Photoelectrochemical photocurrent switching effect on a pristine anodized Ti/TiO2 system as a platform for chemical logic devices

We report here the effect of the photoelectrochemical photocurrent switching (PEPS) observed on highly-ordered pristine anodized Ti/TiO₂ for the first time. At negative potential bias, blue irradiation gives cathodic photocurrent, whereas anodic photocurrent was observed for ultraviolet irradiation. We believe this phenomenon is due to the electron pathway provided by Ti³⁺ defect states.

We report here the effect of the photoelectrochemical photocurrent switching (PEPS) observed on highly-ordered pristine anodized Ti/TiO₂ for the first time.

Titanium dioxide, being one of the most studied materials, still draws much attention from researchers.^1,2 It is considered to be a very promising material due to its high chemical stability, nontoxicity, and its unique properties. Due to stable and robust photoactivity, titania is widely used in the design of solar cells³ and photocatalytic applications.⁴ In addition to the fact that titanium dioxide occurs in several crystalline modifications, it can also be obtained in various forms, such as, for example, nanotubes,⁵ nanofibers,⁶ and nanosheets.⁷ The photocatalytic performance of TiO₂ is highly dependent on crystallinity,⁸ phase content, form, and preparation method.⁹ It was reported that highly ordered arrays of TiO₂ nanotubes are characterized by short charge transport distance and little carrier transport loss.⁵ Therefore, electrochemically fabricated TiO₂ nanotube arrays are preferable compared to random non-oriented titania.¹⁰ Great varieties of photoelectrochemical behaviour can be achieved by doping¹¹ and surface modification.^12,13An interesting feature has recently been demonstrated for highly ordered arrays of TiO₂ nanotubes obtained by double stepwise electrochemical anodization of a titanium foil (Ti/TiO₂). Together with our colleagues observed that localized illumination of Ti/TiO₂ surface in water solution triggers proton flux from irradiated area.¹⁴ The photocatalytic activity of TiO₂ is based on photogenerated electron–hole pairs. Under the electric field of Ti/TiO₂ Schottky junction and due to upward surface band bending, efficient spatial charge separation occurs, and photoexcited holes (h⁺) reach TiO₂ – solution interface. The h⁺, which is a strong oxidizing agent, can react with water, and a pronounced pH gradient arises due to water photolysis. Thus, titanium dioxide can be used to trigger local ion fluxes, and proton release is associated with anodic photocurrent. The use of the light-pH coupling effect to control pH-sensitive soft matter was previously demonstrated.^15,16 Complementary species, H⁺ and OH⁻, annihilating when occurring simultaneously, extend chemical arithmetic with subtraction operation opening way to pure chemical calculations.¹⁷ Ion fluxes consideration as information transducers in solution were proposed¹⁸ and performing simple logic operations was demonstrated.¹⁹ This phenomenon opens perspectives to biomimetic information processing and developing effective human–machine interfaces.²⁰Photoelectrodes using light and potential as inputs and yielding photocurrents are being considered as the basis for logic devices. In this way, optical computing compatible with existing silicon-based devices may be performed.Logic operations are described by Boolean algebra operating with truth values denoted 0 (false) and 1 (true). Elementary logical operations are modelled by logic gates producing single binary output from multiple binary inputs and physically implemented by some switch. As for photoelectrode based information processing, the photoelectrochemical photocurrent switching (PEPS) effect is utilized. This effect is that under appropriate external polarization or/and illumination by light with appropriate photon energy, switching between anodic and cathodic photocurrent may be observed for n-type semiconductors and the opposite for p-type.^21,22Without further modification, this effect was observed for a very limited number of materials, such as bismuth orthovanadate, lead molybdate, V–VI–VII semiconductors, and some others. To show this effect, the majority of semiconductors require electronic structure perturbation creating new electron pathways. A convenient solution is specific modifier adsorption onto the semiconductors'' surface, providing a sufficient level of electronic coupling. Photoelectrodes made of nanocrystalline TiO₂ modified by cyanoferrate,^13,23 and ruthenium²⁴ complexes, thiamine, folic acid,²⁵ and carminic acid²⁶ demonstrated PEPS behavior.Surprisingly, we observed the PEPS effect on non-modified Ti/TiO₂ obtained by anodation of Ti plates.Highly ordered arrays of anatase Ti/TiO₂ were obtained. Crystallinity was proved by XRD (Fig. S1a^†). Fig. 1a shows a SEM image of TiO₂ nanotube arrays obtained as described above. According to SEM image, an average pore diameter is ca. 60 nm. As reported, highly ordered TiO₂ nanotubes possess a short charge transport distance and little carrier transport loss. Therefore, highly ordered TiO₂ nanotube arrays fabricated by electrochemical anodization of titanium may exhibit some enhanced capacity of electron transfer than non-oriented ones of random mixture.¹⁰Open in a separate windowFig. 1(a) SEM image of the TiO₂ nanotubes array. The inset shows cross-section view. (b) Scheme of a cell for photocurrent measurements experiment, CE – counter electrode, RE – reference electrode, WE – working electrode.According to Mott–Schottky analysis, at potential bias more positive than −0,697 V vs. Ag/AgCl reference electrode upwards band bending occurs (Fig. S2^†). Heat treatment in a nonoxidizing atmosphere leads to Ti³⁺ formation. Appearance of Ti³⁺ self-doping was proved by EDX analysis (Fig. S1b^†). It was previously reported that Ti³⁺ introduces gap states which act as recombination centers and pathways for electron transfer.^27–29 Ti³⁺ species in reduced TiO₂ introduce a gap state between valence and conduction bands.^27,28We studied dependence of photocurrent on applied potential. Ultraviolet irradiation (365 nm) gave positive photocurrent for all potentials studied in range from −0.6 V to 0.6 V vs. Ag/AgCl reference electrode (Fig. S3^†). The photocurrent increases as the potential becomes more positive, but eventually saturates. The dependence of the current on the potential under blue irradiation (405 nm) had a different character. Sigmoid function with inflection point at 0–0.2 V was observed for blue light.It should be noticed that photocurrent plotted against time on Fig. 2–4 as well as against potential on Fig. S3^† is ΔI = I_{under illumination} − I_{in darkness}. Steady state current values were used for calculations.Open in a separate windowFig. 2Photocurrent curves under chopped irradiation by (a) 365 nm UV LED, (c) 405 nm blue LED at applied potential bias +300 mV vs. Ag/AgCl, and corresponding scheme of electron pathway at +300 mV polarization under irradiation by (b) 365 nm UV LED and (d) 405 nm blue LED.Open in a separate windowFig. 3Photocurrent curves under chopped irradiation by (a) 365 nm UV LED, (c) 405 nm blue LED at applied potential bias −300 mV vs. Ag/AgCl, and corresponding scheme of electron pathway at −300 mV polarization under irradiation by (b) 365 nm UV LED and (d) 405 nm blue LED.Open in a separate windowFig. 4(a) XOR logic realized on negatively polarized (−0.3 V) pristine Ti/TiO₂ by two source irradiation, input A – UV light (365 nm), input B – blue light (405 nm); blue light gives anodic photocurrent, UV – cathodic photocurrent. The current, significantly different from the dark one, is taken as output 1, otherwise – 0. When irradiated by blue and UV light simultaneously, anodic and cathodic current compensate each other, and no total photocurrent observed. Thus output 0, when both inputs are 1 (b) OR logic realized on positively polarized (+0.3 V) non-modified Ti/TiO₂ by two sources of irradiation. Irradiation by any of them, blue or UV, gives anodic photocurrent.At +300 mV vs. Ag/AgCl irradiation by both blue and ultraviolet light give anodic photocurrent (Fig. 2a and c). The UV-irradiation (λ = 365 nm, 5 mW cm⁻²) excites electron directly to the conduction band (CB) of TiO_2, which is further transferred to conducting titanium support (Fig. 2b). When Ti/TiO₂ electrode in thermodynamic equilibrium with electrolyte, an upward surface band bending occurs at the semiconductor–liquid junction. This phenomenon obstructs electron injection from the conduction band into the electrolyte and forces electron drift to conducting substrate. The fast and steady photocurrent production/extinction upon light on/off indicates efficient charge separation and low recombination.Blue light (λ = 405 nm, 70 mW cm⁻²) is characterized by lower energy than UV-irradiation, which is not sufficient to excite the electron to CB. But electron excited by blue light can be trapped by Ti³⁺ located close to the conduction band and transferred to conduction support from these levels (Fig. 2c). An initial current spike following by an exponential decrease suggesting a fast recombination process. It should be also noticed than when irradiation is switched off photocurrent ‘overshoots’ as the remaining surface holes continue to recombine with electrons.At more negative potential (−300 mV vs. Ag/AgCl, for example) applied to non-modified anodized Ti/TiO₂ photoelectrode, we observed anodic photocurrent during irradiation by UV light (Fig. 3a) whereas blue irradiation gave anodic photocurrent (Fig. 3c). Excitation within bandgap by UV-irradiation leads to cathodic photocurrent (Fig. 3b). In the case of irradiation by blue light, electron trapping by Ti³⁺ occurs in the same manner as at +300 mV polarization. But at negative polarization, the energy landscape is such that electron transport to electron donor in solution is preferable (Fig. 3d). As a result, cathodic current occurs.Thereby, photoelectrode activity of non-modified anodized Ti/TiO₂ can be switched from anodic to cathodic and vice versa by applying various potentials and various photon energies. This is the effect of photoelectrochemical photocurrent switching.Thereby, when Ti/TiO₂ is irradiated simultaneously by blue and UV light being negatively polarized, competition between cathodic and anodic photocurrents occurs. Returning to Boolean logic, the PEPS effect allows us to perform annihilation of two input signals and implement optoelectronic XOR logic gate. XOR logic operation outputs true (1) only when input values are different and yield zero otherwise.It is necessary to assign logic values to input and output signals to analyse the system based on Ti/TiO₂ PEPS effect in terms of Boolean logic. Logical 0 and 1 are assigned to off and on states of the LEDs, respectively. Different wavelengths (365 and 405 nm) correspond to two different inputs of the logic gate. In the same way, we can assign logic 0 to the state when photocurrent is not generated and logic 1 to any nonzero photocurrent intensity irrespectively on its polarization (cathodic or anodic). Fig. 4 demonstrates how different types of Boolean logic are realized by irradiation of Ti/TiO₂. Light sources are denoted here as inputs, UV light – A and blue light – B. If the corresponding light source is switched ON and illuminates photoelectrode Ti/TiO₂, this input is ‘1’, otherwise, it''s ‘0’. The photocurrent is read as output. It''s considered to be ‘1’ if significantly differs from dark value and ‘0’ otherwise.At −300 mV vs. Ag/AgCl, pulsed irradiation with UV diode (365 nm, 5 mW cm⁻²) results in anodic photocurrent, which is consistent with electron excitation to CB and transfer to conducting support. Irradiation with blue LED (405 nm) gives cathodic photocurrent due to electron capture by Ti³⁺ states following by transferring to electron acceptor in solution. Simultaneous irradiation with two LEDs with adjusted intensity yields zero net current as anodic and cathodic photocurrents compensate effectively (Fig. 4a).At positive potentials, pulsed irradiation with UV diode gives anodic photocurrent pulses, as well as the blue one. It is interesting to note that when two sources of light are simultaneously irradiated, the photocurrents created by each of them individually do not summarize. At +300 mV, photocurrent output under the influence of two light inputs (365 nm and 405 nm) follows OR logic giving positive output if at least one of inputs is positive (Fig. 4b). Fig. 5 demonstrates the reconfigurable logic system which characteristics can be changed via an appropriate polarization of the photoelectrode regarded as programming input. Two irradiation sources are considered as inputs. OR/XOR logic is realized depending on programming input.Open in a separate windowFig. 5A reconfigurable logic system based on non-modified Ti/TiO₂. Light sources are inputs. The choice between XOR and OR function is determined by programming input of potential bias. At +300 mV OR logic is realized, at −300 – XOR logic. Corresponding truth table is presented.In summary, PEPS effect on modified nanocrystalline TiO₂ was previously discussed a lot.^13,23–26 In this work we report the same phenomenon for pristine anodized Ti/TiO₂ system. Due to substructure of Ti/TiO₂ system, it shows characteristic response to various range of illumination, including visible range and polarization. The Ti/TiO₂ system is a simple and robust model of chemical logic gates. Suggested mimicking of logic functions in aqueous solutions allows further integration of element into communication with living objects¹⁶vs. intrinsically associated photooxidation and degradation, but rather activation for needed function.³⁰     

Insight into trifluoromethylation – experimental electron density for Togni reagent I

3,3-Dimethyl-1-(trifluoromethyl)-1,3-dihydro-1-λ3,2-benziodoxole represents a popular reagent for trifluoromethylation. The σ hole on the hypervalent iodine atom in this “Togni reagent” is crucial for adduct formation between the reagent and a nucleophilic substrate. The electronic situation may be probed by high resolution X-ray diffraction: the experimental charge density thus derived shows that the short intermolecular contact of 3.0 Å between the iodine and a neighbouring oxygen atom is associated with a local charge depletion on the heavy halogen in the direction of the nucleophile and visible polarization of the O valence shell towards the iodine atom. In agreement with the expectation for λ3-iodanes, the intermolecular O⋯I–C_aryl halogen bond deviates significantly from linearity.

The experimentally observed electron density for the “Togni reagent” explains the interaction of the hypervalent iodine atom with a nucleophile.

A “halogen bond” denotes a short contact between a nucleophile acting as electron density donor and a (mostly heavy) halogen atom as electrophile;^1,2 halogen bonds are a special of σ hole interactions.^3–5 Such interactions do not only play an important role in crystal engineering;^6–11 rather, the concept of a nucleophile approaching the σ hole of a neighbouring atom may also prove helpful for understanding chemical reactivity.The title compound provides an example for such a σ hole based reactivity: 3,3-dimethyl-1-(trifluoromethyl)-1,3-dihydro-1-λ3,2-benziodoxole, 1, (Scheme 1) commonly known as “Togni reagent I”, is used for the electrophilic transfer of a trifluoromethyl group by reductive elimination. The original articles in which the application of 1¹² and other closely related “Togni reagents”¹³ were communicated have been and still are highly cited. Trifluoromethylation is not the only application for hypervalent iodine compounds; they have also been used as alkynylating¹⁴ or azide transfer reagents.^15,16 The syntheses of hypervalent iodanes and their application in organofluorine chemistry have been reviewed,¹⁷ and a special issues of the Journal of Organic Chemistry has been dedicated to Hypervalent Iodine Catalysis and Reagents.¹⁸ Recently, Pietrasiak and Togni have expanded the concept of hypervalent reagents to tellurium.¹⁹Open in a separate windowScheme 1Lewis structure of the Togni reagent, 1.Results from theory link σ hole interactions and chemical properties and indicate that the electron density distribution associated with the hypervalent iodine atom in 1 is essential for the reactivity of the molecule in trifluoromethylation.²⁰ Lüthi and coworkers have studied solvent effects and shown that activation entropy and volume play relevant roles for assigning the correct reaction mechanism to trifluoromethylation via1. These authors have confirmed the dominant role of reductive elimination and hence the relevance of the σ hole interaction for the reactivity of 1 in solution by ab initio molecular dynamics (AIMD) simulations.^21,22 An experimental approach to the electron density may complement theoretical calculations: low temperature X-ray diffraction data of sufficient resolution allow to obtain the experimental charge density and associate it with intra- and inter-molecular interactions.^23–25 Such advanced structure models based on aspherical scattering factors have also been applied in the study of halogen bonds.^26–30 In this contribution, we provide direct experimental information for the electronic situation in Togni reagent I, 1; in particular, we analyze the charge distribution around the hypervalent iodine atom.Excellent single crystals of the title compound were grown by sublimation.^§ The so-called independent atom model (IAM), i.e. the structure model based on conventional spherical scattering factors for neutral atoms, confirms the solid state structure reported by the original authors,¹² albeit with increased accuracy. As depicted in Fig. 1, two molecules of 1 interact via a crystallographic center of inversion. The pair of short intermolecular O⋯I contacts thus generated can be perceived as red areas on the interaction-sensitive Hirshfeld surface.³¹Open in a separate windowFig. 1Two neighbouring molecules of 1, related by a crystallographic center of inversion. The short intermolecular I⋯O contacts show up in red on the Hirshfeld surface³² enclosing the left molecule. (90% probability ellipsoids, H atoms omitted, symmetry operator 1 − x, 1 − y, 1 − z).The high resolution of our diffraction data ^¶ for 1 allowed an atom-centered multipole refinement^33,34 and thus an improved model for the experimental electron density which takes features of chemical bonding and lone pairs into account. Fig. 2 shows the deformation density, i.e. the difference electron density between this advanced multipole model and the IAM in the same orientation as Fig. 1.^|| The orientation of an oxygen lone pair (blue arrow) pointing towards the σ hole of the heavy halogen in the inversion-related molecule and the region of positive charge at this iodine atom (red arrow) are clearly visible. Single-bonded terminal halides are associated with one σ hole opposite to the only σ bond, thus resulting in a linear arrangement about the halogen atom. Different geometries and potentially more than a single σ hole are to be expected for λ3-iodanes such as our target molecule, and as a tendency, the resulting halogen bonds are expected to be weaker than those subtended by single-bonded iodine atoms.³⁵ In agreement with these theoretical considerations, the closest I⋯O contacts in 1 amount to 2.9822(9) Å. This distance is significantly shorter than the sum of the van-der-Waals radii (I, 1.98 Å; O, 1.52 Å (ref. 36)) but cannot compete with the shortest halogen bonds between iodine and oxygen^37,38 or iodine and nitrogen.^39–41Fig. 1 and ​and22 show that the C_aryl–I⋯O contacts are not linear; they subtend an angle C10–I1⋯O1′ of 141.23(3)° at the iodine atom. On the basis of theoretical calculations, Kirshenboim and Kozuch³⁵ have suggested that the split σ holes should be situated in the plane of the three substituents of the hypervalent atom and that halogen and covalent bonds should be coplanar. Fig. 3 shows that the C_aryl–I⋯O interaction in 1 closely matches this expectation, with the next oxygen neighbour O1′ only 0.47 Å out of the least-squares plane through the heavy halogen I1 and its three covalently bonded partners C1, O1 and C10.Open in a separate windowFig. 2Deformation density for the pair of neighbouring molecules in 1; the dashed blue and red arrows indicate regions of opposite charge. (Contour interval 0.10e Å⁻³; blue lines positive, red lines negative, green lines zero contours, symmetry operator 1 − x, 1 − y, 1 − z).Open in a separate windowFig. 3A molecule of 1, shown [platon] along O1⋯C1, and its halogen-bonded neighbour O1ⁱ. Symmetry operator 1 − x, 1 − y, 1 − z.The Laplacian, the scalar derivative of the gradient vector field of the electron density, emphasizes local charge accumulations and depletions and it allows to assess the character of intra- and inter-molecular interactions. A detailed analysis of all bonds in 1 according to Bader''s Atoms In Molecules theory⁴² is provided in the ESI.^‡ We here only mention that the electron density in the bond critical point (bcp) of the short intermolecular I⋯O contact amounts to 0.102(5)e Å⁻³; we are not aware of charge density studies on λ3-iodanes, but both the electron density and its small positive Laplacian match values experimentally observed for halogen bonds involving O and terminal I in the same distance range.⁴³The crystal structure of 1 necessarily implies additional contacts beyond the short halogen bond shown in Fig. 1 and ​and2.2. The shortest among these secondary interactions is depicted in Fig. 4: it involved a non-classical C–H⋯F contact with a H⋯F distance of 2.55 Å.Open in a separate windowFig. 4C–H⋯F contact in 1; additional information has been compiled in the ESI.^‡ Symmetry operators i = 1 − x, 1 − y, 1 − z; ii = 1 − x, 1 − y, −z.The topological analysis of the experimental charge density reveals that this non-classical C–H⋯F hydrogen bond and all other secondary contacts are only associated with very small electron densities in the bcps. Table S8 in the ESI^‡ provides a summary of this analysis and confirms that the I⋯O halogen bond discussed in Fig. 1 and ​and22 represents by far the most relevant intermolecular interaction.The relevance of this halogen bond extends beyond the crystal structure of 1: Insight into the spatial disposition of electrophilic and nucleophilic regions and hence into the expected reactivity of a molecule may be gained from another electron-density derived property, the electrostatic potential (ESP). The ESP for the pair of interacting molecules in 1 is depicted in Fig. 5.Open in a separate windowFig. 5Electrostatic potential for a pair of molecules in 1 mapped on an electron density isosurface (ρ = 0.5e Å⁻³; program MoleCoolQt^44,45). Fig. 5 underlines the complementary electrostatic interactions between the positively charged iodine and the negatively charged oxygen atoms. One can easily imagine to “replace” the inversion-related partner molecule in crystalline 1 by an incoming nucleophile.The ESP tentatively obtained for a single molecule in the structure of 1 did not differ significantly from that derived for the inversion-related pair (Fig. 5), and even the results from theoretical calculations in the gas phase for an isolated molecule²⁰ are in good qualitative agreement with our ESP derived from the crystal structure. In the absence of very short contacts, polarization by neighbouring molecules only has a minor influence on the ESP. The experimentally observed electron density matches the proven reactivity for the title compound, and we consider it rewarding to extend our charge density studies on related hypervalent reagents.     

Proton transfer network with luminescence display controls OFF/ON catalysis that generates a high-speed slider-on-deck

A three-component network for OFF/ON catalysis was built from a protonated nanoswitch and a luminophore. Its activation by addition of silver(i) triggered the proton-catalyzed formation of a biped and the assembly of a fast slider-on-deck (k₂₉₈ = 540 kHz).

Upon addition/removal of silver(i) ions and due to efficient inter-component communication, a supramolecular multicomponent network acts as an OFF/ON proton relay with luminescence display enabling switchable catalysis.

