The relaxation data were collected over a broad frequency range between 4 kHz and 40 MHz, at body temperature. Thanks to the frequency range encompassing four orders of magnitude, the dynamics of water confined in the polymeric matrix was revealed. It is demonstrated that translation diffusion of the confined water molecules exhibits a two‐dimensional character and the diffusion process is slower than diffusion in bulk water by 3–4 orders of magnitude. As far as rotational dynamics of the confined water is concerned, it is shown that in all cases there is a water pool characterized by a rotational correlation time of about 4×10−9 s. In some of the dermal fillers a fraction of the confined water (about 10 %) forms a pool that exhibits considerably slower (by an order of magnitude) rotational dynamics. In addition, the water binding capacity of the dermal fillers was quantitatively compared. Mechanism of water diffusion: Profiting from the potential of NMR relaxometry to reveal mechanism of molecular motion in condensed matter, we demonstrate the 2D character of the translation diffusion of water in hyaluronic acid gels. The diffusion process is slower than diffusion in bulk water by 3–4 orders of magnitude. The unique advantage of NMR relaxometry lies, however, in the ability of revealing not only the time scale of the motion, but also its mechanism;Nuclear magnetic resonance (NMR) relaxometry is one of the most powerful methods of investigating dynamics in condensed matter.1H nuclei) placed in an external magnetic field can assume two quantum states characterized by the magnetic spin quantum number m=±1/2that correspond to the parallel and anti‐parallel orientations of the 1H magnetic moment (spin) with respect to the magnetic field. The energies associated with the two states differ and hence, according to the Boltzmann distribution, the probabilities of occupying the states are different. As a result of the unequal populations of the states, an effective magnetization is created in a system containing many hydrogen atoms (1H nuclei). When the external magnetic field is changed, the magnetization evolves in time to the new equilibrium state determined by the Boltzmann distribution, requiring re‐population of the two quantum states. This is possible due to the fact that the magnetic moments (spins) of protons interact not only with the external magnetic field but also with each other. The mutual magnetic dipole‐dipole interactions fluctuate in time as a result of molecular (atomic) motion. The probability of the spin transitions depends on the strength (amplitude) of the dipole‐dipole interactions, the time scale of their fluctuations and the mechanism of the fluctuations. In consequence, the spin‐lattice relaxation rate, characterizing the speed of the magnetization evolution towards its equilibrium value, is given as a linear combination of spectral density functions that are Fourier transforms of corresponding time‐correlation functions associated with the dynamics causing the fluctuations and hence the relaxation. Mathematical forms of the spectral density functions are significantly different depending on the mechanism of the motion:Protons H nuclei https://en.wikipedia.org/wiki/Extracellular_matrix. It significantly contributes to cell proliferation and migration. Hyaluronic acid is a major component of skin, where it is involved in tissue repair and regeneration.e. g. in tissues) swells, forming gels which can be used, among other applications, for skin treatments in the form of dermal fillers applied to restore volume of skin and subcutaneous tissue lost for various reasons. The water binding involves formation of hydrogen bonds with the hydrophilic groups of hyaluronic acid: N‐acetyl and carboxyl groups.Hyaluronic acid is one of the main components of the In consequence, this work has two intertwined goals. The first one is to reveal the mechanisms of water mobility in dermal filler. The second is to demonstrate the ability of NMR relaxometry to unambiguously identify the mechanism of water diffusion, with the purpose to exploit it for revealing possible differences in water dynamics in healthy and pathological tissues.21H spin‐lattice relaxation originates from 1H‐1H magnetic dipole‐dipole interactions. In the present case, taking into account the prevailing water content, the relaxation reflects dynamics of different pools of water in the system. Generally, one can expect a least two pools of water: This implies that the overall 1H spin‐lattice relaxation rate, R1ω, is given as [Eq. n as Eq. :(1)R1ω=RR1confωand R1bulkωdenote the relaxation contributions associated with the confined and free pools of water, while ωis the 1H resonance frequency in angular frequency units. It is important to stress that the pools of water should not be treated as “isolated fractions” – there are exchange processes at a high rate between the bulk and confined states. As far as the pool of confined water is concerned, the relaxation rate, R1confω, stems from intermolecular and intramolecular 1H‐1H dipole‐dipole interactions modulated by translational and rotational dynamics of water molecules, respectively. Thus, one can rewrite Equation (1) as [Eq. where ) as Eq. :(2)R1ω=RR1conf,interωand R1conf,intraωdenote the intermolecular and intramolecular relaxation