wide spread concern of researchers.^17–20 Furthermore, the special optical properties of diarylethenes enable them to be suitable for long time and real time monitoring in bioimaging.The significance of targeted imaging of fluorescence probe for bio-samples in vitro or in vivo is well-known for their high sensitivity, non-invasiveness, real-time detection and especially selectivity.^21–24 However, the synthesis of these targeting materials is usually complicated. Moreover, for the same type of cells, no selectivity is demonstrated by these complex targeting agents.^21–24 Therefore, easy-prepared and easy-controlled non-labeling fluorescence imaging agents have always been pursued by both the industry and scientific communities. Among the above-mentioned photochromophores, diarylethenes are expected to be promising photochromic imaging candidates due to their favorable characteristics described in preceding paragraph. Prospectively, they can be used for selective long-term tracing if the photostability can be ensured.In our previous work, a series of AIE-active excited-state intramolecular proton transfer (ESIPT) complexes have been demonstrated to be good bio-imaging candidate with many advantages such as simple preparation, good biocompatibility, high quantum yields, fast cell staining as well as long-term anti-photobleaching.^25–27 Sometimes, ESIPT compounds exhibit dual emission, originated from keto and enol state, respectively. This caused extremely fast four-level photophysical cycle (E–E*–K*–K–E), mediated by intramolecular H-bonds immediately after photoexcitation, enables two emissions.^28,29 Herein, we have demonstrated a new type of multifunctional bio-imaging materials based on facile synthesis design concept, by introducing a photochromic diarylethene moiety to enable regional emission turn-on, and introducing an ESIPT moiety to allow good photo stability for long-term tracing. To the best of our knowledge, this is the first example with integrating abilities of non-labeling selectively long-term regional tracing.The new diarylethene derivative (BMBT, see the molecular structure in Scheme 1) has been synthesized (Fig. S1^†) via condensation reaction of 5-amino-2-(benzo[d]thiazol-2-yl)phenol with 4,4-(cyclopentene-1,2-yl)-bis(5-methyl-thiophene-2-formaldehyde), according to a previous reported procedure.^30,31 The molecular structures and purities were confirmed by ¹H NMR spectroscopy and mass spectroscopy (Fig. S6 and S7,^† see the synthesis details and full molecular characterizations in the ESI^†).Open in a separate windowScheme 1Structure and photochromic process of BMBT.The photoirradiation-induced changes in absorption and fluorescence spectra at room temperature are investigated in THF/DMEM mixture (1.0 × 10⁻⁵ mol L⁻¹). The open-ring isomer mainly exhibits two absorption peaked at 304 and 380 nm, respectively (Fig. S2^†). This is ascribed to the internal charge transfer and π–π transition of 2-(2′-hydroxy-phenyl)benzothiazole (HBT),^10,30 coupled with the CT inside the HBT unit from the hydroxyphenyl ring to the benzothiazole ring (Scheme 1). BMBT exhibits two emissions; one blue emission peak around 458 nm and a red emission peak around 600 nm, corresponding to the enol and keto emission, respectively (Fig. 1). Upon irradiation with ultraviolet light at 365 nm, the absorption band at longer wavelength region centered at 595 nm increased obviously with irradiation time (Fig. S2^†). This is caused by the formation of the closed-ring isomer (see the photochromic reaction in Scheme 1). Correspondingly, due to spectroscopic overlap between this longer-wavelength absorbance with the red emission ranged from 560–700 nm, the relative intensity of the red emission substantially decreases with the UV irradiation (Fig. 1), because of efficient energy transfer. Upon further visible light (λ = 520 nm) irradiation, the closed-ring isomer transfers back to the initial open-ring isomer, and thus the longer-wavelength absorption decreases and the red emission restores. This indicates good reversibility of the photochromic reaction.Open in a separate windowFig. 1Fluorescence emission changes of BMBT in THF/DMEM (1 : 200, vol : vol) mixture upon irradiation with 365 nm light (λ_ex = 385 nm) and visible light (λ_ex = 520 nm).Excitingly, BMBT exhibit prominent characteristics of Aggregation-Induced Emission (AIE).^31–33 The fluorescence intensity of BMBT in THF solution was relatively weak, while the powder or nanoparticles of the material dispersed in DMEM buffer exhibited a significantly enhanced emission (Fig. 2). When the DMEM buffer fraction was increased gradually from 0% to 20%, the fluorescence intensity only slightly enhanced. When the DMEM fraction was further increased to 40%, the emission exhibits a significant enhancement. The total intensity at the blue enol emission increased more than 10 times at 100% DMEM fraction as compared with that in THF solution. This AIE property enables high signal-to-noise ratios for favorable bioimaging (Fig. S3^†).Open in a separate windowFig. 2Fluorescence spectra of 1 × 10⁻⁵ M BMBT in the THF/DMEM mixture at different water fractions (λ_ex = 385 nm).Based on the good reversibility of light response and AIE characteristics, the practical application of the BMBT as bioprobe was further investigated. The biological imaging of BMBT was observed by using confocal laser scanning microscopy (CLSM). Blue luminescence in the cytoplasm of HeLa cells was observed after incubation with a THF/DMEM (1 : 200, vol : vol) solution of BMBT (20 μM) for 30 min at 37 °C (Fig. 2 inset). The overlay of luminescent images and bright-field images confirmed that BMBT was located mainly in the cytoplasm of cells rather than the membrane and nucleus (Fig. S3^†). Intense intracellular luminescence with a high signal-to-noise ratio (I₁/I₂ > 7) was detected between the cytoplasm (regions 3 and 1) and nucleus (region 2), also implying weak even few nuclear uptake of BMBT (Fig. S4^†). Besides, BMBT has a low cytotoxicity with the cellular viabilities estimated to be greater than 85% after 24 h incubation with the highest cultural concentration of 50 μM BMBT (Fig. S5^†).The luminescence switching of BMBT can also be achieved while alternating UV and visible light illumination in fixed HeLa cells. Cells (shown in red circle, Fig. 3a) were irradiated with 488 nm light (0.5 mW) for 3 min, the blue fluorescence of the irradiated cells was lamped off while the surrounding cells remained almost unchanged. Such fluorescence quenching is likely ascribed to the intramolecular fluorescence resonance energy transfer of BMBT, due to the intensified short-wavelength absorption band (270–450 nm) of the closed-open form of BMBT with the blue emission band (420–560 nm).³⁴ Upon irradiation with 405 nm light (1.25 mW) the fluorescence of all the selected cell was rapidly recovered within 1 min, caused by decrease of the relative intensity of the 270–450 nm absorption. The fluorescence can be repeatedly erased and recovered many rounds without significant fluorescence quenching, which was hardly achieved by conventional fluorophores (Fig. 3b).^35,36Open in a separate windowFig. 3(a) CLSM image (above) and the overlay image (bottom) of fixed HeLa cells incubated with 20 μM BMBT for 30 min at 37 °C (1) and (5) in original state; (2) and (6) irradiated by 488 nm light (0.5 mW) for a single cell; (3) and (7) all cells, and; (4) and (8) recovered by 405 nm light (1.25 mW). (b) Fluorescence switching of fixed HeLa cells by alternating UV (405 nm, 1.25 mW, 10 s/time) and visible (488 nm, 0.5 mW, 3 min/time) light illumination (λ_ex = 405 nm).This characteristic of selectively opto-marking or de-marking of cells may be used for non-invasive and dynamic tracing of the interested objects in vitro. That is, BMBT can arbitrarily opto-label or de-label interested cells without affecting cell proliferation. As shown in Fig. 4, the cell marked in the red circle was treated with visible light illumination (488 nm, 0.5 mW, 3 min), and its fluorescence was effectively quenched. Cells division of the remaining “bright” cells was observed under the microscope field of vision for a long-term tracing with time up to 36 h. Multiple new cells were produced, as indicated by upper white arrow. Even, as indicated by bottom white arrow, the tacked cells were observed to be doubled. The high brightness of the fluorescence in the proliferated cells indicated the good photo stability of BMBT. This indicated that BMBT can be used as a cell marker for arbitrarily selective erasing the fluorescence of the designated cell. It can also be used for selective lighting their emission by making them as the remaining bright cells after selective photo-erasing or selective photo-recovering after full erasing. This long-term tracking with non-labeling and selective optically marking or de-marking is seldom reported by other photochromophore-based bioimaging agents.³⁷ Herein, the excellent anti-photo bleaching characteristic in the long-term tracing is attributed to high photo stability of the HBT moiety.^26,27Open in a separate windowFig. 4Cells was treated with visible (488 nm, 0.5 mW, 3 min) light illumination for erasing the fluorescence. The remaining bright cells were incubated for another 36 hours, which were observed to be amplified normally (λ_ex = 405 nm).     

Thermally induced fragmentation of nanoscale calcite

Calcite nanorods ∼50 nm wide are thermally separated into nanoblocks. The fragmentation is ascribed to the ion diffusion on metastable crystal surfaces at temperatures (∼400 °C) much lower than the melting point. The presence of water molecules enhances the surface diffusion and induces deformation of the nanorods even at ∼60 °C.

Calcite nanorods ∼50 nm wide are thermally separated into nanoblocks.

Calcium carbonate is a common industrial material that is used as a micrometric filler for papers, rubbers, plastics, and inks. The shape and size of micrometric grains are important parameters that affect the physical and chemical properties of composite materials.^1–5 In recent years, nanometric particles of calcium carbonate have attracted much attention as basal building blocks of biogenic minerals^6–10 and functional materials with high biocompatibility and low environmental load.^11–14 Since various properties are influenced by the miniaturization of crystal grains below 100 nm, characterization of the carbonate particles is necessary for application in practical fields. However, their properties, including the thermal characteristics of nanometric calcium carbonate in the nanometric region, have not been sufficiently clarified because calcium carbonate is easily decomposed above 550 °C.^15,16 In the present work, we studied thermally induced deformation on nanometric calcite at temperatures lower than the melting point (1597 °C at 3 GPa)¹⁷ of the bulky crystal.Since the melting temperature of metals decreases when their size is decreased below ∼50 nm,^18,19 metallic nanowires fragment into nanospheres at temperatures much below the melting point of bulk metal.^20–22 The fragmentation is ascribed to the Rayleigh instability that is known for liquid. These results suggest that the surface diffusion of ions and atoms occurs easily on the nanometric particles when the surface instability is increased. The cleavage of solid nanowires has been observed only for metallic phases. In the current work, we found the morphological transformation of ion crystal nanorods into faceted nanograins through the surface diffusion at relatively low temperatures.The preparation of bulky calcium carbonate consisting of nanocrystals is required for reinforcing materials^23–25 and as a precursor of biomedical materials.^26,27 In general, however, fabricating large, bulky bodies through conventional sintering techniques is difficult because calcium carbonate is thermally decomposed into calcium oxide and carbon dioxide. Several methods, such as sintering with a flux^16,28 or in a carbon dioxide atmosphere²⁹ and hot-pressing under hydrothermal conditions^30,31 were developed to prepare bulky calcium carbonate materials. On the other hand, the thermal behaviors of pure calcium carbonate have not been sufficiently studied at temperatures below the decomposition and melting points.In nature, bulky calcium carbonate crystals are commonly produced as biominerals, such as shells, eggshells, sea urchins, and foraminiferal skeletons, under mild conditions.^{3,9,10,32–34} The bulky biogenic bodies are composed of nanometric grains 10–100 nm in size that are arranged in the same crystallographic direction.³ The formation of bridged architectures through oriented attachment is generally related to the ion diffusion on specific surfaces at relatively low temperatures. A detailed study on the stability of nanometric surfaces at relatively low temperatures is needed to understand the morphological change of calcium carbonate crystals.In the present report, we discuss the morphological change of calcite nanocrystals below the decomposition and melting temperatures by observing metastable crystal surfaces in the nanometer-scale range. Calcite nanorods elongated in the c direction were utilized as a typical nanometric shape covered with metastable surfaces. Here, thermally induced fragmentation was studied with and without water vapor. The surface diffusion was found to occur on the metastable surface at around 400 °C under a dry condition and at around 60 °C with water vapor. Our findings are important for clarification of the surface property of nanometric calcium carbonate and for the fabrication of bulky bodies through the attachment of nanocrystals.Single-crystal calcite nanorods elongated in the c direction were utilized in the present study (Fig. 1 and S1^†). Calcite nanorods up to ∼500 nm were formed through the combination of the carbonation of calcium hydroxide and the subsequent oriented attachment of resultant calcite nanoblocks ∼50 nm in diameter by stirring. The detailed mechanism was described in our previous study.³⁵ As shown in Fig. S1,^† the calcite nanorods were covered with metastable surfaces having a curvature. Moreover, we observed depressed parts originating from the oriented attachment of the original calcite grains. As shown in the Fig. S2,^† the XRD peaks of the nanorods were found to shift to higher angles than those of standard X-ray diffraction data (ICDD 00-005-0586). Thus, the crystal lattice of the nanorods was suggested to be stressed with the coverage of irregular surfaces.Open in a separate windowFig. 1SEM (a) and TEM (b and c₁), HRTEM (c₂) images, and the FFT pattern (c₃) of the lattice in (c₂) of calcite nanorods in aqueous dispersions at pH 12 and at 25 °C with stirring. (c) Reprinted from ref. 35 published by The Royal Society of Chemistry.As shown in Fig. 2, the calcite nanorods were deposited on a silicon substrate for clear observation of the morphological change. We redispersed the calcite nanorods in ethanol and evaporated the dispersion medium to deposit them on the substrate at 25 °C (Fig. 2a and S3a^†).³⁶ The nanorods were arranged on the solid surface through evaporation-driven self-assembly. Specifically, monolayers of the calcite nanorods were obtained by adding poly(acrylic acid) (PAA, MW: 5000 gmol⁻¹) to the ethanol dispersion. The dispersibility of the nanorods in ethanol was improved by the modifying agent. The organic components of the PAA-modified nanorods were removed through oxidation in air at around 300 °C (Fig. S4^†).Open in a separate windowFig. 2SEM images and schematic illustrations of calcite nanorods deposited on a silicon substrate before treatment (a) and heated at 400 °C for 1 (b) and 2 h (c). The PAA-modified calcite nanorods were used to obtain the monolayers (b and c). We used bare nanorods for take the images of original nanorods because the definite surfaces were not observed due to the presence of PAA (a).The calcite nanorods were deposited on the solid surface to study the morphological change when heated in air. Obvious changes were not found in the shape of the nanorods upon heating to a temperature below 350 °C in air for 24 h (Fig. S2^†). On the other hand, we observed significant deformation upon heating to 400 °C (Fig. 2). The depressed parts on the side surfaces enlarged in 1 h. The fragmentation of the nanorods was finally induced after treatment for 2 h. Calcite grains were formed by the thermally induced cleavage (Fig. 2c). The average size of the cleaved grains was ∼100 nm, which was larger than the average width of the nanorods, ∼50 nm. As shown in Fig. 3, we observed the definite surfaces covering nanograins. Most of the definite facets were assigned to the (104) of calcite by FFT analysis of the lattice fringes of nanograins in the HRTEM images. Some (012) planes were found in the nanograins. On the other hand, the surfaces of the deformed nanorods during fragmentation were curved and irregular. According to XRD patterns, the crystal phase was not changed with the fragmentation (Fig. S2^†). The diffraction peaks were shifted to the standard values and sharpened with the treatments. This suggests that the formation of the stable faces with the fragmentation is associated with the lattice relaxation.Open in a separate windowFig. 3TEM (a), HRTEM (b), and FFT images of nanocrystals before and after heating at 500 °C.We observed the stability of rhombohedral grains that were covered with the stable {104} planes. As shown in Fig. 4, the deformation of the rhombohedral grains was not observed at 400 °C in air for 6 h. These results indicate that the ion diffusion is not induced drastically on the stable surfaces at a temperature lower than the decomposition temperature.Open in a separate windowFig. 4SEM (a_1,2 and b_1,2), TEM (a₃), and SAED (a₄) images of calcite nanoblocks before treatment (a) and heated to 400 °C for 6 h (b). (a₃) A schematic illustration of a calcite rhombohedron covered with {104} faces.The fragmentation of the calcite nanorods was enhanced in the presence of water vapor. As shown in Fig. 5a, we found cleavage of the nanorods even at 60 °C in a closed vessel containing water. The formation of unifaceted rhombohedral nanoblocks with {104} faces was clearly observed at 100 °C for 24 h (Fig. 5b–d). Since the morphological change was similar to that under a dry condition, the ion diffusion on the surface is deduced to be assisted by adsorbed water molecules. The X-ray diffraction signals shifted to the standard values with the fragmentation (Fig. S2^†). Thus, the stable faces were formed with the lattice relaxation with the exposure to water vapor.Open in a separate windowFig. 5SEM (a and b), TEM (c), and SAED (d) images of calcite nanorods deposited on a silicon substrate subjected to high humidity at 60 °C (a) and 100 °C (b) for 24 h. (c) A schematic illustration of a calcite rhombohedron covered with {104} faces.The ion diffusion at a relatively low temperature, below the decomposition temperature, has not been reported for calcium carbonate. In the present work, however, we found the fragmentation of calcite nanorods at around 400 °C under a dry condition and at around 60 °C with water molecules. These results suggest that the ion diffusion occurs on the nanoscale calcite crystals. On the other hand, rhombohedral calcite grains covered with the stable {104} planes were not deformed at those temperatures. Thus, the diffusion at low temperatures is induced only on metastable surfaces that are exhibited on the nanoscale calcite. Moreover, the presence of water molecules enhances the ion diffusion on the metastable surfaces.The fragmentation of the calcite nanorods can be explained by Rayleigh instability. The cleavage by Rayleigh instability is ascribed to the enlargement of tiny perturbations on cylindrical liquids,³⁷ polymers, and metals. In general, the cylindrical bodies evolve into several spheres to decrease the total surface energy. Recently, Rayleigh instability was applied to the thermally induced fragmentation of metal^20–22 and organic³⁸ nanowires. The breakup phenomena were attributed to surface oscillations due to the high surface energy induced by increased surface-to-volume ratios.³⁹ Thus, metal nanowires are cleaved and form isotropic nanoblocks at temperatures well below the melting point. In the present work, we found fragmentation of the calcite nanorods at relatively low temperatures. The ion diffusion is induced on the metastable surfaces that are exhibited on nanoscale crystals. The depressed parts exist as perturbations on the side faces of the original nanorods. The cleavage occurs through enlargement of the depressed parts and formation of the stable faces to reduce the surface energy and relax the lattice strain.Polyhedral grains covered with flat planes were formed instead of spherical particles by the fragmentation of calcite nanorods. Formation of the stable {104} plane is achieved to reduce the surface energy. The {012} plane of calcite is not stable under the ambient temperature. However, the surface energy of {012} decreases with increasing temperature.⁴⁰ Thus, the facets are deduced to also be formed at temperatures around 400 °C.     

Separating graphene quantum dots by lateral size through gel column chromatography

Graphene quantum dots (GQDs) prepared through photo-Fenton reaction of graphene oxide are separated via gel column chromatography. The as-separated GQDs were selectively introduced into the active layer of organic solar cells and achieved an enhancement of power conversion efficiency (PCE).

GQDs prepared through a photo-Fenton reaction were separated into eight groups with different sizes and fluorescent colors via gel column chromatography.