Proton relays in living systems¹ are often coupled to bio-machinery with enzymatic catalysis.² Frequently, such proton transfer systems comprise intricate networks with information exchange^3,4 amongst multiple components. Using manmade molecules,⁵ however, it is still arduous to create networks that approach the complexity⁶ and functions of biosystems.Although activation via proton transfer plays an eminent role in molecular logic,⁷ switching,^8–10 and machine operation,^11–15 proton transfer networks in combination with the regulation of catalysis remain scarce.^16,17 Beyond that, we present the in situ catalytic synthesis of a highly dynamic supramolecular device by a new proton transfer network requiring interference-free information handling,¹⁸ self-sorting^19–21 and communication²² among multiple components.^23–25Recently, our group has developed various chemically actuated networks of communicating components and used them for the ON/OFF regulation of transition metal catalysis.^26–28 Herein, we describe a proton relay system that reversibly toggles between two networked states and is useful for regulation of acid catalysis. It catalyzes an amine deprotection which was utilized to fabricate in situ a high-speed slider-on-deck,^29,30 for the first time based on aniline → zinc-porphyrin³¹ (N_an → ZnPor) interactions.The communicating network consists of nanoswitch 1 (ref. 32) and luminophore 2 (ref. 33) as receptors (Fig. 1). In the self-sorted networked state (NetState-I), the added proton is captured in the cavity of nanoswitch 1 resulting in [H(1)]⁺, while luminophore 2 remains free. NetState-I is catalytically inactive as the encapsulated proton is unable to catalyze the deprotection of the trityl protected slider biped 4. Upon addition of AgBF₄ to NetState-I, the silver(i) coordinates to switch 1 and the proton is released onto luminophore 2 thus leading to formation of [Ag(1)]⁺ and [H(2)]⁺ in NetState-II (see Scheme 1), which may be followed by characteristic ratiometric emission changes. If NetState-II is formed in presence of 3 and 4, the catalytic trityl deprotection of 4 produces the free biped 5 that eventually binds to deck 3 furnishing the slider-on-deck [(3)(5)].Open in a separate windowFig. 1Chemical structures and cartoon representations of the ligands used in the study.Open in a separate windowScheme 1Schematic representation of the reversible proton relay network. Green circle represents catalytic site.The concept of the catalyzed biped formation through a proton transfer network relies on the following considerations: (a) the initial networked state should be catalytically inactive as expected from a self-sorting of the proton inside the nanoswitch leaving the luminophore free. (b) Addition of silver(i) ions should release the proton from the protonated switch thus forming the protonated luminophore which we anticipated to act as catalyst. (c) A di-protected biped should undergo acid catalyzed deprotection by the protonated luminophore and eventually form a slider-on-deck by coordination to a deck. (d) Finally, the product of the catalysis should not intervene in any state of the operation with the main system.As proton receptor we chose the known phenanthroline–terpyridine nanoswitch 1 (ref. 32) in order to tightly capture the proton/silver(i) as HETTAP (HETeroleptic Terpyridine And Phenanthroline) complexes.³⁴ Luminophore 2 was based on the 2,4,6-triarylpyridine scaffold which is known to fluoresce strongly upon protonation due to intramolecular charge transfer (ICT).³⁵To investigate the self-sorting within the network, we mixed nanoswitch 1, luminophore 2 and TFA in 1 : 1 : 1 ratio in CD₂Cl₂ which furnished NetState-I = [H(1)]⁺ + 2. ¹H NMR shifts of the 4/7-H proton signals to 8.43 and 8.69 ppm and of 9-H and e-H proton peaks to 6.60 and 8.48 ppm confirmed formation of [H(1)]⁺. Moreover, proton signals of c′-H and d′-H at 7.07 and 7.92 ppm corroborated presence of the free luminophore 2 (Fig. 2). Addition of one equiv. of AgBF₄ to NetState-I instantaneously generated NetState-II = [Ag(1)]⁺ + [H(2)]⁺. In the ¹H NMR of NetState-II, the 9-H and e-H proton signals shifted to 6.55 and 8.23 ppm, respectively, which substantiated the formation of [Ag(1)]⁺ whereas the shift of proton peaks c′-H and d′-H to 7.12 and 8.00 ppm, respectively, established the protonated luminophore [H(2)]⁺ (Fig. 2). Upon excitation at 300 nm, NetState-I exhibited emissions at λ = 346, 372 and 393 nm which represent an overlap of the emission peaks of the free luminophore 2 at λ = 346 and 372 nm and of [H(1)]⁺ at λ = 376 and 393 nm (Fig. 3c). In NetState-II, the fluorescence spectrum showed a single emission at λ = 461 nm, which is attributed to [H(2)]⁺ (Fig. 3). In contrast, [Ag(1)]⁺ in NetState-II is nonfluorescent as advocated by the full quenching of the emission of [H(1)]⁺ at λ = 376 and 393 nm upon titration with silver(i) ions (ESI, Fig. S18^†). When one equiv. of tetra-n-butylammonium iodide (TBAI) was added to NetState-II, complete restoration of NetState-I was achieved as confirmed from ¹H NMR data. TBAI acted as a scavenger for silver(i) by removing AgI through precipitation thus reversing the translocation sequence. Multiple switching of NetState-I ⇆ NetState-II was readily followed by ¹H NMR (Fig. 2) and luminescence changes. Upon successive addition of AgBF₄ to NetState-I, the emission changed to λ = 461 nm (sky blue emission of [H(2)]⁺) in NetState-II (Fig. 3a and b) in a clean ratiometric manner that allowed monitoring of the amount of liberated protons. Alternating switching between the NetStates was followed by emission spectroscopy over three cycles with a small decline in emission intensity (Fig. 3c and d), which may be due to the progressive accumulation of AgI.Open in a separate windowFig. 2Partial ¹H NMR spectra (500 MHz, CD₂Cl₂, 298 K) showing reliable switching between NetState-I and II over 3 cycles. (a) After mixing of TFA, 1, and 2 in 1 : 1 : 1 ratio; (b) after adding 1.0 equiv. of AgBF₄ furnishing [Ag(1)]⁺ and [H(2)]⁺; (c) after addition of 1.0 equiv. of TBAI; (d) after adding 1.0 equiv. of AgBF₄; (e) after addition of 1.0 equiv. of TBAI, (f) after adding 1.0 equiv. of AgBF₄; (g) after addition of 1.0 equiv. of TBAI. Blue asterisk marked proton signals represent the tetra-butylammonium ion.Open in a separate windowFig. 3(a) Ratiometric emission titration of 2 (2.1 × 10⁻⁵ M) with TFA (5.3 × 10⁻³ M) in CH₂Cl₂. (b) Change of emission upon protonation of 2. (c) Reversibility of the network NetState-I ⇆ II over three cycles upon addition of silver(i) and TBAI (see text, blue arrow assigns direction of switching) monitored by luminescence changes at λ = 461 nm. (d) Multiple cycles monitored by the change of the fluorescence intensity.For probing the catalytic efficiency of each state towards a simple amine deprotection reaction we chose the trityl-protected aniline 6 (Fig. 4) performing two NetState-I ⇆ NetState-II cycles (Fig. 4b). We first prepared NetState-I by mixing 1, 2 and TFA (1 : 1 : 1) in CD₂Cl₂ and then added 6 and TMB (1,3,5-trimethoxybenzene as an internal standard) in a 10 : 10 ratio. After heating for 2 h at 40 °C, the ¹H NMR spectrum revealed no product 7. Clearly, the proton in [H(1)]⁺ is catalytically inactive (ESI, Fig. S10b^†). Addition of one equiv. of AgBF₄ and probing the catalysis again for 2 h at 40 °C revealed formation of product 7 in (30 ± 2)% yield (ESI, Fig. S10c^†). When the catalytic experiment was probed with only [H(2)]⁺ under identical reaction conditions, it provided the same yield (ESI, Fig. S16^†). This suggested that the reversible proton relay network (Scheme 1) functioned reliably also in presence of 6 with complex [H(2)]⁺ as the catalytically active species. Addition of one equiv. of TBAI to the above NetState-II mixture and probing for catalysis, revealed no further product formation. Apparently, addition of TBAI translocated back the proton to nanoswitch 1 thus restoring NetState-I, which resulted in the OFF state for catalysis. Adding one equiv. of AgBF₄ to the mixture and probing again the catalytic activity resulted in formation of further (29 ± 2)% of product 7, which was basically identical to the yield produced in the first cycle of NetState-II (ESI, Fig. S10e^†). Adding TBAI and the consumed amount of substrate showed no further deprotection. In sum, the robust catalytic performance of the proton relay network as reflected in two successive catalytic cycles resulted in no significant decline of the catalytic activity (Fig. 4b).Open in a separate windowFig. 4(a) Representation of the OFF/ON regulation of the trityl deprotection reaction in NetState-I and II. (b) Reversible switching between the network states furnished reproducible amounts of product 7 in NetState-II (two independent runs). Consumed amounts of substrates were added (green asterisk).The OFF–ON switching of the catalytic trityl deprotection of aniline 6 by the proton relay network inspired us to utilize the system to catalytically fabricate a slider-on-deck based on the N_an (= aniline) → ZnPor interaction. Deprotection of the protected biped 4 should afford the bis-aniline biped 5 and enable its attachment to the tris-zinc porphyrin deck 3 thus forming the slider-on-deck [(3)(5)].Addition of one equiv. of biped 5 (for synthesis and characterization, see ESI^†) to deck 3 (ref. 25) in CD₂Cl₂ quantitatively generated the slider-on-deck [(3)(5)], which was unambiguously characterized by ¹H NMR, ¹H DOSY NMR and elemental analysis (ESI, Fig. S5 and S6^†). In the ¹H NMR-spectrum, several diagnostic shifts of biped and deck proton signals attested the binding of biped 5 to the deck. Formation of the slider-on-deck was further confirmed from a single set in the ¹H DOSY in CD₂Cl₂ (ESI, Fig. S9^†).The single set of protons (p-H, r-H, s-H, t-H, u-H, v-H) in the ¹H NMR spectrum for all three-zinc porphyrin (ZnPor) stations of deck 3 unmistakably suggested rapid intrasupramolecular exchange^29a of the aniline biped across all three ZnPor sites requiring fast N_an → ZnPor bond cleavage. At low temperature (−65 °C) the ZnPor p-H proton signals separated into two sets (1 : 2) (ESI, Fig. S7 and S8^†). Kinetic analysis over a wide temperature range provided the exchange frequency (k) with k₂₉₈ = 540 kHz at 298 K (Fig. 5a) and the activation data as ΔH^‡ = 41.5 ± 0.7 kJ mol⁻¹, ΔS^‡ = 2.8 ± 0.7 J mol⁻¹ K⁻¹ and ΔG^‡₂₉₈ = 40.7 kJ mol⁻¹ (Fig. 5).Open in a separate windowFig. 5(a) Experimental and theoretical splitting of p-H proton signal of nanoslider [(3)(5)] in VT ¹H-NMR (600 MHz) furnishing rate data in CD₂Cl₂. (b) Eyring plot providing kinetic parameters.After determining the kinetic parameters of the slider-on-deck [(3)(5)], we investigated its in situ catalyzed formation by the proton relay network. Therefore, we mixed nanoswitch 1, luminophore 2, TFA, trityl-protected biped 4, deck 3 and TMB in 1 : 1 : 1 : 5 : 5 : 5 ratio in CD₂Cl₂ and heated the reaction mixture for 12 h at 40 °C. The ¹H NMR spectrum suggested no formation of any slider-on-deck in NetState-I (ESI, Fig. S13^†). Addition of one equiv. of AgBF₄ to the same mixture (converting NetState-I to NetState-II) and heating it at 40 °C for 12 h revealed full formation of the slider-on-deck as indicated by the disappearance of the 2′-H, 3′-H and 4′-H proton signals of biped 4 at 6.60, 6.31 and 6.51 ppm (ESI, Fig. S11^†), respectively, and the quantitative emergence of 4′-H proton signal of biped 5 bound to deck 3 at 5.84 ppm (ESI, Fig. S12^†). The upfield shift of the deck''s meso t-H protons from 10.36 to 10.29 ppm allowed monitoring of the formation of [(3)(5)], e.g. by the gradually increased chemical shift difference of proton signal t-H (Δδ_t-H) (Fig. 6a; ESI, Fig. S11^†). In sum, the protonated [H(2)]⁺ in NetState-II catalyzed the deprotection of the trityl-protected biped 4 with the effect that the resultant bis-aniline biped 5 would quantitatively bind to deck 3 affording the slider-on-deck [(3)(5)]. Furthermore, the OFF/ON switching of catalysis was probed for shortened time periods of 3 h. As illustrated in Fig. 6b, the process could be readily followed using the proton signal t-H (Δδ_t-H) (ESI, Fig. S14^†) and the growth of proton signal 4′-H at 5.84 ppm of bound biped 5 (ESI, Fig. S15^†). One clearly sees recurring OFF/ON regulation in NetState-I/II.Open in a separate windowFig. 6Plot of the shift difference of proton signal t-H (Δδ_t-H) vs. time for (a) the formation of [(3)(5)] from 3 and 4. Red curve: formation of [(3)(5)] over 12 h catalyzed by NetState-II. Black curve: OFF state of catalysis in NetState-I. (b) OFF/ON catalytic cycles by switching between NetState-I and NetState-II. Addition of Ag⁺, see blue asterisk; addition of TBAI: red asterisk.In conclusion, we have designed a supramolecular multicomponent network³⁶ that acts as an OFF/ON proton relay with a luminescence display³⁷ and is useful for switchable catalysis.³⁸ The network is toggled by chemical input and intercomponent communication^15,16 and as a result is a conceptual complement to photoacids driving networks.^7,8The release and capture of protons is demonstrated by the ON/OFF trityl deprotection of anilines. To demonstrate functioning in a complex setting, the network was utilized to catalyze formation of a high-speed slider-on-deck assembly based on N_an → ZnPor interactions (sliding frequency of k₂₉₈ = 540 kHz). The robust operation of the proton relay furnishing dynamic machinery³⁹ through acid catalysis followed by self-assembly is a valuable step mimicking sophisticated biological proton relays.²     

Size effect studies in catalysis: a simple surfactant-free synthesis of sub 3 nm Pd nanocatalysts supported on carbon

Supported Pd nanoparticles are prepared under ambient conditions via a surfactant-free synthesis. Pd(NO₃)₂ is reduced in the presence of a carbon support in alkaline methanol to obtain sub 3 nm nanoparticles. The preparation method is relevant to the study of size effects in catalytic reactions like ethanol electro-oxidation.

A simple surfactant-free synthesis of sub 3 nm carbon-supported Pd nanocatalysts is introduced to study size effects in catalysis.

A key achievement in the design of catalytic materials is to optimise the use of resources. This can be done by designing nanomaterials with high surface area due to their nanometre scale. A second achievement is to control and improve catalytic activity, stability and selectivity. These properties are also strongly influenced by size.^1–3 To investigate ‘size effects’ it is then important to develop synthesis routes that ensure well-defined particle size distribution, especially towards smaller sizes (1–10 nm).Metal nanoparticles are widely studied catalysts. In several wet chemical syntheses, NP size can be controlled using surfactants. These additives are, however, undesirable for many applications^4,5 since they can block active sites and impair the catalytic activity. They need to be removed in ‘activation’ steps which can negatively alter the physical and catalytic properties of the as-produced NPs. Surfactant-free syntheses are well suited to design catalysts with optimal catalytic activity⁶ but their widespread use is limited by a challenging size control.³Palladium (Pd) NPs are important catalysts for a range of chemical transformations like selective hydrogenation reactions and energy applications.^7–9 It is however challenging to obtain sub 3 nm Pd NPs, in particular without using surfactants.² Surfactant-free syntheses are nevertheless attracting a growing interest due to the need for catalysts with higher performances.^10–14Promising surfactant-free syntheses of Pd NPs were recently reported.^8,15 The NPs obtained in these approaches are in the size range of 1–2 nm and show enhanced activity for acetylene hydrogenation⁸ and dehydrogenation of formic acid.¹⁵ Enhanced properties are attributed to the absence of capping agents leading to readily active Pd NPs. The reported syntheses consist in mixing palladium acetate, Pd(OAc)₂, in methanol and the reduction of the metal complex to NPs occurs at room temperature. The synthesis is better controlled in anhydrous conditions to achieve a fast reaction in ca. 1 hour. Another drawback is that the synthesis must be stopped to avoid overgrowth of the particles. Therefore, a support material needs to be added after the synthesis has been initiated and no simple control over the NP size is achieved.^8,15In this communication a more straightforward surfactant-free synthesis leading to sub 3 nm carbon-supported Pd NPs in alkaline methanol at ambient conditions is presented. A solution of Pd(OAc)₂ in methanol undergoes a colour change from orange to dark, indicative of a reduction to metallic Pd, after ca. a day. However, only ca. 1 hour is needed with Pd(NO₃)₂, Fig. 1 and UV-vis data in Fig. S1.^† The fast reduction of the Pd(NO₃)₂ complex in non-anhydrous conditions is a first benefit of the synthesis presented as compared to previous approaches.Open in a separate windowFig. 1Pictures of 4 mM Pd metal complexes in methanol without or with a base (as indicated).For particle suspensions prepared with Pd(OAc)₂ or Pd(NO₃)₂ the NPs agglomerate and quickly sediment leading to large ‘flake-like’ materials. When the reduction of Pd(NO₃)₂ in methanol is performed in presence of a carbon support and after reduction the solution is centrifuged and washed in methanol, a clear supernatant is observed indicating that no significant amount of NPs are left in methanol. Transmission electron microscope (TEM) analysis confirms that NPs are formed and well-dispersed on the carbon support surface and no unsupported NPs are observed, Fig. 2a. Likely, the reduction of the NPs proceeds directly on the carbon support. However, the size of the NPs is in the range 5–25 nm, which is still a relatively large particle size and broad size distribution.Open in a separate windowFig. 2TEM micrographs of Pd NPs obtained by stirring 4 mM Pd(NO₃)₂ in methanol and a carbon support for 3 hours, (a) without NaOH and (b) with 20 mM NaOH. Size distribution histograms are reported in Fig. S4.^† The same samples after electrochemical treatments are characterised in (c) and (d) respectively. Size distribution histograms are reported in Fig. S7.^†Assuming a ‘nucleation and growth’ mechanism, the NPs should become larger over time.¹⁶ But the reaction is so fast that by stopping the reaction before completion, size control is not achieved and unreacted precious metal is observed, Fig. S2.^† To achieve a finer size control and more efficient use of the Pd resources, a base was added to the reaction mixture, e.g. NaOH.³ In alkaline media, the formation of Pd NPs is slower; it takes ca. 60 minutes to observe a dark colour for a 5 mM Pd(NO₃)₂ solution with a base/Pd molar ratio of 10 in absence of a support, Fig. 1.Also in alkaline methanol, the NPs agglomerate over time in absence of a support material. However, if the alkaline solution of Pd(NO₃)₂ is left to stir in presence of a carbon support the desired result is achieved, i.e. Pd NPs with a significantly smaller size and size distribution of ca. 2.5 ± 1.0 nm, Fig. 2b. The NPs homogeneously cover the carbon support and no unsupported NPs are observed by TEM suggesting that the NPs nucleate directly on the carbon surface. Furthermore, the supernatant after centrifugation is clear, indicating an efficient conversion of the Pd(NO₃)₂ complex to NPs, Fig. S3.^† Furthermore, there is no need for an extra reducing agent as in other approaches, for instance in alkaline aqueous solutions.⁹The benefits of surfactant-free syntheses of Pd NPs for achieving improved catalytic activity have been demonstrated for heterogeneous catalysis.^8,15 Surfactant-free syntheses are also well suited for electrochemical applications where fully accessible surfaces are required for fast and efficient electron transfer. Several reactions for energy conversion benefit from Pd NPs. An example is the electro-oxidation of alcohols,⁷ in particular ethanol¹⁷ (see also Table S1^†).Previous studies optimised the activity of Pd electrocatalysts by alloying,^18–20 by using different supports^17,21–23 or crystal structures.^24,25 Investigating NPs with a diameter less than 3 nm was challenging.^2,26,27 The surfactant-free synthesis method presented here allows to further study the size effect on Pd NPs supported on carbon (Pd/C) for electrocatalytic reactions.In Fig. 3, results for ethanol electro-oxidation in 1 M ethanol solution mixed with 1 M KOH aqueous electrolyte are reported based on cyclic voltammetry (CV) and chronoamperometry (CA) with Pd/C catalysts exhibiting 2 significantly different size distributions. The electrode preparation, the measurement procedure and the sequence of electrochemical treatments are detailed in the ESI.^† In order to highlight size effects, we compare geometric and Pd mass normalized currents (Fig. 3a and c) as well as the oxidation currents normalized to the Pd surface area (Fig. 3b).Open in a separate windowFig. 3Electrochemical characterisation of carbon supported Pd NPs with 5–25 nm (grey) and 2.5 nm (dark) size in 1 M KOH + 1 M ethanol aqueous electrolyte. (a) 2^nd CV before chronoamperometry (CA), (b) current normalised by the electrochemically active surface area of Pd, (c) CA recorded at 0.71 V vs. RHE after 50 cycles between 0.27 and 1.27 V.It is clearly seen that based on the geometric current density, the smaller Pd NPs exhibit significantly higher currents for ethanol oxidation than the larger NPs. To differentiate if this observation is a sole consequence of the different surface area, the electrochemically active surface area (ECSA) has been estimated based on “blank” CVs (without ethanol) recorded between 0.27 and 1.27 V vs. RHE in pure 1 M KOH aqueous electrolyte and integrating the area of the reduction peak at ca. 0.68 V, Fig. S5.^† As conversion factor, 424 μC cm⁻² was used.²⁸Using this method, the smaller NPs with a size around 2.5 nm exhibit an ECSA of 92 m² g⁻¹ whereas the larger NPs with a size in the range 5–25 nm exhibit an ECSA of 47 m² g⁻¹, consistent with a larger size. Normalising the ethanol electro-oxidation to these ECSA values instead of the geometric surface area, Fig. 3b, still indicates a size effect. It is clearly seen that the smaller Pd NPs exhibit higher surface specific ethanol oxidation currents, in particular at low electrode potentials. Furthermore, a clear difference in the peak ratios in the CVs is observed. The ratio in current density of the forward anodic peak (j_f, around 0.9 V) and the backward cathodic peak (j_b, around 0.7 V vs. SCE) is around one for the smaller NPs, whereas it is about 0.5 for the larger NPs. The forward scan corresponds to the oxidation of chemisorbed species from ethanol adsorption. The backward scan is related to the removal of carbonaceous species not fully oxidised in the forward scan. The higher j_f/j_b ratio therefore confirms that the smaller NPs are more active for ethanol electro-oxidation and less prone to poisoning, e.g. by formation of carbonaceous species that accumulate on the catalyst surface.^29,30 This observation is further supported by a chronoamperometry (CA) experiment, Fig. 3c, at 0.71 V performed after continuous cycling (50 cycles between 0.27 and 1.27 V at a scan rate of 50 mV s⁻¹). In the CA testing of the thus aged catalysts at 0.71 V, the ethanol oxidation current on the two catalysts starts at around the same values, however, its decay rate is significantly different. The Pd mass related oxidation currents for the smaller NPs are after 30 minutes almost twice as high (ca. 200 A g_Pd⁻¹) as for the larger ones (ca. 130 A g_Pd⁻¹), confirming that the small Pd NPs are less prone to poisoning. In particular a factor up to 4 in the Pd mass related ethanol oxidation currents after 1800 s of continuous operation is achieved compared to a recently characterised commercial Pd catalyst on carbon,²⁰ Table S1.^† Despite different testing procedure reported in the literature, it can be concluded from these investigations that the surfactant-free synthesis presented shows promising properties for electrocatalytic ethanol oxidation even after extended cycling.The extended cycling, however, has different consequences for the two catalysts. For the small (2.5 nm) NPs of the Pd/C catalyst, a massive particle loss, but only moderate particle growth is observed as highlighted in Fig. 2 (see also Fig. S6^†). TEM micrographs of the two Pd/C samples recorded before and after the complete testing (CVs and CAs, for details see Fig. S7^†) show that the catalyst with small Pd NPs exhibits a pronounced particle loss as well as a particle growth to ca. 6 nm probably due to sintering. By comparison, for the Pd/C catalyst with the large Pd NPs, no significant influence of the testing on particle size or particle density is apparent.     

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.     

Highly efficient hydrogen peroxide direct synthesis over a hierarchical TS-1 encapsulated subnano Pd/PdO hybrid

We report a hierarchical TS-1 encapsulated subnano Pd/PdO hybrid catalyst that shows unprecedented activity in H₂O₂ direct synthesis from H₂ and O₂. The macro reaction rate in 30 min is up to 35 010 mmol g_Pd⁻¹ h⁻¹ at ambient temperature. Such high catalytic activity is achieved due to the hierarchical porous structure of TS-1 and the formation of the encapsulated subnano Pd/PdO hybrid after oxidation/reduction/oxidation treatment.

A hierarchical TS-1 encapsulated subnano Pd/PdO hybrid catalyst that shows unprecedented activity in H₂O₂ direct synthesis from H₂ and O₂.