contributions, respectively, while A=R1bulkωdenotes the relaxation rate of bulk water. The last term is frequency independent as a result of the fast dynamics of bulk water. According to spin relaxation theory, the relaxation rates are given as linear combinations of spectral density functions characterizing the dynamical process leading to the fluctuations of the corresponding dipole‐dipole interactions and hence the relaxation, i. e. [Eq. where . e. Eq. :1, 2, 3,Cinter/intradenote the dipolar relaxation constants for the intermolecular and intramolecular relaxation contributions, respectively. The spectral density function, Jintraω, takes the form [Eq. where form Eq. :1, 2, 3,τrotdenotes the rotational correlation time characterizing fluctuations of the 1H‐1H dipole‐dipole coupling in confined water molecules. The form of the spectral density, Jinterω, associated with translation diffusion of the confined water molecules depends on the dimensionality of the diffusion. Anticipating the results, for 2D dynamics the spectral density function yields [Eq. where elds Eq. :8, 10 converges to Jinterω∝-τtranslnωτtranswhich means that the relaxation rate shows a linear dependence on lnω. Eventually, after subtracting the contribution to the overall relaxation originating from bulk water, the relaxation rate R1ωis given (under the assumption of 2D translation diffusion of the confined water) as [Eq. where ) as Eq. :(6)R1ω=C 3Water dynamics in five kinds of dermal filler has been investigated by means of NMR relaxometry. The compounds are labelled as DF1–DF5; their properties are summarized in Table 1H spin‐lattice relaxation studies have been performed for these compounds at body temperature (37 °C) in the frequency range spanning four orders of magnitude, from 4 kHz to 40 MHz. The relaxation process was found to be single‐exponential for all samples at all magnetic fields; examples of 1H magnetization curves recorded versus time are shown in Figure 1H spin‐lattice relaxation dispersion data for the set of dermal fillers, DF1–DF5, is shown in Figure A=0.29s−1, and normalizing the remaining relaxation rates to unity at the lowest frequency, the graphs shown in Figure lnωover two frequency decades; such a dependence is a marker of 2D translation diffusion. One might attribute the diffusion in the reduced dimensionality (2D) to a translation motion along the hyaluronic chains. The 2D diffusion suggests that the chains form locally flat surfaces.The τtransand τrotfor DF1, DF2 and DF3 are very similar, while the dipolar relaxation constants scale by factors which can readily be estimated from the normalization procedure. The analysis of the relaxation data for DF1, DF2 and DF3 and the decomposition into the R1conf,intraωand R1conf,interωcontributions is shown in Figure τtrans=d2Dtrans,d is referred to as distance of closest approach and can be estimated by the diameter of water molecule (2.75 Å).As the relaxation rates overlap, one can easily conclude that the correlation times 1H water relaxation rates in confinement are proportional  to the fraction of bound water molecules, providing the exchange lifetime, τex, between the confined and bulk water pools is sufficiently short, so it can be omitted in the dominator of Equation P and q quantities denote the mole fraction of water protons in the bound position  and the coordination number, respectively.The R1(norm)conf,intra(ω), can be treated as R1(ω)for Pq=1. Comparing Equation (7) with Equation (2) and Equation (3) one can easily see that Equation (2) takes the form of Equation (7), when Cintraincludes the Pq product. Analogously, the intermolecular dipolar relaxation constant, Cinter, includes the number of confined water molecules (number of protons) per unit volume, i. e. it is proportional to the Pq value. The ratios between the dipolar relaxation constants listed in Table CinterCinter(DF1)and CintraCintra(DF1)) are in agreement with the ratios between the hyaluronate contents in DF1, DF2 and DF3, that yields: 1.2 : 1.8 : 2.6 (i. e. 1 : 1.5 : 2.16). This implies that the water binding capacity of this group of dermal fillers is very similar. The translational diffusion coefficient transD is about 1×10−12 m2/s and it is about 3×103 times smaller than the diffusion coefficient of bulk water.transD denotes the relative diffusion coefficient of the confined and confining molecules, but one can expect that the water diffusion is much faster than the diffusion of the polymer. The rotational correlation time of the confined water molecules, τrot, is of the order of 4×10−9 s, yielding the τtransτrotratio of 17–21. The ratio is typical for liquids.A, yielding (0.32–0.33) s−1, is close to the relaxation rate of bulk water.The coordination number (number of protons bounded to a single confining molecule) can be here considered as a measure of water binding capacity of the dermal fillers. The relaxation rate, Equation  and EquaEquation  one can Diffusion of water confined in DF4 and DF 5 is also of 2D character. However, as the translation dynamics in DF4 and DF5 is slower, the linear dependence is limited to lower frequencies. However, the rotational dynamics is more complex. The concept of a single pool of confined water breaks down as the relaxation data cannot be interpreted in terms of Equation (6). The data indicate the presence of two pools of confined water that undergo rotational dynamics on different time scales, although their translation dynamics is the same. This requires extending Equation (6) to the form Eq. :(8)R1(ω)Cintraand Cintra(extra)include the Pq and Pq(extra) factors, respectively. The translation diffusion coefficient for DF4 has a value smaller than the values for DF1, DF2, DF3 by a factor of about 5. The rotational correlation time for the dominating pool of confined water is of about 4×10−9 s, in analogy to the DF1, DF2, DF3 group. The resulting τtransτrotratio (33) is still in the range reported for several liquids.