Graphene quantum dots (GQDs) show potential applications in photovoltaic devices, bio-probers, sensors, and catalysts.^1–6 As the properties of GQDs can be affected severely by their lateral sizes and size distributions,^7,8 to acquire GQDs with controlled size and narrow distribution is prerequisite. However, GQDs prepared directly by the methods developed so far usually assume wide size distribution which limits somehow the practical applications of GQDs.^2,9–12Recently, several protocols have been developed for post separation of GQDs, such as dialysis,¹³ ultrafiltration,¹⁴ gel electrophoresis,⁸ reverse micelle methods,¹⁵ column chromatography on silica¹⁶ or Sephadex G-25 gel,¹⁷ chromatographic separation,¹⁸ and size-selective precipitation,^19,20 but can''t satisfy the bulk production. For examples, Kim et al. successfully obtained GQDs with different sizes using dialysis bags with different interception molecular weights and a 20 nm nanoporous membrane, but with an unacceptable yield.¹³ Fuyuno et al. obtained the GQDs with different fluorescence by the size-exclusion high performance liquid chromatography (HPLC).¹⁸ Jiang et al. separated the single atomic layered GQDs from reaction mixture containing double multilayer allotropes successfully through a Sephadex G-25 gel.¹⁷ In our previous work, the GQDs generated through the photo-Fenton reaction of graphene oxide (GO) have been sorted into three categories with different fluorescence by gel electrophoresis.^2,8,20 Nevertheless, it is still challenging to obtain high quality GQDs with controlled size and size distribution which can satisfy the practical applications.Herein, we describe an efficient GQDs separation procedure via Sephadex G25 gel column. The GQDs prepared through photo-Fenton reaction of GO with wide size distribution are separated into eight groups of GQDs with different size and fluorescent colours.² The size and morphology of as-obtained GQDs were characterized by atomic force microscopy (AFM) and transmission electron microscopes (TEM) measurements. The optoelectronic properties of the GQDs were studied by photoluminescence (PL) and UV-vis absorption spectroscopy techniques. The results showed that this separating technique is very beneficial for obtaining high quality GQDs with a variety of specific sizes and properties. Finally, the as-separated GQDs were introduced into the inverted hybrid solar cells based on the poly(3-hexylthiophene) (P3HT) and poly(3-hexylthiophene)/(6,6)-phenyl-C61 butyric acid methylester (PCBM), and it is found that the solar cell containing the separated GQDs showed a higher performance than that with the raw GQDs, which verified the importance of the size separation for GQDs.The raw GQDs used in the work are first characterized using atomic force microscopy imaging. As shown in Fig. S1a and b,^† their sizes are ranged from 2 to 40 nm with obviously large size distribution, that is further confirmed by PL spectrum and image (Fig. S1c,^† and the inset). Sephadex G25 gel column, one of common size exclusion column, is widely used to purify or separate protein or peptide.²¹ Here, GQDs are separated through Sephadex G25 gel column by size and the as-separated GQDs are named as GQDs 1–8 according to the collection order. The yield of GQDs is close to 80% with this separating technique. Actually, unlike other separating methods such as multiple dialysis¹³ and ultrafiltration,¹⁴ there is almost no significant loss of GQDs in our separation process. Taking the well dispersibility of as-prepared GQDs in water into consideration, we selected water as developing solvent in this work. Their morphologies, size and size distributions are revealed by AFM imaging and the results are shown in Fig. 1. The average size of GQDs 1–8 (calibrated with the parameters of AFM tip deconvolution⁸) are of 27.5, 23.5, 15.5, 12.0, 8.5, 6.3, 5.2 and 3.0 nm, respectively, with narrow size distributions (see the as-inset histograms in Fig. 1). As shown in Fig. S2,^† the sizes of GQDs 1–8 are also measured by TEM imaging, which are in agreement with the AFM images.Open in a separate windowFig. 1Tapping mode AFM images (height) of the separated GQDs samples along the collecting order (a) GQDs 1, (b) GQDs 2, (c) GQDs 3, (d) GQDs 4, (e) GQDs 5, (f) GQDs 6, (g) GQDs 7, (h) GQDs 8. The insets are the histograms of the size.The PL and UV-vis spectra of the GQDs can reflect the size difference, too.^7,8,13,18 The top row of Fig. 2a shows the optical images of the GQDs 1–8 acquired under a daylight lamp, and they were all transparent. The bottom row of Fig. 2a shows the optical images of the corresponding GQDs 1–8 observed under a UV irradiation (302 nm), illustrating that the GQDs 1–8 have fluorescence properties, which are red, orange, yellow, green, cyan, light blue, blue, and purple, respectively. In contrast to the raw GQDs, the result indicates that the GQDs are successfully separated by size through Sephadex G-25 gel column. This should be beneficial for further exploring the relationship between the size and properties of GQDs. As shown in Fig. 2b, the PL spectra of GQDs 1–8 match well with the fluorescence photos. The peak wavelengths of their PL are 587, 565, 554, 483, 462, 452, 385, 384 nm, correspondingly.Open in a separate windowFig. 2(a) The top row is the optical images aqueous suspensions of GQDs 1–8 obtained under daylight lamp; the bottom row is the optical images of the aqueous suspensions GQDs 1–8 acquired under UV irradiation (302 nm). (GQDs-1/red, GQDs-2/orange, GQDs-3/yellow, GQDs-4/green, GQDs-5/cyan, GQDs-6/light blue, GQDs-7/blue, GQDs-8/purple). (b) PL spectra of GQDs 1–8 (the excitation wavelength is 340 nm). (c) UV-vis absorption spectra of GQDs 1–8 (the spectra were normalized at 200 nm for comparison).As shown in Fig. 2c, with the size decreasing, the absorption onset of the GQDs blue-shifted gradually. The absorption around 225 nm corresponding to the π → π* transition of sp² domains in GQDs, and the absorption in the range of 275–325 nm from the n → π* transition of C O groups at the edge of GQDs are also observed clearly, which is similar to the literature.⁸ More specifically, the absorption peak in the range of 275–325 nm becomes more and more obvious with the GQDs size decreasing, indicating the density of carboxylic groups at the edge of GQDs increases from GQDs-1/red to GQDs-8/purple. The reason is that the number of GQD carboxylic groups is directly proportional to its lateral size and the area of GQDs is proportional to the square of its size, as a result the small sized GQDs have higher density of carboxylic groups than the large sized GQDs and present the obvious peak around 275–325 nm.⁸The PL spectra of the as-separated GQDs are shown in Fig. 3 and Fig. S3.^† Comparably, the PL intensities of GQDs 3, GQDs 4, GQDs 5, GQDs 6 (Fig. 3) are stronger than those of others (Fig. S3^†). The PL emission peaks of the GQDs 3 and GQDs 4 shift more obviously than that of GQDs 5 and GQDs 6 with the increase of the excitation wavelength, implying the size distributions of GQDs 3 and GQDs 4 are worse than those of GQDs 5 and GQDs 6. For GQDs 3, as shown in Fig. 3a, there are two peaks in the emission spectra with the excitation wavelength of 360, 380, 400 nm. The left peak is attributed to the π* → n transition of carbonyl or carboxylic and the right peak is attributed to the sp² domains in carbon skeleton. For GQDs 4, as displayed in Fig. 3b, with the excitation wavelength increasing from 300 to 400 nm, the main contribution for the PL is still the sp² domains in carbon skeleton. With the excitation wavelength of 380, 400 nm, two slight shoulders occur in the emission spectra, the left is attributed to the π* → n transition of carbonyl or carboxylic, too. With the decreasing of GQDs size, the PL intensity from the sp² domains gets weak, but the one of π* → n transition increases. The emission peaks of GQDs 5 and GQDs 6 shift slightly (Fig. 3c and d), which is mainly attributed to the π* → n transition of carbonyl or carboxylic, but partly from the sp² domains in carbon skeleton.^8,22Open in a separate windowFig. 3The fluorescence emission spectra of GQDs-3/yellow samples with excitation wavelengths from 360 nm to 520 nm (a), GQDs-4/green with excitation wavelengths from 300 nm to 400 nm (b), GQDs-5/cyan with excitation wavelengths from 280 nm to 380 nm (c), and GQDs-6/light blue samples with excitation wavelengths from 280 nm to 360 nm (d).The PL quantum yields (QYs) of raw and separated GQDs are measured using quinine sulfate as a reference (QY = 57.7%),^3,23 and are summarized in Table S1.^† The QYs of the raw GQDs and GQDs 1–8 are 0.99, 0.611, 0.758, 2.592, 5.905, 1.816, 0.486, 0.259, and 0.199%, respectively. Obviously, the QYs of GQDs 3, GQDs 4, and GQDs 5 are much higher than those of others, but the QYs of GQDs 1, GQDs 2, GQDs 6, GQDs 7, and GQDs 8 are much lower than that of the raw GQDs. This may be resulted from the comprehensive factors from the quantum confinement effect, and the functional groups on the edge of GQDs.In fact, the size and surface functionality of the raw GQDs are the key factors dominating Sephadex G25 gel column separation efficiency. By simply varying the photo-Fenton reaction time, different raw GQDs are prepared. Accordingly, as shown in Fig. S4,^† GQDs assuming different fluorescent colours can be obtained. When the photo-Fenton reaction time was 90 minutes, big sized GQDs with yellow and orange fluorescence can be rarely obtained. Only blue and cyan fluorescence GQDs could be separated using G25 gel chromatography (Fig. S5^†). It can be concluded that the separating extent is seriously depended on the size and surface functionality of GQDs as-obtained via photo-Fenton reaction. Similarly, only the GQDs with blue fluorescence could be separated from the raw material of GQDs prepared by hydrothermal method⁴ using G25 gel chromatography. Recently, various GQDs prepared by reported methods are mono-fluorescence such as blue or green and they are not suitable for the suggested separating technique.^24–26 Thus, the suggested separating technique is not universal for GQDs obtained via different preparing methods.In order to explore the advantages of the as-separated GQDs, raw GQDs and GQDs 4 with the highest QY are used as additivity for the electron acceptor material PCBM, and inverted structure organic ternary hybrid solar cells (Ag/MoO₃/P3HT:PCBM:GQDs/ZnO/ITO) were assembled. The photovoltaic performances of as-fabricated solar cells were characterized, and the results are depicted in Fig. 4 and Table S2.^† The power conversion efficiencies (PCEs) of the solar cells containing raw GQDs and GQDs 4 are of 3.46% and 3.91%, respectively, which are higher than that of the control group (3.07%). Further, the performance of the solar cell with GQDs 4 is even better than that with raw GQDs, which means that the size and size distribution are crucial to the optoelectronic performances. However, the detailed mechanism of the photovoltaic performances of the as-assembled inverted structure organic ternary hybrid solar cells are not clear for us at moment, and will be further addressed in our coming work.Open in a separate windowFig. 4 J–V characteristics of the solar cells based on P3HT:PCBM:GQDs active layers with different GQDs.     

Effect of nanostructured silicon on surface enhanced Raman scattering

Non-metallic materials are often employed in SERS systems by forming composite structures with SERS-active metal materials. However, the role of the non-metallic structures in these composites and the effect of them on the SERS enhancement are still unclear. Herein, we studied the effect of silicon morphology on SERS enhancement on silver nanoparticles-coated different structured silicon surfaces. Our finding will help to further understand the SERS mechanism and pave the way for making more efficient SERS systems.

The surface morphology of non-metallic silicon has a big effect on the SERS enhancement of silver nanoparticle-coated silicon surfaces.

In past decades, increasing attention has been attracted to surface enhanced Raman scattering (SERS) due to the dramatically enhanced detection sensitivity of Raman scattering (down to single-molecule sensitivity). The Raman intensity of the molecules located in the vicinity of SERS nanostructures can be enhanced up to 10¹⁰ to 10¹¹ times,^1,2 largely extending the application of SERS in the fields of physics, chemistry and biology, etc.^3–9 Metal materials are mainly employed in fabricating SERS structures, especially gold or silver ones for visible spectrum excitation.^10,11 To date, many nanostructures have been reported that can enhance Raman scattering enormously, leading to so called Raman hot spots,^12–15 including nanogaps,^16,17 nanostars,¹⁸ nanotriangles and nanorods,¹⁹ mainly due to the introduction of a localized electromagnetic field under illumination. This kind of enhancement is referred to as the electromagnetic mechanism, which dominates the SERS enhancement in most cases.Non-metallic structures can also contribute to Raman enhancement, although the enhancement factor is usually very low. It has been reported that Cu₂O,²⁰ TiO₂ ²¹ and ZnS²² nanoparticles can enhance the Raman intensity of adsorbed molecules. Graphene has also been proved to be an efficient platform for Raman enhancement.²³ Although non-metallic materials can be used directly for SERS applications, they are usually used by forming a composite with SERS-active metal material, where they act as supporting materials or borrow the SERS activity from the metallic Raman hot spots. It has been reported that SERS activity can be borrowed from SERS-active materials through ultrathin SERS-inactive transition metals (e.g., Pt, Ni, Co and Pd)²⁴ or dielectric (e.g., SiO₂, Al₂O₃)²⁵ layer. Tian et al. reported the shell-isolated nanoparticle-enhanced Raman spectroscopy (SHINERS) by using the gold nanoparticles coated with ultra-thin silica or aluminum oxide shell.^26,27 Raman enhancement can be achieved at the silica shell surface by borrowing the SERS activity from the gold core. However, the role of non-metallic structures in enhancing Raman scattering and the interactions between these two kinds of materials are still fuzzy.Silicon nanostructures fabricated by catalytic etching method can be easily metalized with silver or gold by electroless deposition for SERS applications.^28–32 However, up to now, only limited types of silicon nanostructures have been reported for SERS applications.^17,30,31 In addition, the role of the nanostructured silicon surface in Raman enhancement is still unclear.In this report, we fabricated SERS structures on two types of silicon surfaces, flat silicon and nanoporous silicon, by metallizing the silicon structure with silver nanoparticles (AgNPs). Compared to the fabricated SERS structure on flat silicon surface, the one fabricated on nanoporous silicon surface showed obvious enhancement on the Raman spectrum of adsorbed probe molecules. The effect of pore size and depth of nanoporous silicon on Raman enhancement was investigated in detail.To investigate the role of silicon nanostructures in Raman enhancement, we compared the Raman spectra of probe molecules adsorbed on AgNPs-coated flat silicon and nanoporous silicon surfaces, respectively. The nanoporous silicon was fabricated by following a modified reported procedure (see details described in Experimental section and scheme shown in Fig. S1 in (ESI^†)).²⁸ Vertical nanopores were produced on silicon surface, and the pore size and pore depth can be easily tuned by varying the reaction parameters. Then a flat silicon and a nanoporous silicon substrates were both metallized with silver by immersing them into a mixed aqueous solution of AgNO₃ and HF,^29,33 forming AgNPs with size of 60 ± 30 nm. The as-prepared AgNPs-coated silicon surfaces (see scanning electron microscopy (SEM) images shown in Fig. S2 in ESI^†) served as SERS-active substrates. After p-aminothiophenol (PATP) molecules (Raman probe) were adsorbed on the AgNPs-coated silicon structures, both AgNPs-coated substrates showed uniform and strong Raman enhancements (Fig. 1). On AgNPs-coated nanoporous silicon surface, the Raman bands of PATP molecules at 1076 and 1142 cm⁻¹ are 4.2 and 7.4 times stronger compared to those on AgNPs-coated flat silicon surface (Fig. 1), respectively, demonstrating the vital role of silicon morphology in the obtained Raman enhancement. There are two widely accepted mechanism for SERS enhancement, electromagnetic mechanism and charge transfer mechanism.¹ Both flat silicon and nanoporous silicon substrates are composed of same material and the only difference between them is the silicon morphology. Thus the charge transfer mechanism should contribute similar effect in both conditions. Moreover, both flat and nanoporous silicon surfaces were covered with a thin layer of silicon dioxide,^28,34 which limit the charge transfer between AgNPs and silicon surface. This is also confirmed by the XPS measurement on the nanoporous silicon surface (Fig. S3 in ESI^†). Therefore, the observed different enhancement may attribute to the electromagnetic mechanism, which will be discussed latter. In addition, the Raman enhancement is uniform over the whole substrate. This is probably due to the uniform coating of AgNPs on high-density silicon nanopore structures. The enhancement factor (EF) can be calculated by using the following equation, EF = (I_SERS/I_bulk)(N_bulk/N_SERS),³⁵ where I_SERS and I_bulk represent the Raman intensities in SERS and bulk Raman measurements, respectively; N_SERS and N_bulk represent the number of probe molecules located in the excitation volume under these two conditions. For Raman band at 1076 cm⁻¹ (represents a₁ vibration mode of PATP,¹¹ which sits at 1089 cm⁻¹ for bulk,³⁶ Fig. S4 in ESI^†), the average EFs over the whole surface were calculated as 6.7 × 10⁵ and 2.8 × 10⁶ for SERS structures on flat silicon and nanoporous silicon, respectively. The strong Raman band at 1142 cm⁻¹ indicates a chemical conversion from PATP to 4,4′-dimercaptoazobenzene (DMAB) upon light irradiation.¹¹Open in a separate windowFig. 1Raman spectra of PATP molecules adsorbed on the AgNPs-coated (a) flat silicon and (b) nanoporous silicon. A nanoporous silicon with pore depth of 220 nm was used here. The schemes at the bottom right and top right show the structures of the AgNPs-coated flat silicon surface and AgNPs-coated nanoporous silicon surface, respectively. The size of AgNPs was not drawn to scale.As discussed, electromagnetic mechanism dominates the observed SERS enhancement. To confirm the role of silicon morphology, we did numerical simulation using the finite-difference time-domain (FDTD) method to investigate the localized electromagnetic field distributions on AgNPs-coated flat and nanoporous silicon surfaces. Note that, for the AgNPs-coated nanoporous silicon surface, many AgNPs sit on the edge of silicon nanopores (Fig. S2D in ESI^†). In this case, the electromagnetic field around the AgNPs is more localized. Compared with the AgNPs-coated flat silicon, the electromagnetic field is five times more localized on the AgNPs-coated nanoporous silicon surface (Fig. 2), which in principal could introduce 25 times stronger Raman enhancement.³⁷ However, in real case, only a proportional of the AgNPs locates on the edge of silicon nanopores and the shapes of the coated AgNPs are not exactly same with the ones we used in simulation, which explains the smaller SERS enhancement we observed on nanoporous silicon surface.Open in a separate windowFig. 2Schemes (side views) and FDTD simulations on the AgNPs-coated flat silicon (A and B) and nanoporous silicon (C and D). Dashed circles in (D) indicate the positions of silicon pores. The schemes in (A) and (C) were not drawn to scale.As discussed above, the morphology of nanoporous silicon contributes to the enhanced Raman signal. By varying the pore size and pore depth of the nanoporous silicon, different Raman enhancement should be observed.First, we studied the effect of pore depth of nanoporous silicon on Raman enhancement. The depth of silicon nanopores can be easily tuned by varying the period of catalytic etching of silicon. The Raman intensity of the probe molecules increased continuously with increased silicon nanopore depth (from 40 to 220 nm, Fig. 3). When PTAP molecules were adsorbed on the AgNPs-coated nanoporous silicon surface, the Raman intensity measured on silicon with 220 nm pore depth was increased about 2 times compared to that on silicon with 40 nm pore depth. Further increasing the pore depth to 900 nm, the Raman intensity dropped instead (Fig. 3). These results indicate the important role of the pore depth in Raman enhancement. FDTD simulations were carried out to investigate the mechanism behind (Fig. S5 in ESI^†). As the pore depth increases, the electromagnetic field becomes more localized, which is consistent with the experimental data. However, the Raman intensity decreased on surface with very deep silicon nanopores (900 nm), which can be explained by the enhanced light trapping.^38,39 In this case, part of the Raman scattering light cannot escape from the nanopores (confirmed by the dark black color of the sample, Fig. S6 in ESI^†), leading to a weaker Raman signal. This can also be double confirmed by studying the Raman scattering from nanoporous silicon samples with AgNPs located at the bottom of the nanopores (discussed in ESI and Fig. S7^†).Open in a separate windowFig. 3Raman intensity variation (peak at 1076 cm⁻¹) on the AgNPs-coated nanoporous silicon (pore size of ∼40 nm was used) surface with four different depths. The scheme on top was not drawn to scale.Second, the size of silicon nanopores also plays a role in the Raman enhancement. The pore size on silicon surface can be tuned by controlling the size of catalysts (AgNPs) deposited on silicon wafer (Fig. S1B in ESI^†), whose size was replicated by the nanopores in subsequent catalytic etching process (Fig. S1C^†). By varying the deposition time, nanoporous silicon samples with four different pores sizes, 31 ± 10, 41 ± 11, 80 ± 24 and 160 ± 50 (Fig. 4A–D), were fabricated, respectively. For the AgNPs-coated nanoporous silicon samples, the Raman intensity of probe molecules slightly changed while increasing the pore size (Fig. 4E), indicating a weak effect of pore size on Raman enhancement. As aforementioned discussion, the Raman enhancement is mainly contributed by the AgNPs that locate on the nanopore edges. Therefore, the Raman enhancement is strongly dependent on the perimeter of all the nanopores and the number of AgNPs that locate on the edge of silicon nanopores. While increasing the pore size, the perimeter of single pore increases. However, many pores are fused together, compensating the increase of the perimeter of single pore. Thus, the total perimeter of all nanopores does not change much when increasing the pore size. In this case, the amount of the AgNPs locating on the edge of silicon nanopores may not change too much, which may explain the less dependency of the pore size on SERS enhancement. To investigate the structure of the AgNPs-coated nanoporous silicon, we deposited gold nanoparticles (AuNPs) onto it. In this case, however, a stronger Raman scattering was observed due to the formation of AgNP–AuNP nanogaps and the enhancement varied on different sized silicon nanopores. When increasing the pore size from 31 ± 10 to 80 ± 24 nm, the Raman scattering became stronger. Further increasing the pore size to 160 ± 50 nm, the Raman intensity decreased. As known, two particles formed nanogap shows a more localized electromagnetic field when the polarization of incident light is parallel to the center to center axis of the two particles.^16,40 Therefore, horizontal positioned two-particle nanogaps will give much stronger Raman enhancement. If the size of the silicon nanopores is too small, it is difficult for the AuNPs (13 nm in diameter) to enter the pore, limiting the number of AgNP–AuNP nanogaps that are horizontally positioned, and in turn limiting the Raman enhancement. While increasing the nanopore size, we have a better chance to form the ideally positioned AgNP–AuNP nanogaps to improve the Raman enhancement (Fig. S8 in ESI^†). However, when the pore size is too big, only a small part of the nanostructures locates inside the excitation volume during Raman measurement, leading to a weaker Raman signal.Open in a separate windowFig. 4(A–D) Nanoporous silicon with different pore sizes obtained by varying the silver deposition time described in Fig. 1A and B. Scale bars = 200 nm. The pore depth here was set as 220 nm. (E) Raman intensity variation (peak at 1076 cm⁻¹) on the AgNPs-coated nanoporous (four different pore sizes shown in A–D) silicon surface without (magenta bars) and with (cyan bars) the adsorption of AuNPs.The Raman enhancement can also be affected by the size of the AgNPs coated on nanoporous silicon. The AgNP size can be tuned by varying the AgNP deposition time shown in Fig. S1E.^† It has been reported that AgNPs with several tens of nanometers showed optimized plasmon resonance with excitation wavelength of 632.8 nm.⁴¹ In this work, the 40–75 nm AgNPs coated on nanoporous silicon show higher Raman enhancement than those with smaller or bigger AgNPs (Fig. S9^†), since the size of these AgNPs fall into the optimized range for Raman enhancement, which is consistent with the reported work.     

Green synthesis of crystalline bismuth nanoparticles using lemon juice

Lemon juice effectively served as a reducing and capping agent for an easy, cost-effective, and green synthesis of crystalline bismuth nanoparticles (BiNPs) in basic aqueous media. Spherical BiNPs with a rhombohedral crystalline structure are capped by phytochemicals and stably dispersible in aqueous media. The BiNPs effectively catalyze the reduction of 4-nitrophenol to 4-aminophenol by NaBH₄.

Lemon juice effectively served as a reducing and capping agent for an easy, cost-effective, and green synthesis of crystalline bismuth nanoparticles (BiNPs) in basic aqueous media.