Hydrogen peroxide as a clean and strong oxidant is one of the commonly used chemicals in various fields of chemical industry, such as the pulp and paper industry, the textile industry, wastewater treatment, green chemical synthesis metallurgy, electronics manufacture, propulsion and the food industry.¹ Compared to the traditional anthraquinone process (sequential hydrogenation and oxidation of alkyl anthraquinone), the direct synthesis of hydrogen peroxide (DSHP) from hydrogen and oxygen was recognized as an efficient and environmental alternative process owing to its remarkable adherence to green chemistry perspectives, such as low energy consumption, minimized toxicity and infrastructure investment.^2–5Pd supported catalysts were the most extensively and earliest studied catalysts for the DSHP since 1914.⁶ Both DFT and experimental results indicated that subnano Pd particles were most effective for the selective oxygen hydrogenation to hydrogen peroxide,⁷ and the activity and selectivity are also highly dependent upon the oxidation state of the Pd particles.⁸ However, there were limitations in applying Pd nanoparticles catalyst to the reaction due to the thermal vulnerability in a calcination and reduction activation process.⁹ To solve this problem, many preparation methods have been adopted to stabilise Pd nanoparticles and control the particle size and morphology, such as yolk–shell structure,¹⁰ core–shell structure¹¹ and other encapsulation structure supports. But there were still problems that the size of metal particles is larger than 2.5 nm. Encapsulation of Pd species by mercaptosilane-assisted dry gel conversion (DGC) synthesis method can provide a precise control over the nanoparticle size as well as limitating the aggregation under high temperature during activation.¹² However, active sites deep inside the encapsulated nanoparticles were often hardly accessible since the internal configuration diffusion limitations of reactants and products in micropores, leading to low H₂ conversion and decomposition of the long residence time of synthetic H₂O₂.¹³ So, the role of the porous structured catalyst was essential for encapsulated metal nanoparticles.Titanium silicalite-1 (TS-1) has already been used as an excellent catalyst for a variety of selective oxidation reactions employing hydrogen peroxide as oxidant.^14,15 Moreover, in situ H₂O₂ generation coupled with these selective oxidation reactions leading to the desired products such as propylene,^16,17 benzyl alcohol,¹⁸ cyclohexene¹⁹ was a desirable, green and lower cost route. More importantly, the Ti–OOH species formed on the TS-1 during selective oxidation might improve the stability of OOH, which is a key reaction intermediate during the DSHP.²⁰ Hutchings et al. reported that hierarchical titanium silicalite supported Au–Pd catalysts showed high peroxide production rate and benzaldehyde production rate for oxidation of benzyl alcohol by in situ generated H₂O₂.²¹ In this report, the encapsulation of subnano-sized Pd metal particles within conventional (Pd@TS-1) and hierarchical titanium silicalite-1 (Pd@HTS-1) has been achieved (see Scheme 1). The Pd@HTS-1 catalyst after oxidation–reduction–oxidation pre-treatment showed unprecedented activity in direct synthesis of hydrogen peroxide from hydrogen and oxygen under ambient temperature without any promoter.Open in a separate windowScheme 1Schematic diagram of the preparation method for Pd@HTS-1.The TS-1 and HTS-1 encapsulated Pd sub-nanoparticles were first synthesized via solvent evaporation-assisted dry gel conversion method, where the Pd was encapsulated in situ through hydrothermal crystallization in assistance of 3-mercaptopropyl-trimethoxysilane (Scheme 1). The results of ICP analysis confirmed that total Pd contents in Pd@TS-1 and Pd@HTS-1 were 0.094 and 0.106 wt%, respectively. The characteristic diffraction “finger peak” on the X-ray diffraction in Fig. S1^† proved that the TS-1, Pd@TS-1 and Pd@HTS-1 had a well-crystallized MFI structure,²² which was further confirmed by the asymmetric stretching of Si–O–Ti in the spectra of Fourier Transform Infrared Spectroscopy (FT-IR, see Fig. S2^†). For all of the samples, the diffraction peak at 2θ of 25.4° was not observed. Meanwhile, the diffraction peak of crystalline Pd was also not detected for Pd@HTS-1 and Pd@TS-1, indicating that the Pd particles were well dispersed in the zeolite.⁷ Besides, the diffuse reflectance UV-vis spectra of the TS-1, Pd@TS-1 and Pd@HTS-1 were shown in Fig. S3.^† The band at 210 nm in three samples confirmed the tetrahedral structural geometry of Ti in these silicates, and the weak band at 280 nm was assigned to small amounts of penta/hexacoordinated Ti species.²³ Moreover, the absorption band around 300 nm indicated that the three samples contain anatase TiO₂.²⁴The textural properties of the synthesized Pd@TS-1 and Pd@HTS-1 were characterized by N₂ adsorption/desorption and the results were shown in Fig. 1 and Table S1.^† Notably, typical irreversible type IV adsorption isotherms with an H1 hysteresis loop were observed over the Pd@HTS-1 sample (Fig. 1b), indicating the presence of a mesoporous structure. The mesopore size of Pd@HTS-1, obtained through the BJH method, and the obtained graph peaked at about 7.0 nm. Volume of the micropores was around 0.14 cm³ g⁻¹ for both Pd@TS-1 and Pd@HTS-1, but the surface area of Pd@HTS-1 (509.9 m² g⁻¹) was 48.9 m² g⁻¹ larger than that of Pd@TS-1 (461.0 m² g⁻¹) due to its mesoporous structure, which is beneficial for the diffusion of reactants and products through the catalysts.²⁵Open in a separate windowFig. 1Nitrogen adsorption–desorption isotherms of the synthesized TS-1: (a) Pd@TS-1 and (b) Pd@HTS-1.Comparison between the experimentally obtained results from ammonia temperature-programmed desorption (NH₃-TPD) analysis (Fig. S4^†) and the previously reported data showed that the peaks observed were corresponding to weak acid sites, medium acid sites, and strong acid sites of the catalysts.²⁶ Furthermore, pyridine adsorption peak on the FT-IR spectra of these samples (Fig. S5^†) revealed that titanium silicate (TS-1) was an acidic support with a large number of Lewis acid segments and few Brønsted acid segments. As shown in scanning electron microscopy (SEM) image (Fig. 2), Pd@TS-1 particles were crystallites with a morphology close to cuboids and a mean particle size of about 3–5 μm, while the Pd@HTS-1 has spherical morphology with a particle size of about 1.3 μm.Open in a separate windowFig. 2SEM images of the synthesized Pd-modified TS-1: (a) Pd@TS-1 and (b) Pd@HTS-1.The synthesized TS-1 and HTS-1 encapsulated Pd sub-nanoparticles were then subjected to oxidation/reduction/oxidation treatment to adjust the valence states of Pd.²⁷ Such heat treatment cycle can switch off the sequential hydrogenation and decomposition reactions in the DSHP. However, Ostwald ripening, thus the migration and coalescence of metal clusters, will occur at a higher temperature. Therefore, high temperature treatments was used to emulate the conditions used in the literature mentioned before,^28,29 and the thermal stability of the encapsulated Pd@TS-1 catalysts before and after the treatments were also evaluated and compared to investigate the effect of high temperature and the thermal treatments on the catalysts. The Pd@TS-1 and Pd@HTS-1 samples after an air/H₂/air thermal treatments at 500/400/500 °C for 4/2/6 h were denoted as Pd@TS-1-O, Pd@TS-1-OR, Pd@TS-1-ORO, Pd@HTS-1-O, Pd@HTS-1-OR, Pd@HTS-1-ORO respectively with O denoting oxidation and R denoting reduction. The Pd particle size distribution after such treatments was first released by the high-resolution transmission electron microscopy (HRTEM) image in Fig. 3 and S6.^† The Pd particles encapsulated within microporous TS-1 zeolites were well dispersed and uniformly distributed throughout the zeolite crystals. The average sizes of Pd particles encapsulated in the TS-1 and HTS-1 were in the range of 1–2 nm, which, however, was bigger than those of the MFI topology channels (0.53 × 0.56 nm) and intersectional channels (∼0.9 nm). Nevertheless, the successful encapsulation of the Pd particles in the TS-1 zeolites was verified by comparing the hydrogenation rates of a mixture of nitrobenzene and 1-nitronaphthalene. As shown in Fig. S7,^† the reaction rate for the hydrogenation of nitrobenzene and 1-nitronaphthalene was much higher over the Pd@HTS-1-OR compared to the Pd@TS-1-OR. We anticipated that the slightly larger Pd size than the zeolite channels might reflect the local disruption of the crystal structures near the location of the particles during the in situ synthesis. More detailed size distributions of Pd particles encapsulated in the TS-1 and HTS-1 zeolites after air, Ar/H₂ and air treatments were shown in Fig. 3d–f and j–l, respectively. The particle sizes of most of the Pd species still remain below 2 nm on average, which indicated the absence of metal clusters migration and coalescence by Ostwald ripening even after such higher temperature treatments. The high thermal stability of the Pd subnano particles resulted from the embedding confinement.³⁰Open in a separate windowFig. 3HRTEM images and metal particle size distributions of the Pd@TS-1 and Pd@HTS-1 before and after high-temperature oxidation–reduction–oxidation treatments. (a, d Pd@TS-1-O. b, e Pd@TS-1-OR. c, f Pd@TS-1-ORO. g, j Pd@HTS-1-O. h, k Pd@HTS-1-OR. i, l Pd@HTS-1-ORO.)The Pd dispersion and average Pd nanoparticle size for Pd@TS-1 and Pd@HTS-1 after the air/H₂ treatment were further determined by CO chemisorption measurements (see Table S2^†). The dispersions of Pd in Pd@TS-1 and Pd@HTS-1 are 85% and 81%, respectively. The average Pd particle sizes for Pd@TS-1 and Pd@HTS-1 calculated by CO adsorption measurements are 1.06 nm and 1.17 nm, respectively, which was smaller than that estimated from the TEM analysis. This was probably due to the presence of Pd nanocluster or single atoms, which cannot be directly observed by HRTEM.We now turn to the Pd valence states of the catalysts after the oxidation/reduction/oxidation treatment by the XPS (see Fig. S8^†). The Pd3d spectra signals were hardly observed when the concentration of Pd atoms was low, the binding energy peaks for different oxidation states of Pd atoms were collected after peak fitting by prolonging the scanning time.³¹ The XPS results demonstrated the presence of both metallic Pd and PdO. The binding energy of peaks for Pd⁰3d_5/2 and Pd⁰3d_3/2 correspond to 335.5 and 340.6 eV, respectively, while the binding energy for Pd²⁺3d_5/2 and Pd²⁺3d_3/2 were at 337.8 and 341.9 eV, respectively.³¹ The transformation of valence state could be observed in Fig. S8a–c,^† which was derived from XPS measurements. Moreover, the ratios for Pd⁰ and Pd²⁺ atoms in Pd@TS-1 and Pd@HTS-1 were approximately 2 and 1, respectively. On the basis of these results, we proposed a reaction mechanism for the synthetic process of the catalysts, subnano-sized Pd particles might be oxidated from Pd⁰ to Pd²⁺ to form PdO on the surface of the catalysts during reoxidation.The catalytic performance of the TS-1 and HTS-1 encapsulated subnano-sized Pd/PdO hybrid in the direct synthesis of hydrogen peroxide from H₂ and O₂ were tested at ambient temperature without any promoters. Compared to the Pd supported by the active carbon, the selectivity of hydrogen peroxide was higher, the reason might be the formation of Ti–OOH³² and the confinement effect of the Pd encapsulated in the channel of the zeolite (Scheme 2). Both HTS-1 zeolite and Pd@zeolites showed significant amount of O₂ adsorption according to the O₂-TPD (Fig. S9^†), which might be the reason for high activity/selectivity. The selectivity for hydrogen peroxide on Pd@TS-1-OR is lower than that on Pd@TS-1-O, while the degradation rate of hydrogen peroxide on Pd@TS-1-OR are higher than that on Pd@TS-1-O (Fig. 4 and Table S3^†), which was attributed to the change in oxidation state from Pd²⁺ to Pd⁰ after reductive treatment, in agreement with previous reports.^27,33 The selectivity of hydrogen peroxide over Pd@TS-1 increased after an oxidation/reduction/oxidation cycle, the reason might be the weaker adsorption of O₂ and H₂, the intermediate OOH and the production H₂O₂ and the suppression of H₂O₂ decomposition.²⁰Open in a separate windowScheme 2Schematic of the mechanism for DSHP by Pd@TS-1.Open in a separate windowFig. 4H₂O₂ selectivity of DSHP over Pd@TS-1 with different oxidation states for 5 min reaction. Reaction conditions (same as Fig. 5 and ​and6):6): H₂/Ar (2.9 MPa) and air (1.35 MPa), 8.5 g solvent (2.9 g water, 5.6 g MeOH), 0.02 g catalyst, RT, 1200 rpm.The productivity of DSHP over Pd@TS-1 increased with oxidation, reduction and reoxidation treatment in 30 minutes (Fig. 5 and Table S3^†), demonstrated that PdO layer on monometallic Pd catalysts could suppress oxygen dissociation and H₂O₂ degradation,¹² the appropriate PdO formed on the surface of the catalysts after reoxidation can optimize the H₂O₂ production. The hierarchical Pd@TS-1 (35 010 mmol g_Pd⁻¹ h⁻¹) is remarkably higher than those of conventional Pd@TS-1 (3210 mmol g_Pd⁻¹ h⁻¹), the superior hydrogen peroxide production rate of Pd@HTS-1-ORO indicating that the Pd encapsulated by uniformed topology structure of TS-1 highly limited by the effect of pore-diffusion resistance.¹¹ Compared to Pd@TS-1, it was noteworthy that Pd@HTS-1 with only 0.1 wt% Pd content and subnano size after oxidative treatments showed famous reaction activity without any promoters under mild condition, which could be mainly ascribed to the presence of internal diffusion limitation within encapsulated micropore zeolites. The micropore structure limited the use of Pd metal because a part of the Pd crystal surface was blocked by zeolite supports, the hydrogen and oxygen were restricted by the configurational diffusion of zeolite to the Pd surface. Moreover, the formed and desorption H₂O₂ was also constrained by the micropore and thereby resulted in prolonged residence time of the product leading to degradation of H₂O₂. The intracrystal diffusion no longer limited the mass transport process of the hierarchical zeolite due to the presence of additional porosity. Although the physical and structural properties (including the primary particle size, the properties of the external surface and so on) were different between Pd@HTS-1 and Pd@TS-1, we may still draw a conclusion that the excellent catalytic activity is mainly attributed to the presence of mesopore favours diffusion of both reactants and products to and off the active sites in micropores.Open in a separate windowFig. 5Macro reaction rate for H₂O₂ production over Pd@TS-1 and Pd@HTS-1. ^aPd/C#C&Pd/C#Ex from Young-Min Chung;^34bPd–Sn/TiO₂ from Hutchings.²⁹The TON of H₂O₂ production at different reaction time over the six different Pd@TS-1 and Pd@HTS-1 catalysts were shown in Fig. 6. The TON increases with increasing reaction time, however, the slop of the TON–time curves (d_TON/d_t) seems decreased with increasing time, which revealed that the net productivity rate of hydrogen peroxide synthesis declined slightly with increasing time, especially for the Pd@HTS-1-OR at the reaction period of 30–60 min. The accumulative productivity of hydrogen peroxide slowed down, the reason might be the rapid decrease of hydrogen partial pressure in the medium and the ongoing H₂O₂ degradation.Open in a separate windowFig. 6The TON of H₂O₂ production with different reaction time over Pd@TS-1 and Pd@HTS-1 catalysts. TON (turnover number) = mol (H₂O₂)/mol (surface Pd).In summary, successful encapsulation of subnano-size Pd metal particles within titanium silicate (TS-1) voids was achieved via the mercaptosilane-assisted DGC synthesis method. The subnano-size Pd nanoparticles encapsulated in HTS-1 zeolites exhibited superior thermal stability after the oxidation/reduction/oxidation heat treatment process adjusting Pd/PdO hybrid owing to the embedding confinement. The synthesized high-efficiency Pd@HTS-1-ORO showed the famous hydrogen peroxide synthesis productivity, a hydrogen peroxide production rate as high as about 35 010 mmol H₂O₂ g_Pd⁻¹ h⁻¹. Our strategy brings about a finely tailored method to control particle size down to the subnano level and eliminate the diffusion inside metal encapsulated microporous zeolites, which is advantageous for catalytic activity and selectivity in direct synthesis of hydrogen peroxide. Thus, our approach opens up the possibility that the titanium-containing zeolites encapsulated noble metal catalyst can be extended further to selective oxidation reactions with H₂O₂ generated in situ from H₂ and O₂.     

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.     

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.     

Photocatalytic degradation of organic pollutants through conjugated poly(azomethine) networks based on terthiophene–naphthalimide assemblies

A conjugated poly(azomethine) network based on ambipolar terthiophene–naphthalimide assemblies has been synthesized and its electrochemical and UV-vis absorption properties have been investigated. The network has been found to be a promising candidate for the photocatalytic degradation of organic pollutants in aqueous media.

A conjugated two-dimensional poly(azomethine) network based on ambipolar terthiophene–naphthalimide assemblies has been synthesized and its electrochemical and UV-vis absorption properties have been investigated.

Due to the rapid growth of urbanization and intensive industrialization, pollution has evolved into a serious concern that produces a great negative impact on human health and the environment.^1,2 Therefore, many efforts are currently devoted to addressing environmental remediation through the degradation and removal of hazardous contaminants.^3–5 In this regard, photocatalysis has been identified as a suitable approach for environmental remediation given that it is an energy efficient technique that does not require chemical input and does not produce sludge residue.⁶ In recent years, organic semiconducting polymers have evolved into a new type of metal-free and heterogeneous photocatalyst suitable for solar-energy utilization.⁷ The modularity of organic polymers allows the efficient tunning of their electronic and optical properties by bottom-up organic synthesis through the choice of suitable monomeric building blocks.^8–10 Within this context, there is a growing demand for new organic polymeric semiconductors carefully designed to have suitable energy levels of the frontier orbitals, an appropriate bandgap and good intrinsic charge mobility.¹¹For the design of suitable polymeric semiconductors for photocatalysis, it is not only important that the photocatalysts absorb light in the visible light range but also an efficient dissociation of the photogenerated charge carriers is required. The combination of electron-poor acceptor (A) and electron-rich donor (D) moieties in the polymer structure may prevent a fast recombination process following photoexcitation.¹² In addition, it has been found that polymers networks bearing conjugated moieties may exhibit π-stacked columns that can facilitate charge transport.¹³In this respect, molecular and polymeric materials based on the combination of oligothiophene^14,15 and naphthalimide moieties^16,17 connected through conjugated linkers have shown to be very effective in order to efficiently tune their frontier orbital levels and produce tunable organic semiconductors with good charge transport properties.^18–24 As an example, in Fig. 1 is depicted the structure of NIP-3T, an ambipolar organic semiconductor, for which the one-electron HOMO–LUMO excitation consists of the displacement of the electron density from the HOMO, primarily localized on the oligothiophene fragment, to the LUMO, localized on the naphthalimide unit.²⁵Open in a separate windowFig. 1(a) Monomer containing an electron donor terthiophene system directly conjugated with an electron acceptor naphthalimide moiety through a conjugated pyrazine linker (NIP-3T). (b) HOMO and (c) LUMO computed orbital topologies for NIP-3T.²⁵Among conjugated polymers, poly(azomethine)s have found application as organic semiconductors in heterogeneous photocatalysis because of their π-conjugated system and suitable band levels matching the redox window of water.²⁶ The incorporation of D–A monomeric assemblies into poly(azomethine) networks represents an efficient strategy to obtain ambipolar polymeric networks with tunable frontier orbital levels for photocatalytic applications. Thus, in this communication we report the synthesis of a novel donor–acceptor poly(azomethine) network (NIP3T-ANW, Scheme 1) based on NIP-3T monomers. The potential of this system as photodegrading agent for the elimination of contaminant organic dyes in aqueous media is also explored.Open in a separate windowScheme 1Schematic representation of the synthesis of NIP3T-ANW.The synthesis of the macromolecular poly(azomethine) network NIP3T-ANW is acomplished through Schiff-base reactions between trigonal monomers endowed with amine functionalities (TAPB,²⁷Scheme 1) and linear naphthalimide–thiophene-based monomers endowed with complementary aldehyde functional groups (NIP3T2CHO,²⁴Scheme 1). Typically, both monomers were dissolved in an o-dichlorobenzene/n-butanol/acetic acid (1 : 1 : 0.1) mixture, which was then heated at 120 °C under solvothermal reaction conditions for 72 h. A black solid was obtained which was insoluble in common solvents such as water, acetone, THF, toluene or chlorinated solvents like dichloromethane or chloroform. The obtained solid was washed several times with THF to remove the starting materials and low-molecular weight by-products. After drying under vacuum, a black solid was obtained. The yield, as determined by weight, was 98%.To investigate the chemical nature of the material, as well as to determine the conversion of the functional groups after the reaction, we have employed attenuated total reflectance Fourier transform infrared (ATR-FTIR) spectroscopy (Fig. 2). The bands arising from the NH₂ stretching (3000–3400 cm⁻¹) and NH₂ deformation (1650 cm⁻¹) vibrations of the primary amine group of TAPB and the signals from the aldehyde groups of NIP3T2CHO around 2870 (C–H stretching) and 1663 cm⁻¹ (C O stretching) are virtually absent in the NIP3T-ANW spectrum. In addition, a prominent new band is found at 1573 cm⁻¹, which can be assigned to the C N stretching vibration of the imine linkages within the newly formed poly(azomethine) network.^28–30Open in a separate windowFig. 2(a) IR spectra of NIP3T2CHO (blue), TAPB (red) and NIP3T-ANW (black). (b) Solid-state ¹³C CP-MAS NMR spectrum of NIP3T-ANW.Solid-state ¹³C cross-polarization magic angle spinning NMR (¹³C CP-MAS NMR) spectrum (Fig. 2) reveals the characteristic imide signals of the 1,8-naphthalimide moiety at 164.4 ppm, as well as the signal corresponding to the imine carbon at 154 ppm and a signal at 148 ppm which can be assigned to the aromatic carbon neighbouring the nitrogen of the C N group. The absence of the sp² carbons from the NIP3T2CHO ²⁴ aldehyde functionalities above 180 ppm satisfactorily confirms the condensation between the aldehyde and the amine derivatives.Due to the rigidity and geometry of the building blocks, the imine linkers could be ideally generated in such a way that result in a canonical layered hexagonal structure³¹ as predicted by theoretical calculations (Fig. S1 and S2^†). However, in the actual framework, X-ray diffraction (XRD) measurements indicate that the material is mainly amorphous with only some ordered regions, as indicated by the good agreement between the weak and broad diffraction peaks observed at 2θ values larger than 3 and those predicted by calculations for the ideal canonical layered hexagonal structure (Fig. S3^†). In this regard, NIP3T-ANW was submitted to an exfoliation process following a previously described protocol³² for the exfoliation of two-dimensional polymers (see ESI^† for details). The exfoliated material was analysed by dynamic light scattering (DLS) showing a monomodal size distribution of ca. 400 nm (Fig. S4^†) and transmission electron microscopy (TEM) reveals a sheet-like structural aspect (Fig. S5^†).Thermogravimetric analysis of the poly(azomethine) network NIP3T-ANW shows that the degradation starts at around 450 °C and only 40% weight loss is observed at 700 °C (Fig. S6^†). This thermal stability is significantly higher than that observed for NIP3T (Fig. 1), the analogous molecular system based on terthiophene connected with naphthalimide through pyrazine for which the degradation starts at 200 °C (Fig. S6^†).^22,25In photocatalysis mediated by semiconductors, electron–hole pairs (excitons) are generated after light absorption and afterwards dissociate into free charge carriers that can be utilized for redox reactions^33,34 such as CO₂ fixation,³⁵ water splitting³⁶ or organic mineralization.³⁷ Some of the crucial factors that make the photocatalytic process favourable are the levels of conduction and valence as well as the width of the band gap.^38–40 Thus, in order to characterize these parameters, the electrochemical and optical properties of NIP3T-ANW have been analysed.The electrochemical properties of NIP3T and NIP3T-ANW were studied by cyclic voltammetry (Fig. 3). Both materials show ambipolar redox behaviour in which the reversible reduction processes are characteristic of the naphthalimide unit, while the oxidation processes can be ascribed to the conjugated oligothiophene moiety.^19–22,24 For NIP3T-ANW, the first reversible reduction wave (−1.29 V) is shifted to less negative values in comparison with NIP-3T (−1.41 V). On the other hand, the first oxidation half wave potential for NIP3T is observed at +0.41 V and for NIP3T-ANW at +0.44 V. These shifts agree with the electron acceptor ability of the imine linker.Open in a separate windowFig. 3(a) The UV-vis DRS spectra of NIP3T and NIP3T-ANW. (b) Cyclic voltammetry of the NIP3T monomer and the corresponding polymer.The absorption spectrum of NIP3T-ANW as determined by UV-vis diffuse reflectance spectroscopy (UV-vis DRS, Fig. 3) shows a strong absorption in all the UV-vis range, extending even to the near infrared. This broad absorption is red-shifted in comparison with the one observed for NIP3T (Fig. 3), which reflects the formation of the new polymer network with an extended conjugation through the alpha positions of the terthiophenes.Using the corresponding cut-off wavelengths, the optical band gaps E_g found for NIP3T and NIP3T-ANW are 1.59 and 1.42 eV respectively. This optical result suggests that the incorporation of the NIP3T core into an extended conjugated system efficiently harvests photons from the visible range, even extending into the near IR region.To shed some light into the degree of crystallinity of the synthesized NIP3T-ANW network, we have carried out a battery of density functional theory (DFT)-based calculations with the QUANTUM ESPRESSO plane-wave DFT simulation code⁴¹ (see details in ESI^†). We have considered periodic boundary conditions to obtain a fully-relaxed ground-state crystal structure. Optimization of the cell-shape and size, simultaneously to the relaxation of the structure, reveals a hexagonal 2D lattice with an optimized parameter of 48.46 Å. Different interlayer stacking fashions have been tested, with only one yielding a good agreement with the experimental diffractogram from 2θ values >3°. The most favourable stacking predicted by theory consists in an intermediate configuration between the perfectly eclipsed and staggered configurations, with an interlayer distance of 3.42 Å, and permits an adequate accommodation of the layers profiting adjacent pores. Details on the structure can be found in the ESI.^†Additionally, we have computed the electronic band diagram of the obtained crystal structure along the high-symmetry k-path Γ → K → M → Γ, revealing a wide-gap (1.91 eV) semiconducting character, with rather dispersive valence and conduction bands, mainly resembling the molecular HOMO and LUMO of the molecular building blocks (see Fig. S7^†). Besides, computed time-dependent DFT (TDDFT) UV-vis spectrum manifests an excellent agreement with the experimental UV-vis spectra (Fig. S8^†). A broad and pronounced peak-feature is obtained between 600 and 800 nm, centred at around 720 nm (1.7 eV), which agrees with the optical gap of 1.6 eV found for NIP3T from Fig. 3a. This feature corresponds to electronic transitions between the valence and conduction bands, with an energy difference of around 0.2 eV between the optical and the electronic gap, which indicates that charge relaxation in excited states is not much significative. The good agreement between theoretical predictions on the canonically periodic computed system and the experimental evidences seems to justify the presence of some high-crystallinity regions from the synthesis.The band gap of a semiconductor material and the reduction and oxidation potentials are key parameters which determine its light-harvesting properties and types of reaction that can be conducted and therefore the overall photo-catalytic activity. A shift in the adsorption edge of a semiconductor towards longer wavelengths implies a narrower band gap and the efficient harvesting of a wider photons range.⁴²NIP3T-ANW seems to be an appealing material to be utilized as photocatalyst given (i) the optimal light harvesting properties as shown by the optical characterization, (ii) the efficient generation of electron–hole pairs owing to the insertion of terthiophene moieties, and (iii) the right energy band positions for the material.⁷ We therefore evaluated the photocatalytic activity of NIP3T-ANW under white light for the degradation of a model organic pollutant (Rhodamine B dye, RhB) in aqueous solution.⁴³ In the absence of catalyst, RhB remains stable in solution under illumination (Fig. 4a and S9^†). However, in presence of NIP3T-ANW nearly 90% of RhB in an aqueous solution is degraded after 120 min, showing the enhanced catalytic activity of the material (Fig. 4a and S10^†). Furthermore, a good stability is shown upon 4 straight catalytic cycles (Fig. 4b, S11 and S12^†). In contrast, in the presence of the NIP3T moiety, only a 55% degradation of RhB is observed in the same timeframe (Fig. 4a and S11^†).Open in a separate windowFig. 4RhB degradation curves. (a) Comparison between the degradation effect of NIP3T, NIP3T-ANW and without catalyst. (b) NIP3T-ANW stability after four recycling cycles.In the photocatalytic degradation of organic pollutants, they are typically broken down through the attack of superoxide and hydroxyl species, formed when atmospheric oxygen reacts with photogenerated electrons or when water or OH ions are oxidized by holes, respectively.⁴⁴ Additionally, electron–hole pairs (or excitons) can directly reduce or oxidize organic pollutants in aqueous environments. Consequently, the evaluation of the photodegradation mechanism of organic pollutants, despite challenging, can provide meaningful insights about the nature of a semiconductor photocatalyst.⁴⁵ With the aim of evaluating the photodegradation mechanism we performed the measurements in the presence of different scavengers, namely an aqueous solution of AgNO₃ (100 mg L⁻¹), which captures photogenerated electrons, or triethanolamine (TEOA), which traps photogenerated holes (Fig. S13^†).^46,47 In the presence of Ag⁺ we could observe that the photocatalytic efficiency is enhanced, while the addition of TEOA quenched the performance, therefore suggesting that holes are the active specie in the photodegradation mechanism (Fig. S14^†).In summary, we have presented an approach towards the incorporation of D–A π-conjugated monomeric assemblies into poly(azomethine) networks to yield a purely organic semiconductor for the photocatalytic degradation of organic pollutants in aqueous media. The poly(azomethine) network benefits from a straightforward poly(condensation) approach which favourably competes with the elaborate high-temperature protocols applied for the preparation of inorganic materials. This work enriches the family of donor–acceptor organic semiconductor networks and, given its modular nature, paves the way for the development of a promising family of materials for photocatalytic applications.     