τrot(extra), yields about 5×10−8 s). From the ratios CinterCinter(DF1)=1.41 and Cintra+Cintra(extra)Cintra(DF1)=1.44 one might conclude that the water binding capacity is similar than for to the DF1, DF2, DF3 group. It should be noted that the ratios for the intermolecular and intramolecular dipolar relaxation constants should be the same, as the whole fraction of confinement water participates in translational and rotational dynamics; the small differences stem from the fitting procedure keeping both of them independent. Comparing the Cintraand Cintra(extra)dipolar relaxation constants one can determine that that the pool of water molecules undergoing slower rotational motion includes about 10 % of the confined water. Eventually, one should mention that the A term for DF4 is close to the relaxation rate of bulk water.The translation diffusion of water confined in DF5 is an order of magnitude slower than that of water in DF1, DF2, DF3. Again, the rotational correlation time for the dominating pool of confined water is of about 4×10−9 s. However, in this case the rotational dynamics of the second pool is even slower; the ratio τrot(extra)τrotreaches about 20. The CinterCinter(DF1)=0.88 and Cintra+Cintra(extra)Cintra(DF1)=0.91 values show that the water binding capacity for DF5 is considerably lower; following the hyaluronate concentration it should yield 2.92 (i. e. about a factor of 3 more). In analogy to DF4 the pool of water molecules undergoing slower rotational dynamics constitutes about 10 % of the confined water. The A term for DF5 is about a factor of 1.5 larger than the relaxation rate of bulk water, but this does not change the overall scenario of the water dynamics.Figure 41H NMR relaxometry has been exploited to enquire into the mechanism and time scale of water dynamics in selected hyaluronic dermal fillers. It has been found that in all cases the translation diffusion of the confined water is of 2D character and the corresponding diffusion coefficients are smaller than for water in bulk by 3–4 orders of magnitude. Interestingly, for dermal fillers labelled as DF1, DF2 and DF3, the translation diffusion coefficients are almost the same. Following the line of similarities, it has been demonstrated that in all cases there is a pool of the confined water molecules undergoing rotational dynamics characterized (in a very good approximation) by a correlation time of 4×10−9 s. For DF1, DF2 and DF3 the whole fractions of confined water exhibit such rotational dynamics. For DF4 and DF5 about 10 % fractions of the confined water perform rotational motion on a significantly longer (by an order of magnitude) time scale. Moreover, it has been shown that water binding capacity of the dermal fillers DF1, DF2, DF3 and DF4 are very similar, while for DF5 the water binding capacity is about a factor of 3 lower.Independently of the knowledge about water dynamics in hyaluronic dermal fillers gained from the studies, the results demonstrate the unique ability of NMR relaxometry to reveal in a single experiment not only the time scales of molecular motion in a broad time range, but also the underlying mechanisms of the motional processes, rendering this method a powerful tool for studying dynamics in biological systems.1H spin‐lattice relaxation measurements were performed using a commercially‐available NMR Fast Field‐Cycling relaxometer . The sample temperature was controlled with an accuracy of 0.5 °C. Below 10 MHz the pre‐polarized sequence was used, while above 10 MHz the non‐polarized sequence was applied.The The authors declare no conflict of interest.


However, fabrication of high‐quality CsPbI2Br perovskite film normally requires annealing at a high temperature (>250 °C) that is not compatible with the plastic substrate. In this work, a Lewis base adduct‐promoted growth process that makes it possible to fabricate high quality CsPbI2Br perovskite films at low temperature is promoted. The mechanism is attributed to synthesized dimethyl sulfoxide (DMSO) adducts which allow a low activation energy route to form CsPbI2Br perovskite films during the thermal annealing treatment. A power conversion efficiency (PCE) of 13.54% is achieved. As far as it is known, this is the highest efficiency for the CsPbI2Br solar cells fabricated at low temperature (120 °C). In addition, the method enables fabrication of flexible CsPbI2Br PSCs with PCE as high as 11.73%. Surprisingly, the bare devices without any encapsulation maintain 70% of their original PCEs after being stored in ambient air for 700 h. This work provides an approach for preparing other high performance CsPbX3‐based perovskite solar cells (PSCs) at low temperature, particularly for flexible ones.All‐inorganic CsPbX