Nanostructures of bismuth, the heaviest element among the ‘safe ones’ earning the status of a ‘green element’,¹ are particularly interesting due to their large magnetoresistance and excellent thermoelectric properties.^2–4 Bismuth nanoparticles (BiNPs) are the most extensively used nanostructures of bismuth. For example, BiNPs are utilized as contrast agents for computed tomography, photoacoustic imaging and infrared thermal imaging,^5,6 catalysts for the reduction of nitroaromatic compounds,^7,8 and removal of NO from air.⁹ Furthermore, BiNPs can act as intermediates for the synthesis of other nanostructures of bismuth, such as thermoelectric Bi₂Te₃ (ref. 10) and seeds for the solution–liquid–solid growth of nanowires.^11,12 Techniques for the synthesis of pure BiNPs could be categorized into (i) thermal decomposition,¹² (ii) mechanochemical processing,¹³ (iii) photochemical reduction,¹⁴ and (iv) solution-phase chemical reduction methods.^6,15–19 The thermal decomposition method requires harsh preparation conditions, expensive organometallic precursors, high temperature, and long reaction time, while producing high-quality monodispersed BiNPs. The mechanochemical processing technique is advantageous in terms of using inexpensive and nontoxic bulk bismuth pellets as precursors. However, the energy consumption and costly instrumentations are the limitations. The photochemical reactions typically require long time for sufficient conversion to bismuth nanoparticles, while they can also produce highly monodispersed BiNPs. The solution-phase chemical reduction methods are most popular due to the facile procedures and accessible reagents. However, stabilizers and toxic reducing agents often used are the limitations. To resolve the above-mentioned limitations, simple and green alternative methods are highly demanded.The use of abundant plant sources, which has been applied for synthesis of various metal nanoparticles such as Ag and Au, is a promising solution.^20,21 Plant sources contain a wide variety of biomolecules potentially serving as reducing and capping agents. Edible plant sources are obviously safest. To the best of our knowledge, there is no report on the synthesis of crystalline BiNPs using plant sources, while P. Poltronieri et al. reported synthesis of amorphous BiNPs using hydroalcoholic extract of Moringa oleifera.²² We focused on lemon, a very common fruit containing abundant antioxidants such as polyphenols, limonoids, citric acid, ascorbic acid, and vitamins potentially reduce ions with high oxidation states. Some of the phytochemicals, namely carbohydrates and proteins bearing ionic moieties, can be capping agents. Accordingly, lemon juice was successfully applied for the formation and in situ stabilization of silver and gold nanoparticles in aqueous media.^23,24 We presumed that a similar mechanism also works for BiNPs.In this communication, we introduce a greener strategy for the synthesis of crystalline BiNPs using lemon juice as a reducing and capping agent. For example, the synthesis was carried out using Bi(NO₃)₃·5H₂O (0.25 mmol) and freshly prepared lemon juice (25 mL) at 80 °C for 2 h under an aerobic basic condition. The X-ray diffraction (XRD) pattern (Fig. 1) of the obtained product was indexed to the pure rhombohedral phase of elemental bismuth (JCPDS no. 44-1246), indicating that the obtained product is BiNPs without detectable oxide phases. The average crystallite size was calculated to be 20 nm applying the Scherrer''s equation on the peaks of the (012), (104) and (110) plane. This result indicates that the phytochemicals present in lemon juice have ability to reduce bismuth salts to form BiNPs. The plausible phytochemicals for reduction are ascorbic acid, citric acid, and sugars. The reduction of Bi³⁺ with glucose⁸ and ascorbic acid¹⁷ was reported, and we accordingly performed the control experiment using possible reducing agents contained in lemon juice, namely ascorbic acid, glucose, and starch without any stabilizing agents. All of them could form elemental bismuth in highly aggregated forms as confirmed by the XRD patterns (ESI, Fig. S9^†) and SEM images (ESI, Fig. S10^†).Open in a separate windowFig. 1XRD pattern of obtained BiNPs synthesized using lemon juice with that of authentic Bi (JCPDS no. 44-1246).Electron microscopy was employed to confirm the size and the morphology of synthesized BiNPs. The scanning electron micrography (SEM) image (Fig. 2a) shows spherical objects having the size in the range of 50 to 100 nm agglomerated and covered with amorphous substances presumably originating from lemon juice. This agglomeration occurred during drying as confirmed by the stable water dispersibility of BiNPs with an average hydrodynamic diameter (D_h) of 255 nm investigated by dynamic light scattering (DLS) (ESI, Fig. S1a^†). The high colloidal dispersibility of the BiNPs and the coating layer observed in the SEM image suggest that phytochemicals present in lemon juice act also as capping agents. The larger D_h measured by DLS than the size observed by SEM can be attributed to the surrounding hydration layer and swelled phytochemicals attached to the surface of the BiNPs. We then performed transmission electron microscopy (TEM) analysis to get the actual size of the Bi cores. The TEM images of the BiNPs (Fig. 2b and c) indicated the presence of heavy elements in the matrix of light elements. The particles are spherical, and the diameter ranges from 8 to 30 nm. The high-resolution TEM image (Fig. 2d) shows the lattice fringes of 0.298 nm, 0.253 nm, and 0.224 nm of the typical crystallite agreeing well with the distance of the (012), (104) and (110) plane, respectively, of the rhombohedral Bi(0).Open in a separate windowFig. 2(a) SEM and (b–d) TEM images of BiNPs synthesized using lemon juice.The EDX spectrum (ESI, Fig. S2^†) reveals that obtained BiNPs comprise bismuth (Bi), carbon (C), and oxygen (O). The strongest signal of bismuth testifies the successful synthesis of BiNPs, whereas the C and O signals demonstrate the presence of the capping layer on the BiNPs.The FT-IR spectrum of the synthesized BiNPs was analyzed to presume the possible components on the BiNPs surface with the comparison with the solid content of lemon juice (Fig. 3). The FT-IR spectrum of the BiNPs shows the following major absorption bands. The broad absorption band at 3050–3500 cm⁻¹ is assignable to the O–H stretching of alcohol moieties. The absorption band around 2840–2940 cm⁻¹ is assignable to the C–H stretching of alkane, and the absence of the sharp absorption around 2950–3050 cm⁻¹ suggests the absence of trace contents of unsaturated C–H groups. The absence of the adsorptions originating from COO–H around 2400–3100 cm⁻¹ and C O around 1700–1730 cm⁻¹ indicates that citric acid^25,26 and other carboxylic acids contained sufficiently in lemon juice are not the major components on the BiNPs, while the presence of the strong absorption around 1600 cm⁻¹ is assignable to carboxylate moieties. The absorption bands around 930–1130 cm⁻¹ assignable to C–C and C–O stretching and around 1200–1420 cm⁻¹ assignable to O–C–H, C–C–H and C–O–H bending suggest the presence of sugars on BiNPs.^25,26 The content of the organic moieties was estimated approximately 14% by thermogravimetric analysis from weight loss that occurred at the temperature range from 130 to 450 °C (ESI, Fig. S3^†), at which the negligible C and O signals were observed in the EDX spectrum of the residue after TGA (ESI, Fig. S2^†).Open in a separate windowFig. 3FT-IR spectra of solid content of lemon juice and obtained BiNPs synthesized using lemon juice.The ¹H and ¹³C NMR and FTIR spectroscopic analysis (ESI, Fig. S4–S8^†) of the ethanol and chloroform extracts of the BiNPs suggests that the major capping phytochemicals are polysaccharides and fatty acid derivatives. Other minor possible components may include amino acids, terpenes and phenolic compounds contained in lemon juice.The successful formation of BiNPs made us enthusiastic to investigate their catalytic performance in the reduction of nitroaromatic compounds, problematic pollutants arising from explosives, analgesic, and antipyretic drugs and dyes.²⁷ Herein, the catalytic performance of our BiNPs was investigated by selecting the reduction of 4-nitrophenol (4-NP) to 4-aminophenol (4-AP) by NaBH₄ at room temperature as a model reaction.^7,8,28 The progress of the reduction was monitored by UV-vis absorption spectroscopy. The colour of the 4-NP aqueous solution changed from light yellow to deep yellow upon the addition of NaBH₄ to produce 4-nitrophenolate ion and an intense absorption peak at 403 nm appeared instead of the original absorption peak of 4-NP at 316 nm in the UV-vis spectrum.⁷ In the absence of BiNPs, the colour of the solution and the intensity of the peak were retained even after 12 h. In the presence of BiNPs, the deep yellow mixture became almost colourless within 180 min (Fig. 4a). The intensity of the absorption peak at 403 nm decreased overtime, and simultaneously a new absorption peak appeared and grew at 299 nm, indicating the formation of 4-AP (Fig. 4a).^8,27,29 The reaction rate almost follows pseudo-first-order kinetics agreeing with the Langmuir–Hinshelwood mechanism; in which two reactants react after adsorption on the solid surface, and then the product desorbs. The rate constant (k) was determined from the plot of ln(A/A₀) vs. time (Fig. 4b) according to the previous reports.^7,8,28,29 The k value of 0.0134 min⁻¹ (activity factor = 0.1 s⁻¹ g⁻¹) is lower than previously reported PVP-coated bismuth nanodots (6.033 s⁻¹ g⁻¹)⁸ and starch coated BiNPs (0.02751 s⁻¹).⁷ The lower catalytic rate of our BiNPs can be attributed to the electrostatic repulsion of 4-nitrophenolate ions with negatively charged BiNPs (zeta potential value = −31.7 mV) in the reaction mixture. The initial rate is slower than the rate after 50 min. Possible factors are (i) the inductive period for the surface activation of the nanoparticles at the initial stage as reported for bismuth nanodots by Liang et al.⁸ and (ii) partial and gradual release of phytochemicals under the basic conditions. After the catalytic reaction, the D_h value measured by DLS (195 nm) became smaller than the original one maintaining the good water dispersibility, suggesting unravelling of primary particles fused by the capping layer (ESI, Fig. S1b^†). In addition, the weight loss in the TGA curve became lower (ESI, Fig. S3^†), and the SEM image implies the partial loss of the coating (ESI, Fig. S11^†). This decrease of the organic moieties is attributable to the partial removal of coating substances such as fatty acids during the catalytic reaction under the basic and reductive conditions, supported by the disappearance of the absorption band around 1600 cm⁻¹ assignable to carboxylate moieties in the FTIR spectrum of BiNPs after the catalytic reaction (ESI, Fig. S12^†). Carboxylic acids and their esters are reported to be reduced to alcohol by NaBH₄ in the presence of electrophiles.^30–33 The initial presence of the amphiphilic and anionic layer plausibly delays the catalytic reaction.Open in a separate windowFig. 4(a) Optical images and absorption spectra of catalytic reduction of 4-NP ([4-NP] = 15 ppm) by NaBH₄ (4.28 × 10⁻⁴ M) in presence of BiNPs (142 mg L⁻¹); (b) pseudo-first order kinetic plot of catalytic reduction.In conclusion, we have demonstrated a green, cost-effective, and successful approach for synthesis of crystalline BiNPs using lemon juice as a reducing as well as capping agent. The obtained BiNPs effectively catalysed the reduction of 4-NP to 4-AP by NaBH₄. The importance of this method lies in the simple synthetic procedure, uses of safe and low-cost lemon, and good dispersibility over conventional chemical approaches.     

Microwave-assisted synthesis of octahedral Rh nanocrystals and their performance for electrocatalytic oxidation of formic acid

Uniform and well-defined octahedral Rh nanocrystals were rapidly synthesized in a domestic microwave oven for only 140 s of irradiation by reducing Rh(acac)₃ with tetraethylene glycol (TEG) as both a solvent and a reducing agent in the presence of an appropriate amount of KI, didecyl dimethyl ammonium chloride (DDAC), ethylene diamine (EDA) and polyvinylpyrrolidone (PVP). KI, DDAC and EDA were essential for the creation of octahedral Rh nanocrystals. Electrochemical measurements showed a significantly enhanced electrocatalytic activity and stability for the as-prepared octahedral Rh nanocrystals compared with commercial Rh black.

Octahedral Rh nanocrystals were rapidly synthesized in a domestic microwave oven for only 140 s of irradiation by reducing Rh(acac)₃ with tetraethylene glycol as both a solvent and a reducing agent.

To date, platinum group metals play an indispensable role as efficient catalysts for some important reactions in industry. However, due to their limited reserves and high prices, a large number of platinum group metal nanoparticles with different particle sizes, morphologies and surface structures have been synthesized by means of various methods to reduce their cost.¹ As a platinum group metal, Rh has good catalytic activity and stability, and is often used as a typical catalyst for some chemical reactions such as hydrogenation,^2–7 nitrogen oxide reduction,⁸ CO oxidation,^9–11 cross coupling,^12–14 hydroformylation,^15–19 in fuel cells^20,21 and other chemical reactions.²² Therefore, controlled syntheses of Rh nanoparticles with different morphologies have attracted much attention. In recent years, people have successfully prepared Rh nanostructures with various morphologies such as sheet,^23–27 flower,⁶ polyhedron,^28–33 porous ball,⁸ multi branches,^34–39 stars,⁴⁰ nanoframes^13,14,41 and nano nail.⁴² These Rh nanoparticles with unique structures effectively improve the atom utilization as well as their catalytic reaction performances. However, similar to other platinum group metals, the difficulty of large-scale preparation of Rh nanomaterials with single morphology and uniform size still greatly restricts their industrial application.Microwave irradiation has been widely used in chemical synthesis because of its simple, rapid and efficient characteristics as well as special heating mode from the inner. We have synthesized many metallic nanoparticles with different shapes by using microwave irradiation for about 80 to 120 seconds. Herein, we report a simple and fast strategy for the synthesis of octahedral Rh nanocrystals under microwave irradiation with using domestic microwave oven. In a typical synthesis, octahedral Rh nanocrystals with uniform and well-defined morphologies were successfully synthesized with Rh(acac)₃ as the precursor, polyvinyl pyrrolidone (PVP) as the stabilizer, triethylene glycol (TEG) as both a solvent and a reducing agent in the presence of didecyl dimethyl ammonium chloride (DDAC), KI and ethylene diamine (EDA) under microwave irradiation in a very short time. Meanwhile, the electrocatalytical performance of the as-prepared octahedral Rh nanocrystals for the electro-oxidation of formic acid was also investigated with commercial Rh black as a contrast.The TEM and SEM images of the representative Rh nanoparticles obtained under the optimal experimental conditions are shown in Fig. 1, S1 and S2.^† Wherein, the prepared Rh nanoparticles demonstrated uniform and well-defined octahedral structure with sharp edges and corners as well as smooth surfaces (Fig. 1a and b), in which the average side length is about 65 nm. The high-resolution TEM (HRTEM) image (Fig. 1c) shows well-resolved continuous fringes clearly. The corresponding fast Fourier transform (FFT) pattern, as the inset shown in Fig. 1c, shows a lattice distance of 0.194 or 0.216 nm, which can be attributed to the {200} and {111} lattice planes of the octahedral Rh with face-centered cubic structure, respectively, confirming its single-crystal nature. Furthermore, the regular octahedral feature of the as-prepared Rh nanoparticles can be well distinguished from SEM images, as shown in Fig. 1d and S2.^† These results show that the octahedral Rh nanocrystals with a single morphology can be rapidly synthesized in a great quantity by irradiation with domestic microwave oven for only 140 s.Open in a separate windowFig. 1TEM and SEM images of the as-prepared octahedral Rh nanocrystals. (a) and (b) Typical TEM images with different scales. The inset in (b) is the schematic illustration; (c) typical HRTEM image. The inset is the corresponding FFT pattern; (d) SEM image. Fig. 2a shows the XRD pattern of the as-prepared typical octahedral Rh nanocrystals. As can be seen, the diffraction peaks at 2θ values of 41.26, 47.95, 70.18 and 84.33° are observed, which can be well indexed to the diffractions of (111), (200), (220) and (311) lattice facets of metallic Rh referring to the standard powder diffraction card (JCPDS card No. 05-0685), respectively. This observation further confirmed their fcc Rh structure. In addition, the narrow and sharp (111) diffraction peak implied that the typical octahedral Rh nanocrystals exhibited a high purity and crystallinity. The XPS spectrum was taken for the as-prepared octahedral Rh nanocrystals and the result was displayed in Fig. 2b, As it can be seen, two peaks corresponding to the electron binding energies of Rh 3d_3/2 and Rh 3d_5/2 were observed at 311.85 eV and 307.10 eV with an interval of 4.75 eV, respectively, which were consistent with the literature values (311.75 and 307.0 eV),⁴³ revealing Rh(0) metallic state of the octahedral nanocrystals.Open in a separate windowFig. 2XRD pattern (a) and XPS spectrogram (b) of octahedral Rh nanocrystals.The dependence of the morphological evolution of Rh nanocrystals upon irradiation time was investigated. When irradiated for 120 s, the octahedral structural Rh nanocrystals with about 65 nm of the side length produced except for unclear edges and corners as well as a shorter side length, as shown in Fig. 3a. As microwave irradiation progressed to 140 s, uniform and well-defined octahedral Rh nanocrystals with smooth surfaces generated (Fig. 3b). While the irradiation time was extended to 160 s, however, the vertices of some octahedral structures were truncated although with no change of the sizes, as shown in Fig. 3c. As the irradiation time reached 180 s, the octahedral structural feature of most particles disappeared with a further truncation of their vertices (Fig. 3d), which should be ascribed to higher surface free energies for the metallic atoms at the apexes and edges as well as a higher internal temperature due to a longer irradiation time. These results indicated that the optimum microwave irradiation time was 140 s for the creation of regular octahedral Rh nanocrystals.Open in a separate windowFig. 3TEM images of Rh nanoparticles prepared at different reaction time. (a) 120 s; (b) 140 s; (c) 160 s; (d) 180 s.It was noteworthy that KI played a crucial role in controlling synthesis of octahedral Rh nanocrystals. When no KI was used, it would produce irregular Rh nanoparticles, as shown in Fig. 4a. While with addition of 0.6 mmol of KI, octahedral Rh nanostructures with blunt vertices and an average side length of about 50 nm were generated (Fig. 4b), implying an incomplete growth relative to the case of 0.8 mmol of KI as in the typical experimental process (Fig. 1). Nevertheless, the amount of KI was increased to 1.2 mmol, only less octahedral structure features could be observed except for few obscure polyhedral outlines (Fig. 4c). These results indicated that the existence of KI was advantageous to the generation of octahedral Rh nanocrystals. Generally, the eight triangular surfaces of metallic Rh octahedron consists of (111) lattice planes. According to the previous report,^44–49 it can be considered that the preferential adsorption of I⁻ anions on Rh (111) planes is one of the main factors driving the formation of octahedral structure. As a result, a growth along 〈111〉 directions was confined and a growth along 〈100〉 directions was facilitated, which created octahedral structures due to anisotropic growth. However, excessive I⁻ ions would adsorb non-selectively on the surfaces of Rh nanoparticles, which resulted in passivation of the edges and corners of polyhedron. In addition, an equivalent amount of KBr or KCl was used instead of KI, respectively, to clarify the role of I⁻ ions under the same other conditions. As can be seen (Fig. S3, ESI^†), no octahedral Rh nanocrystal except for agglomerated irregular nanosheets was observed in these two contrast experiments. This may be ascribed to the change of the precursor. In the presence of a large number of I⁻ ions, the precursor can be transformed to a more stable [RhI₆]³⁻ complex.^44–47 As a result, the reducing rate of Rh(iii) to Rh atom decreased, which may be favourable for the nucleation of Rh nanoparticles and the oriented growth of Rh octahedra.Open in a separate windowFig. 4TEM images of Rh nanoparticles prepared with different amounts of DDAC or KI under the same other conditions. (a) Absence of KI; (b) 0.6 mmol of KI; (c) 1.2 mmol of KI; (d) absence of DDAC; (e) 0.2 mmol of DDAC; (f) 0.6 mmol of DDAC.Meanwhile, the influence of DDAC on the generation of octahedral Rh nanocrystals was also studied under the same other conditions. In the absence of DDAC, only agglomerated irregular Rh nanoparticles were observed (Fig. 4d). When 0.2 mmol of DDAC was added, octahedral Rh nanostructures with an average side length of about 45 nm, a smaller size relative to the case of 0.4 mmol of DDAC as in the typical experiments (Fig. 1), were generated accompanying with a few irregular nanoparticles (Fig. 4e). With increasing the amount of DDAC to 0.6 mmol, agglomerated irregular polyhedral nanostructures formed (Fig. 4f). Thus, the addition of DDAC was also indispensable for the growth of octahedral Rh nanostructures under microwave irradiation. Whereas an excessive amount of DDAC was also unfavourable for creation of the octahedral Rh nanocrystals. Moreover, no octahedral nanostructures generated except for urchin-like Rh hierarchical superstructures when adding an equivalent amount of cetyltrimethylammonium chloride (CTAC) instead of DDAC (Fig. S4a, ESI^†). While didoctyl dimethyl ammonium bromide (DDAB) was used instead of DDAC, the formation of octahedral Rh structures can be still observed although accompanying with other irregular polyhedral (Fig. S4b, ESI^†). These results suggested that the formation of octahedral Rh nanostructures were strongly dependent upon the hydrophobic chains of DDAC or DDAB but nothing to do with Cl⁻ or Br⁻ anions. The effect of other halide ions can be ignored due to the existence of a large number of I⁻ ions. That is because the strength of adsorption of I⁻ ions on metal surfaces is generally stronger than that of Cl⁻ or Br⁻ ions.⁴⁸Accordingly, the generation of octahedral Rh nanocrystals should be ascribed to the synergistic effect of KI and DDAC under the above experimental conditions. We believe that DDAC could enhance the role of I⁻ ions in generating (111) facets of octahedral by adjusting the adsorption selectivity of I⁻ ions on (111), (100) or (110) facets. On the one hand, the amount of KI would manipulate the reducing kinetics to form octahedral Rh nanostructures under microwave irradiation. A slow reducing rate was favourable for the oriented growth of Rh octahedra due to the formation of a more stable coordinated anion [RhI₆]³⁻. On the other hand, the confinement of DDAC induced the selective adsorption of I⁻ ions on Rh {111} facets which restrained the growth along 〈111〉 directions of Rh nuclei and prompted the growth along 〈100〉 directions. In addition, a proper quantity of DDAC confined the deposition of Rh atoms on {111} facets, which may be beneficial to the growth along 〈100〉 directions. However, an excessive amount of DDAC was unfavourable for the formation of shaped Rh nanoparticles since they disturbed the adsorption of I⁻ anions on Rh {111} facets.Furthermore, it was also found that ethylene diamine (EDA) demonstrated an important effect on the creation of octahedral Rh nanostructures. Under keeping the total volume of the reaction system unchanged, the significantly agglomerated irregular polyhedral nanoparticles with sharp horns were observed in absence of EDA (Fig. S5a^†). When 0.5 mL of EDA was added, a few octahedral nanostructures began to generate though accompanying with agglomerated irregular polyhedra (Fig. S5b^†). While the amount of EDA was increased to 1 mL, uniform and well-defined octahedral Rh nanocrystals with flat and smooth surfaces were produced (Fig. S5c^†). However, a more amount of EDA was added, a part of octahedral nanostructures become deformation as well as agglomeration (Fig. S5d^†). In the reaction system, TEG as a solvent was also served as a reducing agent. As can be seen, even though without adding EDA, the rhodium salt was still reduced completely to produce metal Rh nanoparticles. With the addition of EDA, octahedral Rh nanocrystals began to generate, while an excessive amount of EDA resulted in unclear edges and corners of the octahedral structures. Obviously, EDA demonstrated significant effect on the morphology control of octahedral Rh nanocrystals. It should be ascribed to the coordination adsorption of EDA on the surface of metal particles.⁵⁰ Furthermore, no octahedral nanostructures but irregular nanoparticles or Rh dendrites generated with using an equivalent amount of n-butylamine or n-octylamine instead of EDA (Fig. S6a and b^†). Therefore, we suggest that EDA plays a synergistic role together with DDAC in regulating the rate of atomic packing and nanoparticle growth by coordination adsorption. The growth rate of nanoparticles is faster in absence of EDA, while the growth rate slows down with the increase of EDA dosage. An appropriate amount of EDA facilitates the generation of uniform octahedral Rh nanocrystals by adjusting the balance between nucleation rate and growth rate. Nevertheless, excessive EDA makes a slower growth than nucleation due to their extreme adsorption, resulting in obscure appearances of some octahedral Rh nanoparticles.In addition, PVP was also found to be important but not essential for the formation of octahedral Rh nanocrystals. Either without or with a few amount of PVP, octahedral Rh nanocrystals can also produce except for a little agglomeration (Fig. S7a and b^†). An appropriate amount of PVP contributed to uniform and well dispersed octahedral Rh nanocrystals, while excessive PVP caused aggregation (Fig. S7c and d^†). These results indicated that PVP served mainly as a protecting and dispersing agent for the nanocrystals.The catalytic performance of the synthesized octahedral Rh nanocrystals was tested by cyclic voltammetry (CV) and chronoamperometry (CA) with the formic acid electrooxidation reaction as the model reaction system. Fig. 5a exhibits the representative CV curves obtained for the electrochemical oxidation of 0.5 mol L⁻¹ HCOOH over the octahedral Rh nanocrystals and commercial Rh black in 0.5 mol L⁻¹ HClO₄ solution, respectively. CV measurements showed the peak current density for the octahedral Rh nanocrystals was 3.53 mA cm⁻² at 0.544 V, while it was 1.01 mA cm⁻² at 0.609 V for Rh black. The formic acid electrooxidation indicated that the electrocatalytic activity of octahedral Rh nanocrystals was about 3.5 times that of Rh black, demonstrating an obvious morphological dependence for their electrochemical property. The corresponding CA curves of formic acid electro-oxidation at 0.55 V is shown in Fig. 5b. As can be seen, a higher current retention through the whole measuring range were observed over the as-prepared octahedral Rh nanocrystals than Rh black though both of them showed an equivalent attenuation rate in the initial 20 seconds. The CV curve of continuous cycle scanning for octahedral Rh nanocrystals in 0.5 mol L⁻¹ HClO₄ solution showed a decrease of the electrochemical activity only by 9.6% after 2000 cycles (Fig. S8^†). These results reveal that octahedral Rh nanocrystals exhibit a remarkably enhanced electrochemical activity and stability compared with Rh black. Their enhanced catalytic activity should be attributed to the uniform geometric structure with single surface lattice.Open in a separate windowFig. 5The CV (a) and CA (b) curves for the electrochemical oxidation of 0.5 mol L⁻¹ HCOOH over the octahedral Rh nanocrystals and Rh black in 0.5 mol L⁻¹ HClO₄ solution, respectively.Additionally, CO stripping voltammetry measurements were performed. As shown in Fig. S9a,^† no CO electro-oxidation (CO_ox) was observed for the freshly-prepared octahedral Rh nanocrystals in 0.5 M HClO₄ solution. Subsequently, a current peak for CO_ox appeared at 0.550 V (versus SCE) after adorbing CO for the clean octahedral Rh-modified electrode, as shown in Fig. S9b.^† Then CO_ox peak disappeared in the following second potential scanning, as shown in Fig. S9c.^† These results showed that CO adsorbed on Rh surfaces can be easily removed in the process of electrocatalytic oxidation, showing well CO resistence.In summary, uniform octahedral Rh nanocrystals could be rapidly prepared with domestic microwave oven in only 140 s of irradiation by reducing Rh(acac)₃ with TEG as both a solvent and a reducing agent, PVP as a protecting and dispersing agent in the presence of proper quantities of DDAC, KI and EDA. The formation of octahedral Rh nanocrystals was attributed to the synergism of KI, DDAC and EDA. The electrochemical oxidation of formic acid demonstrated higher electrocatalytic activity and stability for the as-prepared octahedral Rh nanocrystals than Rh black, displaying a significant dependence upon their morphologies.     