Facile one-pot solvothermal preparation of two-dimensional Ni-based metal–organic framework microsheets as a high-performance supercapacitor material

We report a facile one-pot solvothermal way to prepare two-dimensional Ni-based metal–organic framework microsheets (Ni-MOFms) using only Ni precursor and ligand without any surfactant. The Ni-MOFms exhibit good specific capacities (91.4 and 60.0 C g⁻¹ at 2 and 10 A g⁻¹, respectively) and long-term stability in 5000 cycles when used for a supercapacitor electrode.

Two-dimensional Ni-based metal–organic framework microsheets (Ni-MOFms) were synthesized via a facial one-pot solvothermal approach and exhibited good specific capacities and excellent long-term stability when used for a supercapacitor electrode.

With the continuous growth of energy demand worldwide, high-performance, environmental-friendly, and low-cost energy storage devices have attracted extensive research interest.^1–3 Among them, supercapacitors are considered most promising because of their high power density, long lifespan, and fast charging/discharging speed.^4–6 To date, numerous materials have been explored for fabricating supercapacitors. Carbon materials have been usually used for electrical double-layer capacitors (EDLCs), including carbon fibers, carbon nanotubes, carbon spheres, carbon aerogels, and graphene,^7–12 while conducting redox polymers and transition metal oxides/hydroxides are widely explored as active materials for pseudocapacitance and battery-type electrodes.^13–16Metal–organic frameworks (MOFs), a porous crystalline material composed of metal nodes and organic linkers, have been widely applied in versatile fields including chemical sensors, catalysis, separation, biomedicine, and gained more and more attention in the area of energy storage.^17–25 Recently, two-dimensional (2D) MOFs have aroused great interest as a new kind of 2D materials.^26,27 Compared with traditional bulk MOFs, 2D MOFs possess distinctive properties, such as short ion transport distances, abundant active sites, and high aspect ratios, making them exhibit better performance than their bulk counterparts.^28–32 Bottom-up methods are generally adopted to prepare 2D MOFs with the addition of surfactants to control the growth of MOFs in a specific direction.^33–35 However, the use of surfactants inevitably blocks part of the active sites at the expense of the performance of materials. Therefore, it is highly necessary to explore and develop a direct solvothermal synthesis of 2D MOFs with the advantages of additive-free, simple operation, and easy scale-up.Herein, we report a facile one-pot solvothermal method to synthesize 2D Ni-based MOF microsheets (denoted as Ni-MOFms) by treating nickel chloride hexahydrate (NiCl₂·6H₂O, the metal precursor) together with the trimesic acid (H₃BTC, the ligand) in a mixed solvent of N,N-dimethylformamide (DMF), ethanol (EtOH) and H₂O. During the whole preparation process, only Ni precursor and the ligand are used while no surfactant is added. When used as active materials for a supercapacitor electrode, the obtained Ni-MOFms displayed excellent reversibility and rate performance. It also exhibited specific capacities of 91.4 and 60.0 C g⁻¹ at 2 and 10 A g⁻¹, respectively. Besides, they showed a good cycling performance in 5000 cycles with about 70% of the specific capacity and almost 100% of the coulombic efficiency maintained.Morphologies of the Ni-MOFms were characterized by scanning electron microscopy (SEM) and transmission electron microscopy (TEM). As shown in Fig. 1a and b, the Ni-MOFms were successfully fabricated via the facile one-pot solvothermal method with varying lateral sizes on the micron scale. Energy dispersive spectroscopy (EDS) mapping indicated that the obtained microsheets were mainly composed of C, O, and Ni. A trace amount of N was also observed, which could be attributed to the residual DMF in the mixed solvent (Fig. 1c). These elements were uniformly distributed throughout the whole microsheet. To measure the exact thickness of the Ni-MOFms, atomic force microscopy (AFM) was used. Fig. 1d showed that the thickness of the microsheet was about 58 nm. Considering the large lateral size, even such thickness could produce a relatively high aspect ratio, which is beneficial to the electrochemical performance.Open in a separate windowFig. 1(a) TEM image, (b) SEM image, (c) EDS mapping, and (d) AFM image and the corresponding height profile of the Ni-MOFms.The composition information of the Ni-MOFms was analyzed by X-ray diffraction (XRD) and the resulting diffraction pattern was shown in Fig. 2a. It was clear that the sample was a crystalline material. However, the exact structure was difficult to determine because no matching MOF structure has been found. Therefore, the structure of the Ni-MOFms was further confirmed by Fourier transform infrared spectroscopy (FT-IR). As shown in Fig. 2b, there was a sharp peak at 1721 cm⁻¹ for H₃BTC, which could be ascribed to the stretching vibration of C O in the nonionized carboxyl group.³⁶ For the Ni-MOFms, the peak at this location disappeared while four new peaks appeared. Bands at 1634 and 1557 cm⁻¹ were related to the asymmetric stretching vibration of carboxylate ions (–COO⁻) and peaks at 1433 and 1371 cm⁻¹ were the characteristic peaks of the symmetric stretching vibration of –COO⁻.^37,38 All these changes indicate that the ligand interacted well with the metal precursor.Open in a separate windowFig. 2(a) XRD pattern of the Ni-MOFms. (b) FT-IR spectra of H₃BTC and the Ni-MOFms.The chemical status and surface composition of the Ni-MOFms were further examined by X-ray photoelectron spectroscopy (XPS). From Fig. S1a^† we could see that the Ni-MOFms were composed of C, O, Ni, and N, which was consistent with the result of EDS mapping. High-resolution spectra of C 1s, Ni 2p, O 1s, and N 1s were shown in Fig. S1b–e.^† Characteristic peaks of C 1s at 288.27, 286.50, 285.85, and 284.80 eV were related to O C–OH, C–O, C–C, and C C, respectively, suggesting the presence of H₃BTC (Fig. S1b^†).³⁹ The Ni 2p spectrum showed two peaks at 873.32 and 855.77 eV, which could be ascribed to Ni 2p_1/2 and Ni 2p_3/2, respectively, together with two satellite peaks at 879.26 and 861.05 eV, verifying the existence of Ni²⁺ (Fig. S1c^†).⁴⁰ In the O 1s region, bands positioned at 532.94 and 531.40 eV could be ascribed to the adsorbed H₂O molecules on the surface of Ni-MOFms and typical metal–oxygen bonds, respectively, further corroborating the coordination between H₃BTC and Ni²⁺ (Fig. S1d^†).³⁹ Finally, the high-resolution spectrum of N 1s was also analyzed (Fig. S1e^†). There were two main peaks at 400.18 and 402.21 eV that could be ascribed to neutral amine and charged nitrogen, respectively,⁴¹ further proving the residual DMF on the Ni-MOFms surface.To explore the crucial factors in the formation process of the Ni-MOFms, the reaction time and temperature, the solvent, the ligand addition amount, and the ligand type were studied. As shown in Fig. S2,^† different crystalline materials were obtained at different reaction times. With the increase of reaction time, the material gradually changed from sphere to sheet. The reaction temperature is another crucial factor. At 120 °C, the material was amorphous and spherical. When the temperature rose, the crystal formed and appeared as microsheets (Fig. S3^†). The effect of solvent was illustrated in Fig. S4.^† Microsheets could not be synthesized in DMF or DMF with a small amount of EtOH. In the mixed solvent of DMF and H₂O, crystals could be prepared, indicating the vital role of H₂O. However, spheres existed in the products. Only when a mixture of DMF, EtOH, and H₂O with a certain proportion was used as the solvent, the Ni-MOFms could be obtained. Furthermore, we investigated the effect of the ligand addition amount. From Fig. 1 and S5^† we can see that the Ni-MOFms crystals formed when the molar ratio of Ni precursor and H₃BTC was 1 : 2 (Fig. 1). We speculated that ligands could simultaneously act as regulators to adjust the morphology of materials, avoiding the use of additional surfactants. When the ligand was replaced with 2-methylimidazole (2-MI) or terephthalic acid (H₂BDC), flower-like crystals rather than microsheets were obtained (Fig. S6^†), indicating the importance of the ligand type. Taking the above factors into account, we could finally determine the suitable conditions for preparing the Ni-MOFms (see the experimental section in ESI^†).The potential application of the Ni-MOFms in supercapacitors was first tested by cyclic voltammetry (CV) in 3 M KOH between 0 and 0.4 V (vs. saturated calomel electrode, SCE). As can be seen from Fig. 3a, all CV curves had similar shapes and the peak currents improved gradually as the scan rate increased, suggesting the good capacitive behavior of the Ni-MOFms electrode.⁴² When the scan rate was as high as 150 mV s⁻¹, redox peaks could still be observed, which indicated the excellent rate performance and kinetic reversibility.⁴³ Besides, as the scan rate went up from 10 to 150 mV s⁻¹, the reduction and oxidation peaks moved towards negative and positive potential, respectively, demonstrating the electrode polarization at large scan rates.⁴⁴Open in a separate windowFig. 3Electrochemical measurements of the Ni-MOFms. (a) CV curves at different scan rates. (b) GCD curves at various current densities and (c) corresponding specific capacities. (d) The EIS Nyquist plot at the bias potential of 0.4 V and the equivalent circuit model with the fitted plots (the red dots).The galvanostatic charge–discharge (GCD) behavior was further investigated to assess the coulombic efficiency and the specific capacity of the Ni-MOFms (see the ESI^† for detailed calculation method).^45,46 As shown in Fig. 3b, the shape of GCD curves was highly symmetric during charging and discharging, indicating that the coulombic efficiency of Ni-MOFms was almost 100% at various current densities. The specific capacities of 91.4, 78.4, 71.4, 64.0, and 60.0 C g⁻¹ were achieved at current densities of 2, 4, 6, 8, and 10 A g⁻¹ (Fig. 3c), respectively, demonstrating the excellent rate capability with about 65.6% of the specific capacity maintained from 2 to 10 A g⁻¹. The specific capacity at 2 A g⁻¹ was comparable with or even superior to that of some MOF materials reported in the literatures (Table S1^†).^47–50The kinetics of the electroanalytical process was then investigated by electrochemical impedance spectroscopy (EIS). Fig. 3d showed the Nyquist plot of Ni-MOFms from 0.01 to 100000 Hz and the corresponding equivalent circuit model (inset) with the fitted plots. CPE was the constant phase element related to the double layer capacity.⁵¹ The equivalent series resistance was denoted by R_s and its value obtained from the x-axis intercept was about 2.1 Ω, indicating the low resistance of the solution.⁴³R_ct represented the charge-transfer resistance at the interface of the electrode and electrolyte.⁵² For Ni-MOFms, the value of R_ct was up to 147.1 Ω, which could be attributed to the poor conductivity of MOF materials.The long-term stability of Ni-MOFms was also explored by charging–discharging at 10 A g⁻¹ for 5000 consecutive cycles. From Fig. 4 we could see that the specific capacity retention remained about 70% after 5000 cycles and the coulombic efficiency was maintained at almost 100% throughout the whole process. Furthermore, the inset in Fig. 4 exhibited that the GCD curves of the last 10 cycles were the same as the first 10 cycles, indicating excellent cycling stability.Open in a separate windowFig. 4Cycle property of Ni-MOFms at 10 A g⁻¹. Inset: GCD curves of the first 10 cycles (left) and the last 10 cycles (right).     

Electrochemical preparation system for unique mesoporous hemisphere gold nanoparticles using block copolymer micelles

Gold nanoparticles (AuNPs) are widely used in various applications, such as biological delivery, catalysis, and others. In this report, we present a novel synthetic method to prepare mesoporous hemisphere gold nanoparticles (MHAuNPs) via electrochemical reduction reaction with the aid of polymeric micelle assembly as a pore-directing agent.

Mesoporous hemisphere Au nanoparticles using self-assembled micelles, for the first time, are demonstrated by using electrochemical reduction on a Ti substrate.

Gold (Au) is one of the most stable and versatile elements utilized in various fields, including catalysis, optics, and industrial purposes. Consequently, various shapes and sizes of AuNPs have been intensively studied to improve the performance of Au in different applications.^1–6 Previously, nanoporous or dendritic metal nanostructures, including Au nanostructures, have been synthesized by employing different reagents and conditions such as SH-terminated amphiphilic surfactant,⁷ pH controlling,⁸ and hard-templates.^9,10 The reported porous and dendritic Au nanostructures possess high surface areas and rich active sites, which in turn lead to highly enhanced catalytic activities.Recently, a soft-template method using self-assembled micelles or lyotropic liquid crystals as pore-directing agents has allowed the successful synthesis of mesoporous nanoparticles^11–13 and films^14–17 with different metal compositions. The metals with mesoporous structures demonstrate superior catalytic activity per weight or surface area over their nonporous bulk forms. Previously, our group reported a several-fold increase in the catalytic activity of mesoporous metals in reactions such as the methanol oxidation reaction (MOR),^14,15 ethanol oxidation reaction (EOR),^13,15–17 and nitric oxide reduction¹² as compared to their bulk nanoparticles and films. Such improvement in the catalytic activity of mesoporous structures is mainly attributed to their significantly larger surface areas, more exposed catalytically active sites, and increased durability against aggregation.Interestingly, nanoporous or mesoporous Au structures had been successfully synthesized by using a dealloying method¹⁸ and a hard templating method.⁹ Such methods, however, are a little complicated, and pore-directing templates often remain within the pores, thus leading to severe contamination. Using a thiol group is an alternative way to synthesize mesoporous Au nanospheres.⁷ A significant drawback of using a thiol group, however, is its strong chemical bonding with Au, thus becoming unable to be removed. The synthesis of mesoporous structures using self-assembled polymeric micelles as soft-templates, on the other hand, is a more facile method with fewer synthetic steps, and it is also known to be free of contaminations within the pores. Although a soft-templating method using polymeric micelles has been utilized for the preparation of mesoporous Au and Au-based alloy films towards surface-enhanced Raman scattering (SERS) signals,¹⁹ glucose sensing,^20,21 and MOR,²² the obtained morphologies have been limited to only films.Despite such apparent benefits arising from mesoporous structures and their synthesis using soft-templates, the synthesis of mesoporous AuNPs using soft-templates has not been achieved yet. It is mainly due to the physical and chemical properties of Au which make it extremely hard to form mesoporous structures. Herein, we adopt an electrochemical approach and the soft-template method to synthesize MHAuNPs successfully. As discussed above, we expect MHAuNPs to be highly efficient in various applications in medical diagnosis,²³ optical sensing,²⁴etc.In this report, MHAuNPs with different shapes and sizes are for the first time reported by changing various electrochemical deposition conditions such as applied potentials between electrodes and deposition times. Scheme 1 shows the schematic illustrations of the entire process of precursor preparation (Scheme 1a) and the MHAuNPs fabrication process (Scheme 1b), including the deposition and the detachment of the nanoparticles. The characterization methods implemented in this paper are mentioned in ESI.^†Open in a separate windowScheme 1(a) The process of Au precursor solution preparation and (b) fabricating MHAuNPs by electrochemical reduction.In a typical experiment, a p-doped silicon (Si) wafer was cleaned by using acetone, isopropyl alcohol, and deionized water (DIW) with sonication for 5 minutes, followed by nitrogen (N₂) gas blowing to dry the Si wafer. After the wet cleaning process, the Si surface was treated by oxygen (O₂) plasma for 5 minutes (Oxford Instruments PlasmaPro 80 Reactive Ion Etcher) to remove residual organic impurities. Then, 10 nm of titanium (Ti) layer and 100 nm of Au layer were deposited sequentially by electron beam evaporation (Temescal FC-2000 e-beam evaporator) at 10⁻⁶ torr. Commercially available Au etchant (Sigma-Aldrich) was used to etch the Au film to expose the Ti area (the left image in Scheme 1b). During etching, about 20 percent of Au area was left to be connected to the electrochemical work station, as drawn in Scheme 1b. In preparation of the Au precursor solution, 5 mg of poly(styrene)-block-poly(ethylene oxide) (PS-b-PEO, the number of average molecular weight (M_w) for each block is 18 000 for PS and 7500 for PEO, respectively) was mixed in 1.5 ml of tetrahydrofuran (THF) followed by stirring at 300 rpm for 8 hours. Then, 0.75 ml of ethanol, 0.5 ml of HAuCl₄ aqueous solution (40 mM), and 1.25 ml of DIW were added sequentially. The solution was stirred for another 30 minutes at 200 rpm. The existing block copolymer micelles can be confirmed by TEM observation, and the average diameter is 25 nm, as shown in Fig. S1.^† For the electrochemical deposition, an electrochemical workstation (CH Instruments Inc. 660e) with three electrode system was used to deposit MHAuNPs on the Ti/Si substrate. After the deposition, the particles were carefully washed by chloroform, followed by a rinse using DIW to remove the residual micelles completely. To detach and collect MHAuNPs from the Ti/Si substrate, the substrate was soaked in ethanol and strongly sonicated for a few minutes (Scheme 1b).Fig. S2^† shows the details of the growth mechanism of MHAuNPs by different deposition times. At the initial stage (Fig. S2a^†), small nanoparticles are generated by reducing Au ions in the precursor solution throughout the substrate. Then, the seed starts growing and forming MHAuNPs as the deposition time increases (Fig. S2b–e^†). This similar growth mechanism is the same as the previous report.¹⁹ The high-angle annular detector dark-field scanning transmission electron microscopy (HAADF-STEM) image (Fig. S2f^†) shows the mesopores inside the MHAuNPs are homogeneously generated. As-obtained MHAuNPs consist of a pure Au element without any impurities, as shown in Fig. S3.^† Fig. 1 and S4^† show scanning electron microscope (SEM) images of MHAuNPs deposited at different voltages from −0.2 V to −0.9 V vs. Ag/AgCl at high magnification and low magnification, respectively. Different deposition voltages lead to significant changes in the particle sizes but slight differences in the particle shapes. The size distributions of MHAuNPs and the plots of the average diameters of MHAuNPs by different deposition voltages are described in Fig. 2. The distribution graphs show the large sizes of particles, such as more than 1 μm in diameter, when the high voltage (−0.2 V vs. Ag/AgCl) is applied (Fig. 2a). The distribution becomes narrower upon the lower applied voltage. The average diameter-applied voltage plots in Fig. 2b show that the average particle size decreases from around 1.1 μm at −0.2 V to about 300 nm at −0.9 V. Thus when the lower deposition voltages are applied (i.e., the deposition rate is higher) (Fig. 1g–h), the smaller particles with a higher degree of size uniformity are obtained. The opposite trend is observed at higher deposition voltages (i.e., the deposition rate is lower) (Fig. 1a and b), at which the particles become larger and their size uniformity decreases. This trend is because the higher voltage allows only a limited number of seed particles to be deposited on the Ti/Si substrate, and each seed individually grows with no additional seed formation. Whereas the lower voltage can allow a higher number of seeds, leading to a uniform supply of electrons from the working Ti/Si electrode (Fig. S5^†). In addition, the lower deposition voltages make the particle shape more hemispherical in Fig. 1f–h.Open in a separate windowFig. 1The SEM images of MHAuNPs electrochemically deposited at (a) −0.2 V, (b) −0.3 V, (c) −0.4 V, (d) −0.5 V, (e) −0.6 V, (f) −0.7 V, (g) −0.8 V, and (h) −0.9 V for 500 s. The scale bars indicate 200 nm.Open in a separate windowFig. 2(a) Size distributions of MHAuNPs generated by different voltages and (b) the average diameter–the applied voltage plots. Fig. 3 shows the SEM images of MHAuNPs deposited at −0.2 V and for different deposition times from 250 s to 1000 s. Although longer deposition time does not change the number of MHAuNPs, it leads to the growth of MHAuNPs in lateral and vertical directions. Although the MHAuNPs grow more than about two or three times larger at long deposition time, the mesoporous formation does not seem to be changed, as shown in insets in Fig. 3. This point indicates that the deposition time is not the main factor affecting the formation of mesoporous structures as well as the number of particles (seeds), but it affects the sizes of particles.Open in a separate windowFig. 3The SEM images of MHAuNPs deposited at −0.2 V (vs. Ag/AgCl) for (a) 250 s, (b) 500 s, and (c) 1000 s. The scale bars indicate 10 μm. The insets in each figure are magnified SEM images of each condition (The scale bars in insets indicate 500 nm).In this report, 10 nm Ti layer on Si wafer plays an important key role in the formation of MHAuNPs, as previously mentioned in the experimental procedure. The use of the Ti substrate with low conductivity (ca. 2.38 × 10⁶ S m⁻¹), which is about only 5.8% in comparison with that of Au (ca. 4.10 × 10⁷ S m⁻¹), is not common in the electrochemical plating research field.^25–30 Most of the papers on mesoporous metal structures synthesized by electrochemical deposition have utilized Au or Pt substrates due to its chemical stability and high electrical conductivity.^14–17 Fig. S6^† shows the amperometry (i–t) curves during the deposition of MHAuNPs (black dots) on a Ti/Si substrate and mesoporous gold films (red dots) on an Au substrate at the same deposition condition. As shown in Fig. S6,^† around 1/7 times less current flows on the Ti/Si substrate throughout the deposition time. This low current density on the Ti/Si substrate is one of the factors for fabricating MHAuNPs. Low current density causes the formation of a few particles (i.e., seeds) at the initial stage of the deposition and leads to seed growth in a few places, as explained in Fig. S2.^† Furthermore, the use of Ti/Si substrates affects the bottom parts of MHAuNPs to become an arch shape. Only edges of MHAuNPs attach onto the Ti/Si substrates, as shown in Fig. 4. This attachment is because the interaction between the deposited MHAuNPs and the Ti substrate surface (probably, the Ti surface can be partially oxidized, forming TiO_x) is very weak. Therefore, the deposited MHAuNPs can be easily detached from the Ti/Si substrates by sonicating the substrates in solvents (Scheme 1b). The collected MHAuNPs in a solvent are obtained as colloidal particles as shown in Fig. S7.^† Such interesting hemispherical mesoporous nanoparticles have advantages to electrocatalytic activities in comparison to spherical mesoporous metals.³¹ The method using a Ti/Si substrate as a working electrode can be repeatedly implemented with one substrate and without change of the precursor solution, thus it can be effective for mass production in the future.Open in a separate windowFig. 4The SEM image of MHAuNPs deposited at −0.6 V. The arrow shows that the bottom of the MHAuNPs is an arch.Finally, surface-enhanced Raman scattering (SERS) effects on MHAuNPs were investigated by using an adsorbate called rhodamine 6G (R6G), as shown in Fig. S8.^† The resulting MHAuNPs at all conditions (−0.3 V, −0.6 V, and −0.9 V) show substantially strong SERS intensity (Fig. S8a^†), while Ti/Si and Si substrates without MHAuNPs show noise level of intensity. To further investigate enhancement factor (EF) and limit of detection (LoD), various concentrations of R6G with MHAuNPs fabricated at −0.9 V were used for the SERS studies (Fig. S8b^†). The main peak of SERS is 1363 cm⁻¹, and it disappears from less than 10⁻⁶ M concentration, while the 1183 cm⁻¹ peak still exists at 10⁻⁸ M (Fig. S8c^†). The maximum EFs at 1363 cm⁻¹ (10⁻⁶ M) and 1183 cm⁻¹ (10⁻⁸ M) are 1.5 × 10⁴ and 3.1 × 10⁶, respectively (Fig. S8d^†). Transmission electron microscope (TEM) images in Fig. S9^† show the detailed particle structures and the electron diffraction (ED) pattern confirmed the crystal structure is the face-center cubic (FCC) structure. The sharp surface structures and the pores on MHAuNPs provide abundant hot spots that have been reported as the origin that enhances SERS intensity owing to the plasmon resonances.^19,32 Besides, the high density of small-sized MHAuNPs (Fig. S5 and S10^†) boosted higher SERS intensity.In conclusion, we have synthesized MHAuNPs by using 10 nm Ti-coated Si substrates as a working electrode on a Si wafer and electrochemical deposition using self-assembled polymeric micelles as pore-directing agents. The low current generates Au seeds at only a few places, and it acts as the points that MHAuNPs start growing. The particle shapes and sizes can be controlled by changed applied voltages and deposition times. The lower voltages make small particles and the great hemispherical AuNPs with mesoporous architecture. The long-time deposition does not affect any mesoporous formation, but the particle shape and size. Besides, the low affinity between Au and Ti (probably, oxidized layer) results in the arch on the bottom of MHAuNPs, which helps the particles detached from the substrates easily. These results indicate that different thicknesses and compositions of working electrodes can provide different metal deposition phenomena, which can bring out unique shaped particles with mesoporous architectures in the future.     

The synthesis of highly active carbon dot-coated gold nanoparticles via the room-temperature in situ carbonization of organic ligands for 4-nitrophenol reduction

Due to the serious pollution issue caused by 4-nitrophenol (4-NP), it is of great importance to design effective catalysts for its reduction. Here, a novel and simple strategy was developed for the synthesis of carbon dot-decorated gold nanoparticles (AuNPs/CDs) via the in situ carbonization of organic ligands on AuNPs at room temperature. The enhanced adsorption of 4-NP on CDs via π–π stacking interactions provided a high concentration of 4-NP near AuNPs, leading to a more effective reduction of 4-NP.