Therefore, it is highly desirable to develop low temperature processed methods for preparing high quality CsPbX3 inorganic perovskite films which can not only simplify the complicated device process, but also promote emerging flexible device technologies.Recently, all‐inorganic cesium lead halide perovskites  over 14%, but also exhibited good thermal and illuminated stability.2Br with superior photovoltaic performance at low temperature. Until now, there have been few reports on the preparation of CsPbI2Br PSCs at low temperature. For example, Wang et al. applied HPbI3 to replace PbI2 as a precursor for preparing CsPbI2Br film.2+HPbI3) can decrease the formation energy barrier for the α‐CsPbI2Br phase to help form a compact and pinhole‐free film. The champion CsPbI2Br PSC achieves a PCE of 10.56% and exhibits long‐term phase stability at 130 °C. Lau et al. incorporated strontium into CsPbI2Br to prepare PSCs at low temperature (100 °C).0.98Sr0.02I2Br PSCs delivered a highest PCE of 11.3% and better thermal stability. Most recently, Rao et al. controlled the morphology of CsPbI2Br film by dimethyl sulfoxide (DMSO) solvent and obtain a pure α‐phase CsPbI2Br at low temperature (120 °C).2Br PSCs prepared at low temperature (≤150 °C) are still lower than 12% and 8%, respectively.Among CsPbX2(DMSO), PbBr2(DMSO) adducts are obtained via treating lead halides with DMSO under low temperature conditions (60 °C), which can efficiently prevent rapid reaction of precursors and slow down the crystal growth. On the other hand, DMSO adducts can also reduce the formation energy of CsPbI2Br perovskite that needs to be overcome by conventional precursors (PbI2 and PbBr2). By using the as‐prepared DMSO adducts instead of commercial PbI2 and PbBr2 for the perovskite precursor solution, a high‐quality perovskite film can be successfully prepared under low temperature conditions (120 °C), which is free of pinholes and impurities, and has high crystallinity and stability. Using the low temperature crystallized perovskite films, a high performance CsPbI2Br PSC with PCE of 13.54% is fabricated (the current best CsPbI2Br PSCs fabricated at low temperature (≤150 °C) have a PCE of <11%).2Br PSCs with a high PCE of 11.73%. Furthermore, the flexible CsPbI2Br PSCs are very stable, which can maintain 90% and 70% of their original PCEs after being bent under a curvature radius of 5 mm and stored in an ambient environment for 700 h, respectively. This work not only provides a novel route for preparing high‐quality CsPbI2Br perovskites at low temperature but also represents an important step for their application in high‐performance and low‐cost flexible electronics.In this work, we propose a low temperature scheme of DMSO‐adduct promoted process (DAPP) for high quality perovskite films. The PbInditions 0 °C, whiFigure2 and PbBr2, the formation of all‐inorganic perovskite materials needs to overcome a large energy barrier (Ea). Interestingly, it is found that the formation of all‐inorganic perovskite materials undergoes two‐stage reactions in a Lewis base adduct‐promoted pathway, PbI2(Lewis base)X formation and perovskite formation. The first step needs to overcome a small energy barrier (Ea1) from PbI2 to PbI2(Lewis base)X, meaning that PbI2(Lewis base)X can be obtained at low temperature for a short time. Ea2 is the Lewis base adduct‐promoted pathway from PbI2(Lewis base)X to perovskite, which is substantially lower than that of the conventional pathway (Ea). Therefore, a high‐quality perovskite film can be readily obtained at low temperature for a short time through the Lewis base adduct‐promoted growth.In thermodynamics, activation energy or formation energy represents the energy required for a chemical reaction. According to the recent literature, the reaction coordinate diagram of the all‐inorganic perovskite material formation via different pathways is presented in 2Br. The DAPP is a two‐stage reaction pathway with a lower activation energy and different from the conventional single pathway with a higher activation energy. First, the DMSO adducts, PbI2(DMSO), and PbBr2(DMSO), are synthesized at 60 °C as precursors to replace PbI2 and PbBr2 in conventional all‐inorganic perovskite solution. On the one hand, the DMSO adducts can effectively retard the fast reaction between PbI2, PbBr2, and CsI during the evaporation of solvent. On the other hand, the DMSO adducts can reduce the formation energy of CsPbI2Br perovskite as discussed before.Therefore, in order to obtain high quality inorganic perovskite materials at low temperature, in this work, we propose a DMSO adduct‐promoted process (DAPP) for formation of all‐inorganic perovskite CsPbI2(DMSO) and PbBr2(DMSO) adducts contains two steps, namely, I) adding toluene as a nonsolvent into PbI2 or PbBr2 solution dissolved in DMSO to obtain white precipitation and II) annealing the white precipitation at 60 °C for 24 h.2(DMSO) and PbBr2(DMSO) phase reported in literature.2 powders, the absorption edge of PbI2(DMSO) precipitation is removed to short wavelength corresponding to the change from yellow‐colored PbI2 powered to colorless PbI2(DMSO) adducts, which is in agreement with the reported results.The fabrication of PbI2Br precursor films are deposited on the rigid substrate by coating precursor solution via antisolvent ways.2Br perovskite, the