Resistive non-volatile memories fabricated with poly(vinylidene fluoride)/poly(thiophene) blend nanosheets

Ferroelectric poly(vinylidene fluoride)/semiconductive polythiophene (P3CPenT) blend monolayers were developed at varying blend ratios using the Langmuir–Blodgett technique. The multilayered blend nanosheets show much improved surface roughness that is more applicable for electronics applications than spin-cast films. Because of the precisely controllable bottom-up construction, semiconductive P3CPenT were well dispersed into the ferroelectric PVDF matrix. Moreover, the ferroelectric matrix contains almost 100% β crystals: a polar crystal phase responsible for the ferroelectricity of PVDF. Both the good dispersion of semiconductive P3CPenT and the outstanding ferroelectricity of the PVDF matrix in the blend nanosheets guaranteed the success of ferroelectric organic non-volatile memories based on ferroelectricity-manipulated resistive switching with a fresh high ON/OFF ratio and long endurance to 30 days.

Ferroelectric poly(vinylidene fluoride)/semiconductive polythiophene blend nanosheets show good resistive non-volatile memory performance with a fresh high ON/OFF ratio and long endurance to 30 days.

Ferroelectric capacitors memorize information by the two antiparallel polarization states coded as “0” or “1”, even after removing the applied power, serving as non-volatile memory.^1,2 However, the reading out of the stored information in such memories will change the polarization state, which is known as destructive read-out. To realize non-destructive read-out and low-energy writing in such devices, conjugated polymers were introduced into a ferroelectric polymer matrix to form phase-separated blends, which contributed to novel ferroelectric resistive switches with very simple two-terminal sandwiched devices superior to the complicated transistor type electronic elements.^3,4 In the ferroelectric nonvolatile memories, the bi-stable polarization states in ferroelectric polymers such as poly(vinylidene fluoride) (PVDF) derivatives were used to manipulate the charge transfer and injection in conjugated polymers (CPs).^3–7 This manipulation results in different resistance at on-state and off-state of the devices controlled by the ferroelectric polarization directions. Consequently, non-destructive low-voltage operation was realized in which the ferroelectric PVDF provides switching functionality. The read-out operation is performed by conduction through a nearby semiconducting layer.⁵ Thereafter, the operational mechanism of ferroelectric-driven organic resistive switches was clarified by Kemerink et al.⁸ They stressed that the stray field of the polarized ferroelectric phase can modulate the charge injection from a metallic electrode into the organic semiconductor, switching the diode from injection-limited to space-charge-limited.Although the devices show promising applications, the random phase-separation morphology gives them uncertain properties.^9,10 The most important obstacle is to prepare well-dispersive PVDF–CP phase-separated blend nanofilms because a well-dispersive hybrid nanofilm is indispensable for higher information storage density. In other words, embedding semiconductor polymers into ferroelectric polymer matrixes with domain size of CPs smaller than a few tens of nanometers (e.g. 50 nm) has remained extremely challenging to date.^6,9 Phase-separated films prepared using traditional methods such as spin-coating show a much disordered surface with root-mean-square (RMS) roughness exceeding 39.97 nm.¹¹ Such a highly rough surface tends to induce large leakage current. Especially, the annealing treatment for PVDF materials to increase crystallinity and ferroelectric phase crystals also brings a significant increment to the phase-separated domain size. From another perspective, to decrease the writing and read-out operational voltages, ultrathin blend films (thinner than 100 nm) with high-content ferroelectric crystals are also necessary. However, the reported methods to prepare polymer ferroelectric ultrathin films, e.g. spin-coating are usually performed at the cost of decreasing the ferroelectric phase or requiring high energy consumption.Previous reports have proposed a versatile method to achieve highly oriented ferroelectric polymer Langmuir–Blodgett (LB) nanofilms consisting of ferroelectric poly(vinylidene fluoride) (PVDF) and tiny amphiphilic poly(N-dodecylacrylamide) (pDDA) supplying film stability at the air–water interface and transfer properties (Fig. 1a).^12,13 The LB nanofilms, with a controllable film thickness at several nanometers'' scale and a smooth film surface, can be transferred onto solid and flexible substrates irrespective of their surface wettability. The contents of ferroelectric β crystals in the as-prepared nanofilms with no post-treatment can be up to 95%, which is the best content in such thin films ever reported.^2,14 Therefore, we investigated a strongly enhanced ferroelectricity of the PVDF LB nanofilms with no electrical poling showing one of the highest ferroelectric remanent polarization values, 6.6 μC cm⁻², for nanometer-scale films, in addition to long endurance because of the high orientation of β crystals.² The PVDF–pDDA LB nanofilms benefit both from the good film-formation properties of amphiphilic pDDA and the functionality of ferroelectric PVDF. To extend their applications, the PVDF–pDDA LB nanofilms are promising for application as a platform and matrix to combine with other functional materials such as organic functional materials and inorganic nanoparticles. Similarly, it is promising for production of well-dispersive ferroelectric/semiconductive polymer blends by addition of a third composition: semiconductive polymers at the air–water interface. Moreover, we have reported that semiconductive poly(3-hexylthiophene) (P3HT) monolayers were stabilized by mixing with the amphiphilic poly(N-dodecylacrylamide) (pDDA) at the air–water interface for field-effect transistors.¹⁵Open in a separate windowFig. 1(a) Chemical structures of poly(vinylidene fluoride) (PVDF), poly[3-(5-carboxypentyl) thiophene-2,5-diyl] (P3CPenT) and poly(N-dodecylacrylamide) (pDDA), (b) surface pressure (π)–area (A) isotherms, and (c) compressibility modulus (C_s⁻¹)–surface area (A) isotherms of PVDF–P3CPenT–pDDA Langmuir films for various weight contents of P3CPenT relative to PVDF. The straight arrow in (b) indicates the line extrapolated to determine the limiting surface area.This study examines the facile preparation of blend nanosheets of ferroelectric PVDF and semiconductive polymer, by incorporating poly[3-(5-carboxypentyl) thiophene-2,5-diyl] (P3CPenT) into PVDF–pDDA Langmuir–Blodgett (LB) nanofilms. P3CPenT is a conjugated polythiophene carboxylate (Fig. 1a) that occupies superior charge transport performance resulting from a delicate balance between its solubility and solid-state film morphologies.^16,17 The blend LB nanosheets can be transferred readily and regularly from the air–water interface onto various substrates. Phase-separation morphologies were studied based on the varying P3CPenT contents. Resistive switching properties were demonstrated in a simple sandwiched device of Al/PVDF–P3CPenT–pDDA/Au at the opposite polarization direction of the PVDF matrix.The PVDF–P3CPenT–pDDA monolayers were prepared by spreading PVDF–P3CPenT mix solution and pDDA solution onto water surface in sequence at a molar ratio of 50 : 1 (see ESI^† for details). Fig. 1b shows surface pressure (π)–area (A) isotherms of PVDF–P3CPenT–pDDA Langmuir films at varying P3CPenT contents (relative to PVDF) at 20 °C: they are designated as P3CPenT 6 wt%, P3CPenT 9 wt% and P3CPenT 23 wt%. The sharply rising curves with high collapse surface pressures up to 50 mN m⁻¹ are indicative of the formation of stable and compact Langmuir blend monolayers at the air–water interface. These pressure values are much higher than those of pure PVDF (25 mN m⁻¹) or P3CPenT (15 mN m⁻¹) (Fig. S1^†). Results suggest that amphiphilic pDDA can greatly improve film stability at the air–water interface because of excellent hydrogen bonding interactions that occur among the amide groups in pDDA backbones.¹³ The compressibility modulus (or elasticity) C_s⁻¹ was calculated from isotherm data based on eqn (1) to obtain the monolayer elastic behavior.¹⁸C_s⁻¹ = −A(dπ/dA)1In the C_s⁻¹–A curves shown in Fig. 1c, each compressed Langmuir film shows a peak value of C_s⁻¹*, manifesting the formation of a critical concentrated solid state.^19–22 Table S1^† shows the limiting surface area (S_aver), collapse surface pressure (π_c), peak compressibility modulus (C_s⁻¹*), corresponding surface pressure (π*), and corresponding occupied area (A*) for the respective P3CPenT contents. The change in the limiting surface area from 0.0224 (6 wt%) to 0.0250 nm² (23 wt%) shows the successful introduction of P3CPenT into the PVDF matrix.After confirming the film formation properties, the monolayers were transferred onto hydrophobic silicon substrates regularly (Fig. S2^†). The transfer ratios for both downstrokes and upstrokes are almost unity, based on eqn (S1).^† Low surface roughness is a crucially important parameter for electronic applications. In earlier reports, lowering the root-mean-square (RMS) roughness below 30 nm is quite a challenge for the ferroelectric/semiconductive polymer blend films because of the crystalline properties of the two polymers. The AFM images (Fig. 2a–d) of the as-prepared PVDF–P3CPenT blend LB nanosheets manifest much smoother surfaces, even for the 80-layer nanosheets by an RMS value of 10.8 nm. This value is smaller than that of the reported spin-coating films.^6,11 In Fig. 2c, the nanofiber structures in the P3CPenT 23 wt% monolayer orientate at one direction as the black arrow indicates. All the nanofibers show a regular size of 7.9 nm (Fig. 2e). In contrast, such nanofiber structures do not appear in the monolayers at low contents of P3CPenT in Fig. 2a and b, which proves that the nanofiber structures originate from P3CPenT molecules. In Fig. 2d and S3,^† the interlaced nanofiber networks in the 80-layer LB blend nanosheet caused by repeated deposition of Langmuir monolayers will provide the nodal points for electric conductivity. EDX mapping images show homogeneous dispersion of P3CPenT in PVDF, which also ensured the continuous conducting path (Fig. S4^†), while the nano-domains were not clear seen. It may originate from the similar solubility of PVDF and P3CPenT in NMP, which made the unclear phase-separation structure.Open in a separate windowFig. 2AFM images of (a)–(c) PVDF–P3CPenT–pDDA blend LB monolayer at 6, 9, and 23 wt% P3CPenT, (d) an 80-layer blend LB nanosheet (P3CPenT 23 wt%), and (e) height profile of the fiber structures in (c).The blend monolayers were readily transferred onto various substrates such as rigid silicon and flexible PET substrates (Fig. 3). The uniform interference colours of the macroscopic films at each layer number are indicative of the regular layer structures and homogeneous film transfer from water surface to the substrates. The result also evidenced the controllable preparation of functional polymer blends at several nanometres'' scale for the present work, which cannot be achieved by general film preparation methods such as drop casting and spin-coating. UV-vis spectra in Fig. 4a show a broad absorbance peak near 462 nm for P3CPenT and PVDF–P3CPenT solutions in NMP. The detected absorbance peaks at 442, 428, and 415 nm at 23 wt%, 9 wt%, and 6 wt%, respectively, evidenced the successful introduction of P3CPenT into the PVDF LB nanofilms (Fig. 4b). The tunable domain size and controllable film thickness are not the only virtues of the blend nanosheets. In Fig. 5a, the FT-IR spectra of the blend nanosheets show significant peaks at 508, 840, 1276, and 1400 cm⁻¹ at each content, which are assigned for the ferroelectric β crystals (Form I).^23–25 Characteristic peaks of paraelectric α crystals (763 and 976 cm⁻¹) were not detected. The rich β phase formation was also confirmed using XRD patterns, which showed the appearance of 110/200 reflection at 20.1° (Fig. S5^†).^12,25 The PVDF–pDDA LB films were crystalline films with a crystallinity of 50%.¹² Results confirmed that the blend nanosheets contain almost 100% ferroelectric β crystal, which is consistent with the reported value for PVDF–pDDA LB nanofilms,¹² indicating that the addition of P3CPenT did not affect the crystal structures of PVDF. The result ensures that the ferroelectric matrix in the blend nanosheets can supply a stable and large polarization electric field to modulate charge transfer and accumulation. Monolayer thickness was confirmed as 2.3, 2.6, and 2.6 nm respective for 6 wt%, 9 wt%, and 23 wt% P3CPenT containing LB nanosheets by XRD spectra (Fig. S6 and Table S2^†). All the monolayers can be uniformly and repeatedly transferred to substrates to adjust the film thickness. Therefore, the film thickness of the nanosheets can be finely modulated at several nanometres'' scale.² This report is the first of the relevant literature to describe success on the preparation of blend LB nanosheets with highly ferroelectric matrix and well-dispersive semiconductive P3CPenT.Open in a separate windowFig. 3Photographs of blend LB nanosheets with 23 wt% P3CPenT at various layers on Si substrate (left) and a 30-layer blend nanosheet on flexible PET membrane (right).Open in a separate windowFig. 4UV-vis absorption spectra of (a) P3CPenT and PVDF–P3CPenT solutions in NMP and (b) the blend LB nanosheets for various weight contents of P3CPenT relative to PVDF.Open in a separate windowFig. 5FT-IR spectra of 80-layer PVDF–P3CPenT–pDDA blend nanosheets for various P3CPenT contents.Memory devices were constructed with a simple sandwiched structure: glass substrate/gold (Au)/blend LB nanofilms/aluminum (Al) (see the inset of Fig. S7 and ESI^†).^26,27 Fig. S7a^† shows typical current (I)–voltage (V) characteristics of 30-layer blend nanosheets (78 nm), which demonstrated a reversal current hysteresis with voltage sweeping between ±20 V.^9,11 The cross-sectional SEM image of the memory device is shown in Fig. S7b,^† indicating a well-defined layer structure. The current can only flow through the P3CPenT domains because of the insulation characteristic of PVDF matrix, however, no hysteresis loop was obtained for the blend LB nanosheets. This is due to the conducting path formed by the interlaced P3CPenT nanofibers even with the presence of rich ferroelectric β phase. So our device will not follow Asadi''s mechanism as introduced in the Introduction part.Local surface potential patterns were visualized using Kelvin probe microscopy to investigate the polarization state of the devices.¹² A two-step measurement was conducted (Fig. 6a). First, a positive/negative bias voltage was applied to the conductive AFM cantilever under ambient conditions, making the sample surface polarized with a square shape. Then, the KFM mode was conducted to obtain images of the polarized patterns, which is called the reading procedure. Fig. 6b and c portray the surface potential images before and after polarization. Clearly, the inverse profile of the positive and negative polarized area is seen in Fig. 6b, which indicates the LB films are ferroelectric.Open in a separate windowFig. 6(a) Schematic illustration of Kelvin probe force microscopy measurement and surface potential images (b) before polarization and (c) after polarization.To investigate whether the device is resistive switching or not, we applied a polarizing pulse for information writing with pulse width of 2.5 s and voltage of 30 V to induce polarization states of the ferroelectric matrix in a 40-layer blend nanosheet (104 nm) as shown in Fig. 7a. The pristine state refers to as-fabricated devices that have not undergone any electrical poling. The writing voltage magnitude for both ON and OFF states is set as 30 V for 2.5 s higher than its coercive voltage: 20 V (195 MV m⁻¹ × 104 nm)² of the PVDF LB nanosheet, as we previously reported, to ensure the high polarization value in the ferroelectric matrix. The read-out operation was operated at quite low voltage by detecting the current flow in the devices. The blend nanosheet after polarizing at +30 V exerts ten-fold improved current density (ON-state) in comparison with the pristine nanosheet before polarization (Fig. 7b). In contrast, the current density decreased after polarization at −30 V (OFF-state). The results verified the resistive switching characteristics and the non-volatility nature of the memory devices fabricated with the blend nanosheets. For comparison, a control experiment for pure pDDA LB nanofilms showed no ferroelectric polarization-reversal-induced resistive switching (Fig. S8^†). It is noteworthy that by combining the ferroelectric PVDF nanosheet with semiconductive polymer, the read-out voltage is optimized even to lower than 1 V. Results show that it is highly competitive over the ferroelectric capacitors without semiconductive components, which require a reading-out voltage higher than coercive voltage, about 20 V for the 100 nm-thick films.² The ON/OFF ratio values were recorded at varying reading voltages (Fig. 7c). The maximum value of the ON/OFF ratio was approximately 891 with increasing read-out voltage to 15 V, which hit one of the highest records over the previous ferroelectric polymer – thiophene-based semiconductor polymer blend memories.^3,5,11,28 This result derives from the synergetic effects of uniform dispersion of semiconductive domains and improved ferroelectricity of the PVDF matrix. The ultrathin PVDF matrix can survive at a strong electric field up to 500 MV m⁻¹ as our previous work reported, which ensured a large operation window for reading out the stored information as well as high ON/OFF ratio values. In other words, the bottom-up construction method for functional polymer blends affords us a novel platform for materials design and micro/nanostructures manipulation. At the first stage of the work design, we expected a band bending mechanism to explain the resistive switching. However, the I–V curve in Fig. S7a^† indicated Asadi''s mechanism is not applicable for our devices. The resistive switching toward ferroelectric polarization shows similar behaviour to the reported devices.²⁸ They used a symmetric device of Au–PVDF/P3HT–Au, however, the layer structure was asymmetric because they spin-coated P3HT on the patterned P(VDF–TrFE) layer. It seems that P3HT has a continuous phase through the layer, which is different from the clearly phase-separated structures.^3,4 As mentioned above, P3CPenT took a conductive path in the nanosheet sandwiched with two different electrodes, leading to no band bending formation at the interfaces. A lot of work is necessary in the future to optimize the device structure and materials combination to understand the operation mechanism in terms of structure–property relationship, e.g. using other semiconductive polymers with different solubility from PVDF. We also investigated the device cycle endurance (fatigue property) and retention properties for the blend-nanosheet-based ferroelectric nonvolatile memories (Fig. 7d). The devices show reversible switching over many cycles up to 1000 cycles in spite of a little linear decrease of ON-current and increase of OFF-current as a function of cycle numbers, which is consistent with the reported values.²⁹ The retention of current density was demonstrated with consistent current density at the ON state up to 4 days (Fig. 7d, bottom) and the OFF state up to 30 days (Fig. S9^†). A slight degradation of current density at the ON state was observed after 4 days, which might be ascribed to the interfacial charge induced depolarization of ferroelectric domains.³⁰Open in a separate windowFig. 7(a) Schematic illustration of the operation program for the ferroelectric nonvolatile memories, (b) current density (J)–voltage (V) characteristics of the blend LB nanosheets (P3CPenT 23 wt%, 40 layers) before and after polarization, (c) ON/OFF ratios at varying read voltage, and (d) (top) the device cycle endurance at ON- and OFF-states reading at 1 V and (bottom) the retention property at ON-state reading at 3 V. Dashed lines represent the linear fits at different states.In conclusion, a bottom-up construction of ferroelectric PVDF/semiconductive P3CPenT nanoblends was achieved simply by in situ mixing at the air–water interface. The PVDF matrix manifests ultrahigh ferroelectric crystal content, approaching 100% and guaranteeing the ferroelectricity necessary for the modulation of charge transfer and accumulation. Semiconductive P3CPenT components uniformly dispersed in PVDF matrix with some nanostructures such as particles and nanofibers. Ferroelectric nonvolatile memories based on the blend nanosheets were demonstrated with resistive switching properties along with the reversal of ferroelectric polarization direction. The ON/OFF ratio values and retention time are superior to most of the previously reported values, indicative of a great potential for application of such blend nanosheets to flexible electronics.     

Coumarin-based fluorescent probe for the rapid detection of peroxynitrite ‘AND’ biological thiols

A coumarin-based novel ‘AND’ logic fluorescent probe ROS-AHC has been developed for the simultaneous detection of ONOO⁻ and biological thiols. ROS-AHC was shown to exhibit only a very small fluorescence response upon addition of a single GSH or ONOO⁻ analyte. Exposure to both analytes, however, resulted in a significant fluorescence enhancement.

A coumarin-based novel ‘AND’ logic fluorescent probe ROS-AHC has been developed for the simultaneous detection of ONOO⁻ and biological thiols.