Due to the serious pollution issue caused by 4-nitrophenol (4-NP), it is of great importance to design effective catalysts for its reduction.

With the rapid development of the global economy and the continuous progress of the society, environmental problems, especially water pollution caused by nitroaromatic compounds, have become increasingly serious during the last decade.^1,2 As one of the toxic phenolic pollutants, 4-nitrophenol (4-NP) is usually found in the wastewater discharged from many chemical industries, which causes severe environmental issues and also a serious toxic effect on the living organisms.^3–6 Hence, 4-NP has been classified as a priority toxic pollutant by the US Environmental Protection Agency.⁷ In contrast, 4-aminophenol (4-AP), a product of 4-NP, is less toxic and an important intermediate that can be applied in many fields.⁸ Therefore, developing a novel method to effectively convert hazardous 4-NP to nontoxic 4-AP is urgent for solving the environmental issues.⁹It has been reported that 4-NP can be reduced to 4-AP by borohydride ions (BH₄⁻) enabled by catalysts, which plays an important role in this process. Among various catalysts, noble-metal materials (especially Au) have been demonstrated to be effective catalysts and are still the main catalysts used for 4-NP reduction.^10–13 In any heterogeneous catalysis process, the catalytic reaction must occur on the surface of the catalyst.^14–16 Therefore, the adsorption of reagents on the catalyst surface is a vital process before the chemical reaction. Recently, the efficient adsorption of reagents with aromatic rings on π-rich supports has been proved based on π–π interactions.^17–19 As an aromatic compound, 4-NP is also a π-rich molecule in nature, and its adsorption on a carbon-based catalyst by these π–π stacking interactions has also been demonstrated.^20,21 By decorating metal nanoparticles (MNPs) on these supports, the catalytic activity of MNPs toward the reduction of 4-NP to 4-AP by NaBH₄ has also been effectively enhanced via a synergistic effect.^22,23 Based on this, various π-rich carbon materials like carbon fibers,²⁴ carbon nanotubes,²⁵ graphene oxide,²⁶ reduced graphene oxide,²⁷ and graphitic carbon nitride²⁸ have been utilized as supports to prepare metal/carbon-based catalysts with enhanced catalytic activity. However, these methods suffer from drawbacks such as a multiple-step preparation process, prior functionalization, high cost and low yields, limiting their wide practical applications.^29,30 More importantly, as described previously, the reduction of 4-NP mainly occurs on the AuNP surface;³¹ thus, we can speculate that anchoring smaller π-rich carbon materials on the AuNP surface may lead to higher catalytic activity for the 4-NP reduction. Carbon dots (CDs) with the structure of sp² carbons^32,33 can adsorb 4-NP via π–π stacking, with the additional advantages of excellent stability and smaller size.³⁴ Thus, the catalytic activity of AuNPs for 4-NP reduction can be improved by the surface decoration of CDs, which, however, has not been reported before.Herein, a convenient and simple method was proposed to synthesize CD-decorated AuNPs (AuNPs/CDs) via the room-temperature in situ carbonization of cetylpyridinium chloride monohydrate (CPC) pre-adsorbed on AuNPs as an organic ligand. This AuNP/CD hybrid was demonstrated to have higher catalytic activity for the reduction of 4-NP, which was about 2.7-fold higher than that of AuNPs. The improved catalytic activity of AuNPs/CDs can be attributed to the synergistic effect between AuNPs and CDs.According to previous research, CPC can be carbonized to carbon dots in the presence of NaOH.^35,36 Inspired by this, we used CPC as both the capping and reducing agent to synthesize AuNPs in the presence of NaOH, and the synthesis process is displayed in Fig. 1A. After adding HAuCl₄ into an aqueous solution of CPC, yellow suspended solids were generated quickly (Fig. 1A(a) and (b)), which should be attributed to the formation of CPC-Au(i) between AuCl₄⁻ and CPC. Fig. 1B shows the UV-vis spectra of the aqueous solution of CPC before and after the addition of AuCl₄⁻. The obvious absorption peak at 340 nm (green) should be due to the generated CPC-Au(i). The generated Au(i), confirmed by the XPS analysis results (Fig. 1D), can originate from the reduction of Au(iii) by CPC because there are no extra reducing agents present in this system. After adding NaOH into the above-mentioned CPC-Au(i) solution at room temperature (Fig. 1A(b) and (c)), the color of the solution gradually changed to dark red (Fig. 1A(c)). The UV-vis spectrum of the above-mentioned mixture solution was then recorded. The absorbance between 280 and 450 nm proved the formation of CDs due to the carbonization of CPC under alkaline conditions;³⁵ moreover, the obvious absorption peak at 520 nm (Fig. 1C (red)), corresponding to the surface plasmon resonance (SPR) of AuNPs,³⁷ indicated the formation of AuNPs in this process. The formation of AuNPs could be attributed to the further reduction of CPC-Au(i) after the addition of NaOH. XPS analysis was also performed to investigate the oxidation state of the Au species in the process of AuNP synthesis. The high-resolution Au 4f XPS spectra shown in Fig. 1D and E indicate the presence of different Au species in these two samples. For CPC-Au(i) (Fig. 1D), the four evident peaks at 84.8, 88.4 eV and 87.4, 91.1 eV were assigned to the Au 4f_7/2 and Au 4f_5/2 signals,³⁸ respectively, indicating the co-existence of nonmetallic Au⁺ and Au³⁺ in this sample.³⁹ However, for AuNPs/CDs (Fig. 1E), two typical binding energies at 86.8 eV and 83.1 eV, attributed to the binding energies of metallic Au(0),⁴⁰ were clearly observed, indicating the successful reduction of Au(iii) and/or Au(i) to Au(0) in the presence of NaOH.Open in a separate windowFig. 1(A) Diagram for the synthesis of AuNP/CD nanocomposites at room temperature. (B and C) The UV-vis absorption spectra of CPC (black), CPC-Au(i) (green), CDs (blue) and AuNPs/CDs (red) in aqueous solutions, respectively. The high-resolution XPS spectra of Au 4f in CPC-Au(i) (D) and AuNPs/CDs (E). Fig. 2A displays the TEM image of the synthesized AuNPs/CDs. From the image, it can be seen that the monodispersed AuNPs/CDs are spherical, with an average size of 3.5 ± 0.8 nm and a standard deviation of the particle sizes of 22.9% (Fig. S1^†). This relatively small size should be attributed to the use of CPC as a capping agent, which can effectively control the growth of AuNPs. Fig. 2B shows the HRTEM image of AuNPs/CDs; the characteristic lattice spacing of 2.06 Å can be attributed to the (200) plane of face-centered cubic (fcc) gold,⁴¹ which also indicates the successful synthesis of Au nanocrystals under these conditions. In addition, the presence of CDs on the AuNP surface can be clearly observed (Fig. 2B and C), and the Raman peak of AuNPs/CDs between 1100 and 1800 cm⁻¹ (Fig. S2^†) also indicates the presence of carbon dots on AuNPs.⁴²Fig. 2D shows the XRD patterns of CDs (black) and AuNPs/CDs (red). There is no diffraction peak for CDs, indicating the amorphous structure of CDs. However, the obvious peaks at 38.9°, 44.8°, 65.1° and 78.1° (red) for AuNPs/CDs, assigned to the diffraction from the (111), (200), (220) and (311) planes, respectively, of fcc Au crystals,⁴³ also prove the successful formation of Au nanocrystals and agree well with the results of TEM. Fig. 2E displays the FT-IR spectra of AuNPs/CDs (red) and CDs (black). Similar groups can also be found on both CDs and AuNPs/CDs, indicating the presence of CDs on the AuNP surface. In a word, the results described above not only prove the successful synthesis of AuNPs, but also the structure of CDs decorated on the AuNP surface based on the in situ carbonization of CPC at room temperature.Open in a separate windowFig. 2(A) TEM and (B) HRTEM images of AuNPs/CDs. (C) The structure of AuNPs/CDs. (D) XRD profiles and (E) FT-IR spectra of CDs (black) and AuNPs/CDs (red).The reduction of 4-NP to 4-AP with an excess amount of NaBH₄ was carried out to quantitatively evaluate the catalytic properties of AuNPs/CDs. In the absence of the catalyst, a strong absorbance peak at 400 nm was observed for the mixture of 4-NP and NaBH₄, which was attributed to the 4-NP ions under alkaline conditions.⁴⁴ After adding the catalysts CDs and AuNPs/CDs into the reaction system, the reduction process was monitored by measuring the time-dependent absorption spectra of the mixed reaction solution. As shown in Fig. 3A, the nearly unchanged absorption even after one hour when only CDs were present as the catalyst indicates the non-active nature of CDs for 4-NP reduction. However, when we used AuNPs/CDs as the catalyst, the absorbance intensity of 4-NP at 400 nm decreased quickly as the reaction time was extended, and this was accompanied by the appearance of an absorbance peak of 4-AP at 300 nm (Fig. 3B), indicating the higher catalytic activity of AuNPs/CDs for the reduction of 4-NP to 4-AP by NaBH₄. This result also indicates that the reduction of 4-NP should be attributed to the catalysis of AuNPs in AuNPs/CDs due to the non-active CDs. It should be noted that the reduction of 4-NP by NaBH₄ could be completed within 10 minutes, with the observation of fading and ultimate leaching of the yellow-green color of the reaction mixture in the aqueous solution. However, a longer reaction time (more than 18 minutes) was required to achieve the full reduction of 4-NP under similar conditions with AuNPs alone (synthesized with trisodium citrate; see ESI^† for preparation details) as the catalyst (Fig. 3C). Fig. 3D shows the absorption changes of the solutions at 400 nm in the presence of different catalysts, and the quicker decrease in absorption indicates the higher catalytic activity of AuNPs/CDs than that of AuNPs. Since the concentration of BH₄⁻ was constant during the catalytic reaction and much higher than that of 4-NP, the rate constants could be evaluated by the pseudo-first-order kinetics using ln(C_t/C₀) = −kt, where k is the apparent first-order rate constant; its value estimated directly from the slope of the straight line can be used to evaluate the catalytic activity of a catalyst.⁴⁵ As shown in Fig. 3E, a good linear relationship of ln(C_t/C₀) versus the reaction time (t) is observed for the two catalysts; the rate constant (k) values were calculated as 1.83 × 10⁻⁴ min⁻¹, 0.11 min⁻¹ and 0.30 min⁻¹ for CDs, AuNPs and AuNPs/CDs, respectively. These results clearly demonstrated the higher catalytic activity of the AuNPs/CDs composites, which was about 2.7-fold higher than that of AuNPs. In addition, the AuNPs/CDs catalyst exhibited better/comparable catalytic activity compared to/to that of previously reported modified gold catalysts in terms of the time needed for the reduction reaction and rate constant (Table S1^†), demonstrating its potential applications in catalysis.Open in a separate windowFig. 3The UV-vis absorption spectra for the reduction of 4-NP by CDs (A), AuNPs/CDs (B) and AuNPs (C). The plots of absorbance (D) and ln(C_t/C₀) (E) versus reaction time for the catalytic reduction of 4-NP based on different catalysts. (F) The mechanism for the enhanced catalytic activity of AuNPs/CDs for 4-NP reduction.Based on the experimental results and analysis mentioned above, the enhanced catalytic activity of AuNPs/CDs for 4-NP reduction should be attributed to the presence of CDs on AuNPs, which leads to a synergistic effect between AuNPs and CDs. This synergistic effect plays an active part in catalysis and therefore enhances the catalytic activity of AuNPs/CDs, which can be explained as follows: as a π-rich molecule, 4-NP can be adsorbed onto the surface of CDs via π–π stacking interactions.^20,23 In addition, the adsorption of 4-NP on AuNPs/CDs can be demonstrated by the decreased absorbance of the 4-NP solution after adding AuNPs/CDs in the presence of NaOH instead of NaBH₄ (Fig. S3^†). This endows AuNPs/CDs with anchoring sites for the adsorption of 4-NP, and such strong adsorption results in a higher concentration of 4-NP near the surface of AuNPs,²⁰ leading to efficient contact between them and therefore speeding up the catalytic process (Fig. 3F(I)). In contrast, for AuNPs without the presence of CDs, 4-NP must collide with AuNPs by chance and must remain in contact for the catalysis to proceed; or else, 4-NP will pass back into the solution and the reaction cannot occur (Fig. 3F(II)) until it reaches the AuNP surface again.⁴⁶ This can also be demonstrated by the increased reaction rate constant for the 4-NP reduction with more catalysts present in the reaction system (Fig. S4^†). In addition, due to the excellent electron acceptor and donor properties of CDs,^47,48 which can obtain/give excess electrons from/to AuNPs (Fig. 3F), like a reservoir of electrons, the electron transfer from CDs to AuNPs increases the local electron concentration and maintains AuNPs in an electron-enriched state,⁴⁹ facilitating the uptake of electrons by the 4-NP molecules.In summary, a nanocomposite of AuNPs/CDs was successfully synthesized via the room-temperature in situ carbonization of CPC as a capping agent for AuNPs. As a catalyst for 4-NP reduction, this composite showed superior catalytic performances to AuNPs, which could be rationally attributed to the enhanced adsorption of 4-NP on CDs providing a high concentration of 4-NP near AuNPs for a more effective reduction of 4-NP. This work not only offers an attractive catalyst material for 4-NP reduction, but also opens an exciting new avenue for the rational design and development of CD/metal nanostructure hybrids for various applications.     

NiFe LDH/CuO nanosheet: a sheet-on-sheet strategy to boost the active site density towards oxygen evolution reaction

The amount of active sites of a catalyst is of great importance to enhance the oxygen evolution reaction (OER) activity. Here, the sheet-on-sheet strategy is proposed to effectively increase the density of active sites of NiFe layered double hydroxide (NiFe LDH) catalyst in terms of structural engineering. As a non-precious electrocatalyst for the OER, NiFe LDH is grown directly on CuO nanosheets. As a result, the received NiFe LDH/CuO nanosheet catalyst with sheet-on-sheet structure shows an ultralow overpotential of 270 mV at 20 mA cm⁻², much lower than that of RuO₂ as a benchmark. The CuO nanosheets, as substrate, play the vital role in downsizing the NiFe LDH, leading to the raised active site density.

The amount of active sites of a catalyst is of great importance to enhance the oxygen evolution reaction (OER) activity.

Electrochemical water splitting is an important technique for facile energy storage and transfer.¹ The oxygen evolution reaction (OER), as a half reaction of water splitting, suffers from sluggish kinetics due to the complicated charge transfer process.³ As a result, a larger overpotential is required to drive the OER process. Relative to the state-of-the-art noble metal catalysts, such as RuO₂ and IrO₂, reliable non-precious electrocatalysts are urgently needed to boost the OER activity and reduce the cost.⁴Currently, mixed NiFe layered double hydroxide (NiFe LDH) electrocatalysts exhibit favourable intrinsic OER activities.^1–3 However, restriction on applications of NiFe LDH lies in the limited electrochemically active surface area and low conductivity.⁴ Some strategies, such as defect engineering,^5–7 and NiFe based ternary LDH,⁸ have been proposed for enhancing the OER activity. Recent studies have shown that the OER activity of NiFe LDH can be further enhanced by structure engineering.^8,9 Especially, NiFe LDH with 3D architecture provides large specific surface area and sufficient active catalytic reaction sties.^10–12 The substrate used for growing the NiFe LDH is important to construct the 3D architecture. However, the substrates, such Ni/Cu foam,^13,14 and porous carbon,^15,16 are commonly in the size scale of micro- or even centi-meter, and leads to the large size of NiFe LDH, which undermines the activity. Hence, downsizing the substrate is essential for decreasing the size of NiFe LDH and improving the catalytic activity. Bin et al. fabricated the NiFe LDH on AgNWs with heterointerfaces which optimized the electronic structures and boost the OER activity.¹⁷ Zhou et al. deposited the NiFe LDH on CuO nanorods to construct a core–shell heterostructure, which demonstrated the overpotential of 290 mV at 50 mA cm⁻² in 1 M KOH.¹⁸ Zhang et al. deposited the NiFe LDH on NiCoP nanoarray, the resulting electrode required the overpotential of 220 mV to deliver 10 mA in 1 M KOH.¹⁹ 1D nano-structure with NiFe LDH directly growing on nanowires seems to be effective for support engineering.Herein, we explore a sheet-on-sheet structure of NiFe LDH catalyst for enhancing the OER activity. We first grow CuO nanosheets (CuO NSs) on Cu nanowires (Cu NWs), then deposit NiFe LDH on CuO NSs, constructing a NiFe LDH/CuO NS structure. The CuO NS offers the nano-sized substrate for NiFe LDH growth, and enhance the number of active sites, thus greatly improving the OER performance of catalysts.The as-prepared Cu NWs with the smooth surface (Fig. S1^†) were used as templates. XRD shows the typical Cu diffraction with peaks at 2θ of 43.3°, 50.4°, and 74.1°, corresponding to (111), (200), and (220) facets (JCPDS 04-0836), respectively (Fig. 1). After treated in NaOH solutions, the leaf-like nanosheets were grown perpendicularly and homogeneously on the surface of CuNWs at high density (Fig. S1^†). These nanosheets are in the triangle shape and have the 2D size of around 200 nm. The diffraction peaks at 35.5°, 38.7°, 48.7°, 58.3°, 61.5°, 66.2° and 67.9° belong to (002), (111), (−202), (202), (−113), (311), and (113) facets of CuO (JCPDS 45-0937), respectively, illustrating the nanosheets consists of CuO. With the formation of NiFe LDH, Cu diffraction signals disappear, due to the reaction of Cu with the ammonia and final dissolution. The peaks at 11.4°, 23.4°, and 59.8° correspond to the (003), (006), and (113) facets of NiFe LDH (JCPDS 40-0215). The XRD results show that the final composition of the as-prepared catalyst is NiFe LDH/CuO NS.Open in a separate windowFig. 1XRD pattern of CuNWs, CuO/CuNWs, and NiFe LDH/CuO NS.The SEM image reveals the remaining of nanowire shape even after removal of CuNWs (Fig. 2). High density of very thin NiFe LDH can be observed. NiFe LDH has the 2D size of around 100 nm. The homogeneous dispersion of O, Cu, Ni and Fe elements can be detected in NiFe LDH/CuO NS. The mass ratio of Cu, O, Ni, and Fe detected by EDS is shown in Table S1,^† the Ni : Fe atomic ratio is close to 4 : 1. NiFe LDH directly grown on CuNWs with different dosages of Ni and Fe precursors (NiFe LDH/CuNWs-L, NiFe LDH/CuNWs, and NiFe LDH/CuNWs-H) were also prepared for comparison. NiFe LDH is found on NiFe LDH/CuNWs-L and NiFe LDH/CuNWs with a low density (Fig. S2^†). NiFe LDH/CuNWs-L shows NiFe LDH with very small size. As the dosage of Ni and Fe precursors rises (LDH/CuNWs-H), both the density and the size of NiFe LDH increase. The size of NiFe LDH becomes large and even reaches one micrometer. The difference in the morphologies of NiFe LDH/CuO NS and NiFe LDH/CuNWs demonstrates that the CuO nanosheets, as nano-sized substrate, are crucial for the morphology engineering of NiFe LDH. A control experiment with different NiFe dosages but the same amount of support was also carried out (Fig. S3^†). The density of NiFe LDH becomes bushier as NiFe dosage increases.Open in a separate windowFig. 2High (A) and low (B) SEM images of NiFe LDH/CuO NS, and the EDS mapping of Cu, Ni, O, and Fe elements.The composition and valence information were surveyed by XPS (Fig. 3). Fig. 3A shows the Cu 2p orbital. The peaks at around 934 and 953 eV are ascribed to Cu 2p_3/2 and 2p_1/2, respectively. The satellite peaks in the range of 938–947 eV relate to Cu²⁺.²⁰ The peaks of Fe 2p at 712.5 and 724.6 eV (Fig. 3B) are assigned to Fe 2p_3/2 and Fe 2p_1/2, respectively, which are typical values of Fe³⁺. The peaks of the Ni 2p at 856 and 873 eV are attributed to Ni p_3/2 and Ni 2p_1/2 (Fig. 3C), respectively, corresponding to Ni²⁺.¹⁷ As for the O 1s region (Fig. 3D), the peaks locating at 532.6 and 531.1 eV are corresponding to the oxygen atoms from adsorbed water on the surface and the OH groups, another peak at 529.4 eV relates to the oxygen in metal–O bonds.²¹Open in a separate windowFig. 3XPS analysis of Cu 2p (A), Fe 2p (B), Ni 2p (C), and O 1s (D) orbitals of NiFe LDH/CuO NS.The linear scanning voltammetries were carried out to evaluate the OER activity in alkaline solutions. As presented in Fig. 4, NiFe LDH/CuO NS behaves much better OER activity than others (Fig. 4A). NiFe LDH/CuO NS merely requires an overpotential of 270 mV to achieve the current density of 20 mA cm⁻², which is much lower than NiFe LDH/CuNWs (310 mV) and commercial RuO₂ (330 mV), respectively. CuO NS/CuNWs show negligible activity, implying that CuO NS has very low inherent OER activity, and NiFe LDH is mainly responsible for the excellent performance. The high OER activity of NiFe LDH/CuO NS continues with the potential going up. The overpotential at 50 mA cm⁻² is 340 mV, still lower than 410 mV of RuO₂. In contrast, the OER activity of NiFe LDH/CuNWs increases slowly at potentials above 1.55 V. The high activity of NiFe LDH/CuO NS lies on the high density of NiFe LDH relative to NiFe LDH/CuNWs. CuO nanosheets, as nano-sized substrates, greatly increase the density of NiFe LDH. NiFe LDH/CuNWs with different NiFe dosages were also prepared and used for activity comparison. As presented in Fig. S4,^† NiFe LDH/CuNWs performs higher activity than NiFe LDH/CuNWs-L and NiFe LDH/CuNWs-H, but they are still better than RuO₂. For low dosages of Ni and Fe, both the size and the density of NiFe LDH decrease, leading to the low OER activity. While for the large size of NiFe LDH and NiFe LDH/CuNWs-H, it decreases the active sites as well, hence leading to the poor OER activity.²² The peaks located at ≈1.43 V (vs. RHE) corresponded to the oxidation of M(ii) to M(iii or iv) (M = Ni and Fe), which is common observed during the electrochemical process of Ni-based electrocatalysts.^23,24Open in a separate windowFig. 4Polarization curves (A) and Tafel plots (B) of CuO NS/CuNWs, NiFe LDH/CuNWs, NiFe LDH/CuO NS, and RuO₂ in 1 M KOH at the scanning rate of 5 mV s⁻¹; Nyquist plots of CuO NS/CuNWs, NiFe LDH/CuNWs, and NiFe LDH/CuO NS (C); the polarization curves of NiFe LDH/CuO NS before and after 1000 cycles of CV (D).The Tafel plots were used to unravel the intrinsic OER activity of NiFe LDF/CuO NS catalysts. As exhibited in Fig. 4B, the Tafel slope of NiFe LDH/CuO NS is 34.7 mV dec⁻¹, which is smaller than that of RuO₂ (142.6 mV dec⁻¹), indicating the much more favourable catalytic kinetics.²⁵ NiFe LDH/CuNWs and CuO NS/CuNWs show the slopes of 82.9 and 256.7 mV dec⁻¹, respectively. The great gap in the slop between NiFe based catalysts and CuO NS/CuNWs confirms the main responsibility of NiFe LDH for the OER activity. NiFe LDF/CuO NS shows the good activity even compared to catalysts reported in literatures (Table S2^†).Nyquist plots obtained by electrochemical impedance spectroscopy (EIS) were used to investigate the charge-transfer process. As shown in Fig. 4C, the smallest R_ct is found on NiFe LDH/CuNWs, due to the excellent conductivity of CuNWs. Both NiFe LDH/CuNWs and NiFe LDFs/CuO NS have smaller R_ct than CuO NS/CuNWs, underlying the better conductivity of NiFe LDH than CuO NS. Generally, the small R_ct of NiFe LDF/CuO NS indicates the fast charge transfer, attributing to the improved conductivity and surface adsorption ability.²⁶ NiFe LDF/CuO NS shows the highest OER activity but lower conductivity than NiFe LDH/CuNWs, which evidences that increasing the active sites is more important than improving the conductivity.The durability was also examined by evaluating the HER activity after 1000 cycles of continuous CV. As shown in Fig. 4D, small decays are shown at high current densities. The overpotential at 50 mA cm⁻² is 340 mV, only a 10 mV decay is present.In summary, NiFe LDFs/CuO NS performs high activity towards OER. The CuO nanosheets, as nano-sized substrates, are crucial to downsize NiFe LDH. As a result, the amount of active sites for OER is greatly increased. Compared with the effect of conductivity, the structure engineering plays a larger role in the OER activity. This sheet-on-sheets strategy is proven to be effective for boosting the active sites and hence, enhancing the OER activity. Undoubtedly, this work provides a facile method to synthesize the nanoscale substrate, which will greatly reduce the size of LDH.     

Transition metals enhance prebiotic depsipeptide oligomerization reactions involving histidine

Biochemistry exhibits an intense dependence on metals. Here we show that during dry-down reactions, zinc and a few other transition metals increase the yield of long histidine-containing depsipeptides, which contain both ester and amide linkages. Our results suggest that interactions of proto-peptides with metal ions influenced early chemical evolution.

Transition metals enhance prebiotic proto-peptide oligomerization reactions through direct association with histidine.