CsPbI2Br precursor films without annealing are stored in a glove box for 24 h . It is found that the CsPbI2Br precursor film without DAPP turns black, indicating crystallization of CsPbI2Br film. Conversely, the CsPbI2Br precursor film with DAPP shows dark brown, meaning that the DAPP can slow down the crystallization process of CsPbI2Br. In order to further confirm the influence of the DAPP, the CsPbI2Br precursor films are annealed at low temperature (35 °C) for different times , the CsPbI2Br films without and with DAPP turn black and light brown after being annealed at 35 °C for 35 min, respectively. Therefore, it is concluded that the DAPP can control the crystallization kinetic of CsPbI2Br perovskite, which is crucial for high quality CsPbI2Br film.Then the CsPbIFigure2Br perovskite films with and without DAPP annealed at 65 °C or 120 °C for 10 min on a hotplate. Many studies have found that the cubic α‐CsPbI2Br phase is often obtained when the annealing temperature is higher than 250 °C.2Br phase transition easily occurs by using DMSO adducts as precursor even at 65 °C for 10 min. The characteristic Bragg peaks of CsPbI2Br thin films at 2θ = 14.6°, 29.5° show that both films are well‐oriented in the cubic (100) direction. The peak splitting of the (100) and (200) planes is observed in the control film  annealed at 65 °C for 10 min, indicating the separate growth of I‐rich  crystals.2Br films is enhanced after annealing at high temperature (120 °C), which benefits from the improved crystallinity of the CsPbI2Br film. The intensity ratio of CsPbI2Br films with and without DAPP is 1.2, which means that the quality of CsPbI2Br films is improved after DAPP optimizing.2Br films with and without DAPP are shown in Figure S6 (Supporting Information). The CsPbI2Br film with DAPP exhibits an increased absorption over the entire absorption range compared to CsPbI2Br film without DAPP, which is due to its higher crystallinity  and time‐resolved PL (TRPL) measurements are performed . The PL intensity of the CsPbI2Br film with DAPP is much higher than that of the film without DAPP, suggesting a significantly reduced nonradiative recombination rate within the perovskite film. The PL decay lifetimes of the CsPbI2Br films are determined by TRPL measurements . The CsPbI2Br film with DAPP shows a much longer lifetime (11.57 ns) than that of the film without DAPP (3.41 ns), indicating that the improved crystalline quality of CsPbI2Br film can significantly prolong the lifetimes of carriers.The UV–Vis absorption spectra of CsPbI2Br PSC on the rigid substrate. The device has a structure of (FTO/TiO2/CsPbI2Br/Spiro‐OMeTAD/Au). The energy band levels of the CsPbI2Br PSC are shown in Figure S8 (Supporting Information). In this device, the 300 nm CsPbI2Br perovskite layer is prepared by PbI2(DMSO), PbBr2(DMSO) and CsI, which is free of pinholes and has a compact surface  that is critical for the photovoltaic performance of CsPbI2Br PSC. The DMSO in DMSO adducts, PbI2(DMSO) and PbBr2(DMSO), could not easily escape from the precursor film after spin‐coating, which could enhance the mass transport and diffusion, slow down the rate of crystallization, and eventually improve the film quality.2 layer, some pinholes can be observed on the film probably due to the rapid crystallization , which could lead to the lower PCEs and instability for inorganic PSCs. This result indicates that high quality CsPbI2Br perovskite film can be obtained by using DMSO adducts through a low temperature solution process. The representative current density–voltage (J–V) curve of the CsPbI2Br PSCs is given in Figure VOC of 1.18 V, a FF of 77.23%, a JSC of 14.89 mA cm−2, and a PCE as high as 13.54%, which is one of the highest PCEs among the all‐inorganic perovskite PSCs to date and much higher than the PCEs of the reported all‐inorganic perovskite PSCs that prepared at low temperature . Figure 2Br device with an integrated photocurrent density of 14.20 mA cm−2, which is close to the JSC derived from the J–V measurement.Figure 2Br PSCs have been successfully fabricated on rigid substrate under low temperature (120 °C), the high performance flexible CsPbI2Br PSCs on flexible substrate are expected. The flexible CsPbI2Br PSCs are further fabricated by utilizing ITO/polyethylene terephthalate (PET) as the conductive transparent electrode/substrate. Nb2O5 film is e‐beam evaporated at room temperature as an electron transport layer.2O5 film on flexible PET substrate is shown in Figure2O5 film has a thickness of 60 nm and shows a smooth surface morphology with uniform grain structure. As shown in Figures 2Br film on flexible substrate with DAPP has a dense and smooth surface, and passivated grain boundary. In contrast, the control CsPbI2Br film without DAPP has a rough surface and some cracks, which is probably due to the rapid crystallization process.Considering that the high performance CsPbI2Br films on the flexible substrates, the CsPbI2Br solution with DMSO adducts are spin‐coated on Nb2O5/ITO/PET substrate and then annealed at different temperatures  on a hotplate to form a 300 nm CsPbI2Br perovskite film. The XRD results of CsPbI2Br films are shown in Figure 2Br thin films show the Bragg peaks at 14.6° and 29.5°. As the annealing temperature increases, the intensities of the diffraction (100) peaks increase