Peroxynitrite (ONOO⁻) is a short-lived reactive oxygen and reactive nitrogen species (ROS and RNS) produced intracellularly by the diffusion-controlled reaction of nitric oxide (NO˙) with superoxide (O₂˙⁻).^1–3 Despite playing a key role as a physiological regulator,⁴ it is commonly known for its high reactivity towards most types of biomolecules, causing deleterious effects and irreversible damage to proteins, nucleic acids, and cell membranes.^5,6 ONOO⁻ is therefore a central biological pathogenic factor in a variety of health conditions such as strokes, reperfusion injuries or inflammatory and neurodegenerative diseases (Parkinson''s disease, Alzheimer''s disease).^7–9 Conversely, biothiols such as glutathione and cysteine are endogenous reducing agents, playing a central role in the intracellular antioxidant defence systems.^10–12 Glutathione (GSH), in particular, is the most abundant biothiol in mammalian cells, and exists as both its reduced GSH form, and as the oxidised disulphide form GSSG.^13–15 Peroxynitrite and biothiols such as GSH are intimately linked, as abnormal levels of highly oxidising ONOO⁻ can perturb the delicate GSH/GSSG balance, causing irreversible damage to key processes such as mitochondrial respiration.¹⁶ Thus, abnormal levels of GSH are common in cells undergoing oxidative stress, in which the regulation of and interplay between ONOO⁻ and GSH is closely associated with physiological and pathological processes.^17,18 One such example is drug-induced liver injury (DILI), in which upregulation of ONOO⁻ occurs in hepatotoxicity. Treatment with GSH could be used to remediate this type of cell injury by depletion of ONOO⁻.^19–22One of our core research interests lies in the development of dual analyte chemosensors capable of detecting two distinct analytes such as biological reactive oxygen species and biothiols.^23–26 Although a wide range of single-analyte probes exist for the detection of ROS and thiols separately,^27–30 ‘AND’ logic sensors for their simultaneous detection are still rare.^31–33 We are therefore interested in developing such probes, containing two distinct sensing units, one for each analyte, working simultaneously or in tandem to elicit a fluorescence response.³⁴ This approach allows the monitoring of multiple biomolecular events and factors involved in specific disease pathologies, in order to achieve optimal predictive accuracy for diagnosis and prognostication.³⁵Using these principles, our group has recently focused on developing a range of ‘AND’ logic based sensors exploiting a variety of sensing units and mechanisms of fluorescence. Two such probes are shown below: GSH-ABAH (Fig. 1a), an ESIPT probe with a 4-amino-2-(benzo[d]thiazol-2-yl)phenol (ABAH) core, employing a maleic anhydride thiol-acceptor group;³¹ and JEG-CAB (Fig. 1b), a coumarin-based probe, this time with a salicylaldehyde homocysteine-reactive unit.²⁴ Both of these sensors employ a benzyl boronate ester as their peroxynitrite-reactive unit.Open in a separate windowFig. 1(a) GSH-ABAH, previously reported probe for simultaneous detection of ONOO⁻ and GSH. (b) JEG-CAB, previously reported probe for simultaneous detection of ONOO⁻ and GSH. (c) AHC – a core fluorescent unit that enables the synthesis of ‘AND’ based fluorescent probe for the detection of ONOO⁻ and GSH (d) ROS-AHC, a novel probe detailed in this work for simultaneous detection of ONOO⁻ and GSH.Herein, we set out to develop an ‘AND’ logic gate based fluorescence probe for simultaneous detection of ONOO⁻ and GSH. 3-Amino-7-hydroxy-2H-chromen-2-one (AHC) was chosen as a suitable coumarin fluorophore core for the development of an ‘AND’ logic based sensor, as its free phenol and amine functional groups provided a good opportunity for independent derivatization (Fig. 1).^36–39Previous literature reports show that protection of AHC with a maleic anhydride group results in quenching of the coumarin''s fluorescence intensity due to photoinduced electron transfer (PeT) processes. This fluorescence is rapidly restored in the presence of biological thiols, however, due to their fast addition to this functional group.⁴⁰ Therefore, we suggested that functionalization of the free phenol of this sensor using a benzyl boronic ester should further block the fluorescence, whilst serving as reporter unit for ONOO⁻. The greatly increased reactivity of peroxynitrite over other ROS towards boronate esters^41–43 should allow this functionality to act as a peroxynitrite-selective reporter, leading to an ‘AND’ logic based probe for the detection of ONOO⁻ and biological thiols (Fig. 1, Scheme 1). ROS-AHC was synthesized in 5 steps, starting with a 4-step synthesis of compound 1 adapted from literature procedures,^40,44 followed by the addition of the benzyl boronic pinacol ester (see Scheme S1 ESI^†).Open in a separate windowScheme 1Fluorescence ‘turn on’ mechanism of ROS-AHC in the presence of ONOO⁻ and GSH.The UV-Vis behaviour of ROS-AHC before and after exposure to both GSH and ONOO⁻ was evaluated in pH 7.40 buffer solution, showing a maximum absorption peak at 340 nm for both the unreacted probe and the probe following exposure to GSH, shifting to 350 nm with the addition of ONOO⁻ to the probe and 365 nm after sequential additions of GSH and ONOO⁻ to the probe (Fig. S1 ESI^†). Fluorescence experiments were then carried out. As expected, ROS-AHC was initially non-fluorescent, with a small fluorescence increase upon addition of ONOO⁻ (6 µM) (Fig. 2 and S2 ESI^†). Incremental additions of GSH (0–4.5 µM) resulted in a much larger increase in fluorescence intensity (>69-fold, see Fig. 2 and S3 ESI^†), demonstrating the need for both GSH and ONOO⁻ in order to achieve a significant ‘turn on’ fluorescence response.Open in a separate windowFig. 2Fluorescence spectra of ROS-AHC (5 µM) with addition of ONOO⁻ (6 µM), wait 5 min then incremental addition of GSH (0–4.5 µM), 5 min incubation before measurements in PBS buffer solution (10 mM, pH = 7.40). Fluorescence intensities were measured with λ_ex = 400 nm (bandwidth 8 nm). The green line represents the highest intensity after addition of GSH (4 µM).Similar fluorescence experiments were then carried out in reverse order, with the addition of GSH (6 µM) to ROS-AHC resulting in only a small increase in fluorescence intensity (Fig. 3 and S4 ESI^†). As before, incremental addition of the second analyte, in this case ONOO⁻ (0–5.5 µM), resulted in a large increase in fluorescence intensity (>46-fold, Fig. 3 and S5 ESI^†), confirming that ROS-AHC requires both GSH and ONOO⁻ for a full fluorescence ‘turn on’ response.Open in a separate windowFig. 3Fluorescence spectra of ROS-AHC (5 µM) with addition of GSH (6 µM), wait 5 min then incremental addition of ONOO⁻ (0–5.5 µM) with 5 min incubation before measurements in PBS buffer solution (10 mM, pH = 7.40). Fluorescence intensities were measured with λ_ex = 400 nm (bandwidth 8 nm). The orange line shows the highest intensity after addition of ONOO⁻ (5 µM).Subsequently, the selectivity of this probe towards both analytes was evaluated. A range of amino acids were evaluated (Fig. S6 ESI^†), with only thiol-containing analytes (glutathione, cysteine and homocysteine) eliciting significant fluorescence response, whilst non-thiol amino acids led to no changes in fluorescence intensity. A broad screen of ROS analytes was also carried out, demonstrating excellent selectivity for ONOO⁻, even over H₂O₂ (Fig. S7 ESI^†).The time-dependent response of ROS-AHC with both ONOO⁻ and GSH was also examined (Fig. S8 and S9 ESI^†). After initial addition of GSH or ONOO⁻ to the probe, subsequent addition of the second analyte triggered a rapid and significant increase in fluorescence, achieving maximum fluorescence intensity within 78 s in both cases. Furthermore, LC-MS experiments confirmed the formation of the suggested non-fluorescent intermediates, as well as the final fluorescent species shown in Scheme 1 (Fig S10, S11 and S12^†).In summary, we have developed a coumarin-based dual-analyte ‘AND’ logic fluorescent sensor, ROS-AHC, for the simultaneous detection of ONOO⁻ and biological thiols. ROS-AHC has shown high sensitivity and selectivity towards both ONOO⁻ and biological thiols.     

Micro-supercapacitors based on oriented coordination polymer thin films for AC line-filtering

Reported herein is a facile solution-processed substrate-independent approach for preparation of oriented coordination polymer (Co-BTA) thin-film electrodes for on-chip micro-supercapacitors (MSCs). The Co-BTA-MSCs exhibited excellent AC line-filtering performance with an extremely short resistance–capacitance constant, making it capable of replacing aluminum electrolytic capacitors for AC line-filtering applications.

Micro-supercapacitors exhibiting excellent AC line-filtering with oriented coordination polymer thin-film electrodes are fabricated based on a substrate-independent electrode fabrication strategy.

Micro-supercapacitors (MSCs), as important Si-compatible on-chip electrochemical energy storage devices, have attracted rapidly growing attention due to their rapid energy-harvesting features and burst-mode power delivery.^1,2 In the past few years, a variety of materials including carbon nanotubes,³ graphene,⁴ graphene oxide and mesoporous conducting polymers,^5,6 have already been explored to fabricate the electrodes of MSCs for improving their electrochemical performance. Unfortunately, fabrication procedures of most of these active materials suffer from high cost, harsh and complicated processing conditions, as well as easy cracking and delamination of active films,^1,4 extremely limiting their commercial applications. Moreover, their performances are unsatisfactory for alternating current (AC) line-filtering, which is a key parameter to implement high-frequency operation in most line-powered devices.^7–9For AC line-filtering, capacitors need to respond harmonically at 120 Hz to attenuate the leftover AC ripples on direct current voltage busses.¹⁰ Notably, the development of more compact and miniaturized capacitors to replace traditional aluminum electrolytic capacitors (AECs) for AC line-filtering has become one of the major tasks for future electronics.¹¹ However, typical supercapacitors are incapable for AC line-filtering at this frequency due to their limited ion diffusion and charge transfer efficiency, corresponding to the unsuitable architectures and low conductivity of electrode materials.^10–12 Therefore, the design and fabrication of highly conductive electrodes with optimized architectures for facial electron/ion transportation is crucial for improving the performance of MSCs in AC line-filtering.^12,13 It is worth mentioning that great advancements have been achieved by utilizing vertically oriented graphene sheets as well as 3-dimensional graphene/carbon nanotube carpets prepared by chemical vapor deposition (CVD),^7,8 yielding efficient filtering of 120 Hz AC with short resistance–capacitance (RC) time constants of less than 0.2 ms, which is competitive with those of porous carbon-based supercapacitors (RC time constant = 1 s) as well as AECs (RC time constant = 8.3 ms).⁸ However, the CVD method necessitated in the fabrication of graphene/carbon nanotube electrodes suffers from high cost and complicate procedures.Coordination polymers with an unrivalled degree of structural and property tunability which could be realized by facial procedures, are promising candidates for energy storage.^14,15 Recently, a remarkable achievement which demonstrated a facile and low-cost solution-processed method towards on-chip MSCs based on an azulene-bridged coordination polymer framework (PiCBA) on a Si wafer-supported Au surface was reported.¹⁴ Nevertheless, the reported preparation of coordination polymer film exhibited strong dependence on the surface chemistry (functionality) of the substrate and further improvement of their electrochemical stability was needed. Therefore, the development of substrate-independent fabrication strategies of large-scale and uniform coordination polymer films is in great need not only for fundamental studies, but also for technological applications especially in electronics.Herein, we demonstrate a facial solution-based substrate-independent approach to fabricate oriented coordination polymer (Co-BTA) thin-film electrodes. Remarkably, rigid and flexible Co-BTA-based MSCs with excellent electrochemical stability and AC line-filtering performance were realized, indicating great application potential in micro-supercapacitors.As demonstrated in Fig. 1a–c, a large scale and continuous Co-BTA coordination polymer film composed of one dimensional (1D) molecules ([Co(1,2,4,5-bta)]_n) was prepared at the air–liquid interface through a coordination reaction between 1,2,4,5-benzenetetramine tetrahydrochloride (BTA) and cobalt acetate tetrahydrate (Co(CH₃COO)₂·4H₂O). Notably, the preparation of Co-BTA film is basing on mild conditions and independent of any substrates. The resulting film can be transferred onto any supports including rigid silicon (Si) wafer, glass, as well as flexible PET substrate, indicating great substrate-independence and making it practically applicable for various applications. Besides of a brown film formed at the air–liquid interface, a powder product is also obtained at the bottom of the reaction bottle.Open in a separate windowFig. 1(a) Synthesis of Co-BTA through the coordination reaction between BTA and cobalt ions. (b) Illustration of the gas–liquid interface growth of Co-BTA film. (c) Photographs of the reaction system before and after the coordination reaction.To study the morphology of the resulting Co-BTA film, the brown film was transferred onto a SiO₂/Si wafer by immersing the wafer down to the reaction mixture and subsequently lifting the film up. The scanning electron microscopy (SEM) image reveals a highly uniform and large-scale distribution of the obtained film without cracks or wrinkles (Fig. 2a), which is superior to other reported coordination polymer films obtained via a similar method.¹⁶ An average thickness of approximately 60 nm of the Co-BTA film is observed from the cross-sectional SEM image as shown in Fig. 2b. Interestingly, thickness of the obtained coordination polymer film could be well controlled and Co-BTA films with thicknesses up to several hundred nanometers could be well prepared by adjusting the ratio of raw materials (Fig. 2c and d).Open in a separate windowFig. 2(a) Planar SEM image of Co-BTA film. Cross-sectional SEM images of Co-BTA films with a thickness of (b) 60 nm, (c) 160 nm and (d) 260 nm.To investigate the structure information of the resulting Co-BTA and further explore the coordination reaction, characterizations including powder X-ray diffraction measurements (PXRD), X-ray photoelectron spectroscopy (XPS) and Fourier transform infrared spectroscopy (FTIR) were carried out. The PXRD pattern of Co-BTA powder shown in Fig. S1a^† is in great agreement with that simulated from the crystal structure of Ni(dhbq)·nH₂O (Fig. S1b^†), suggesting that Co-BTA and Ni(dhbq)·nH₂O is isostructural and forms 1D structures with straight infinite chain extends.¹⁷ More interestingly, PXRD measurements employing two different scattering geometries (Fig. S2^†) on the Co-BTA thin-film demonstrate two quite different diffraction patterns. As observed in Fig. 3a, the diffraction pattern observed for the out-of-plane scattering geometry exhibits three characteristic peaks of the Co-BTA film at ∼12°, 24° and 36°, which are corresponding to (001), (002) and (003), respectively. In contrast, the in-plane PXRD profile employing grazing-incidence XRD (GIXRD) technique at an incident angle (α) of 0.2° demonstrates a main peak at ∼18°, which is assigned to the (110) diffraction peak. Different diffraction peaks observed through these two XRD scattering geometries indicate an orientation nature of the as-prepared Co-BTA film,¹⁸ which exhibits better crystallinity compared with the powder Co-BTA product. In addition, the N 1s core level spectrum for Co-BTA film exhibit one typical peak at 399.1 eV, which is corresponding to the amido coordinated with Co^II, indicating the strong coordination between Co^II and BTA (Fig. 3b). The weak peak at ∼401 eV is assigned to N–O due to the oxidation of ligand BTA in ambient environment before reaction. The atomic ratio of N : Co is calculated to be 3.53 : 1 for Co-BTA film and 3.71 : 1 for Co-BTA powder respectively (Fig. S3 and Table S1^†), which is close to the theoretical stoichiometric ratio (4 : 1) for Co-BTA structure, suggesting a high degree of coordination in the resulting product through one Co cation and two benzenetetramine groups. Moreover, the disappearance of two characteristic N–H stretching modes from –NH₂ after the coordination reaction whereas the phenyl-related vibration still exists, further confirms the existence of –NH– in the product through the loss of one H per –NH₂ (Fig. S4^†).¹⁹Open in a separate windowFig. 3(a) PXRD profiles of out-of-plane XRD, in-plane XRD and simulated PXRD pattern of Ni(dhbq)·nH₂O,¹⁷ respectively. *SiO₂/Si substrate. (b) N 1s core level spectra of the Co-BTA film.On the basis of facial fabrication, substrate independence, highly orientation nature, low band gap (1.68 eV, calculated from Fig. S5^†) and excellent stability in acid environment (Fig. S6^†), the resulting Co-BTA film is considered as a promising candidate for MSCs application. Fig. 4a schematically depicts the stepwise fabrication of a planar Co-BTA film based MSC on a SiO₂/Si wafer and its electrochemical performance is first examined by cyclic voltammetry (CV) with scan rates ranging from 50 mV s⁻¹ to 1000 V s⁻¹ (Fig. 4b and c). At a low scan rate of 50 mV s⁻¹, the 60 nm-thick Co-BTA film based MSC exhibited a pronounced pseudocapacitive effect, implying the occurance of faradaic reaction.²⁰ With the increase of scan rate, a gradual transition of the CV curves from the pseudocapacitive to the typical electrical double-layer capacitive behavior with a nearly rectangular CV shape was observed. Remarkably, the device exhibited a maximum volumetric capacitance of 23.1 F cm⁻³ at 50 mV s⁻¹, which is comparable with those of reported carbon- or graphene-based MSCs (Table S2^†), e.g., onion-like carbon,²¹ vertically oriented graphene,⁸ and carbon nanotubes/graphene.⁷ Even though a trend that C_V decreased gradually with increasing scan rate was observed, the Co-BTA-based electrode still delivered a C_V of 2.7 F cm⁻³ even at a high scan rate of 1000 V s⁻¹, suggesting an excellent capacitive performance of this Co-BTA-based MSC device.⁷Open in a separate windowFig. 4(a) Schematic illustration of the fabrication of MSC device with the Co-BTA film electrode. (b) CV curves of Co-BTA-based MSCs in the H₂SO₄–PVA gel electrolyte at different scan rates. (c) C_V evolution of the MSCs at different scan rates.Electrochemical impedance spectroscopy (EIS) measurements were performed to evaluate the charge transport properties within the Co-BTA-based MSCs. The Nyquist plot shown in Fig. S7^† indicated the kinetic features of electron transfer/ion diffusion at the electrode, from which an almost straight line especially in the low frequency region is observed. Notably, the plot shows a closed 90° slope without a charge transport semicircle at high frequency which is corresponding to an almost ideal capacitive ion diffusion mechanism, due to the excellent charge transfer property of the oriented Co-BTA electrode film. Moreover, this microdevice exhibited a low equivalent series resistance of 13.48 Ω (Fig. S7^† (inset)), further suggesting the ultrafast ion diffusion characteristic in such a Co-BTA-based-MSC.²² It''s suggested that the unique kinetic feature of fast ion diffusion and charge transfer benefits from the intrinsic characteristics of the oriented polymer film composed of 1D molecules, which can not only facilitate rapid ionic diffusion but also facilitate the interfacial charge transfer and faradaic redox reaction between the electrode material and electrolyte.What''s more, the dependence of the phase angle on frequency shown in Fig. 5a delivered a high characteristic frequency f₀ of 6812 Hz at the phase angle of −45° (the resistance and reactance of the device have equal magnitudes),¹⁰ which is much higher than that of an active carbon supercapacitor (5 Hz),²³ sulfur-doped graphene MSCs (3836 Hz),²² or an azulene-bridged coordination polymer framework based MSCs (PiCBA-MSCs) (3620 Hz) and so on,¹⁴ as summarized in Table S2.^† Moreover, a max phase angle of −80° at a frequency of 18 Hz is observed, indicating the performance of this Co-BTA based MSCs is 89% of that of an ideal capacitor. Importantly, a large impedance phase angle of −78.6° was obtained at a frequency of 120 Hz, which is the largest reported value for coordination polymer based MSCs (Table S2^†), suggesting an excellent AC line-filtering performance of our microdevice.⁷ To further conform the ultrahigh fast ion diffusion in Co-BTA-based-MSCs, the relaxation time constant τ₀ (τ₀ = 1/f₀, the minimum time needed to discharge all the energy from the device with an efficiency of greater than 50% of its max. value) was calculated to be only 0.15 ms (6812 Hz), which is orders of magnitude higher than that of conventional electrical double-layer capacitors (1 s),⁸ activated or onion-like carbon MSCs (<200 ms, <10 ms),^21,23 and much shorter than those of MSCs based on carbon nanotubes/reduced graphene oxide (4.8 ms) as well as azulene-bridged PiCBA coordination polymer framework film (0.27 ms).^14,24 Moreover, a short RC time constant (τ_RC) of 0.32 ms was obtained (Fig. 5b) through a series-RC circuit model, making it capable of replacing AECs for AC line-filtering application. To the best of our knowledge, this is the first report of coordination polymer-based MSCs exhibiting such a small relaxation time constant and RC time constant.Open in a separate windowFig. 5(a) Impedance phase angle on the frequency for the Co-BTA-based microdevices. (b) Plot of capacitance (C_V′ = volumetric real capacitance and C_V′′ = imaginary capacitance) versus the frequency of Co-BTA-based microdevices. (c) Cycling stability of Co-BTA film with 10 000 cycles at the scan rate of 50 V s⁻¹. Inset displays the CV curves at the first, five thousandth and ten thousandth cycle, respectively. (d) Ragone plots for Co-BTA film, compared with commercially applied Li-thin-film batteries,²¹ electrolytic capacitors,² CNT-graphene carpets,²⁴ PiCBA coordination polymer and MXene-reduced graphene oxide.^14,25Impressively, this oriented electrode structure exhibits excellent long-term electrochemical stability with ∼96.3% capacitance retention even after 10 000 cycles of charging/discharging at a scan rate of 50 V s⁻¹ (Fig. 5c), which has also been confirmed by comparing the CV curves before and after testing for 10 000 cycles (inset of Fig. 5c). It''s worth pointing out that the as-made Co-BTA-based MSCs exhibit the best electrochemical stability among reported MSCs with coordination polymer electrodes.¹⁴ On the basis of the above discussion, it is reasonable to conclude that the ultrahigh fast ion diffusion/charge transfer in Co-BTA-based-MSCs attributed to the oriented architecture of Co-BTA thin-film electrodes, the excellent AC line-filtering performance, as well as remarkable electrochemical stability contributes to the excellent performances of Co-BTA-based-MSCs. Moreover, the power density and energy density of the as-made device is calculated and compared with that of MSCs based on other electrode materials to evaluate the energy storage performance of the Co-BTA based MSCs. The Ragone plot in Fig. 5d reveals a high power density of 1056 W cm⁻³ for our device, which is at least five orders of magnitude higher than that of commercial thin-film lithium batteries. What''s more, our device exhibits an energy density of up to 1.6 mW h cm⁻³ at 50 mV s⁻¹, which is at least one order of magnitude higher than that obtained for MSCs based on CNTs-graphene carpet and high-power electrolytic capacitors.^2,24To further demonstrate the substrate independence of this fabrication strategy, flexible Co-BTA-based-MSC device was fabricated and investigated basing on a flexible polyethylene terephthalate (PET) substrate instead of rigid Si substrate in the same way (Fig. S8–S10^†). The as-fabricated device exhibited a maximum volumetric capacitance of 22.0 F cm⁻³ at 50 mV s⁻¹, a short relaxation time constant τ₀ of 0.15 ms and a RC time constant (τ_RC) of 0.42 ms, which are close to the properties of devices with a Si substrate, confirming the substrate independence of this fabrication scheme. As a proof-of-concept application, bending tests were carried out and the bended device (radius = 1 cm) exhibited a small relaxation time constant τ₀ of 0.21 ms and RC time constant (τ_RC) of 0.42 ms, suggesting that the Co-BTA-based MSC with PET substrate in a bended state still delivers a good ion diffusion and AC line-filtering performance.In conclusion, we have demonstrated a facile method that can be used to construct large scale and highly oriented uniform Co-BTA coordination polymer thin films using a very convenient and fast process. With this method, Co-BTA-based MSCs are fabricated without any dependence of the substrate. The as-fabricated MSCs on Si substrate exhibit high specific capacitance, energy density as well as excellent electrochemical stability. Particularly, the fabricated Co-BTA based MSCs deliver excellent AC line-filtering performance with an extremely short RC time of 0.32 ms, attributed to the facilitated ion diffusion beneficial from the oriented architecture of Co-BTA thin film. The high-performance electrochemical properties of Co-BTA-MSCs makes Co-BTA films promising materials to provide more compact AC filtering units for future electronic devices.     