Around half of all known proteins associate with metal ions or organometallic cofactors.^1–4 Metals stabilize protein structure (e.g., zinc fingers) and mediate catalysis (nitrogenases), electron transfer (cytochromes), ligand transport (hemoglobin), and signalling (calmodulin). Metals stabilize nucleic acids; around 500 divalent metal cations associate with a single ribosome.⁵ The dependence of biochemistry on metals appears ancient and rooted in chemical evolution, preceding emergence of life on Earth.^6,7 Long-standing questions about the origins of life might be answered by understanding how metals interact with ancestral biopolymers. For example, metals could have affected oligomerization reactions of early proto-polymers, and could have conferred stability and function, in analogy with their roles in extant biochemistry.Depsipeptides, containing mixtures of ester and amide linkages, are plausible ancestors of polypeptides.^8–14 During drying of hydroxy acids and amino acids at mild temperatures (i.e., 65 °C) ester bonds form first and are converted to peptide bonds by ester–amide exchange. Hydroxy acids and amino acids are thought to have been abundant on the prebiotic earth.^15–19Here we tested the hypothesis that interaction with Zn²⁺ or other metal ions with the amino acid histidine (His) might affect rates and extents of depsipeptide formation in dry-down reactions. We observed that Zn²⁺ and other selected transition metals increase yields of long His-containing depsipeptides. The increase in yield is specific for imidazole-containing amino and hydroxy acids and metals that interact directly with them.We previously showed that His is readily incorporated into depsipeptides during dry-down reactions under mildly acidic conditions (pH ∼3) at 85 °C.¹² These conditions are known generally to promote oligomerization in mixtures of hydroxy acids and amino acids into depsipeptides.^8–11 Here, mixtures of glycolic acid (glc) and l-His were reacted under these known reaction conditions in the presence or absence of various metals.Addition of Zn²⁺ led to an increase in the average lengths of His-containing depsipeptides. Depsipeptide oligomers are readily detected by high-performance liquid chromatography (HPLC) using a C18 column (Fig. 1a). Longer oligomers exhibit longer retention times²⁰ allowing us to compare various dry-down reactions. The yield of products with a higher retention times on HPLC is increased following addition of Zn²⁺ to the reaction mixture (Fig. 1a). ¹H NMR analysis confirmed that the distribution of oligomeric species changes with addition of Zn²⁺ to the dry-down reaction (Fig. 1b–d). MS analysis verified the increased abundance of longer His-containing depsipeptides following dry-down with Zn²⁺ (Fig. S1^†). On the other hand, Zn²⁺ decreases the extent of conversion of His into oligomers. While 42% of His monomer converted into oligomers in the absence of Zn²⁺, only 22% of His monomer converted into oligomers in its presence (Fig. 1b–d). Oligomerization is indicated by shifts in the β-proton resonance envelope centered at ∼3.37 ppm and in the imidazole proton envelopes at ∼7.40 ppm and ∼8.68 ppm (Fig. 1b–d). Extent of conversion to oligomers was calculated based on the integrals of residual nonreacted imidazole proton resonances (Fig. 1b–d).Open in a separate windowFig. 1Zinc increases the yield of long His-containing depsipeptides in dry-down reactions. (a) His monomer was dried with glycolic acid (glc) at a 1 : 1 molar ratio at 85 °C for seven days in the presence or absence of Zn²⁺, at a 1 : 1 molar ratio (His : Zn²⁺). Analysis of samples via C18-HPLC showed a dramatic increase in the yield of longer oligomers in the presence of Zn²⁺. (b and d) ¹H NMR spectrum of a mixture of glc and His in D₂O before (b) and after (c and d) dry-down at 85 °C for seven days in the absence (c) or presence (d) of 1 eq. of Zn²⁺. (e) A possible coordination complex between Zn²⁺ and two His monomers.The observed increase in depsipeptide length upon addition of Zn²⁺ might result from direct association between Zn²⁺ and His monomers (Fig. 1e)²¹ because the effect is Zn²⁺ dose-dependent and is maximal at a 1 : 1 molar ratio of Zn²⁺ to His (Fig. S2^†). Specific interaction of Zn²⁺ with His monomer has been reported previously.²² The increased production of long His-containing depsipeptides with increasing Zn²⁺ is reversed when the number of Zn²⁺ equivalents exceeds that of His. Three equivalents of Zn²⁺ completely inhibited oligomerization reactions (Fig. S3^†). The effect of Zn²⁺ on reactivity of His during the dry-down conditions might arise from either a locked chelate conformation or increased electrophilicity of the carbonyl. Importantly, Zn²⁺ did not cause polymerization of His in the absence of glycolic acid (Fig. S4^†).The effect of Zn²⁺ on depsipeptide formation is not a generic effect and is specific to His. We dried-down mixtures of glc with either alanine (Ala) or lysine (Lys). The addition of Zn²⁺ to a 1 : 1 molar ratio with these amino acids does not increase the yield of products in any length range. In fact, Zn²⁺ inhibited depsipeptide formation for both Ala and Lys (Fig. 2).Open in a separate windowFig. 2Zinc does not increase extent of polymerization or length of oligomers of Ala- or Lys-containing depsipeptides in dry-down reactions. Alanine (a) or lysine (b) were dried with glc at a 1 : 1 molar ratio, for one week at 85 °C under unbuffered conditions, in the presence or absence of Zn²⁺, at a 1 : 1 molar ratio of amino acid to Zn²⁺. The addition of Zn²⁺ hindered depsipeptide formation for both amino acids tested, as opposed to the reverse trend observed with His.In addition to Zn²⁺, we investigated the effects of Na⁺, K⁺, Ca²⁺, Mg²⁺, Cu²⁺, and Co²⁺ on the oligomerization of His in depsipeptides. We dried-down mixtures of glc and His in the presence or absence of various metals (1 : 1 molar ratio of M⁺ or M²⁺ to amino acid). Analysis by HPLC showed that Zn²⁺, Cu²⁺, and Co²⁺, but not Na⁺, K⁺, or Mg²⁺, increased the production of longer His-containing depsipeptides (Fig. S5 and S6^†). Ca²⁺ decreased the production of His-containing depsipeptides (Fig. S5^†).Circular dichroism spectroscopy (CD) supported our hypothesis that enhancement of oligomerization of His in the presence of Zn²⁺ results from direct association between monomeric His and Zn²⁺. We added various concentrations of Zn²⁺ to monomeric glc plus His. The CD spectrum inverts at equal molar ratio of Zn²⁺ and His (Fig. 3a). We attribute the inversion to formation of His–Zn²⁺ complex in concert with a change in the conformation of His. An example of a possible complex is shown in Fig. 1e.Open in a separate windowFig. 3Circular dichroism analysis confirms that zinc and several other transition metals interact with His monomer. Glycolic acid (glc) and His were mixed, at a 5 : 1 molar ratio, in 50 mM Tris buffer (pH 7.2). Circular dichroism (CD) spectra were collected for the mixture in the absence or presence of 1 eq. (referring to the amount of His monomer) of (a) Zn²⁺ or (b) various other metals. CD spectra showed a clear shift in the signal of His following addition of Zn²⁺, Ni²⁺, Cu²⁺, and Co²⁺, but not for the other metals tested.CD spectra of glc plus His only report conformational changes of His because glc is achiral. The conformational change upon Zn²⁺ binding is dose dependent; the change in the CD spectra increases with increasing concentrations of Zn²⁺ (Fig. S7^†). Addition of Zn²⁺ to the dry-down product mixture of His-containing depsipeptides resulted in far more subtle changes in the CD spectra, which appears to arise from binding of Zn²⁺ to small amounts of remnant His monomer that was not converted into polymers during the dry-down reaction (Fig. S8^†).¹² In accordance with the observed species-specific impact of metals on dry-down reactions (Fig. S5^†), the inversion of the CD signal of His monomer was also observed for Co²⁺, Cu²⁺, and Ni²⁺, but not for Na⁺, K⁺, Li⁺, Mg²⁺, or Ca²⁺. Thus, it appears that low-lying d-orbitals of metals are important for interaction with His. The differences in the electron configuration between the different metals affect their metal coordination properties. Transition metals are more electronegative and have more oxidation states than alkaline and alkaline-Earth metals, and their valence electrons in the d-shell tend to promote stable coordination complexes. By contrast, Zn²⁺ inhibited oligomerization of Ala or Lys in dry-down reactions. This inhibition is consistent with recent thermodynamic calculations²³ that examined effects of metals on the monomer–oligomer equilibria of glycine. Metals shift the equilibria toward the monomer, particularly at neutral and alkaline pH.²³To determine if oligomerization of imidazole-containing monomers other than His is promoted by Zn²⁺, we dried l-β-imidazole lactic acid (the hydroxy acid analog of His, herein termed his) with glc for one week at 85 °C. The reaction produced polyesters, co-polymers of his and glc (Fig. S9^†). Addition of Zn²⁺ to his and glc dry-down reaction mixtures increased the yield of longer polyester oligomers (Fig. 4a). ¹H NMR analysis of these polyesters indicate that Zn²⁺ did not change the extent of conversion of his monomer into oligomers: 39% of his converted into oligomers in the absence of Zn²⁺ and 38% of his converted into oligomers in its presence (Fig. S10^†). Therefore, Zn²⁺ does not increase the overall oligomeric yield but rather the distribution of product oligomers, increasing the yield of longer oligomers (Fig. 4a and Fig. S10^†). These results imply that a terminal alcohol can support a chelation complex with Zn²⁺, in analogy with the suggested chelation complex of Zn²⁺ by His (compare Fig. 4b to Fig. 1e).Open in a separate windowFig. 4Zinc increases lengths of his-containing polyesters in dry-down reactions. (a) l-β-Imidazole lactic acid (His) was dried with glycolic acid (glc) at a 1 : 1 molar ratio at 85 °C for seven days in the presence or absence of Zn²⁺, at a 1 : 1 molar ratio of his to Zn²⁺. Analysis of samples via C18-HPLC showed a dramatic increase in the yield of longer oligomers in the presence of Zn²⁺. (b) Possible coordination complex between Zn²⁺ and his monomers.Several distinct non-exclusive mechanisms can explain why Zn²⁺ promotes formation of longer His-containing depsipeptides. Dry-down reactions are conducted under mildly acidic conditions (pH ∼ 3), in which the imidazole side chain (pK_a of ∼6) and the α-amino group (pK_a of ∼9) of monomeric His are protonated and the carboxylic acid (pK_a of ∼2) is deprotonated. Deprotonation of α-amino group would be promoted by His coordination of Zn²⁺ (Fig. 1), favoring an intermediate in formation of depsipeptides through ester–amide exchange. Zn²⁺ coordination by His might also lock His in a specific reactive conformation and/or increase the electrophilicity of the His carbonyl group. Zn²⁺ is expected to pull electron density from His and expose the carbonyl to nucleophilic attack (Fig. 1e). Dehydration would promote His coordination with Zn²⁺ by depleting competing water molecules.^21,24 In principle, it is possible that a complex is formed in which glc and His simultaneously chelate Zn²⁺, or in which His and glc-based oligomers chelate Zn²⁺ such that a favored configuration for a nucleophilic attack is reached.The effects of metals on oligomerization of amino acids by methods other than dry-down reactions has been investigated previously.^23,25–30 Concentrated sodium chloride (>3 M) promotes oligomerization in the presence of Cu²⁺, to increase the yield of glycine (Gly)- and Ala-based peptides.²⁵ Various metals can affect oligomerization of chemically activated amino acids (N-carboxyanhydrides).^31,32 Chemical activation studies focused on Gly, the simplest and most reactive amino acid, but resulted in only low yields of very short peptides.^{23,25–28,33–39} It has been proposed that minerals might catalyze dry-down oligomerization of amino acids.^33–39 McKee et al. observed that silica hinders the amidation of Gly in the presence of lactic acid, the hydroxy acid analog of Ala.⁴⁰ However, silica did lead to an enrichment of amide bonds over ester bonds.⁴⁰His may not be a prebiotic amino acid. It has been proposed that the prebiotic chemical ancestor of His might be imidazole-4-acetaldehyde,^41–44 which is produced by Strecker synthesis.^45,46 The His-containing depsipeptides produced here do not appear to bind to Zn²⁺ (Fig. S8^†). This absence of chelation is consistent with the low number and density of His side chains, and the absence of backbone amines at ester linkages. Longer depsipeptides with greater number and density of His residues may bind Zn²⁺ and might lead to emergence of small metalloenzymes.The importance of small proto-metalloproteins on the prebiotic Earth is supported by the cooperative interactions of metals and proteins in extant biology. Mulkidjanian proposed the zinc world theory, according to which the first metabolism was driven by zinc sulfide minerals that catalyzed photochemical reactions.^47–49 Primordial cooperation may have existed between metals and proto-peptides prior to the emergence of coded protein synthesis.^50–53 For instance, amyloidogenic heptapeptides can function as Zn²⁺-dependent esterases.⁵¹ Zn²⁺ promotes fibril formation by His-containing peptides, acting as a catalytic cofactor. Short peptides with acidic residues, such as aspartic acid and glutamic acid,^52,53 could have protected short RNA molecules against Mg²⁺-induced degradation.^54,55 Coordination of metal ions induces peptide conformational changes and supramolecular assembly.^56–60 In addition to effects on peptide self-assembly and function, metal–peptide interactions are utilized for fabrication of nanofiber materials for various applications, including three-dimensional cell culture and tissue engineering.^56,61,62It is generally accepted that Zn²⁺ concentration has remained fairly constant in seawater through time, whereas the concentrations of Co²⁺ and Ni²⁺ were higher in earlier stages of Earth history than in modern seawater.^63–72 If accumulation of metals occurred at shallow lakes or similar environments that were subjected to dry-wet cycling, they might have affected distribution of polymers that formed within these environments in a specific manner. Our results suggest that the close relationship between metals and biopolymers has roots in prebiotic chemistry and shaped their co-evolution.     

Viscosity effect on the strategic kinetic overgrowth of molecular crystals in various morphologies: concave and octapod fullerene crystals

A kinetic overgrowth allowing organic molecular crystals in various morphologies is induced by temperature-dependent viscosity change of crystallization solution. By this strategy, concave cube and octapod fullerene C₇₀ crystals were successfully obtained by antisolvent crystallization (ASC). The structural analysis of fullerene C₇₀ crystals indicates that the morphological difference is the result of kinetic processes, which reveals that viscosity, the only variable that can change dynamics of solutes, has a significant influence on determining the morphology of crystals. The effect of solvent viscosity in the stage of crystal growth was investigated through time-dependent control experiments, which led to the proposal of a diffusion rate-based mechanism. Our findings suggest morphology control of organic crystals by diffusion rate control, which is scarcely known compared to inorganic crystals. This strategic method will promote the morphology controls of various organic molecular crystals, and boost the morphology–property relationship study.

A kinetic overgrowth allowing organic molecular crystals in various morphologies is induced by temperature-dependent viscosity change of crystallization solution.

Because morphology directly affects the properties of crystals, morphology control of crystals has been one of the major research subjects in chemistry. While the strategies for the morphology control of inorganic metal crystals have been established quite well in the solution phase, primarily through surface chemistry guiding crystallization to occur only on non-passivated crystal planes or through kinetic overgrowth,^1–3 the absence of such strategic methods controlling the morphology of organic molecular crystals restricts their potential. Thus, enormous efforts have been devoted to exploring new approaches to control the morphology of molecular crystals through solvent^4,5 and temperature controls.^6,7As frequently demonstrated from inorganic metal crystals, kinetic overgrowth can be a prominent method to obtain molecular crystals in various shapes. The kinetic overgrowth induces preferred growth at specific sites resulting in morphologies that are not available thermodynamically, as demonstrated from concave cubes and octapods.^8–13 The kinetic overgrowth usually occurs with a concentration gradient around the seed crystals, which can be achieved by regulating the diffusion rate (V_diff) of solutes and crystal growth rate (V_grow). Therefore, it is important to understand and develop kinetic overgrowth process for organic molecular crystals, which can correspond to the well-established surface chemistry for inorganic metal crystals.Considering the importance of concentration gradient near seed crystals for kinetic overgrowth, among various crystallization methods, antisolvent crystallization (ASC) process in which supersaturation and nucleation are induced by the injection of antisolvent to which target molecules have a low solubility (Scheme 1),^14–18 is an ideal method for kinetic overgrowth of molecular crystals since it has many variables that can change micro-environment near seed crystals. In addition, fullerenes are good target molecules because obvious results are expected when kinetic product is successfully contrasted to well-known thermodynamic morphologies such as tubes, rods, and even polyhedrons,^19–23 as Yang et al. and Ariga et al. showed.^24–26 In this case, kinetically favorable fullerene concave cubes could be formed by involving sonication²⁴ and controlling solvent ratio,^25,26 and clearly contrasted with fullerene cube crystals. However, the origin and detailed mechanism of these kinetic overgrowths are still veiled, which prevents further application.Open in a separate windowScheme 1Schematic illustration of traditional ASC process. Target molecules are effectively crystallized via solvation shell mechanism.In this regard, the influence of temperature on kinetic overgrowth of organic molecular crystals should be investigated not only because of its contribution for thermodynamic versus kinetic reaction control in many branches of chemistry,^27–29 but also because it can alter the behavior of molecules significantly during ASC, especially by regulating solution viscosity. Therefore, in this study, we aim to investigate the effect of viscosity upon temperature change for kinetic overgrowth of fullerene crystals. Herein, we show that kinetic overgrowth can be induced by controlling temperature in ASC. Kinetically overgrown fullerene C₇₀ crystals, in concave cube and octapod shapes, are successfully obtained at low temperature, while only cube crystals are obtained at higher temperature. These results originate from increased solution viscosity, which causes slow diffusion of C₇₀ molecules to seed crystals, and consequent kinetic overgrowth. Diffusion rate-based mechanism of inorganic metal nanocrystals is successfully applied to this strategic morphology control.All the fullerene crystallization has been performed by ASC method. Isopropanol (IPA), an antisolvent, is added to C₇₀ solution in mesitylene to obtain cube-shaped C₇₀ crystals at room temperature. After 3 h, black precipitates have been separated from the solution by filtration for characterization. Well-defined faces, edges, and vertices of C₇₀ cubes are confirmed using a scanning electron microscope (SEM), which agrees well with previous report (Fig. 1a).¹⁹ To investigate the effect of growth temperature on the morphology, ASC of C₇₀ has been performed at lower temperature. When the growth temperature is lowered from RT to −16 °C using a refrigerator and −78 °C using dry ice bath, concave cube-shaped and unprecedented octapod C₇₀ crystals are obtained, respectively (Fig. 1b and c). The average sizes of the C₇₀ cube, concave cube, and octapod are 2.0 μm, 2.5 μm, and 1.7 μm, respectively (Fig. S1^†). The resulting crystals show a high degree of homogeneity in their morhpologies.Open in a separate windowFig. 1C₇₀ crystals prepared by ASC at different conditions. (a) Cubes at 25 °C, (b) concave cubes at −16 °C, and (c) octapods at −78 °C. (d) Crystallization condition and obtained morphology of product (scale bar: 2 μm).Because these morphologies are well-known kinetically overgrown products for inorganic metal nanocrystals,^8–13 we have checked the viscosity of crystallization solution at each temperature, which is directly related to V_diff of C₇₀ molecules by the Stokes–Einstein equation.³⁰ The apparent viscosity of the solution at the shear rate of 1000 Hz increases from 2.05 cp to 6.40 cp at −16 °C and further to 218.36 cp at −78 °C (Fig. 1d), which indicates dramatic decrease of V_diff of C₇₀ molecules may occur. From these results, it can be assumed that this morphology difference comes from the slow diffusion of C₇₀ molecules, which is induced by high viscosity at low temperature.The crystal structures of each product have been examined by powder X-ray diffraction (Fig. 2a). For C₇₀ cube crystals, the overall diffraction pattern including the intense peak from (100) plane with d-spacing of 10.56 Å indicates a simple cubic structure as known from previous results.¹⁹ Importantly, the XRD patterns of C₇₀ concave cube and octapod crystals are the same as cube crystals, which implies that the concave cube and octapod morphologies are the results of kinetic crystal overgrowth from cube crystals.²⁵Open in a separate windowFig. 2(a) X-ray diffraction patterns and (b) Fourier transform infrared spectra of C₇₀ cubes (black), concave cubes (blue), octapods (red), and pristine powder (green).Furthermore, to examine the effect of solvent inclusion on the crystal morphologies,²² the intercalated molecules have been analyzed using Fourier-transform infrared spectroscopy (Fig. 2b). Four representative peaks of mesitylene (2717, 2841, 2901, 3019 cm⁻¹) are observed equally from all three types of C₇₀ crystals.²⁵ This result suggests that the intercalation of solvent is irrelevant to the morphology decision. The influence of intercalated mesitylene on the rotational motion of fullerene is also investigated by Raman spectroscopy, and no peak shift also indicates that the intercalated mesitylene does not affect the motion of fullerene molecules in crystal lattice (Fig. S2^†).²⁵ Therefore, we conclude that morphological diversity of C₇₀ crystals is not originated from common structural difference, but from the crystal overgrowth induced by diffusion rate change induced by viscosity change.The overall mechanism of morphology control during ASC via viscosity control can be proposed as follows (Fig. 3). When antisolvent is injected into C₇₀ solution of mesitylene, emulsion droplets containing C₇₀ and mesitylene are formed instantaneously, followed by the diffusion of mesitylene out to continuous phase consisted of antisolvent (Fig. 3a).³¹ As a result, supersaturation and nucleation occur, and seed crystals are formed. To examine if the morphology decision is made at the nucleation stage or growth stage, the seed crystals obtained at 25 °C, −16 °C, and −78 °C are identified by SEM to confirm their cube-shaped morphology (Fig. 3b–d and Fig. S3^†). Therefore, the morphological difference must be induced at the crystal growth step, rather than the nucleation step, which is also on the line of seed-mediated kinetic overgrowth mechanism.Open in a separate windowFig. 3(a) Schematic illustration of temperature-independent nucleation of C₇₀ immediately after the injection of antisolvent, and the SEM images of seed crystals prepared at (b) 25 °C, (c) −16 °C, and (d) −78 °C (scale bar: 500 nm). (e) Schematic illustration of diffusion rate-dependent growth from cubic seed crystals.After nucleation, the reaction system goes through metastable states, where crystals continue to grow and kinetic overgrowth starts to play. All the mesitylene emulsion droplets are broken and release C₇₀ nucleates. Then, C₇₀ molecules remained in the continuous phase at the stage of nucleation are attached to the nucleates. In this stage, the competition between V_diff of C₇₀ molecules and V_grow plays a key role in the determination of crystal morphology, where V_diff is controlled effectively by viscosity (Fig. 3e). At 25 °C, the viscosity of solution is quite small, so C₇₀ molecules can easily diffuse from the bulk solution to the seed crystals (V_diff ≫ V_grow). In this condition, the concentration of C₇₀ around seed crystals is maintained almost same, hence no preferential overgrowth occurs, resulting in seed crystal morphology-retained cube crystals.¹⁹ Whereas, C₇₀ molecules cannot diffuse rapidly at −78 °C due to the high viscosity of the solution (V_diff ≪ V_grow). In this regime, the concentration gradient of solute around seed crystals is generated because of the fast consumption of C₇₀ molecules near seed crystals with slow refill from the bulk solution. Such a concentration gradient induces preferential attachment of distant solutes to the sites possessing the highest reactivity, vertices for cubic crystals, as known for the inorganic metal crystals.^8–13 Eventually, octapod-shaped C₇₀ crystals are formed. The kinetic overgrowth that frequently results in anisotropic crystal growth at specific sites^32–34 rather than thermodynamic isotropic growth resulting in simply bigger crystals was supported from time-dependent SEM images of C₇₀ crystals at −78 °C (Fig. 4) showing petals that are preferentially and continuously grown out of cube crystals. In contrast, an isotropic growth only with a size increase from cubic seed has been observed in the case of the growth at 25 °C (Fig. S4 and S5^†), which indicates thermodynamic crystal growth at low viscosity.Open in a separate windowFig. 4Time-dependent SEM images of C₇₀ octapod crystals obtained at −78 °C. Growth time for each image is (a) 0 min, (b) 1 min, (c) 5 min, and (d) 30 min, respectively (scale bar: 1 μm).To verify if this morphology control is directly related with viscosity change rather than temperature change itself, control experiments using less viscous solvent at the same temperature have been conducted, finding no morphology changes. When acetone is used as antisolvent instead of IPA, no morphological change is observed even at the growth temperature of −78 °C, (Fig. 5 and S6^†) and this result owes to the low viscosity of acetone even at low temperature.³⁵ In other words, the addition of acetone to mesitylene solution does not cause a dramatic viscosity change, which implies the importance of the selection of good solvent and antisolvent for the successful morphology control by ASC process. On the other hand, the use of other alcohols (ethanol, 1-propanol, and 1-butanol having viscosity of 5.26, 6.52, and 7.53 cp at −16 °C and 30.0, 79.8, and 132.0 cp at −78 °C, respectively) show clear kinetic overgrowth at low temperature (Fig. S7^†). Other than the viscosity change inducing morphology control, there is another important property change, i.e. solubility change upon temperature change (Fig. S8^†) to be considered, which requires further studies in the future.Open in a separate windowFig. 5C₇₀ cubes prepared using acetone as antisolvent at (a) 25 °C, (b) −16 °C, and (c) −78 °C (scale bar: 2 μm).     

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.     

Preparation and characterization of an edible metal–organic framework/rice wine residue composite

In this communication, using rice wine residue (RWR) as the support, an edible γ-cyclodextrin-metal–organic framework/RWR (γ-CD-MOF/RWR) composite with a macroscopic morphology was synthesized. The obtained edible composite is promising for applications in drug delivery, adsorption, food processing, and others.

An edible metal–organic framework/rice wine residue composite was made with large surface area for potential applications in drug delivery, adsorption, food processing, and others.