while the full width at half maxima (FWHM) decreases due to the increased crystallinity of perovskite film. However, with further increases of the temperature to 140 °C, the intensity of the diffraction peak decreases while the FWHM increases, which is probably due to the influence of thermal deformation behavior of PET substrates.2Br films at the optimal annealing temperature is also investigated . The intensities of (100) peak increase from 2 to 30 min due to the increased crystallinity of perovskite film. With further annealing up to 40 min, the change of the (100) peak intensity is not obvious.In order to investigate the effect of temperature on the crystallization of CsPbI2Br PSCs, as shown in Figure2O5/CsPbI2Br/Spiro‐OMeTAD/Au) is adopted. The energy band levels of the flexible CsPbI2Br PSC are shown in Figure S12 (Supporting Information). The effect of different annealing temperatures on the photovoltaic performance of the CsPbI2Br PSCs is investigated. The detailed photovoltaic parameters are summarized in Table2Br PSCs are increased with increasing temperature due to the increased crystallinity of perovskite film, and the CsPbI2Br PSC exhibits the best performance when the temperature is 130 °C. When annealing temperature up to 140 °C, the performance of device decreases, which may be due to the influence of thermal deformation behavior of PET substrates.2Br PSCs are also investigated. The detailed photovoltaic parameters and their statistical distributions are summarized in Table S1 and Figure S14 (Supporting Information), respectively. The optimized annealing time at 130 °C is 10 min. The J–V curve of the champion flexible PSC annealed at optimized condition (130 °C‐10 min) is shown in Figure 2Br PSC can reach up to 11.73% with a VOC of 1.19 V, a FF of 67.25%, and a JSC of 14.61 mA cm−2. In contrast, the control flexible CsPbI2Br PSC without DAPP only exhibits a PCE of 9.48% with a VOC of 1.08 V, a FF of 65.0% and a JSC of 13.55 mA cm−2 . Obviously, all the photovoltaic parameters of CsPbI2Br PSCs with DAPP have higher values than the control samples. Compared with the control device, the PCE of CsPbI2Br PSCs with DAPP is increased about 24%. Figure 2Br device with an integrated JSC value of 14.28 mA cm−2, which agrees well with the J–V measurements. To confirm the reliability of the J–V measurements, the current density and PCE at the maximum power point (0.94 V) are recorded as functions of time, as presented in Figure JSC of 12.37 mA cm−2 is obtained, which is very close to the efficiency value obtained by the J–V measurements.To build up flexible CsPbI2Br PSCs, the mechanical bending tests of flexible CsPbI2Br PSCs are carried out with different curvature radii. The J–V curves of the flexible devices before and after bending at curvature radii of 12, 7, and 5 mm for 300 cycles are shown in Figures2Br PSC exhibits no significant degradation after bending at a curvature radius of 12 mm for 300 cycles. In addition, the flexible CsPbI2Br PSC retained more than 90% of its initial PCE value even after bending at a curvature radius of 5 mm for 300 cycles, proving excellent mechanical stability.To demonstrate the flexible property of CsPbI2Br films and PSCs are stored without encapsulation at room temperature in ambient air at ≈30% relative humidity (RH). As shown in Figure2Br films annealed at 65 °C with and without DAPP undergo a complete degradation after storing in air for 4 h. When the annealing temperature is increased to 120 °C, CsPbI2Br film without DAPP shows a significant degradation, whereas the CsPbI2Br film with DAPP remains in its initial state, proving that DAPP method can enhance the stability of CsPbI2Br film that annealed at low temperature. Figure 2Br PSCs with or without DAPP, the flexible PSCs with DAPP still retain ≈70% of the initial efficiencies after 700 h aging in air. Such a long‐term stability of flexible CsPbI2Br PSCs is closely related to the high‐quality of CsPbI2Br film. In contrast, the efficiencies of the control device without DAPP rapidly decrease to less than 60% of their initial values within 5 h.To examine the effect of the DAPP method upon the device air stability, CsPbI2Br PSCs with or without DAPP are also investigated, respectively . It turns out that the CsPbI2Br PSCs with DAPP have better thermal and humidity stabilities than their counterparts without DAPP. However, the stability performance of CsPbI2Br PSCs is still inferior to the CsPbBr3 and CsPbIBr2 PSCs with the lower iodide content.Meanwhile, the thermal (100 °C) or high‐humidity  stabilities of flexible CsPbI2Br film through reducing the reaction energy of precursors by using PbI2(DMSO) and PbBr2(DMSO) to replace PbI2 and PbBr2. This method can obtain a high quality and stable CsPbI2Br film even at 120 °C. A CsPbI2Br PSC with PCE of 13.54% is fabricated on rigid substrate, showing negligible hysteresis. In addition, the low temperature DAPP method based on DMSO adducts enables the fabrication of the flexible CsPbI2Br PSCs with a highest reported PCE of 11.73%, the flexible CsPbI2Br PSCs can maintain 90% and 70% of their original PCE after being bent under a curvature radius of 5 mm and stored in an ambient environment for 700 h, respectively. Therefore, this work not only proves the feasibility of producing highly efficient CsPbI2Br PSCs on rigid