Coumarin-based fluorescent ‘AND’ logic gate probes for the detection of homocysteine and a chosen biological analyte

With this research we set out to develop a number of coumarin-based ‘AND’ logic fluorescence probes that were capable of detecting a chosen analyte in the presence of HCys. Probe JEG-CAB was constructed by attaching the ONOO⁻ reactive unit, benzyl boronate ester, to a HCys/Cys reactive fluorescent probe, CAH. Similarly, the core unit CAH was functionalised with the nitroreductase (NTR) reactive p-nitrobenzyl unit to produce probe JEG-CAN. Both, JEG-CAB and JEG-CAN exhibited a significant fluorescence increase when exposed to either HCys and ONOO⁻ (JEG-CAB) or HCys and NTR (JEG-CAN) thus demonstrating their effectiveness to function as AND logic gates for HCys and a chosen analyte.

With this research we set out to develop of a number of coumarin-based ‘AND’ logic fluorescence probes that were capable of detecting a chosen analyte in the presence of HCys.

Homocysteine (HCys) is a non-proteinogenic amino acid, formed from the de-methylation of methionine,¹ which is then converted into cysteine (Cys) via a vitamin B₆ cofactor. Typical physiological concentrations of HCys range between 5–15 μmol L⁻¹.² However, elevated levels of HCys (>15 μmol L⁻¹), which is known as hyperhomocysteinemia (hHCys),³ have been associated with pregnancy disorders, Alzheimer''s disease, cardiovascular disease and neurodegenerative diseases (NDs).^4–6 It is believed that the main cause of HCys induced toxicity is through the non-enzymatic modification of proteins. This is achieved through irreversible covalent attachment of the predominant metabolite of HCys, homocysteine thiolactone (HTL), to lysine residues; a phenomenon known as ‘protein N-homocysteinylation’ that results in the loss of a proteins structural integrity leading to loss of enzymatic function and aggregation.⁷A number of fluorescent sensors have been developed for the detection of HCys to help improve our understanding of its role in biological systems.^8–11 However, these fluorescent probes have focused on the detection of a single biomarker (HCys), however, processes associated with HCys induced toxicity often involve more than one biochemical species. For example, it has been reported that peroxynitrite (ONOO⁻) and nitric oxide (NO˙) play a significant role in HCys-mediated apoptosis in trigeminal sensory neurons¹² and HCys has been reported to induce cardiomyocytes cell death through the generation of ONOO⁻.¹³ The production of ONOO⁻ is believed to be the result of an increased production of superoxide (O₂˙⁻) by HCys activating the enzyme NADPH oxidase.^14–16 This increased production of O₂˙⁻ leads to a reduction in the bioavailability of NO˙ by increasing the formation of ONOO⁻ (NO˙ + O₂˙⁻ → ONOO⁻).¹⁷ The reported ONOO⁻ concentrations in vivo are believed to be approximately 50 μM but, higher concentrations of 500 μM have been found in macrophages.^18,19 Furthermore, hypoxia has been reported to facilitate HCys production in vitamin-deficient diets²⁰ where hypoxia leads to an upregulation of nitroreductase (NTR) – a reductive enzyme upregulated in cells under hypoxic stress.^21,22 Therefore, the development of tools that enable an understanding of the relationship of HCys with these biologically important species would be highly desirable.To achieve this, a number of fluorescent probes have been developed that are capable of detecting two or more analytes.²³ These include AND logic gate based-fluorescence probes, which require both analytes to work in tandem to produce a measurable optical output.^24–28 In our group, we have developed several ‘AND’ reaction-based probes for the detection of ROS/RNS and a second analyte.^29–32 These ‘AND’ logic scaffolds have been used to detect two analytes within the same biological system.^24,33Owing to the pathological role of HCys, we set out to develop the first example of a fluorescent probe for the detection of HCys and biological related analyte. Aiming towards that target, we became interested in a previously reported coumarin-based fluorescent probe developed by Hong et al.CAH, with a salicylaldehyde (Fig. 1).³⁴ Salicylaldehyde is a known reactive unit towards HCys/Cys, therefore we believed CAH could be used as a scaffold for the development of ‘AND’-based systems for the detection of HCys/Cys and a second analyte.³⁴ In the presence of HCys, CAH exhibited a ‘turn-on’ fluorescence response which is attributed to the nucleophilic nature of the nitrogen and sulfur atoms resulting in thiazine ring formation (Scheme S1, Fig. S1 and S2^†).^34–36Open in a separate windowFig. 1(a) CAH – a core fluorescent unit that enables the synthesis of ‘AND’ based fluorescent probe for the detection of HCys/Cys and a second analyte. (b) JEG-CAB enables the detection of HCys/Cys and (ROS/RNS) while (c) JEG-CAN enables the detection of HCys/Cys and NTR.We believed that CAH was a useful core unit that can be used to introduce the chosen reactive chemical trigger on the phenol for the detection of the corresponding analyte with HCys/Cys. Owing to the relationship between HCys and ONOO⁻/NTR, we set out on the development of a HCys AND ONOO⁻ probe and a HCys AND NTR probe.Therefore, we set out to prepare JEG-CAB and JEG-CAN, which are able to detect HCys/Cys and peroxynitrite (ONOO⁻) or nitroreductase (NTR), respectively (Scheme 1). For JEG-CAB, a benzyl boronate ester was introduced as a ONOO⁻ reactive unit.³⁷ For JEG-CAN, a p-nitrobenzyl group was installed as it is known to be an effective substrate for NTR.^38–40Open in a separate windowScheme 1Synthesis of target probe JEG-CAB and JEG-CAN.To afford CAH, compound 2 was synthesized by refluxing umbelliferone and acetic anhydride at 140 °C. Compound 2 was then dissolved in trifluoroacetic acid at 0 °C followed by the addition of hexamethylenetetramine (HMTA). The mixture was heated to reflux overnight and the solvent was then removed. The intermediate was then hydrolyzed in H₂O for 30 min at 60 °C. Upon isolating CAH, it was then alkylated using 4-bromomethylphenylboronic acid pinacol ester and K₂CO₃ in DMF at r.t. to afford JEG-CAB in 51% yield. JEG-CAN was produced by alkylating CAH using 4-nitrobenzyl bromide and K₂CO₃ in DMF at r.t. to give 49% yield (Scheme 1).We then evaluated the ability of JEG-CAB to detect ONOO⁻ ‘AND’ HCys in PBS buffer solution (10 mM, pH 7.40). The maximum absorption of JEG-CAB at 336 nm shifted to 323 nm with the addition of HCys and then slightly shifted to 328 nm following the addition of ONOO⁻ (Fig. S3^†). As shown in Fig. 2, JEG-CAB was initially non-fluorescent, but the addition of HCys (1 mM) led to a small increase in fluorescence intensity, the subsequent additions of ONOO⁻ (0–24 μM) led to a significant increase in fluorescence intensity (>17-fold, see Fig. S5^†). These results demonstrated the requirement for both ONOO⁻ ‘AND’ HCys to obtain a significant turn ‘‘on’’ fluorescence response.Open in a separate windowFig. 2Fluorescence spectra of JEG-CAB (15 μM) with addition of HCys (1 mM) and incubated for 40 min then measured. Followed by incremental additions of ONOO⁻ (0–24 μM). The data was obtained in PBS buffer solution (pH 7.40, 10 mM) at 25 °C. λ_ex = 371 (bandwidth 20) nm. Dashed line represents JEG-CAB and Hcys addition only. Blue line represents highest intensity after addition of ONOO⁻.The addition of HCys and ONOO⁻ were then performed in reverse where JEG-CAB exhibited a negligible increase in fluorescence intensity upon addition of ONOO⁻ (16 μM). However, in a similar manner to that shown in Fig. 2, a large increase in fluorescence intensity was produced after the subsequent addition of HCys (0–5.5 mM) (Fig. 3 and S6^†). LC-MS experiments were carried out to ascertain the reaction mechanism and the results confirmed the sequential formation of the thiazine ring in the presence of HCys followed by boronate ester cleavage in the presence of ONOO⁻ or vice versa (Scheme S2 and Fig. S19–S21^†).Open in a separate windowFig. 3Fluorescence spectra of JEG-CAB (15 μM) with addition of ONOO⁻ (16 μM) and followed by incremental additions of HCys (0–5.5 mM) measurements were taken after 40 min of both additions. The data was obtained in PBS buffer solution (pH 7.40, 10 mM) at 25 °C. λ_ex = 371 (bandwidth 20) nm. Dashed line represents JEG-CAB and ONOO⁻ addition only. Blue line represents highest intensity after addition of HCys.As expected, probe JEG-CAB was shown to have excellent selectivity with ONOO⁻ against other ROS in the presence of HCys (1 mM) (Fig. S9 and S10^†). Furthermore, JEG-CAB exhibited a high degree of selectivity towards a series of amino acids where only HCys and Cys led to a fluorescence response in the presence of ONOO⁻. This is due to the formation of stable five or six-membered thiazine rings (Fig. S7 and S8^†).³⁴We then evaluated the changes in the fluorescence of JEG-CAN with both HCys and NTR in PBS buffer solution (10 mM, pH 7.40, containing 1% DMSO). As shown in Fig. 4, addition of HCys led to a small increase in fluorescence intensity. However, subsequent addition of NTR (4 μg mL⁻¹) led to a large time dependant increase in fluorescence intensity. To ensure both analytes were required, NTR and NADPH was kept constant (4 μg mL⁻¹ and 400 μM respectively) resulting in a 3.4 fold fluorescence increase (Fig. 5). We attribute the large initial increase to background fluorescence of NADPH.⁴¹ NTR then facilitates reduction of the nitro group of JEG-CAN releasing the core probe CAHvia a fragmentation cascade (Scheme S3^†).^38,42 Subsequent addition of HCys (2.0 mM) led to a 2 fold increase in fluorescence intensity. Again, LC-MS experiments confirmed the proposed reaction mechanism (Fig. S22^†).Open in a separate windowFig. 4Fluorescence spectra of JEG-CAN (15 μM) with initial addition of HCys (2 mM) and incubated for 60 min. Followed by addition of nitroreductase (4 μg mL⁻¹) and NADPH (400 μM) and measured over 90 min in PBS buffer solution (pH = 7.40, 10 mM, containing 1% DMSO). λ_ex = 363 nm. Ex slit: 5 nm and em slit: 5 nm. Dashed line represents JEG-CAN and HCys addition only. Blue line represents highest intensity after addition of NTR.Open in a separate windowFig. 5Fluorescence spectra of JEG-CAN (15 μM) with initial addition of nitroreductase (4 μg mL⁻¹) and NADPH (400 μM) and incubated for 60 min. Followed by addition HCys (2 mM) and measured over 90 min in PBS buffer solution (pH = 7.40, 10 mM, containing 1% DMSO). λ_ex = 363 nm. Ex slit: 5 nm and em slit: 5 nm. Dashed line represents JEG-CAN and NTR addition only. Blue line represents highest intensity after addition of HCys.Kinetic studies for JEG-CAN with both NTR and HCys were carried out (Fig. S11–S18^†) where it is clear that JEG-CAN exhibits a dose dependant fluorescence increase in response of both HCys and NTR.Unfortunately, the probes failed to give good data in cells, which could be due to their short excitation wavelengths or the extremely low intracellular HCys concentrations (5–15 μM). We are now pursuing the development ‘AND’ logic fluorescence probes with longer excitation and emission wavelengths suitable for in vitro and in vivo applications.In summary, we have developed two coumarin-based ‘AND’ logic fluorescence probes (JEG-CAB and JEG-CAN) for the detection of HCys/Cys and ONOO⁻ or NTR, respectively. CAH is a useful platform that enables easy modification for the development of ‘AND’-based fluorescent probes for the detection of HCys/Cys and a second analyte. Both JEG-CAB and JEG-CAN were shown to be ‘AND’-based fluorescent probes.     

A simple two-photon turn-on fluorescent probe for the selective detection of cysteine based on a dual PeT/ICT mechanism

Herein, a simple two-photon turn-on fluorescent probe, N-(6-acyl-2-naphthayl)-maleimide (1), based on a dual PeT/ICT quenching mechanism is reported for the highly sensitive and selective detection of cysteine (Cys) over other biothiols. The probe was applied in the two-photon imaging of Cys in cultured HeLa cells, excited by a near-infrared laser at 690 nm.

N-(6-acyl-2-naphthayl)-maleimide (1) is a simple two-photon fluorescent probe with selectivity for cysteine, based on a thiol-Michael-addition-transcyclization cascade and dual PeT/ICT quenching mechanism.

Cysteine (Cys), homocysteine (Hcy), and glutathione (GSH) are structurally similar biothiols, but their biological functions are quite different from one another.^1–6 Among these biothiols, Cys functions as one of the twenty-one amino acids for peptide and protein synthesis, and Cys deficiency is also associated with certain disease symptoms.^7–10 Methods for the selective detection and differentiation of Cys among different biothiols, including high performance liquid chromatography (HPLC),¹¹ capillary electrophoresis,¹² electrochemical assay,¹³ UV-vis spectroscopy,¹⁴ and fluorescence-based methods,^15–17 are important for its biological studies. Recently, fluorescent probes have attracted much attention as vital chemical biology tools due to their high sensitivity, convenient operation, and real-time imaging capabilities.^18–20 A number of Cys-selective fluorescent probes have been reported,²¹ which utilize Cys-selective recognition groups such as aldehydes,^11,22 acrylates,²³ thioesters,²⁴ and electron-deficient aromatic halides^25–27 in their structures. However, many of them have relatively long response times and low sensitivity due to a slow cyclization process. In addition, fluorescent probes with high selectivity for Cys over Hcy are difficult to achieve because they differ by only one methylene group.²⁸ Recently, we reported that N-(N′-butyl-1,8-naphthalimide-4-yl)-maleimide, containing a single maleimide group as the recognition group, is a fast, sensitive, and selective fluorescent probe for Cys based on a dual photo-induced electron transfer (PeT) and photo-induced intramolecular charge transfer (ICT) quenching mechanism.²⁸ Different from many other maleimide-based fluorescent probes that only undergo a PeT mechanism,¹⁵ the additional ICT quenching mechanism keeps the 1,8-naphthalimide (NAP) fluorophore in the thiol-Michael adduct in a low fluorescence emission state due to the strong electron-withdrawing effects of the succinimide group at its 4-position. Then, a subsequent transcyclization step, involving the formation of a six-membered thiomorpholinone ring and cleavage of a five-membered succinimide ring, converts the non-fluorescent thiol-Michael adduct into the major fluorescent product, in which the ICT quenching is removed, resulting in a drastic fluorescence turn-on response.²⁸ A similar transcyclization process and the simultaneous removal of ICT quenching allowed us to design a NAP-based turn-on fluorescent probe for γ-glutamyltranspeptidase²⁹ and a coumarin-based turn-on fluorescent probe with dual recognition groups and dual cyclization for the selective detection of Cys.³⁰ In addition, another NAP-based dual PeT/ICT probe was recently reported by Meka and Heagy for the detection of hydrogen sulfide, although two recognition groups instead of one were adopted in their probe to achieve the dual quenching mechanism.³¹Our previous work and that of other groups has demonstrated that the combination of PeT and ICT mechanisms is particularly suitable for the design of fluorescent probes with a significant fluorescence turn-on response.^30–33 However, many of these probes have a short excitation wavelength in the UV or visible range, which is not optimal for biological applications due to enhanced phototoxicity and/or autofluorescence.^34,35 Considering that two-photon fluorescence imaging has advantages such as the excitation process being carried out by a near-infrared (NIR) laser that has a reduced cell toxicity and low fluorescence background,³⁶ in this work, we aimed to introduce a similar dual PeT/ICT quenching mechanism to the known two-photon fluorophore 6-acyl-2-naphthylamine^37–39 in order to design a simple maleimide-based two-photon fluorescent probe, 1, for the selective detection of Cys over Hcy and GSH. It was also tested to determine whether it is a turn-on fluorescent probe with high sensitivity and selectivity, which reacts with Cys via a fast two-step thiol-Michael addition and transcyclization cascade reaction.²⁸ The structure of probe 1 is shown in Fig. 1. It has a maleimide group at its 2-position, which promotes the PeT quenching effect. It also has an additional electron-withdrawing methylcarbonyl group at its 6-position to ensure a pull–pull ICT quenching effect.Open in a separate windowFig. 1Design rationale of the fluorescent probe 1 for the selective turn-on detection of Cys over Hcy and GSH.Probe 1 was conveniently synthesized from 6-acyl-2-naphthylamine (3)³⁹ in a two-step process with a total yield of 38% (see Scheme S1 in the ESI^†). First, the amine 3 was reacted with maleic anhydride to form the maleic amide acid 4. Then, the amide acid 4 was cyclized to afford the maleimide 1 in the presence of acetic anhydride (see the ESI^† for more details).We then investigated the absorption and fluorescence emission response of the probe towards just 1 equiv. of Cys. The time-dependent absorption spectra upon the addition of 1 equiv. of Cys are shown in Fig. 2a. Probe 1 has a maximum absorption peak at 292 nm. Upon addition of Cys, the maximum absorption peak shifts to 314 nm, a red-shift of 22 nm. Notably, an isosbestic point can be seen at 295 nm after 2 min, indicating the formation of an intermediate within 2 min, which is then converted into the final product. The UV spectral changes supported the presence of a proposed cascade reaction sequence for the fast formation of a thiol-Michael adduct intermediate, which then underwent a relatively slow intramolecular transcyclization process to give the final product. From time-dependent fluorescence emission studies (Fig. 2b), probe 1 was found to have almost no fluorescence emission due to dual PeT and ICT quenching effects. Upon the addition of 1 equiv. of Cys, a drastic turn-on fluorescence response (a >3000 fold increase) was observed at 446 nm (see Fig. S1b in the ESI^†). The fluorescence intensity at 446 nm reached its maximum value after around 30 min indicating that the cascade reaction finished in about 30 min (Fig. 2b, and S2a in the ESI^†). The pseudo-first-order reaction kinetic constant based on the fluorescence enhancement was calculated as 0.123 min⁻¹ (half-time = 5.64 min, Fig. S2b in the ESI^†), indicating an overall fast cascade reaction. Fluorescence titration experiments using an increasing amount of Cys from 0 to 4.0 equiv. over 30 min showed a steady increase in the fluorescence intensity and the maximum intensity was reached at exactly 1.0 equiv. of Cys. Further Cys addition did not increase the fluorescence intensity, indicating that probe 1 reacts with Cys in a 1 : 1 molar ratio (Fig. 2c and S3 in the ESI^†), which was also supported by the Job plot (see Fig. S4 in the ESI^†). From the linear relationship of the fluorescence intensity at 446 nm versus the Cys concentrations, the detection limit of probe 1 (2 μM) for Cys was calculated as 1.4 nM (S/N = 3, Fig. 3d), indicating that 1 is a highly sensitive probe for Cys. Moreover, the probe showed excellent selectivity for the detection of Cys over many other species (Fig. 2e, and S5 in the ESI^†), including the structurally similar thiols Hcy, GSH, and N-acetylcysteine (NAC). The fluorescence intensity at 446 nm for 1 equiv. of Cys was significantly higher (12.2-fold, 9.1-fold, and 17.7-fold, respectively) than that of 10 equiv. of Hcy, GSH, or NAC. To further confirm the reaction mechanism, the reaction product, 2, from the reaction of probe 1 with Cys, was isolated and its structure was confirmed using ¹H nuclear magnetic resonance (NMR) spectroscopy, ¹³C NMR spectroscopy, 2D-rotating-frame nuclear Overhauser effect spectroscopy (ROESY), and high-resolution mass spectrometry (HRMS) (see the ESI for more details^†). The fluorescence quantum yields of probe 1 and product 2 were measured as 0.002 and 0.782, respectively (see the ESI for more details^†). Therefore, the formation of the transcyclization product 2 was determined to be responsible for the observed fluorescence turn-on response. For the other thiols, the transcyclization steps of the thiol-Michael adducts were much slower, resulting in the observed high selectivity. Overall, we have shown here that probe 1 is a highly sensitive and selective turn-on fluorescent probe for Cys.Open in a separate windowFig. 2(a) Time-dependent UV-vis spectra of probe 1 (10 μM) upon the addition of 1 equiv. of Cys (a spectrum was recorded every 2 minutes); (b) time-dependent fluorescence emission spectra of probe 1 (2 μM) upon the addition of 1 equiv. of Cys (a spectrum was recorded every 3 minutes); (c) time-dependent fluorescence emission intensity at 446 nm of probe 1 (2 μM) upon addition of Cys (0 to 4 equiv.); (d) a linear relationship of the fluorescence intensity at 446 nm versus the Cys concentration (0.2–2.0 μM); (e) fluorescence response of probe 1 (2 μM) at 446 nm toward various species in PBS buffer (10 mM, pH 7.4): (1) blank; (2) Cys; (3) Hcy; (4) GSH; (5) NAC; (6) valine; (7) glycine; (8) isoleucine; (9) lysine; (10) leucine; (11) histidine; (12) asparagine; (13) methionine; (14) proline; (15) serine; (16) alanine; (17) threonine; (18) arginine; (19) glutamine; (20) aspartic acid; (21) glutamic acid; (22) tyrosine; (23) tryptophan; (24) phenylalanine; (25) glucose; (26) H₂O₂; (27) Na⁺; (28) K⁺; (29) Ca²⁺; (30) Mg²⁺; (31) Fe³⁺; (32) Fe²⁺; (33); Cu²⁺; (34) Zn²⁺ (All measurements were made in 10 mM PBS buffer, pH 7.4, 25 °C, and λ_ex = 314 nm).Open in a separate windowFig. 3Two-photon fluorescence images (b, e, h, k) of HeLa cells collected at 410–510 nm (blue to cyan-blue, λ_ex = 690 nm), the corresponding bright field view (a, d, g, j), and overlap of the fluorescence channel and the bright field view (c, f, i, l) after different treatments: (a–c) the cells were pretreated with 0.5 mM of N-ethylmaleimide (NEM) for 30 min and then incubated with 10 μM of probe 1 for 30 min; (d–f) cells were first pretreated with 0.5 mM of NEM for 30 min, then after addition of 1 mM of Cys were incubated for 30 min, and finally, incubated with 10 μM of probe 1 for 30 min (scale bar = 10 μm); the conditions for (g–i) and (j–l) were similar to those of (d–f), except that 10 μM of Hcy and 10 μM of GSH were used instead of 10 μM of Cys.Encouraged by the fast, selective, and sensitive in vitro fluorescence response of probe 1 for the detection of Cys, we further evaluated its potential use as a two-photon imaging agent for Cys in biological systems, such as in living cells. The fluorescence response of probe 1 towards Cys at different pH values was evaluated and a suitable pH range for Cys detection was determined to be 7.0 to 10.0, which is a good range for cell imaging applications because physiological conditions have a pH of around 7.4 (see Fig. S7 in the ESI^†). HeLa cells were then pretreated with N-ethylmaleimide (NEM, 0.5 mM) for 30 min to remove the endogenous cellular thiols, and incubated with Cys (1 mM), Hcy (1 mM), or GSH (1 mM), respectively for 30 min to increase the specific thiol levels. The samples were then further incubated with probe 1 (10 μM) for 30 min and were then washed with PBS buffer before two-photon fluorescence cell images and the corresponding bright-field view images were taken (Fig. 3(d–l)). Control images were also taken for samples pretreated with NEM (0.5 mM) and then incubated with probe 1 (10 μM) (Fig. 3a–c). Only cells pretreated with NEM and then Cys showed a distinctive blue fluorescence (Fig. 3e). The above cell imaging studies clearly demonstrated that probe 1 is capable of the selective detection and imaging of intracellular Cys over Hcy and GSH in living cells by two-photon fluorescence imaging with low background fluorescence interference.     