As a typical class of porous materials, metal–organic frameworks (MOFs) have attracted increasing attention since being first proposed by Yaghi and co-workers.¹ Over the past two decades, owing to their large surface area, ultrahigh porosity and tunable pore size,² MOFs have exhibited great prospects for gas storage and separation,^3,4 catalysis,^5–8 sensors,⁹ drug delivery,^10–12etc. Among numerous reported MOFs, γ-cyclodextrin-MOF (γ-CD-MOF), which is connected by the (γ-CD)₆ units of alkaline earth metal ions, was initially synthesized and reported by Stoddart et al.^13,14 in the 2010s. Owing to the –OCCO– groups derived from γ-CD, this kind of MOF is edible and therefore opens a new path for preparing green, biocompatible and edible MOF materials.^13,15,16 For example, Stoddart et al.¹¹ reported a co-crystallization approach to trap ibuprofen and lansoprazole inside γ-CD-MOF, and the resultant composite microspheres can be used for sustained drug delivery. Zhang et al.¹⁷ proposed a strategy to graft cholesterol over the surface of γ-CD-MOF to form a protective hydrophobic layer to improve its water stability. Many researchers succeeded in preparing oral delivery medicine with high drug loading and an enhanced therapeutic effect by combining the drug molecules with γ-CD-MOF.^16,18–20 These works present the excellent application prospects of γ-CD-MOF in the medical field.Since MOFs possess so many attractive advantages, extensive studies have focused on combining MOFs with many other functional materials (metal nanoparticles, quantum dots, carbon matrices and polyoxometalates, etc.) by means of the synergistic effect, leading to the formation of novel composites designed for targeted applications.^21–28 However, these reported composites were still presented as loose powders, which may not be convenient for the applications. Therefore, the question of how to prepare MOFs-based composites for larger particles at low cost is of great significance. On the other hand, as a traditional alcoholic beverage, rice wine has been popular in southern China and some other Asian nations for thousands of years.²⁹ The rice wine lees or rice wine residue (RWR) is a by-product of the fermentation process of rice wine. It is a mixture of proteins, amino acids and polysaccharides. It is traditionally a health food in some Asian nations.³⁰ The edibility, extensive source, low cost and specific macroscopic shape make RWR a potential functional material for further use of MOFs.Herein, a facile and environmental-friendly strategy has been developed to realize the growth of γ-CD-MOF on rice wine residue, resulting in the formation of an edible MOF/RWR composite in the shape of rice grains. The material characterization confirmed the obtained composite possesses the characteristics of MOF. Except for the edible γ-CD-MOF/RWR, other MOF/RWR composites (HKUST-1, ZIF-67 and MIL-100(Fe)/RWR composites; shown in Fig. S1^†) were prepared to demonstrate the universality of this synthesis strategy.The synthesis procedure of the γ-CD-MOF/RWR composite is schematically illustrated in Fig. 1. The rice wine residue was soaked in deionized water for 12 h and then washed with deionized water three times before vacuum freeze-drying. Similar to the synthesis of γ-CD-MOF powder,¹⁵ KOH was dissolved into water. Then certain amounts of the aforementioned dry rice wine residue were soaked into the K⁺-containing solution for 2 h in order to absorb the sufficient potassium ions. K⁺ was then linked by the coordination of –OCCO– units in γ-CD and RWR with the three-dimensional interconnected network. After vapor diffusion of MeOH and some other procedures described in the synthesis of γ-CD-MOF powder (seen in ESI), the γ-CD-MOF/RWR composite (Fig. 2) was obtained. This method is convenient as no extra binders are needed during the whole process. The same procedure was employed to prepare the RWR composites with other MOFs (HKUST-1, ZIF-67 and MIL-100(Fe)). And the syntheses are briefly described in the ESI. The images of the obtained composites are shown in Fig. S1.^†Open in a separate windowFig. 1Schematic illustration of the synthesis procedure of γ-CD-MOF/RWR composite.Open in a separate windowFig. 2Digital photo of the γ-CD-MOF/RWR composite.The rice wine residue, of which the elemental analysis is shown in Table S1,^† is mainly composed of polysaccharides and proteins. Thus, a broad peak at around 22.2° in the XRD patterns of rice wine residue can be observed (Fig. S2^†), which is due to its poor crystallinity.³¹ The XRD patterns of γ-CD-MOF and γ-CD-MOF/RWR composite samples are shown in Fig. 3a. The characteristic peaks at 5.6°, 6.9°, 13.3°, 16.6°, 20.6° and 23.2°, observed from the XRD patterns of γ-CD-MOF, agree with the previously reported works.^32,33 Meanwhile, compared with γ-CD-MOF, the γ-CD-MOF/RWR composite shows similar characteristic peaks with lower intensity, indicating a lower crystallinity of the MOF within the composite. Fig. 3b shows the FT-IR spectra of different samples. Compared with the rice wine residue, the peaks in regions 1 and 2 of γ-CD-MOF and γ-CD-MOF/RWR can be ascribed to the stretching vibration of –CH₂ and –C–O–C– of the MOF, respectively.^15,34 These results further confirm the formation of the γ-CD-MOF in the γ-CD-MOF/RWR composite.Open in a separate windowFig. 3XRD patterns (a) and FT-IR spectra (b) of γ-CD-MOF/RWR composite, γ-CD-MOF and RWR.The SEM images were collected to further investigate the micromorphology of the as-prepared samples. As shown in Fig. 4a, a three-dimensional layered network structure and rich macropores of the rice wine residue rough surface can be seen. γ-CD-MOF (Fig. 4b) exhibits a uniform body-centered cubic shape with an average size of 4.27 μm, which is in accordance with the reported works.^15,35,36 Meanwhile, the images of the γ-CD-MOF/RWR composite (Fig. 4c and d) show that the cubic γ-CD-MOF crystals are well dispersed on the surface of the rice wine residue and even partially integrated into the framework of the rice wine residue. Compared with the pristine γ-CD-MOF, some γ-CD-MOF in γ-CD-MOF/RWR is not an intact cubic structure, exhibiting a significantly different morphology. This suggests a synergistic effect between the MOF crystals and the rice wine residue during the growth of MOF crystals, rather than a simple physical mixture of the two materials. The thermal stability of the γ-CD-MOF/RWR composite was investigated via TGA analysis. As shown in Fig. S3,^† the decomposition temperature of γ-CD-MOF/RWR composite slightly increased compared with those of pristine γ-CD-MOF and rice wine residue. Moreover, the γ-CD-MOF/RWR composite was stable in water, methanol and ethanol (shown in Fig. S4^†) even under mild stirring. These results indicate an improved physiochemical stability of γ-CD-MOF after the incorporation of rice wine residue. This finding further confirms the synergistic effect between them.Open in a separate windowFig. 4SEM images of rice wine residue (a), γ-CD-MOF (b) and γ-CD-MOF/RWR composite (c and d). Fig. 5a shows the nitrogen sorption isotherms of the γ-CD-MOF and γ-CD-MOF/RWR composite. Both pristine γ-CD-MOF and γ-CD-MOF/RWR exhibit typical type-I isotherms, demonstrating their microporous structures. The pore size distributions of pure γ-CD-MOF and γ-CD-MOF/RWR (Fig. 5b) confirm the existence of micropores (between 1 and 2 nm). The calculated Brunauer–Emmett–Teller (BET) surface areas, micropore volume and total pore volume are listed in 35,37 The specific surface area of the γ-CD-MOF/RWR composite is 651 m² g⁻¹, which is significantly higher than that of the pure rice wine residue (10.8 m² g⁻¹). Thus, the increase in the specific surface area of γ-CD-MOF/RWR composite can be attributed to the growth of γ-CD-MOF on the RWR support. Therefore, γ-CD-MOF/RWR composite inherits both the high porosity of γ-CD-MOF and the macroscopic morphology of rice wine residue, which should contribute to its practical applications.Open in a separate windowFig. 5N₂ adsorption and desorption isotherms (a) and pore size distributions (b) of γ-CD-MOF/RWR composite and corresponding comparative samples.Summary of the BET areas (S_BET), micropore volume (V_micro) and total pore volume (V_tot) of γ-CD-MOF, γ-CD-MOF/RWR composite and pure rice wine residue

Experimental review of PEI electrodeposition onto copper substrates for insulation of complex geometries

Polyetherimide (PEI) was used for coating copper substrates via electrophoretic deposition (EPD) for electrical insulation. Different substrate preparation and electrical field application techniques were compared, demonstrating that the use of a pulsed voltage of 20 V allowed for the best formation of insulating coatings in the 2–6 μm thickness range. The results indicate that pulsed EPD is the best technique to effectively coat conductive substrates with superior surface finish coatings that could pass a dielectric withstand test at 10 kV mm⁻¹, which is of importance within the EV automotive industry.

Electrophoretic deposition relying on electrodeposition of charged polymers via modulated electrical fields is reported. Superior surface finishes that could pass a dielectric withstand test at 10 kV mm⁻¹ were obtained for pulsed potentials at 20 V.

Electrophoretic deposition (EPD) is a useful and scalable technique for the production of electrically insulating coatings on conductive substrates having irregular shapes, while at the same time allowing a precise control of the coating thickness.^1,2 EPD relies on the preparation of an aqueous colloidal suspensions based on charged polymers that are deposited as particles onto an electrically conductive substrate by an induced electric field.³ In contrast to other electrical field induced deposition methods such as electroplating, EPD does not require the suspension to have high electrical conductivity, allowing for small power losses due to the total current being used for the coating formation.⁴ Advantages also include that more complex shapes/parts can be coated due to the emulsion being able to reach into difficult geometries, consequently covering areas in hard to reach segments that otherwise would remain uncoated in a dip-coating procedure.^1–4To promote high electrophoretic movement of the polymer particles and homogenous formation of the polymer coatings, both high polymer charge and high polymer colloidal suspension concentration are required in the coating system.^5–7 These factors are considered the most important in the selection of polymer for the EPD systems.⁴ The control over the coating quality has shown, however, to rely on several more parameters, including deposition time, electric field strength, suspension viscosity, applied voltage, etc.⁷ To enhance the polymer charge and solubility, functional groups have been added to the polymer structure and/or a mixture of solvents have been utilized, respectively.^8,9 Furthermore, the thermal stability of the selected polymers needs to be considered to withstand the temperature build-up in the conductive substrate (as an applied coating) during regular operations, e.g., in the automotive parts. Here, polyether ether ketone (PEEK) and polyetherimide (PEI) have shown to be useful in terms of producing homogenous coatings and adequately allowing for thermal post-treatments performed. These post treatments are made to ensure uniform concealing and reformation of the insulating polymer characteristics, i.e., post-coating deposition.^10–12 Among the two, the PEI is the most interesting as it is the least expensive.The EPD of PEI requires a quaternization reaction resulting in protonated PEI (qPEI) for the formation of a charged polymer emulsion suspension and subsequent effective electrophoretic deposition of the PEI onto the conductive substrate.^13,14 The PEI''s imide group undergoes a ring-opening reaction by the addition of 1-methylpiperazine, leading to the formation of an amide group in the PEI structure (Fig. 1a and S1^†). The process is followed by acid protonation of lactic acid, and charge neutralization of the amine group, which in turn allow for the preparation of the polymer suspensions (emulsion) required for the EPD process (Fig. 1a). The inset in Fig. 1a shows the resulting polymer emulsion in the round bottom flask. After the EPD emulsion coating was carried out on the substrate, the coated substrate was treated at high temperature to allow the re-imidization of the PEI, producing the insulating coating (Fig. 1a). A challenge in this context is the possible H₂ gas formation on the cathode electrode, from one of the half reactions in the electrolysis of water, and occasional formation of bubbles/voids in the formed PEI coatings.¹⁵ Therefore, alternating currents (AC) and pulsed direct currents (DC) were evaluated as process alternatives to improve the coating properties.¹⁶ However, to our best knowledge, enlightening comparisons between different voltages, current set-ups, surface substrate preparations, and charged polymer suspension formation have previously not been cross-examined, and compared, in terms of finding optimal conditions for generating useful EPD conditions and formation of uniform insulating coatings.Open in a separate windowFig. 1The electrophoretic deposition process from PEI to qPEI followed by re-imidization to PEI (a), and FTIR transmission spectra of the different samples (b). The coated PEI material was scrapped off prior to the FTIR analysis.In this work, copper substrates with different surface treatments (polishing and sandblasting, Fig. S2^†) and acidic cleaning (HCl and HNO₃), were used for the EPD of qPEI under different potentials (2–20 V), see Fig. 2 and S3.^† The selected potentials were obtained from screening useful conditions with associated electrical field strengths, as well as motivated from previous works.^17,18 The qPEI suspension formation was evaluated in terms of lactic acid/water content and mixing temperature of the solution after the addition of water (Table S1^†). The coated copper substrates with qPEI were always identically heat treated to re-imidize the qPEI forming the electrically insulating PEI layer (Fig. S4 and S5^†). To prepare the qPEI, PEI (20 g) was dissolved in an NMP/acetophenone solution and 1-methylpiperazine was added. The solution was treated at 110 °C (2 h), which promoted the ring-opening of the imide ring in the PEI, from the solvated state (sPEI) to the fully protonated state (qPEI) (Fig. 1a and S1^†). The changes in the FTIR spectra for the corresponding sPEI and qPEI is shown in Fig. 1b. The PEI stretching bands at 1780 and 1720 cm⁻¹ (imide carbonyl of the five-member ring) and 1360 cm⁻¹ (carbon–nitrogen vibration of imide five-member ring) were considerably reduced in the sPEI sample, and almost absent in the qPEI sample. Additionally, the strong band observed at 1670 cm⁻¹ (–C O vibration from amide) demonstrated the successful formation of qPEI (Fig. 1b).¹⁴ The band at 1360 cm⁻¹ and the broad shoulder observed for the band at 1670 cm⁻¹ was ascribed to the characteristics bands of acetophenone (Fig. S6^†). The presence of the bands at 1780, 1720 and 1548 cm⁻¹ confirmed the re-imidation of the qPEI forming PEI coating, see Fig. 1b.Open in a separate windowFig. 2Current density decay on the copper substrate during the EPD process using different constant voltages and substrate pretreatments (a), compassion between constant and pulse voltage 20 V (b), and initial current density response when qPEI suspensions of different storage time were used (c). The result in “c” are the average of triplicates and the error bars the standard deviation.The formation of a qPEI polymer suspension was strongly relying on the combined composition of lactic acid, acetophenone, and water, as well as the suspension temperature. The suspensions formed by 0.35 wt% lactic acid, 79.65 wt% acetophenone, and 20 wt% water, mixing temperature = 90 °C, (Suspension 1, Table S1^†), resulted in solid lumps being formed at the bottom of the reactor (Fig. S7a^†). Increasing the water content to 46 wt% on the expense of the combined amount of lactic acid and acetophenone (Suspension 2, Table S1^†) resulted in a more homogenous suspension without apparent particle agglomeration at room temperature if rapid stirring was applied (Fig. S7b^†). An increase of the lactic acid content to 1.3 wt% and water content to 67.5 wt% (Suspension 3, Table S1^†) led to an unstable suspension that immediately phase separated, unless the suspension was stirred rapidly using mechanical stirring (Fig. S7c and d^†). To stabilize the emulsion, the mixing temperature was decreased to 65 °C and the emulsion turned into a milky, highly viscous liquid (Suspension 4). At lower temperatures (23 °C or lower), a milky suspension with lower viscosity was formed (Suspension 5) that was stable for up to 48 h before any phase separation was observed (Fig. S8^†). The pH of this optimized suspension was 4.6 and showed an electrical conductivity of 126 mS m⁻¹. The addition of water to the qPEI was also found necessary to carry out as slow as possible to avoid excessive foam formation in the making of the optimized suspension, see further Fig. S9.^†The EPD process on the copper substrate using Suspensions 1 and 2 did not produce a coating layer on the substrate. Suspension 1 yielded a suspension with lumps formed whereas Suspension 2 appeared more homogenous (Table S1^†) but no coating was formed, possibly due to limited ionic movement, which is a key factor/requirement for producing homogenous EPD coatings.²Trials using Suspensions 3 and 4 for the EPD process revealed that an increased lactic acid content facilitated the formation of the coating, probably due to improved ionic movement, however none of these solutions produced homogenous coatings. Suspension 5 resulted in evenly covered copper substrates (Fig. S5^†). The pH of the qPEI Suspension 3 was 8.6 ± 0.1 while for Suspension 5 was 4.6 ± 0.1. The difference in the pH for the aforementioned suspensions (both having the same lactic acid/water composition, Table S1^†) could result from different zeta potential values, where very low zeta potential values have been previously showed to have a negative effect on the ionic stability/mobility in colloidal dispersions.² Suspension 5 was therefore used for the further EPD of the copper substrates due to its low viscosity, low mixing temperature, and ability to form homogenous qPEI coatings on the substrate (Table S1 and Fig. S5^†). Fig. 2a shows the current density decay of the qPEI suspension during the EPD process for representative samples. The decrease in conductivity shown in Fig. 2a represents the coverage of the copper substrate (electrode) with the insulating qPEI layer. It should be noted that the decrease in current density was not ascribed to a depletion of qPEI in the solution/suspension. The same suspension could be used several times without an evident decrease in the initial current density. Fig. 2a shows that a sandblasted (Sbl) substrate coated at 7 V (constant voltage) resulted in double initial current density compared to the non-sandblasted substrate, independently of the pre-acid treatment used (7 V HCl Sbl – 7 V HNO₃ Sbl vs. 7 V HCl – 7 V HNO₃, respectively). The sandblasting process increased the surface area of the substrate, which was suggested to have favoured a higher exposure of the substrate in the qPEI suspension. Despite the increase in surface for the sandblasted samples, the current density equilibrium was always reached within less than 15 s, similarly to the non-sandblasted samples (Fig. 2a). The lower voltages in Fig. 2a were not further explored due to inefficient formation of the coatings and lack of electric field gradient in the suspension. A higher voltage of 20 V was instead considered but due to the unwanted half reaction, the pulsed voltage setting was explored. Although the use of 20 V pulsed voltage showed similar initial current density, the pulsed case presented a lower current density decay as compared to the constant 20 V settings, and the current density equilibrium was reached beyond 30 s for the pulsed sample (Fig. 2b). Fig. 2c shows the effect of keeping the qPEI suspension stored for extensive time (2 weeks) on the initial current density (prior to initiating the EPD process). The observed decrease in the initial current density for both 7 V and 20 V EPD treatment suggests that the number of charged particles in the suspensions decreased with longer storage times, which was a consequence of the qPEI emulsion aggregation. Accordingly, the initial current density obtained after storing the suspension for 2 weeks and using constant 20 V was similar to 7 V when used fresh (ca. 0.012 mA cm⁻², Fig. 2c). The result show that for a production of qPEI substrate coating by EPD, the quality/thickness of the polymer coating is affected by the qPEI suspension storage time, but this may not be problematic as long as the voltage of the EPD is used to compensate for the reduced total charge of the emulsion. Fig. 3a shows the result of an unsuccessful EPD coating when using the heterogeneous Suspension 3 (Table S1 and Fig. S7c, d^†). The uneven coverage was revealed in the SEM micrographs, and the oxidation of the copper during the thermal treatment (re-imidization step) is visually observed in the image as a yellow/green coat (Fig. 3a). The sandblasting of the copper substrate in combination with coating from Suspension 5 (7 V HCl Sbl) is shown in Fig. 3b. The micrograph of the surface shows a homogenous and complete PEI coverage. No defects from the thermal treatment of the qPEI coated layer (to produce PEI by re-imidization, Fig. 1a) were observed (Fig. 3b). It is noteworthy that the PEI layer coated the rough sandblasted copper substrate perfectly around the periphery of the copper substrate, demonstrating the ability of the EPD process to evenly coat substrates of complex profiles.Open in a separate windowFig. 3Copper substrate after the EPD process and thermal treatment (qPEI re-imidization) and SEM micrographs of the PEI coatings of samples coated (constant voltage) at 7 V HCl – Suspension 3 (a), 7 V HCl Sbl (b), 7 V HCl (c), and 7 V HCl coating layer after being scratched with a scalpel (d). The coatings in (b) and (c) firmly attached to the copper substrates. Fig. 3c shows a copper substrate that was coated identically as Fig. 3b with Suspension 5, but in absence of sandblasting (7 V HCl). The microscopy revealed the smooth and evenly coated substrate obtained, with copper imperfections revealed through the coating (see photo Fig. 3c, left). A micrograph of the same coated 7 V HCl sample display how a scratch on the surface (with a scalpel after the coating procedure) appeared (Fig. 3d, left). The PEI coating layer had a thickness of about 3–4 μm on the copper surface and the coatings was revealed as firmly attached to the copper substrates, as seen from the encircled area shown with higher resolution imaging (Fig. 3d, right). Nano-sized cracks are observed on the coating close to the cut edge (Fig. 3d, right), which resulted from the plastic deformation of the PEI film coating. The results demonstrate the importance of using a homogenous suspension for obtaining well-distributed polymer and properly coated conductive substrates.A typical defect observed in the EPD coated samples with the post-thermal treatment converting the qPEI to PEI (re-imidized) was the formation of bubbles in the coating film, see Fig. 4a. The effect was ascribed to volatiles formed at the surface of the copper as a consequence of hydrolysis occurring during the EPD process, leaving remains of gas (H₂/O₂) and/or hydroxyl groups that during the thermal treatment allowed condensation of water and resulted in poor consolidation of the coating.^15,19 The formation of these bubbles was observed for all the samples after the drying of the coated substrate when using constant voltage for the EPD (60 °C oven for 17 h). On the contrary, the use of pulsed voltage (20 V) favoured the preparation of a smooth coatings on the substrate under the same drying conditions (60 °C oven for 17 h, Fig. 4b). Here, the presence of defects was limited to the edge of the sample (Fig. 4b). It is suggested that the pulsed deposition limited the hydrolysis reactions in the qPEI suspension due to the overall shorter time with applied potential, reducing the formation of hydrolysed fractions trapped between the coating and the substrate.^17,18,20 However, more extensive drying of the coated copper substrate (60 °C oven for 44 h), also resulted in smooth coatings with no apparent defects, both for the constant 20 V and pulsed 20 V depositions (Fig. 4c and d). The longer drying times accordingly favoured the evaporation of the formed hydrolysed species and facilitated the release of entrapped volatiles. In industrial production, long drying times often represents excessive energy costs. The pulsed DC voltage is therefore suggested as a useful condition for generating smoothly coated substrates in short drying times. A stepwise build-up of the coating also enables otherwise entrapped volatiles to dissipate between each pulse when voltage is absent, resulting in a more homogenous morphology. The cross-section of the coated substrates (prepared via constant or pulsed voltage deposition) are shown in Fig. 4c and d, respectively. The EDS mapping of the coatings display the carbon base of the polymer, whereas the thickness could be established to around 4–6 μm.Open in a separate windowFig. 4Shows the copper substrate coated without (a) and with (b) pulsed EPD conditions at 20 V, after for 17 hours of thermal post treatment at 250 °C. Images (c) and (d) show the same substrates after 44 hours 250 °C heat treatment, with accompanying SEM and EDS mapping insets; without pulsed 20 V deposition and with pulsed 20 V deposition, respectively. (e) Schematic representation of the electrical insulation resistance test setup. Note: it is not drawn to scale, as the coating is about 1/1000 the thickness compared to the size of electrode.The DSC and TGA analysis performed on the removed PEI from the copper substrates showed that no significant difference in thermal properties resulted from carrying out the EPD process, i.e. when comparing the thermal characteristics of the virgin PEI before the quaternization reaction and the final PEI from the coating (Fig. S10a and b^†). A slight decrease in the degradation temperature for the PEI retrieved from the coating layer as compared to the raw PEI (528 °C and 541 °C, respectively, Fig. S11 and S12^†) was however observed. Although this could indicate some molecular changes in the PEI polymer on the coating and/or traces of low molecular weight fractions such as lactic acid (Fig. S12a and b^†), these temperatures are well above the service temperature of electrically conductive copper cables (even considering heat build-up).As an example of the insulating characteristics of the PEI coating, the electrical insulation of the coated and re-imidized samples, i.e. 7 V HCl sample (constant voltage), were evaluated using an insulation resistance test at a temperature of 20 °C (see Fig. 4e). The top electrode had a polished surface with a radius of 2 mm and the electrode was kept at a positive potential of 50 V with reference to the negative sample. For automotive applications, this voltage level may fit for VCA-systems (below 60 VDC).²¹ Considering that the coating thickness was on average ca. 5 μm, these conditions yielded a substantial electric stress of 10 kV mm⁻¹ during the test. A Keithley 2450 source-measure-unit (SMU) was used to apply the voltage and to measure the corresponding current and resistance. Non-coated copper substrates showed a resistance of less than 0.1 ohm, essentially a short circuit (the SMU was current limited). The coated samples on the other hand showed an insulation resistance greater than 10⁹ ohm, illustrating the effectiveness of the EDP method and the re-imidization process in forming an insulating polymeric coating.     

Zein film functionalized with gold nanoparticles and the factors affecting its mechanical properties

In this article, we report a simple method to synthesize biodegradable zein films functionalized with gold nanoparticles (AuNPs) with significantly improved mechanical properties, as an environmentally benign substitute to biologically hazardous polymers. Zein-coated AuNPs were synthesized using the zein protein as a reducing agent and characterized with IR, UV, CD, ζ-potential, and TEM measurements. The zein protein interaction with the negatively charged surface of AuNPs provides excellent strength to the zein thin film. For the first time, FT-IR spectral studies suggested the strong interaction between AuNPs and zein protein, which was further supported by the higher binding constant (K_b) value. The films were characterized for mechanical properties with spectroscopic and physical experimental investigations. The surface morphology of AuNP-doped zein film was explored by AFM and SEM, which suggested that the AuNPs prevent the buckling of zein film and increase the strength as well as flexibility of the film.