substrates at low temperature, but also opens a new avenue to realize the high performance flexible CsPbI2Br PSCs on flexible substrates.In summary, we have demonstrated a low temperature DAPP method to fabricate high quality CsPbIMaterials: The PbI2(DMSO), PbBr2(DMSO) adducts were prepared by dissolving the appropriate amount of the commercial PbI2  and PbBr2  in DMSO with heating (60 °C) and stirring, respectively. Then 40 mL of antisolvent  was slowly added into the PbI2 or PbBr2 solution to precipitate the DMSO adducts. After the precipitation was completely produced, it was filtered and dried in a vacuum oven for 24 h to get white PbI2(DMSO)2 and PbBr2(DMSO)2 powders. The PbI2(DMSO) and PbBr2(DMSO) adducts were obtained by annealing PbI2(DMSO)2 and PbBr2(DMSO)2 in a vacuum oven at 60 °C for 24 h, respectively.Device Fabrication: The rigid all‐inorganic CsPbI2Br PSCs were fabricated with a structure of FTO/TiO2/CsPbI2Br/spiro‐OMeTAD/Au. FTO/glass substrates were successively cleaned in an ultrasonic bath of water, acetone and isopropanol alcohol for 10 min. Then the surface of substrates was treated under UVO for 15 min to make a hydrophilic surface. TiO2 layers (40 nm) were prepared by a chemical bath deposition method as electron transport layers. CsPbI2Br perovskite layers were prepared by spin coating the perovskite precursors , 0.200 g PbBr2(DMSO), 0.234 g CsI) in mixed solvent of N,N‐dimethylformamide (DMF) and DMSO  with a concentration of 0.9 m at 1000 rpm for 10 s and 4000 rpm for 40 s, and the chlorobenzene was coated on the film between 14 and 16 s during the second spinning step as an antisolvent. After the spin coating, the precursor films were annealed on a hotplate at 65 or 120 °C for 10 min. Then bis(trifluoromethylsulfonyl)imidelithium salt  and 4‐tert‐butylpyridine  codoped 2,2′,7,7′‐tetrakis‐9,9′‐spirobifluorene (spiroOMeTAD) solution in chlorobenzene (90 mg mL−1) was coated on inorganic perovskite layer at 5000 rpm for 30 s to form the hole transport layer. Finally, a gold film (≈80 nm) was thermally evaporated on spiro‐OMeTAD by using a shadow mask to form a device active area of 9 mm2 as the top electrode. The flexible inorganic PSCs were fabricated with a structure of PET/ITO/Nb2O5/CsPbI2Br/spiro‐OMeTAD/Au. The ITO film was deposited on the PET substrates by a magnetron sputtering method. The MgF2 was deposited on the back of PET by electron beam evaporation as an antireflection coating. The MgF2/PET/ITO flexible substrate has an average transmittance of ≈86% in the spectrum region of 400–800 nm and a sheet resistance of ≈8 Ω sq−1. The flexible substrate was cleaned with deionized water and ethanol in an ultrasonic bath. The Nb2O5 films (60 nm) were deposited on the flexible substrates by an e‐beam evaporation method as the electron transport layer. The CsPbI2Br perovskite film, spiro‐OMeTAD layer and Au electrode were prepared with the same procedure as on the rigid substrates.Device Characterization: Top‐view and cross sectional images of samples were analyzed by field‐emission SEM . XRD spectra were obtained using a D/MAX 2400 diffractometer with Cu Kα radiation (Rigaku). The absorption spectra of DMSO adducts were obtained by a UV/Vis NIR spectrophotometer . The current density–voltage (J–V) curves of PSCs were collected using a Keithley 2400 series sourceMeter under the illumination of an AM 1.5G at 100 mW cm−2 irradiance generated by SAN‐EIELECTRIC XES‐40S2‐CE solar simulator, with the intensity calibrated by a NREL‐traceable KG5 filtered silicon reference cell. Both forward and reverse scans were recorded at a scan rate of 0.03 V s−1. The EQE spectra of the PSCs were recorded by a QTest Station 2000ADI system. The sheet resistance of ITO film deposited on PET was measured by an SB118 four‐probe meter.The authors declare no conflict of interest.SupplementaryClick here for additional data file.


An improved understanding of neuropsychiatric disorders associated with dysfunctions in the dopamine system and the underlying mechanism is a necessity in order to promote the development of new potential therapeutic drugs. In contrast to other broadly applied 11C-radiopharmaceuticals, the production of this radiotracer requires a challenging four-step radiosynthesis involving harsh reaction conditions and reactants as well as an inert atmosphere. Consequently, the production is prone to errors and troubleshooting after failed radiosyntheses remains time consuming. Hence, we aimed to optimize the radiosynthesis of (+)-[11C]PHNO for achieving better activity yields without loss of product quality. Therefore, we synthesized (+)-[11C]PHNO and omitted all heating and cooling steps leading to higher activity yields. As a result, radiosynthesis fully conducted at room temperature led to a time-reduced production procedure that saves about 5 min, which is an appreciable decay-prevention of around 15% of the activity yield. Additionally, we established a troubleshooting protocol by investigating reaction intermediates, byproducts, and impurities. Indeed, partial runs enabled the assignment of byproducts to their associated error source. Finally, we were able to generate a decision tree facilitating error detection in (+)-[11C]PHNO radiosynthesis.