Ni(ii)-based coordination polymers for efficient electrocatalytic oxygen evolution reaction

The exploration of highly efficient, stable and cheap water oxidation electrocatalysts using earth-abundant elements is still a great challenge. Herein, alkaline-stable cationic Ni(ii) coordination polymers (Ni-CPs) were successfully obtained under hydrothermal conditions, which could stabilize the incorporation of Fe(iii) to form Fe-immobilized Fe@Ni-CPs. The newly developed Ni-based CPs were used for the first time as an effective electrocatalyst for the oxygen evolution reaction in strong alkaline media.

An alkaline-stable cationic Ni(ii) coordination polymer showed remarkable oxygen evolution reaction (OER) catalytic activity due to capturing Fe ions.

The oxygen evolution reaction (OER) plays a vital role in energy storage and conversion applications due to energy issues and the need for sustainable development.^1–4 Because of the sluggish kinetics of the OER, excellent electrocatalysts are required to work in acidic or strong alkaline environments.^5,6 Noble metal-based catalysts like IrO₂ and RuO₂ with high efficiency showed excellent OER catalytic activity.⁷ However, noble metal with high-cost and scarcity are impractical for scale-up applications. Currently, to substitute these precious metal-based materials, transition-metal-based (Fe, Co, Ni and so on) and metal-free (e.g., N, P and S) hybrid materials have been extensively developed.^8–10 For example, metal–organic frameworks (MOFs) such as ZIF-8/67 derived metal–carbon composite materials exhibit promising electrocatalytic performance.^11,12 Unfortunately, the pyrolysis process destroys the framework completely and causes agglomeration of metals, resulting in a decreased number of active sites. Therefore, to explore highly efficient and low-cost OER catalysts that can be directly used in the OER without calcination are desired, including complexes and MOFs.Recently, much of transition bimetallic materials showed excellent electrocatalytic activity.^13–15 However, a handful of examples such as Fe–Co-MOFs or Co–Ni-MOFs have been explored due to instability, poor conductivity and harsh synthesis conditions.^16–20 Notably, the disadvantages of MOFs have limited their usage in the potential OER. Therefore, it is urgent to develop cheap, stable and active OER catalysts to replace the precious metals. However, this is still a great challenge. Coordination polymers (CPs) with a low-dimensional framework similar to MOFs are constructed by metal ion and organic ligands with potential active sites and functional groups, exhibiting wide applications in sensing, photoluminescence and photocatalysis.^21–26 However, rare examples of CPs have been directly explored in the OER. For Fe/Ni-based bimetal electrocatalysts, a novel strategy involves doping Fe(iii) into a functional Ni-based CPs, which could enhance the electrocatalytic OER activities.Herein, we report the hydrothermal synthesis of a Ni-based CPs as a high-performance OER electrocatalyst in strong alkaline solutions (Scheme 1). The blue crystals of [Ni(bp)₃·(H₂O)₂]·(bp)·(ClO₄)₂ (bp = 4,4′-biprydine) were obtained upon the reaction of bp ligands with NiClO₄·6H₂O under hydrothermal systems. Chemical stability tests displayed that Ni-CPs could retain their original framework in water or even a strong alkaline (pH = 14) solution after 12 hours (Fig. S1^†), which is rarely reported for most transition metal CPs. The thermogravimetric analysis (TGA) showed that there is a significant change at about 110 °C due to the weight loss of the partial guest (Fig. S2^†). Interestingly, the cationic Ni-CPs successfully captured the Mohr''s salt (ammonium iron(ii) sulfate) by taking advantage of the post-synthetic strategy. The obtained Fe@Ni-CPs exhibited a high-efficient OER activity under strong alkaline conditions.Open in a separate windowScheme 1Illustration of the synthesis process for Fe doped Ni coordination polymers for the OER.Single-crystal X-ray diffraction analysis revealed that the Ni-CPs crystallized in the C2/c space group (Table S1^†). The obtained Ni coordination polymer was the isostructural compound reported by Talham,²⁷ but their packing modes were distinctly different (Fig. S3–S5^†). It also had a similar coordination environment to the railroad-like double chains synthesized by Yaghi.²⁸ The parallel chains were occupied by 4,4′-bpy, perchlorate and water molecules. In the Ni-CPs, most of the phenyl rings adopted the face-to-face mode. There were evident π⋯π interactions between the adjacent 4,4′-biprydine (Fig. 1a). In addition, there were strong intermolecular hydrogen bonds between 4,4′-bpy and perchlorate (strong Cl–O⋯C amongst adjacent layers) (Fig. 1a). The weak reaction increased the high density of the framework and protected the coordination bonds against external guest attacking. These chains and guests were further packed with a three-dimensional structure along the c-axis (Fig. 1b).Open in a separate windowFig. 1(a) The weak reaction between chains in Ni CPs, showing the Cl–O⋯C (green and black dashed line) and π⋯π interactions (pink dashed line) between the adjacent perchlorate and 4,4′-biprydine, respectively; (b) the packing view of Ni CPs along the c-axis. Cl in green, Ni in pale blue, N in blue, O in red and C in black. H atoms and partial guest molecules were omitted for clarity.The unique cationic Ni-CPs framework has the potential to immobilize some counterpart ions. To demonstrate this, Mohr''s salts were investigated. Most strikingly, the color slowly changed from blue to green in an aqueous solution, given by the optical image (Fig. 2a and b), which not only indicated that Ni-CPs captured Mohr''s salts via ion-exchange, but also suggested an alteration in the valence of Fe ions. It was possible that Fe²⁺ may have been further oxidized to Fe³⁺ under the O₂ and water environment when we prepared the Fe@Ni-CPs (4Fe²⁺ + 2H₂O + O₂ = 4Fe³⁺ + 4OH⁻). The PXRD pattern showed that the frameworks remained unchanged after doping with Fe ions (Fig. S1^†). From the transmission electron microscopy (TEM) images (Fig. 2c), after immobilization, the morphology of the Fe@Ni-CPs was still level and smooth; no ring-like patterns arose corresponding to the selected area for electron diffraction (SAED) (Fig. 2d), indicating that no bulk Fe particles formed during ion-exchange. This was further demonstrated using high resolution TEM (HRTEM) (Fig. 2e), in which there was no lattice fringe of crystallized Fe. The well distribution of C, N, O, Cl, Fe, and Ni in Fe@Ni-CPs was demonstrated by elemental mapping (Fig. 2f–l). Energy-dispersive X-ray spectroscopy (EDX) also agreed well with the above mapping data (Fig. S6^†). In addition, the Fe³⁺ uptake was 4.1 wt%, as determined by inductively coupled plasma atomic emission spectroscopy (ICP). These results showed Fe ions to have been successfully immobilized by the Ni-CPs.Open in a separate windowFig. 2(a–b) The optical image of Ni-CPs and Fe@Ni-CPs; (c) TEM images of Fe@ Ni-CPs; (d–e) the corresponding SAED and HRTEM pattern. (f–l) Element mapping of C, N, O, Cl, Fe, and Ni in Fe@Ni-CPs.The X-ray photoelectron spectroscopy (XPS) survey spectrum of the Fe@Ni-CPs also showed the presence of C, N, O, Cl, Fe, and Ni elements (Fig. 3a). The Fe 2p high resolution XPS spectrum exhibited peaks at 725 eV and 711 eV (Fig. 3b), further indicating the presence of the Fe³⁺ oxidation state. This could be explained by the transformations of Fe²⁺ to Fe³⁺ during the ion-exchange process. Similarly, in the Ni 2p spectra (Fig. 3c), two main peaks located at 855.8 eV and 873.5 eV can be ascribed to Ni²⁺ 2p_3/2 and Ni²⁺ 2p_1/2, respectively. These peaks are associated with two shakeup satellite peaks, indicating that Ni still remained in a divalent state. The Cl 2p and N 1s region could be corresponded to the ClO₄⁻ and biprydine, respectively (Fig. S7^†). The O 1s spectrum (Fig. 3d) was divided into two peaks at 531.7 eV and 533.2 eV, which could be assigned to the OH group from filled H₂O molecules and partial ClO₄⁻, respectively.Open in a separate windowFig. 3(a) XPS survey spectrum of the Ni-CPs; XPS spectra of the Ni-CPs in the (b) Fe 2p, (c) Ni 2p, and (d) O 1s regions.The above Ni-CPs with Fe doping encouraged us to investigate its electrocatalytic application in oxygen evolution reaction. To study the electrocatalytic activity of Fe@Ni-CPs for the OER, linear sweep voltammetry (LSV) was performed in a strong alkaline solution (pH = 14) for Fe@Ni-CPs@GC (fresh samples coated on glassy carbon electrode with Nafion binder). Fe@Ni-CPs@GC directly acted as working electrodes and showed good OER activity with an onset potential of 1.52 V (Fig. 4a), overpotential of 368 mV at 10 mA cm⁻², and a Tafel slope of 59.3 mV dec⁻¹ (Fig. 4b). These OER performances are close to some reported MOFs catalysts (Table S2^†) and even better than commercial benchmark OER catalysts like RuO₂ working at the same condition (Fig. S8^†). In contrast, the electrocatalytic OER activities of the pristine Ni-CPs@GC without Fe incorporation displayed much worse activity. The onset potential, overpotential (at 10 mA cm⁻²), and the Tafel slope reached 1.62 V, 458 mV, and 96.8 mV dec⁻¹, respectively (Fig. 4a and b). In addition, Fe@Ni-CPs showed a strong durability during the OER process. The chronoamperometric response of Fe@Ni-CPs displayed a slight anodic current attenuation within 12 h due to the peeling of samples during the evolution of a large amount of O₂ gas (Fig. S9^†). Furthermore, LSV of Fe@Ni-CPs showed negligible changes after OER tests for 12 h (Fig. S10^†). These results indicate that the Fe-doped Ni-CPs with more active sites could serve as an excellent candidate for OER in strong alkaline conditions.Open in a separate windowFig. 4(a) OER polarization curves and (b) Tafel plots of various electrocatalysts in a 1 M KOH aqueous solution; (c) linear relationship of the current density at 1.1 V (vs. RHE) vs. scan rates for Fe@Ni-CPs@GC and Ni-CPs@GC; (d) EIS of Fe@Ni-CPs and Ni-CPs electrode.When Fe(iii) was introduced, the resulting Fe@Ni-CPs catalysts greatly improved OER catalytic performance. There were dynamic collisions between Fe³⁺ ions and Ni-CPs, which allowed for more accessible catalytic active sites compared to the Ni-CPs. Particularly, Fe(iii) doping can contribute to the adsorption and reaction of OH⁻ groups in OER process.²⁹ As a result, Fe@Ni-CPs enhanced charge transfer under an apt electronic environment of the mixed Fe⋯Ni systems. In addition, electrochemical impedance spectrum and double-layer capacitance (C_dl) of the Fe@Ni-CP were also studied. The C_dl of Fe@Ni-CPs was confirmed to be 269.7 μF cm⁻² (Fig. 4c and S11^†), which is higher than that of Ni-CPs (C_dl = 174.2 μF cm⁻²) (Fig. 4c and S12^†). The semicircular diameter in EIS of Fe@Ni-PCP was smaller than that of Ni-CPS (Fig. 4d). These results further showed that Fe@Ni-CPs were more effective in enlarging the catalytically active surface area, conductivity and synergistic effects between Fe and Ni in comparison to Ni-CPs coated on electrodes.In conclusion, a new alkaline-stable cationic Ni(ii) coordinated polymers was synthesized under hydrothermal conditions. The Ni CPs could quickly interact with Mohr''s salt. Interestingly, the Ni CPs could act as a unique oxidation matrix to realize the transformation of Fe²⁺ to Fe³⁺ during the ion-exchange process. Furthermore, the resulting Fe@Ni-CPs electrode, for the first time, showed an excellent electrocatalytic activity for OER in strong alkaline media. This study provides a new avenue to explore stable coordinated polymers by incorporating the low-cost and high-activity transition metal, Fe, which will substitute the rare noble metals used in energy-related research.     

Identification of different carbenium ion intermediates in zeolites with identical chabazite topology via13C–13C through-bond NMR correlations

¹³C–¹³C through-bond NMR correlation experiments reveal the stabilization of different carbenium ion intermediates in two zeolites possessing identical CHA topology (H-SAPO-34 and H-SSZ-13) during the methanol to olefins reaction.

Multidimensional NMR experiments explicitly distinguish differences in carbenium ion intermediates stabilized in two zeolites with identical topology.

The production of light olefins via the methanol-to-olefins (MTO) reaction is an important chemical process that links non-oil resources such as coal and natural gas with olefin-based petrochemicals.^1–5 The catalysts used for the MTO reaction are mainly microporous acidic zeolites amongst which H-SAPO-34, a silicoaluminophosphate zeolite with the chabazite (CHA) topology, is of particular importance due to its high selectivity to ethylene and propene, and is of commercial use.^1,3 H-SSZ-13 is a CHA silicoaluminate analogue of H-SAPO-34 which has been shown to be a potential alternative in the MTO process.⁶Despite the successful industrialization of this process with methanol conversion higher than 99%,³ further improving of the catalyst performances in terms of selectivity has been an important scientific endeavor. For example, the selectivity to ethylene and propene has increased from 79.2% in the first generation industrial DMTO process (“D” refers to Dalian Institute of Chemical Physics) to 85.7% in the second generation DMTO-II process in China.³ However, there is still significant room to improve catalysts performances which can be informed by providing a deeper understanding of the catalytic reaction intermediates and reaction mechanism.We and others have previously investigated the MTO mechanism on H-ZSM-5,^7–10 H-SAPO-34 ^11–14 β zeolite^15,16 using a range of experimental and computational approaches including solid-state Nuclear Magnetic Resonance (NMR). The hydrocarbon pool (HCP) mechanism has been generally accepted for the formation of hydrocarbons from methanol^{5,11,12,17,18} and suggests that for the aromatic cycle routes the organic species (mainly cyclic carbenium ions and neutral aromatic species) confined in the pores of zeolites act as co-catalysts with the inorganic framework. Based on the species observed, a side-chain and a paring reaction pathways have been proposed. While the former involves olefins released through methylation of six-membered ring cations (alkylbenzenium) and elimination of the side chain groups, the later route produces the olefins via expansion of alkylcyclopentenyl cations followed by contraction of the formed six-membered ring cations.^14,15,19 These cyclic carbenium ions are key nodes along the reaction routes and their identification plays an important role in determining the mechanism acting for a given zeolite.Solid-state NMR has played a critical role in achieving this understanding^8,14,20,21 and our recent works^10,16 deploying a range of multidimensional and multinuclear NMR approaches have enabled the unequivocal structural identification of a range of five- and six-membered ring cations (as well as neutral compounds) in H-ZSM-5 and β-zeolites (MFI and BEA topologies, respectively), without the need for previous knowledge or assumption of the carbenium ions structures. In particular, this approach permits to distinguish carbenium ions with closely related structures (e.g., cyclopentenyl cations with various methyl groups) and previously unidentified carbenium ions (e.g., 1,5-dimethyl-3-sec-butyl cyclopentenyl cation and methylnaphthalenium cations), all offering a complete understanding of the reaction routes.Here, we probe the carbenium ions formed during the MTO reaction on two different CHA zeolites (H-SSZ-13 and H-SAPO-34) by utilizing the 2D ¹³C–¹³C refocused INADEQUATE (Incredible Natural Abundance DoublE QUAntum Transfer Experiment)²² NMR experiment (pulse sequence in Fig. S1 in the ESI^†). We experimentally identified 1,2,3,4-tetramethylcyclopentenyl(I) and 1,2,3-trimethylcyclopentenyl(II) cations as the major retained cation species on H-SSZ-13 and H-SAPO-34 CHA zeolites, respectively.Activated zeolites were prepared by flowing ¹³C enriched CH₃OH on H-SSZ-13 and H-SAPO-34 catalyst beds at 275 °C for 25 min and 300 °C for 20 min, respectively, followed by quenching into liquid N₂ to capture the carbenium ion intermediates (full experimental details are given in the ESI^†). The ¹³C CP (Cross Polarisation) MAS (Magic Angle Spinning) NMR spectra of these two activated catalysts are given in Fig. 1 and show multiple signals in the 0–260 ppm region, highlighting the complexity of the retained carbon species. Three general type of species (carbenium ions, aromatics/dienes, adamantane derivatives, see ESI^†) are present, the downfield signals above 150 ppm being characteristic of carbenium ions.^{8,10,14,16,21} It is worth pointing out that the ¹³C spectra of activated H-SSZ-13 and H-SAPO-34 are very similar in this high frequency region; the two main peaks at about 153 and 243 ppm are representative signals for polymethylcyclopentenyl cations.^8,14,20,21 These 1D NMR spectra would suggest that the same cations are retained on these two zeolites with identical CHA topology, however the more informative 2D through-bond ¹³C–¹³C correlation NMR experiment²² reveals that this is not the case (see below).Open in a separate windowFig. 1 ¹³C CP MAS spectra of activated (a) H-SSZ-13 and (b) H-SAPO-34. Spectra were recorded at 9.4 T and at a MAS of 14 kHz. Only characteristic signals for carbenium ions are labelled with chemical shifts. Asterisks (*) denote spinning sidebands (see Fig. S2 and S3^†).The corresponding 2D ¹³C–¹³C refocused INADEQUATE spectra of both activated zeolites are shown in Fig. 2. These experiments are based on through-bond scalar J coupling²² (rather than through-space dipolar-based experiments)^23,24 and the correlation maps directly yield C–C bonds information, unambiguously enabling assignment of the carbon resonances. In this experiment, two peaks resonating at the same frequency in the double quantum (vertical) dimension arise from the sum frequency of the two individual ¹³C peaks in the single quantum (horizontal) dimension that correspond to chemically bonded carbons. The 2D spectra of the two activated zeolites show that the ¹³C–¹³C correlation maps are distinct as evidenced by correlations of the 63 and 243 ppm signals in H-SSZ-13 and of the 47 and 244 ppm peaks in H-SAPO-34 (see red and blue traces in Fig. 2), demonstrating that the main carbenium ions formed are different in both zeolites with the same CHA topology. The full maps are also shown in the 2D spectra and reveal experimental observation of the characteristic correlations C1(I) (243 ppm)–C2(I) (153 ppm) and C2(I) (153 ppm)–C3(I) (251 ppm) identifying cation I as the main retained cation species in H-SSZ-13 (Fig. 2a) and C1,3(II) (244 ppm)–C2(II) (155 ppm) and C1,3(II) (244 ppm)–C4,5(II) (47 ppm) for cation II in H-SAPO-34 (Fig. 2b and S4–S6^† for complete set of correlations for each cation), allowing the structure of the carbenium ions and their ¹³C chemical shifts to be explicitly obtained (Open in a separate windowFig. 22D ¹³C–¹³C refocused INADEQUATE spectra of activated (a) H-SSZ-13 and (b) H-SAPO-34. Data were recorded at 9.4 T and at a MAS rate of 14 kHz. Signals of positive intensities and Fourier Transform (FT) wiggles of negative intensities are coded in black and olive, respectively. Asterisks (*) denote spinning sidebands. The assignments of the different carbenium ions and their corresponding structures are coloured-coded. Some representative traces extracted along the horizontal dimension are also shown. The complete set of traces for carbenium ions I and II is given in Fig. S4–S6.^† Partial correlations for other carbenium ions are coded in maroon and displayed in Fig. S7.^† The correlations coded in green and purple belong to the neutral species (aromatics, dienes and adamantane derivatives) (Fig. S8^†). Signals off the carrier frequency in black dashed box correspond to small artefacts caused by direct current (DC) offset. Numbers in parenthesis are the chemical shifts of the correlated ¹³C sites. ¹³C chemical shifts for carbenium ions I and II (structures are shown in Fig. 2)

Halogen-sensitive solvatochromism based on a phenolic polymer of tetraphenylethene

Herein, we describe the successful preparation of a methylene-bonded tetraphenylethene polymer using a phenolic-resin synthesis protocol. Our novel phenolic polymer showed solvatochromism in response to halogenated organic solvents. Solvatochromism is induced by halogen/π interactions between the polymer and the organic halide.

Herein, we describe novel phenolic polymer showed solvatochromism in response to halogenated organic solvents through halogen/π interactions.