A green chemical approach to substitute biologically hazardous polymer.

During the last few years, the use of polyethylene has become a major environmental threat due to its non-degradable nature.^1,2 It has been estimated that ∼8.3 billion metric tons of plastic has been produced since its invention, out of which 40% has been used for packaging.³ Owing to its non-degradable nature, ∼79% of the total amount of plastic produced remains as waste.^4,5 To overcome this, we need an alternative to plastics which should be bio-degradable and economical. Scientists have developed a number of biologically driven thin films using cellulose, lactic acid, and starch, however, these face issues related to tensile strength, hydrophobicity, and flexibility.^6–9 Thin films derived from naturally occurring zein, a predominant corn protein, can serve as a cost-effective substitute due to its intrinsic hydrophobic nature.¹⁰ Zein is a mixture of proteins, mainly consisting of alpha-zein, a 19 kDa protein which is specifically rich in glutamine, proline and alanine amino acids.¹¹ The natural hydrophobicity of zein arises due to its secondary structure where nine anti-parallel helices arranged in a distorted cylinder which is stabilized by hydrogen bonds.¹² Some major physical properties of the film like its moisture barrier property, tensile strength, flexibility, and gas permeability need to be regulated in order to transform zein into an alternative to polythene. The inherent brittleness of zein restricts its application as a structural material.¹³ This can be improved by treating it with plasticizers, such as polyols, oligosaccharides, lipids, and fatty acids prior to casting in order to make zein films more useful.^14,15 Addition of plasticizers increase the strength and flexibility of the zein films, however, excess will make them too weak and sticky to be handled.¹⁶Several studies have been conducted with the aim to improve the functional properties of zein films using various physical and chemical modification techniques.^17–19 Incorporation of nanoparticles (NPs) into the thin films of zein protein can be used as a strategy to impart the desired tensile strength and flexibility to the films by enhancing fibrillation.^20–22 Nevertheless, the mechanism for protein–NP interaction leading to improved mechanical properties is underexplored and still is a topic of great interest due to its wide range of applications.^23,24Here, we report the synthesis and characterization of zein protein film doped with gold-nanoparticles (AuNPs) having enhanced mechanical properties as an alternative to the synthetic polymer. In this regards, first, the AuNPs functionalized with zein protein were synthesized (vide infra) and characterized.^25–27 Binding constant (K_b) and Fourier Transform Infrared (FT-IR) spectral measurement suggest a strong interaction between the AuNPs and zein protein. These interactions provide the ground to develop a method to prepare AuNP doped zein film where the reaction conditions were most optimum and economic (i.e., NTP). AFM, SEM and spectral characterization of zein films suggested that the strong binding of AuNPs with zein enhances the fibrillation and therefore declines the buckling or otherwise provides extraordinary strength to the film. Physical parameters such as tensile strength and strain also demonstrate explicitly that the AuNP doped zein film is an ideal substitute for the biologically hazardous polymer. The present study, for the first time, explains how the AuNP interaction with zein provides enhanced strength and strain to the zein film.AuNPs, doping agent for the zein thin film, was prepared by heating an aqueous solution of HAuCl₄ with zein protein in the presence of SDS (ESI,^† ES).²⁶ The solution color changed from yellowish to pink, indicating the conversion of tri-cationic gold (Au³⁺) to the AuNPs. UV-visible spectral measurement of the above reaction showed the formation of a band at 550 nm (Fig. 1a). Surface Plasmon Resonance (SPR) absorption band at 550 nm is a characteristic feature for AuNPs. The reaction followed the first order rate kinetics, and the pseudo-first order rate constant determined to be 1.25 × 10⁻³ s⁻¹ (ESI, Fig. S1a^†). To determine the best suitable temperature for uniform AuNP formation with the highest yield, we followed the formation of NPs by UV-visible spectral measurements as a function of temperature (20–90 °C) and found it to be 65 °C (Fig. 1b; ESI, ES, and Fig. S1b^†). Also, we determined the optimal zein amount for AuNP synthesis by following the 550 nm peak in UV-visible spectral measurements and observed to be 0.15% with 0.20 mM HAuCl₄ (ESI, ES, and Fig. S2^†) (Scheme 1).Open in a separate windowFig. 1(a) UV-visible spectral changes of zein protein (0.20%; black line) upon addition of HAuCl₄ (0.20 mM, red line) in 10 mM SDS aqueous solution at 65 °C. (b) UV-visible spectral titration monitored at 550 nm for the formation of AuNPs as a function of temperature (from 20 °C to 90 °C). (c) CD spectra of zein (blue line) and AuNPs (red line) recorded in aqueous medium at 25 °C (d) FT-IR spectra of zein (blue line) and AuNPs (red line) recorded in KBr at 25 °C.Open in a separate windowScheme 1To gain insight into the zein coated AuNP formation and their interaction with zein protein, which provides strength to the zein film, circular dichroism (CD) and FT-IR spectroscopic measurements were carried out. The CD spectra of zein exhibited a band at 222 nm due to peptide n–π* transition, characteristic of typical alpha-helical structure and the band at 208 nm represents exciton splitting of lowest peptide π–π* transition (Fig. 1c).^28–30 The change in CD spectra of zein (red) is attributed to the unfolding of zein or immobilization of zein on AuNPs, reflecting the surface functionalization of the AuNPs. The unfolding of zein in aqueous HAuCl₄ solution in the presence of surfactant leads to an exposure of cysteine which initiate the reduction of Au³⁺ to Au⁰.^31–33 This result can be correlated to the change in the secondary structure of zein protein when functionalized on AuNPs.³⁴ To elaborate further, the CD measurement of zein by varying the temperature from 20 °C to 90 °C suggested that zein denatured in the range of 65 °C to 80 °C (Fig. S3a^†). Interestingly, the CD spectrum of a mixture of zein plus AuNPs (45 °C), as otherwise in case of AuNP doped zein film, exactly matches with the CD spectrum of only zein at 80 °C (Fig. S3b^†). This result suggests that the presence of Au³⁺/AuNPs enhances the zein unfolding/denaturation and hence explain the clear temperature-dependent transitions of AuNP formation as or else shown in the UV-visible titration curve (Fig. 1b). The FT-IR spectrum of zein exhibited a characteristic peak at 1650 cm⁻¹,³⁵ corresponding to an amide C O stretching of peptide chain shifted to 1638 cm⁻¹ in zein-AuNPs due to the interaction of amide functional groups with Au⁰ of AuNPs (Fig. 1d; ESI, Fig. S4^†).³⁶ Disappearance of the peak at 1540 cm⁻¹ (N–H), suggests the removal of H⁺ from N–H and therefore create a negative charge on amide N-atom. The distribution of charge on –N–C O moiety is responsible for shifting/decreasing of C O peak to 1636 cm⁻¹. This was additionally supported from the ζ-potential measurement, which suggested a negative charge on the surface of AuNPs (−35.8 mV) (ESI Fig. S5^†), which was further confirmed by gel electrophoresis (ESI Fig. S6^†).³⁷ Furthermore, Transmission Electron Microscopy (TEM) images (Fig. S7a^†) of AuNPs prepared by zein (0.15% w/v) with HAuCl₄ (0.20 mM) showed that the particles were mostly spherical (Fig. S7a^†) with an size around ∼27 ± 5 nm. TEM images of AuNPs of the same sample show clearly that the zein adheres to the surface of AuNPs (Fig. S7b^†).We then attempted to prepare and characterize the biodegradable AuNP functionalized zein-protein film. AuNP doped zein protein film was prepared by mixing the solution of zein (0.75 g, 7.5% w/v) in aqueous ethanol (80% v/v), glycerol (0.25 g, 2.5% w/v) as plasticizer and a freshly prepared AuNP suspension (15% v/v, vide supra) (ESI,^† ES). Different experiments were performed to evaluate various physical parameters of the AuNP doped zein film. In this regards, first, we calculated the binding constant (K_b) for AuNPs with zein protein to understand the strength of the interaction between the two components. We have calculated the K_b by the fluorescence spectroscopic measurements as zein protein is known to show fluorescence due to the presence of tyrosine (5.25%, major component), rather tryptophan (0.16%, minor component).^38,39 The binding constant was found to be 9.22 ± 1.195 × 10¹⁰ M⁻¹, using Stern–Volmer equation, suggesting the reasonably good binding of AuNPs with zein protein (Fig. 2a; ESI, ES and Fig. S8, ESI^†).^40,41 The Hill coefficient (n) determined to be 1.028 (±0.1391), which is due to the enhanced degree of cooperativity in binding of zein with AuNP surface, suggesting a strong interaction between the immobilized AuNPs'' surface ligand and zein.⁴² Interaction of zein with AuNPs and change in the conformation of protein was confirmed using CD.⁴³ The spectra of zein protein with or without AuNPs are different from each other and showed two negative peaks at about ∼207–210 nm and 220 nm with a positive one around 192 nm, suggesting a typical protein with α + β structure, in which intensity of α-helix is more than β-sheet structure (Fig. 2b; ESI, ES, ESI^†).^44–46 This indicates a major conformational change arising from the interaction of AuNPs with zein. Also, we recorded FT-IR spectrum of zein-AuNP thin film showing a broad peak at 1644 cm⁻¹,²³ an additive spectrum of amide C O stretching of zein peptide chain and the AuNP caped zein ligand [(1650 cm⁻¹ + 1638 cm⁻¹)/2 = 1644 cm⁻¹], clearly indicating a very strong binding between zein and AuNP caped zein ligand (Fig. 2c; ESI, Fig. S9^†). This strong interaction contributes to the mechanical strength of the zein-AuNP thin film, which was further confirmed by the estimation of the tensile strength (TS) and strain at failure (E) (ESI,^† ES, and Fig. 4a and b).^47,48Open in a separate windowFig. 2(a) Fluorescence spectral changes of zein protein (0.25 μM; blue line) upon addition of AuNPs (0.14 nM, red line) in 10 mM SDS aqueous solution at 70 °C. (b) CD spectra of AuNP doped zein film (magenta line) and zein (red line) recorded in aqueous medium at 25 °C (c) FT-IR spectra of film (magenta line) and AuNPs (red line) recorded in KBr at 25 °C. (d) TGA graph of the AuNP doped film recorded from room temperature to 500 °C.Open in a separate windowFig. 4(a) Tensile strength and (b) strain at failure of various films made from different amounts of AuNPs vs. zein (15% v/v) (c) AFM image of zein protein film doped with AuNPs. (d) SEM image of AuNPs doped zein film showing the random distribution of NPs (yellow circles) (e) images of the AuNP functionalized film showing the bending at different angles.The morphology of films in the presence and absence of AuNPs were probed with the help of Optical Microscope, Atomic Force Microscopy (AFM) and Scanning Electron Microscopy (SEM). Optical microscope and AFM images confirmed the change in the morphology of these films (Fig. 3a and ESI, Fig. S10^†). The RMS roughness (Sq), amplitude (A), and wavelength (λ) of the buckled zein film were determined by the optical microscope and found to be 120.5 nm, 42 μm, 137 μm (Fig. 3b), respectively. Optical microscope and AFM images of the AuNP doped film showed that there is almost no buckling (Sq = 1.6 nm, Fig. 4c). However, the AFM resolution was not enough to observe small size AuNPs (Fig. 3a); yet the clear difference in surface morphology was observed for films with or without AuNPs (Fig. 4c, ESI, and Fig. S11^†). The internal compressive stress causes buckling formation due to the difference in Young''s modulus at the adjacent region.^49,50 Doping of AuNPs enhances the effective modulus and decreases critical stress, which regulates the disparities in the tensile stress hence eliminate buckling.⁵¹ Surface topography of the film is a key factor in controlling the extent and alignment of the buckling. Alteration of surface topography as shown in the AFM image induced by NPs diminishes the buckling which in turn increase the mechanical strength as well as the smoothness of the film surface (Fig. 4c).⁵¹ Additionally, SEM imaging was used to determine the surface as well as the distribution of AuNPs in the protein film. AuNPs embedded in the zein film appeared to be bigger than their independent size determined by TEM (Fig. 4d). This result also correlates with the shifting as well as the broadening of the SPR band in UV-visible spectra of the film from 550 to 560 nm (ESI, Fig. S11^†).Open in a separate windowFig. 3(a) Optical microscope image of buckled zein protein film. The average wavelength was found to be 150 μm (b) Optical microscope image of AuNPs doped zein protein film. The addition of AuNPs alleviates the compressive stress and prevents buckling.These results are in good agreement with our observation (vide supra) of higher tensile strength for the films incorporated with AuNPs. Also, we have determined the thermal stability of AuNPs functionalized zein thin film from the TG-DSC experiment by following the temperature range from 27 °C to 500 °C. TG-DSC studies showed that the film could be easily used over a wide range of temperature (0 to 200 °C). However, the protein starts decomposing and completely burns out close to 328 °C (Fig. 2d). The thin film prepared from zein plus AuNPs showed the essential physical parameters such as flexibility, tensile strength, and strain required for a wide range of applications. The AuNPs functionalized film showed great flexibility, the bending at different angles suggesting high pliability, which is essential to adopt various from without breaking (Fig. 4e). Also, we have calculated the parameters (TS and E) by varying the ratio of AuNPs to zein. The titration of AuNPs with zein indicates the most suitable ratio for zein and AuNPs is 15% v/v to 7.5 × 10⁻² nM (TS = 7.5 MPa, S = 8.5 × 10⁻³), respectively. However, an increase in the amount of AuNPs to zein significantly decreases both TS and E, suggesting that an optimum amount of AuNPs required to improve the mechanical properties of zein films (ESI, Fig. S10^†). Additionally, Young''s modulus and hardness of AuNPs doped zein films were obtained by fitting the Derjaguin–Müller–Toporov (DMT) model to the section of the force–distance curve where zein and the tip are in contact, and by measuring the adhesion forces between the tip and the zein film.^52,53 Force–distance curves were obtained and was compared with standard (polystyrene). Young''s Modulus was found to be 3.6 MPa (ESI ES and Fig. S12^†).In summary, we have demonstrated the synthesis of a novel zein thin film functionalized with AuNPs with increased mechanical properties. The characterization of AuNPs, as well as their interaction with zein protein, were studied using a combination of spectroscopic and microscopic techniques. We have also confirmed by FT-IR spectral data, for the first time, that the zein protein strongly binds on AuNPs'' surface and also establishes that the free zein protein has a strong affinity towards the AuNPs. Also, physical parameters (tensile strength, strain, and Young''s modulus) supports that the AuNPs'' interactions provide the strength to zein film, which is further explained by the prevention of zein film buckling by the AuNPs interaction with zein protein. We believe that this is because of the decrease in compressive stress due to the uniform distribution of the AuNPs. We envisage that the further development of such proteins films with cheap, readily available metal oxides NPs will allow us to develop the materials for various industrial applications to a large extent.     

Luminometric dosimetry of X-ray radiation by a zwitterionic uranium coordination polymer

A novel X-ray dosimeter based on a uranium coordination polymer U-Cbdcp was obtained by the judicious synergy between the luminescent uranyl centres and zwitterionic tritopic ligands. Notably, U-Cbdcp exhibits luminescence quenching upon increasing X-ray dose, which in combination with its excellent radiolytic stability, makes it suitable for X-ray dosimetry.

A novel X-ray dosimeter based on a uranium coordination polymer has been developed by the judicious synergy between the luminescent uranyl centres and zwitterionic tritopic ligands.

X-ray radiation has been extensively used in medical diagnosis and treatment, security screening, quality control inspection, scientific instrumentation, etc.^1–3 Overexposure to X-ray radiation cause damage to human cells, which could result in skin burn, tissue damage, and increased incidence of cancer.^4,5 Moreover, X-ray dosimetry is required in many industrial fields, including food irradiation, sterilization, and material modification.⁶ Thus, different types of radiation dosimeters, including ionization chamber, scintillator, semiconductor, thermoluminescence dosimeter, chemical dosimeter, and so on, have been commercialized to quantify the incident X-ray dose.⁶ The former three types of dosimeters are more frequently applied to record the dose-rate of incident radiation.^7–10 Thermoluminescence dosimeters and chemical dosimeters are suitable for dosimetry of accumulated dose, but they suffer from critical drawbacks such as cumbersome reading processing, instrument-demand, or cost-ineffectiveness.^11,12 Therefore, further development of new types of X-ray dosimeters remains essential.Coordination polymers, which are assembled from metal ions and organic ligands, have been met with great interest in diverse fields including catalysis, sensing, sorption, separation, and luminescence.^13–18 Their tunability in terms of chemical composition, structure, and more importantly photophysical property, makes them promising for radiation detection. Indeed, pioneering works by Allendorf and co-workers have demonstrated that scintillating metal–organic frameworks (MOFs) assembled from metal cations and radioluminescent organic ligands can function as a new type of radiation detection materials.¹³ Furthermore, coordination polymers or cluster species showing radiochromism, radio-photoluminescence, fluorochromism, and photoluminescence quenching upon accumulated doses of ionizing radiation have been documented, making them as promising candidates of radiation dosimeters.^19–26 Notably, the abundance of luminescent centers or radio-responsive moieties in some of these materials renders higher saturation point in response to radiation dose. This attribute allows for wider operation ranges or higher upper limits of detection compared with those of traditional metal-ion-doped inorganic dosimeters, e.g. Ag-doped phosphate glass and Mg²⁺-doped LiF (LiF:Mg).²⁷We have recently undertaken a study focused on developing actinide-based coordination polymers or cluster materials for their promising applications in ionizing radiation detection.^19,21,28 The large coordination numbers and diverse coordination geometries of actinide cations engender a myriad of topologies of these materials.^29–31 Moreover, the intrinsically intense green emission from the uranyl cation can be utilized as a radio-luminescent center.^23,26 In addition, the slight radioactivity of ²³⁸U (t_1/2 = 4.47 billion years) can be neglected in the course of high dose radiation detection.^32,33 Herein, a novel zwitterionic uranium coordination polymer is reported, showing rather unique fluorescence quenching response to X-ray radiation. This radio-responsive feature, in combination with its high radiolytic stability, points to the potential implementation of uranium-bearing materials for radiation dosimetry.Solvothermal reaction between UO₂(NO₃)₂·6H₂O, zwitterionic N-(4-carboxybenzyl)-(3,5-dicarboxyl)pyridinium bromide (H₃CbdcpBr), and CH₃COOH in DMF/H₂O mixed solution at 100 °C afforded yellow crystals of UO₂(OH)(H₂Cbdcp)(HCbdcp)·4H₂O (U-Cbdcp) with a yield of 63% based on U.Single crystal X-ray diffraction (SCXRD) analysis revealed that U-Cbdcp crystalizes in the monoclinic P2₁/n space group (Table S1^†). The asymmetric unit of U-Cbdcp network consists of one crystallographically independent UO₂²⁺ cation, two Cbdcp ligands, and one hydroxide group (Fig. S1^†). The coordination geometry of uranyl cation can be best described as a typical pentagonal bipyramid, of which four O atoms on the pentagonal plane are donated from three Cbdcp ligand and the rest one is from a hydroxide group (Fig. 1a).^34–38 One of the organic linkers coordinates with one uranyl cation in a μ₁–η₁ bridging mode, while the other one interconnects with two uranyl cations in a μ₂–η₁:η₂ manner (Fig. 1b). Therefore, these two crystallographically unique ligands can be assigned as H₂Cbdcp⁻ and HCbdcp²⁻ with one and two carboxylate group being deprotonated, respectively. The torsion angles between the carboxybenzyl and (3,5-dicarboxyl)pyridinium moieties are measured to be 110.909° and 112.447° for H₂Cbdcp⁻ and HCbdcp²⁻, respectively, as defined by ∠N–C–C. The assembly of uranyl cations, H₂Cbdcp⁻, and HCbdcp²⁻ ligands results in the formation of a one-dimensional infinite chain extending along the c axis (Fig. 1c). The afforded chains are further extended into a 3D supramolecular network via π–π interactions and hydrogen bonds between the ligands (Fig. 1d). The phase purity of bulky U-Cbdcp sample was confirmed by powder X-ray diffraction (PXRD), showing that the measured pattern matches well with the simulated one (Fig. S2^†).Open in a separate windowFig. 1(a) The coordination environment of UO₂²⁺ cation. (b) The coordination modes of two crystallographically independent ligands. (c) The 1D chain of U-Cbdcp extending along the c axis. (d) Representation showing the network of U-Cbdcp. In figure (a)–(c), U atoms are in green, O atoms are in red, N atoms are in blue, and C atoms are in grey.The solid-state luminescence spectrum (λ_ex = 365 nm) was collected on a tablet of U-Cbdcp, that was fabricated from finely ground powder. As expected, U-Cbdcp exhibits five characteristic bands of uranyl cation centring at 488, 508, 531, 556, and 583 nm (Fig. 2a). This intense green emission can be attributed to the HOMO-LUMO transition occurring in the uranyl bonds upon UV excitation.^38,39 Strikingly, the uranyl-based luminescence is strongly quenched after X-ray radiation (4.7 kGy) as shown by the photographs of U-Cbdcp tablet (Fig. 2a inset). Concomitantly, the intensities of characteristic UO₂²⁺ emission bands, which were measured from the tablet exposed to specific interval of X-ray dose, gradually diminished upon continuous X-ray irradiation (Cu-K_α, 120 Gy min⁻¹). More specifically, approximately 44% luminescence intensity was retained after being exposed to 260 Gy X-ray radiation (Fig. 2a). Further increasing the dose to 4.7 kGy resulted in nearly 90% emission quenching. Interestingly, I₀/I as a function of radiation dose can be well fitted with a linear correlation with R² of 0.9988, where I₀ and I are the luminescence intensities monitored at 508 nm before and after irradiation, respectively. This excellent linearity allows for quantifying X-ray dose in a wide dynamic range spanning from 10 to 4700 Gy via a luminescence “turn-off” manner. To obtain limit of detection (LOD), the calibration curve was established by plotting the quenching rate (I₀ − I)/I₀ as a function of dose at the low dose range (0–30 Gy) (Fig. S3^†). The limit of detection (LOD) is calculated to be 0.093 Gy based on the method reported by Zang and coworkers.⁴⁰ Markedly, this LOD is comparable to 0.047 Gy of the most sensitive photochromic sensor Htpbz@Th-SINAP-2.²¹Open in a separate windowFig. 2(a) X-ray dose-dependent fluorescence spectra and optical micrographs (inset) of a U-Cbdcp tablet. (b) The plot showing the linear correlation between I₀/I and X-ray dose.To decipher the quenching mechanism, the structures of U-Cbdcp before and after X-ray irradiation (5 kGy) were thoroughly characterized by combined techniques including PXRD and SCXRD. The PXRD patterns of U-Cbdcp remained approximately unchanged upon irradiation, ruling out our initial speculation of radiation induced damage to the bulky sample (Fig. S4^†). This supposition is additionally supported by the nearly identical FTIR spectrum of irradiated U-Cbdcp with the nonirradiated one (Fig. S5^†). Furthermore, SCXRD analysis before and after X-ray radiation was conducted on the same single crystal of U-Cbdcp and revealed that the overall network derived from these two datasets retain unchanged as well (Table S1^†). In detail, the local structure as represented by the bond length and bond angle of U-Cbdcp changes slightly, which can be attributed to the standard deviations of these parameters obtained from SCXRD (Table S2^†). This observation further excludes the quenching mechanism via decomposition of U-Cbdcp crystal.There is precedence in literature that the luminescence quenching can be associated with the generation of radicals via radio-induced bond break or electron transfer.^23,26,41,42 Therefore, electron paramagnetic resonance (EPR) spectrum of irradiated U-Cbdcp was collected and indeed shows an intense EPR signal with a g-tensor of 2.0197, corresponding to the value (g = 2.0023) of a free electron (Fig. 3).⁴³ The freshly synthesized sample, however, is EPR silent for comparison. To identify the location of radical species in the coordination polymer, EPR spectra of H₃CbdcpBr ligand before and after irradiation were recorded as well. As shown in Fig. S6,^† the irradiated H₃CbdcpBr exhibits a relatively weak resonance with a g factor of 2.0198, which is comparable with that of U-Cbdcp. In the light of aforementioned results, we may conclude that continuous X-ray radiation generates ligand-based radical species, which functions as a quencher via a nonradiative energy transfer pathway.^44–46Open in a separate windowFig. 3EPR spectra of U-Cbdcp before and after 5 kGy X-ray radiation.Encouraged by the structural integrity of U-Cbdcp upon 5 kGy X-ray irradiation, we further investigated its radiolytic stability by irradiating the sample with high dose β-ray and γ-ray radiations. The radiations were provided by a custom-built electron cyclotron (1.2 MeV) and a ⁶⁰Co irradiation source (2.22 × 10¹⁵ Bq) with dose rates of 150 and 11.8 kGy per h, respectively. PXRD study indicated that no obvious changes in long-range order or loss of crystallinity of U-Cbdcp were observed after radiations, implying excellent radiation resistance of U-Cbdcp (Fig. S7^†).In summary, a new 1D uranium coordination polymer built from uranyl cations and zwitterionic Cbdcp ligands were obtained solvothermally. One of the most intriguing properties of U-Cbdcp is the occurrence of luminescence quenching upon X-ray radiation. This unique radio-induced luminometric response can be utilized as a strategy for X-ray dosimetry. Notably, the quenching response can be well fitted with a linear correlation and the detection limit was calculated to be 0.093 Gy. This finding, in conjunction with the excellent radiation resistance of U-Cbdcp, point to potential applications of uranium bearing materials for radiation detection.