Organic halides, such as chloroform,^1–5 dioxin,⁶ and iodomethane⁷ critically influence biological systems and the environment. In the case of chloroform, inhalation of the vapor can cause cancer. Therefore, for decontamination purposes, halogenated compounds need to be detected, no matter how low the concentration. Existing halogenated compounds have generally been determined using gas chromatography.⁸ In addition, liquid chromatography is a useful method for separating these compounds from pollutants.⁹ However, these methods require pollutant samples to be processed by gel-permeation chromatography and/or silica-gel column chromatography, which can be arduous. Hence, new methodology for the detection of organic halides is required.Recently, chromism has attracted significant attention and has been used to understand a variety of conditions and to engineer products. Numerous types of chromic behavior exist, such as photochromism,¹⁰ mechanochromism,^11,12 electrochromism,¹³ piezochromism,¹⁴ vapochromism,^15–17 and solvatochromism.^18–20 Colors emerging from most forms of chromism are caused by the metamorphoses of compounds with conjugated donor and acceptor units. 1,2-Dithienylethene, which turns red when irradiated with ultraviolet (UV) light, is a famous photochromic material reported by Irie and coworkers.^21–23 The intramolecular conjugation length of this compound increases by a ring-closing photoreaction. Solvatochromism in solvents of different polarity is due to a variation in its absorption or emission spectrum in each solvent. Because organic halide solvents are of low polarity, they can be used to investigate new methods for detecting organic halides because a color transition can occur without an external stimulus, such as light and heat. Previously, yellow-to-red solvatochromism through intramolecular charge transfer in halogenated solvents has been achieved.^24–26The chemical structures of the donor and acceptor units are important when designing and preparing solvatochromic materials. However, it is often difficult to prepare and control the solubilities of these compounds. In addition, crystallinity is an important factor for single-molecule solvatochromism. From these viewpoints, polymer structures are attractive platforms for inducing chromic properties because the solubility of the polymer and its binding sites for organic halides can be controlled. In this study, a novel polymer constructed using 1,1,2,2-tetra(4-methoxyphenyl)ethene (TPE-4MeO) and methylene linkers was prepared by phenolic-polymer synthesis. Phenolic polymers are a class of artificial polymer prepared from phenols and formaldehyde.²⁷ Their high mechanical strengths and chemical and thermal stabilities make these products useful for engineering plastics. Our novel phenolic polymer showed solvatochromism in response to halogenated organic solvents through color changes; in addition, it showed guest and halide selectivity.The phenolic polymer was constructed from TPE units bonded through methylene linkers, and was synthesized according to a previously reported procedure.²⁸ The TPE-based phenolic polymer (TPE-P) was synthesized by polymerizing TPE-4MeO and paraformaldehyde with sulfuric acid in a mixed solution of acetic acid and chloroform (Fig. 1a). The ¹H NMR spectrum exhibited peak broadening with increasing molecular weight, and the ¹³C NMR spectrum exhibited a signal for a methylene unit after polymerization (Fig. S1–S3^†). GPC (Fig. S4^†) and MALDI-Tof-MS of TPE-P revealed: M_n = 1.3 × 10³, M_w = 2.7 × 10³, Đ = 2.1. Other brunching and molecular weights polymers could not be obtained using other substituted TPE monomers, which had its reactive positions blocked by methylene (Fig. S5–S7^†). In addition, a single ¹³C NMR methylene peak was observed, which suggests that TPE-P is composed of a single, linear polymer structure.Open in a separate windowFig. 1(a) Synthesis of TPE-P from TPE-4MeO and paraformaldehyde. (b) Photographic images of solvatochromism in various solvents.The solubilities of TPE-P (40 g L⁻¹) were investigated in chloroform (CHCl₃), dichloromethane (CH₂Cl₂), tetrahydrofuran (THF), and dimethyl sulfoxide (DMSO) (Fig. 1b). TPE-P in the halogenated and non-halogenated solvents exhibited different colors; the former were red but the latter were yellow. Fig. 2a shows the UV-Vis spectra of each TPE-P solution. Although 510 nm absorption peaks were observed in the spectra of TPE-P in the halogenated solvents, such peaks were absent in the spectra in the non-halogen solvents. The observed color generally corresponds to the wavelength of the complementary visible color, as given by the color circle. In addition, the concentration of the visible color is related to the absorption intensity. The halogenated-solvent solutions were red because the 510 nm absorption wavelength is complementary to red. The CH₂Cl₂ solution absorbed more strongly than the CHCl₃, THF + HCl, and THF solutions. Similarly, the color concentration of the CH₂Cl₂ solution was higher. To investigate the effects of halogenated compounds on solvatochromism, 1.1 × 10⁻⁵ mol HCl was added to the solution of TPE-P in THF (Fig. 2a, purple spectrum); this mixture was also red and exhibited a similar absorption peak at 510 nm. Because other halogenated compounds such as carbon tetrachloride, bromoform and bromoethane were also exhibited red, these results suggest that yellow-to-red solvatochromism is induced by halogenated compounds.Open in a separate windowFig. 2(a) UV-Vis absorption spectra of TPE-P in various solvents; 510 nm is the complementary wavelength of visible red. (b) ¹H NMR spectra in various deuterated solvents. The aromatic peak is downfield-shifted in halogenated solvents compared to non-halogenated solvents. (c) FT-IR spectra of TPE-P in CHCl₃ and CH₂Cl₂. (d) UV-Vis absorbance at 510 nm as a function of the chain length of the alkyl halide. (e) DFT ground-state optimized structure of a model TPE-P trimer calculated by B3LYP/3-21G showing the distances between aromatic units, which provide good fits for 1,2-DCE and 1,4-DCB.As we speculated that the solvatochromism behavior of TPE-P requires interaction with a halogenated compound, we next investigated the interactions between halogenated compounds and TPE-P by ¹H NMR spectroscopy. Fig. 2b shows the ¹H NMR spectra of TPE-P in various deuterated solvents. NMR peaks corresponding to aromatic and alkoxy groups were observed in both non-halogenated and halogenated solvents; however, the chemical shifts of these peaks were downfield-shifted only in the halogenated solvents. These downfield shifts are due to lower electron densities in these compounds; hence, these shifts indicate that halogenated compounds and TPE-P interact through halogen/π interactions, which are non-covalent bonds previously reported in computational studies.^29–31 Electrostatic and dispersion energies were reported to be the main contributors to halogen/π interactions between aromatic compounds and halogen atoms. As a result of these halogen/π interactions, TPE-P exhibits solvatochromism through changes in its ground state that results from conformational changes to the polymer structure in the presence of an organic halide. In addition, halogen/π interactions are achieved in aromatic compounds bearing halogen atoms in which the halogen atom binds to the same halogen or carbon. With this in mind, we investigated the solution states by FT-IR spectroscopy, the results of which are shown in Fig. 2c. The IR spectrum of TPE-P shows C–O stretching vibrations at 1250 cm⁻¹ and 1050 cm⁻¹. Both vibrational peaks dramatically shift in dichloromethane and chloroform compared to the non-halogenated solvents. The asymmetric stretching peak at 1250 cm⁻¹ separates into two peaks, and the symmetric stretching peak at 1050 cm⁻¹ shifts to a lower wavenumber. A shift to a lower wavenumber is generally associated with a decrease in electron density. Therefore, the symmetric stretching peak at 1050 cm⁻¹ was low-shifted through interactions with halogenated compounds. Similarly, the asymmetric stretching peak at 1250 cm⁻¹ was also affected by halogenated compounds.In addition, 2D NOESY NMR spectroscopy showed a correlation peak between TPE-P and 1,2-dichloroethane (Fig. S8^†), which indicates that TPE-P solvatochromism is induced through halogen/π interaction with halogenated compounds. On the other hand, we evaluated the effects of free anions as the solvatochromism targets (Fig. S9^†). Various anions, including chloride, bromide, tetrafluoroborate, hypochlorite, and hexafluorophosphate, were dissolved with TPE-P in 1,4-dioxane. These solutions were yellow and did not show any color change. The relationships between dielectric constant, E_T (30), of organic halides and UV-vis absorption maxima were also investigated in Fig. S14.^† Every point was appeared at different wavelength and absorption without systematically lined up. The number of organic halides in TPE-P was different in THF solution because our solvatochromism had dependency on molecular size of organic halides. Moreover, TPE-4MeO did not show any solvent dependency, which indicates that TPE-P can selectively detect organic halides solvatochromatically through halogen/π interactions.To investigate the solvatochromism mechanism, other organic halides were examined by UV-Vis absorption spectroscopy (Fig. S10^†). TPE-P was dissolved in 1,2-dichloroethane (1,2-DCE), 1,3-dichloropropane (1,3-DCP), 1,4-dichlorobutane (1,4-DCB), and 1,5-dichloropentane (1,5-DCP). Not only did the dichloromethane solution exhibit solvatochromism, but the other halogenated solutions did so too. The absorbances of the various halogenated compounds were compared (Fig. 2d); the 510 nm absorption in 1,3-DCP was weaker than in the other halogenated solvents. Similar trends were observed when dibromoalkanes were used as alternative halogenated compounds (Fig. S11^†), which suggests that the halogen/π interactions are influenced by the size of the halogen and/or the compound. To compare the effect of the halogen, we determined the binding constants of 1,2-DCE and 1,2-dibromoethane (1,2-DBE) with TPE-P using Hill''s plots (Fig. S12 and S13^†). The binding constant (K_a) of 1,2-DCE was found to be 1.4 × 10¹ M⁻¹, while that of 1,2-DBE was determined to be 7.1 × 10⁰ M⁻¹; i.e., K_a for 1,2-DCE is 10-times higher than that of 1,2-DCB. Because halogen/π interactions strongly depend on electronegativity, the chloride in 1,2-DCE interacts more strongly with TPE-P.We next determined the distances between TPE-P units by optimizing the geometry of a model trimer by DFT at B3LYP/3-21G. Distances of 46 nm and 73 nm were determined between the aromatic units of TPE-P (Fig. 2e); 1,2-DCE and 1,4-DCB fit well into the spaces created by these distances, respectively. On the other hand, there is insufficient space to accommodate 1,3-DCP. These computational results were confirmed by absorption-intensity data, as shown in Fig. 2d. Halogenated compounds with sizes similar to the above-mentioned distances exhibit higher absorption intensities than compounds of other size. These results indicate that TPE-P shows solvatochromism by interacting with halogen compounds at the unit-molecule level.Finally, we investigated the ability of TPE-P to detect aromatic halides (Fig. 3a). TPE-P solutions in chlorobenzene and bromobenzene were red; this color was also produced using benzene as the solvent, which indicates that this solvatochromism occurs through interactions between TPE-P and the aromatic compounds. Although the benzene solution absorbed less than the other aromatic halides, it nevertheless exhibited solvatochromism (Fig. 3b). In contrast, the 1,4-diiodo-2,3,5,6-tetrafluorobenzene solution was yellow and absorbed poorly, despite being halogenated because the concentration was low due to low solubility in THF. Clearly, dihaloalkane solvatochromatic behavior depends on molecular size, which indicates that TPE-P requires a guest compound with interaction sites and a molecular size matched to its available space to exhibit solvatochromism.Open in a separate windowFig. 3(a) Photographic images of TPE-P solutions in aromatic compounds and (b) effect of aromatic compounds on absorbance. Aromatic compounds exhibited solvatochromism.In conclusion, we successfully prepared a methylene linked tetraphenylethene polymer by phenolic-resin synthesis. The polymer showed solvatochromism and turned red in solutions of halogenated compounds. Experiments with various alkyl halides and DFT calculations at B3LYP/3-21G reveal that halogen/π and π–π interactions between aromatic and halogenated compound induce this phenomenon. Although the polymer was bonded through methylene linkers and was not conjugated, solvatochromism occurs through interactions with guest molecules. Our results suggest that this phenolic polymer may be applied in industrial fields as a sensing, biological, or aerospace material.     

Synthesis of gold nanorod/neodymium oxide yolk/shell composite with plasmon-enhanced near-infrared luminescence

A yolk/shell composite consisting of an AuNR core and an Nd₂O₃ shell with a 19 nm gap is synthesized by a multi-step over-growth method. The near-infrared luminescence of AuNR@Nd₂O₃ is up to 4.6 times higher than that of Nd₂O₃ hollow nanoparticles. The underlying mechanism of plasmon-induced luminescence enhancement is further investigated.

AuNR@Nd₂O₃ yolk/shell nanocomposites are synthesized by a hydrothermal method; the luminescence of Nd³⁺ is enhanced 4.6 times by AuNRs.

Rare earth (RE)-based nanostructures have attracted a lot of attention for their promising applications ranging from photonics to biomedicines.^1–4 The RE-based nanostructures have shown many advantages over the conventional luminescent materials such as semiconductor quantum dots (QDs) and organic dyes, as the luminescence of RE shows high purity, large stock-shifts and excellent stability.^5–7 On the other hand, the RE material is also bio-compatible, which suggests that it has great potential in bio-imaging and therapy.^8–10 Among the lanthanide elements, neodymium (Nd) has drawn a lot of interest for its potentials in sub-tissue imaging and bio-sensing as its luminescence is in the first biological window.^4,11,12 However, the absorption cross-section of Nd is smaller than those of semiconductors or dyes, which seriously affects its fluorescence efficiency, and this prevents its practical applications.^13,14Great efforts are being made to improve the fluorescence of RE,^15–17 in particular the combination between plasmonic noble metal structures and RE is an efficient approach.^18–20 When the light excites noble metal nanostructures, the electron gas collectively oscillates and generates plasmons near the metal surface.^21,22 The large absorption cross-section and strong local electromagnetic field dramatically improve the fluorescence efficiency of the nearby emitters.^23–25 Thus, various hybrids composed of RE nanoparticles and metal nanostructures are designed.^26–31 For instance, J. R. Lakowicz et al. encapsulated lanthanides with silver nanoshells, and the emissions were significantly enhanced by about 10 times.²⁶ For the case of Ag@SiO₂@Y₂O₃:Er synthesized by F. Zhang et al., the up-conversion luminescence (UCL) of Y₂O₃:Er was enhanced 4 times by the inner Ag nanoparticles.²⁷ A. Priyam et al.²⁸ found that the fluorescence of NaYF₄:Yb, Er NPs can be improved by a gold-shell. The metal- and particle-size-dependent enhancements are both investigated.^29,30 In addition, various 3D metamaterials and photonic crystals have been designed to adjust or enhance the emissions of RE.³¹The luminescence of RE can be improved by the coupled plasmons, because the plasmons provide strong electromagnetic field to enhance the excitation/emission process; also, there may be energy transfer between the plasmons and emitters.^32,33 Gold nanorod (AuNR) is a typical plasmonic structure used to enhance the fluorescence of emitters, and it has tunable longitudinal surface plasmon resonance (LSPR) ranging from visible to near-infrared.^34–36 Since the excitation and emission frequencies of Nd are both in the near-infrared region, it is possible to design a resonance structure between the AuNR and Nd structures to achieve luminescence enhancement of Nd³⁺.³⁷ For obtaining luminescence enhancement, isolation between the plasmonic structure and the emitters is very important; silica, alumina, polymers, or DNAs have been applied to adjust the distance to obtain the largest enhancement.^27,38,39 However, there are a few reports on the structure consisting of a plasmonic core and an RE shell with a natural isolation layer; the plasmon-induced RE down-conversion luminescence enhancement is also not a popular topic.^40–42 In this study, we developed a facile method to prepare AuNR@Nd₂O₃ yolk/shell composites containing a 19 nm gap between an AuNR core and an Nd₂O₃ shell. The effects of AuNRs on the down-conversion luminescence (DCL) properties of Nd₂O₃ shells were studied by comparing the luminescence intensities of the AuNR@Nd₂O₃ yolk/shell composites and the corresponding Nd₂O₃ hollow nanoparticles. It was found that the 873 nm emission of Nd³⁺ was enhanced by AuNRs up to 4.6 times. The LSPR-dependent enhancement was also investigated further.Cetyl-trimethyl ammonium bromide (CTAB)-capped AuNRs were first synthesized using the seed-mediated growth method.^43,44 Then, CTAB was replaced by oleate with a ligand exchange approach.⁴² Au@Nd₂O₃ yolk/shell composites were prepared by an oleate-assisted hydrothermal method. In brief, for the 1^st step of the growth procedure, 5 mL aqueous solution of oleate-AuNRs was diluted with 14 mL of ultrapure water. Nd(NO₃)₃ and HMT solutions were injected with stirring to form a well-dispersed solution; the mixture was incubated at 85 °C for 3 h, in which Nd(NO₃)₃ and HMT served as the cation and anion reagents, respectively. Then, the resultant solution was centrifuged; the precipitate was re-dispersed in ultrapure water, and it was used as seeds in the next step. This process was repeated three times to achieve the final yolk/shell composites (Fig. 1(a)). The transmission electron microscopy (TEM) images in Fig. 1(b–e) indicate the morphology evolution in the whole growth process. The length and diameter of the original AuNRs were about 60 nm and 15 nm, respectively, thus suggesting an aspect ratio of 4 (Fig. ESI-1a^†). After the 3 hour 1^st step growth, the Nd₂O₃ nanoparticles loosely surrounded AuNRs (Fig. 1(b)). In the 2^nd growth step, the outer Nd₂O₃ shell became thicker and more compact (Fig. 1(c)). A gap formed between the AuNR core and the Nd₂O₃ shell, and the thickness of the Nd₂O₃ shell decreased from 28 nm to 16 nm after the 3^rd growth step (Fig. 1(d)), which was probably due to the Ostwald ripening.^45,46 When the 4^th growth step was completed, the final products were collected. Fig. 1(e) and Fig. ESI-1b^† demonstrate that the as-prepared hybrids were monodispersed hollow quasi-spheres consisting of AuNR cores and Nd₂O₃ shells. Interestingly, AuNRs were completely separated from Nd₂O₃, and the gap was about 19 nm. During the growth process, LSPR of AuNRs gradually red-shifted from 761 nm to 808 nm as the surrounding Nd₂O₃ increased the refractive index (Fig. 1(f)).⁴⁷ The energy-dispersive X-ray (EDX) spectrum in Fig. 1(g) displays the relative element contents of the final AuNR@Nd₂O₃ composites; the Au/Nd ratio was about 4 : 6.Open in a separate windowFig. 1(a) Schematic illustration of the growth procedure of the AuNR@Nd₂O₃ yolk/shell composites. (b–e) TEM images of AuNR@Nd₂O₃ composites obtained at different growth steps. (f) Normalized extinction spectra of AuNRs and AuNR@Nd₂O₃ composites obtained at different growth steps. (g) EDX spectrum of the final AuNR@Nd₂O₃ yolk/shell composites.To reveal the plasmonic effect on the luminescence of Nd₂O₃, iodide/triiodide redox couple is used to corrode the inner AuNRs and to achieve the Nd₂O₃ hollow nanoparticles (Fig. 2(a)). Fig. 2(b–e) show the morphology evolution against the etching time. After etching for 1 h, AuNRs decrease in size and after 2 h, they transform into quasi-spheres about 10 nm in diameter. Then, AuNRs transform into 2–3 nm spheres after 3 h and finally disappear after 4 h. The completely etched Nd₂O₃ nanoparticles are uniform hollow spheres with inner cavities (Fig. 2(e) and Fig. ESI-1c^†). With respect to size distributions (insets of Fig. ESI-1b and c^†), the average diameters of AuNR@Nd₂O₃ composites and Nd₂O₃ hollow nanoparticles are similar, thus indicating that the iodide/triiodide electrolyte has not destroyed the Nd₂O₃ structure.⁴⁸ The extinction spectra obtained at different etching times are displayed in Fig. 2(f). The LSPR location is blue-shifted, and the intensity is systematically decreased. When AuNRs are completely etched, the absorption of AuNRs completely disappears. As shown in Fig. 2(g), the corresponding EDX spectrum also demonstrates that AuNRs have been etched thoroughly.Open in a separate windowFig. 2(a) Schematic illustration of the etching formation process of the Nd₂O₃ hollow nanoparticles. (b–e) TEM images of the as-synthesised AuNR@Nd₂O₃ nanocomposites after (b) 1 h, (c) 2 h, (d) 3 h and (e) 4 h of etching. (f) Time evolution of the extinction spectra during the etching process. (g) EDX spectrum of the Nd₂O₃ hollow nanoparticles.The plasmon-enhanced near-infrared luminescence of AuNR@Nd₂O₃ is further studied by comparing the composites'' emission spectra before and after etching (Fig. 3(a)). For the measurement, a 730 nm continuous wave laser is used for excitation; the luminescence spectra are recorded by a spectrometer with a liquid nitrogen-cooled CCD. The emission bands for the ⁴F_3/2–⁴I_9/2 transition of Nd³⁺ are located at 873 nm. As AuNRs are corroded, the emission peak position and shape remain unchanged. To calculate the enhancement factors of 873 nm emission against the etching time, the emission spectra are decomposed by Gaussian fitting (Fig. ESI-2^†); the emission intensities at 873 nm are displayed in Fig. 3(b). The corresponding enhancement factors are also calculated in Fig. 3(b); the largest enhancement of about 3.5 is obtained at the beginning. As the AuNRs are corroded, the emission intensity is decreased. In another words, the results suggest that as the Au content increases, the luminescence becomes brighter.Open in a separate windowFig. 3(a) Emission spectra of AuNR@Nd₂O₃ composites during the etching process. (b) Emission enhancement factors and the emission intensities of 873 nm against the etching time.To further reveal the relationship between plasmon wavelength and luminescence enhancement, we prepared three kinds of AuNR@Nd₂O₃ with LSPRs at 740 nm (pink), 808 nm (green) and 880 nm (gray) (Fig. 4(a)). The emission spectra of the AuNR@Nd₂O₃ composites before and after etching were measured, and the enhancement factors at 873 nm emission are calculated in Fig. 4(b). The enhancement factors of Nd³⁺ at 873 nm were 4.6, 3.5 and 2.7 for the AuNR@Nd₂O₃ samples with LSPRs at 740 nm, 808 nm and 880 nm, respectively. When the LSPRs wavelength of AuNR@Nd₂O₃ is closer to the excitation wavelength, the enhancement factor was larger. The schematic of the plasmon-enhanced luminescence is shown in Fig. 5. Due to the gap between AuNR and Nd₂O₃, the non-radiative energy transfer from Nd₂O₃ to AuNR was negligible (dashed arrows).⁴⁹ As the incident light excited AuNR and Nd₂O₃, the strong local electromagnetic field around AuNR enhanced the excitation and the emission processes of the Nd₂O₃ shells. There may be a resonance energy transfer between the AuNR and the Nd₂O₃, which was favorable for the improvement in Nd³⁺ excitation efficiency. As the LSPRs of AuNR@Nd₂O₃ were tuned from 740 nm to 880 nm, the enhancement factors decreased. All the results suggested that plasmons influenced the excitation process more efficiently than the emission process.Open in a separate windowFig. 4(a) Extinction spectra of AuNR@Nd₂O₃ yolk/shell composites (solid lines) and the corresponding AuNRs (dashed lines) with different LSPRs. (b) Enhancement factors at 873 nm of AuNR@Nd₂O₃ with different LSPRs.Open in a separate windowFig. 5Schematic mechanism of the energy transfer between an AuNR core and an Nd₂O₃ shell.In summary, AuNR@Nd₂O₃ yolk/shell composites have been synthesized using a simple hydrothermal method. The reference samples (Nd₂O₃ hollow nanoparticles) are prepared by etching Au cores. The plasmon enhancement factor is also proved to be highly LSPR-dependent; the highest enhancement factor up to 4.6 is obtained when LSPR is resonant with the excitation. These findings provide a general pathway to modulate DCL of RE materials. The composites can be applied in photonics, bio-imaging, energy conversion and other fields.