(+)-[ Recently, the application of this tracer in healthy subjects and patients suffering from schizophrenia showed an enhanced dopamine release in patients with schizophrenia PHNO radiosynthesis has been published by several working groups describing synthetic routes via carbon-11 radiolabeling either on the 1-position or the 3-position of the propyl moiety (11C]-(+)-PHNO radiosynthesis has not become widespread, whereas [1-11C]-(+)-PHNO is widely applied in clinical studies [11C]PHNO (further stated as (+)-[11C]PHNO to be in accordance with the literature) starts with generation of [11C]propionic acid chloride, which can be obtained by a Grignard reaction from [11C]CO2 and ethylmagnesium bromide and subsequent conversion with thionyl chloride (SOCl2) (1 and 2) PHNO synthesis. This method facilitates the establishment of (+)-[11C]PHNO in other PET facilities .–14.2/3 r studies –17. Radi1 and 2) –20. Thiscilities .11C]PHNO, the susceptibility to humidity as well as oxygen contamination remains an unsolved problem. As a consequence, (+)-[11C]PHNO synthesis has the highest failure rate within our facility. After a failed production run, an extensive validation process starts for investigating which of the numerous steps led to the radiosynthetic failure. During troubleshooting, the impact of the reaction temperature needs special considerations. Wilson et al. reported that the reaction temperature regulated the formation of byproducts: the reduction of the carboxyl-group at temperatures below −30°C results in a lower number of side products compared to higher temperatures CO2 was produced by a GE PETtrace cyclotron 860  via the 14N11C nuclear reaction by irradiation of a gas target  filled with high-purity N2 (+1% O2) . Automated syntheses were performed on a TRACERlab™ FX C Prosynthesizer  including a semipreparative HPLC system featuring a Linear Instruments Model 200 UV/Vis detector and a LaPrep HPLC pump . For analytical HPLC, an Agilent 1260 infinity system  incorporating a quaternary pump (G1311B), a column oven (G1316A), a manual injector (G1328C), a multi wavelength UV-detector (G1365D), and a NaI(Tl)-detector from Berthold Technologies  was used. The HPLC device is controlled by GINA Star v5.9 SP17 controlling software .Ethylmagnesium bromide , thionyl chloride  to a PE tube, which was beforehand impregnated with a solution of ethylmagnesium bromide  in THF. The impregnation was performed by diluting 500 μL EtMgBr with 1000 μL THF and pushing the solution trough the loop, and then the loop was flushed with helium for around 5 sec to remove the excess of the impregnation solution. After the [11C]CO2 reaction with EtMgBr that last for around 5 min, SOCl2 was pushed through the tube to obtain the acid chloride and the solution was simultaneously transferred to a reactor containing a solution of (+)-HNO (1.9–2.4 mg) in TEA (50 μL) and THF (400 μL). The intermediate amide species was obtained after 5-6 min stirring at room temperature. Afterwards, a solution of 120 μL lithium aluminum hydride (LAH) in THF (400 μL) was added to reduce the intermediate amide. The reaction was quenched by addition of aqueous HCl  and neutralized with aqueous NaOH . The resulting suspension was filtered over cotton wool and purified by semipreparative HPLC. The collected product peak was diluted in 80 mL water and trapped on a C18-SPE cartridge. Afterwards, the product was eluted with 1.5 mL ethanol, diluted with phosphate buffered saline, and sterile filtered.A reduced reaction time was achieved by omitting all heating and cooling steps with exception of CO11C]PHNO, the synthesis was carried out using the following conditions:−40°C −15°C (as previously published by our group )22°C (shortened synthesis time and no heating and cooling)In order to evaluate the influence of the reaction temperature on the efficacy of the reduction of the intermediate amide with lithium aluminum hydride (LAH) to (+)-CO2 in EtMgBr was investigated as follows:2O) in 1 mL THF, and then the activity was released through the loop. To quench the reaction and hydrolyse the radioactive intermediate, 0.5 mL of water was pushed through the impregnated loop and the reaction mixture was analyzed (F).The PE tube was loaded with a mixture of 0.5 mL Grignard solution (3 M in EtμL of water were added to the Grignard reagent solution and the reaction was stopped after amide formation (G) or after the reduction step (H).5 The PE tube was impregnated with the Grignard reagent solution and the loop was flushed with 0.5 mL of water. The reaction was stopped after amide formation (G1) or after the reduction step (H1).The effect of moisture before and after trapping of [2O (20 μL) to a solution of (+)-HNO (1-2 mg) in THF (400 μL) and TEA (50 μL) prior to the addition of [11C]propionic acid chloride (I). Additionally, the same synthesis was performed and the resulting reaction mixture further treated with LAH (J).The impact of moisture on the amide formation was investigated by adding Hμm, 130 Å column . The mobile phase consisted of solvent A (ammonia phosphate buffer (100 mM), sodium-1-octasulfonate (5 mM), pH 2.1 adjusted with H3PO4); solvent B (90% acetonitrile (ACN)/10% water); solvent C (water); and solvent D (ammonium phosphate buffer (50 mM), pH 9.3 adjusted with NaOH). The gradient started with a composition of 33% A, 17% B, 17% C, and 33% D. Subsequently, B is increased over 2 min from 17% to 34%, whereas C is reduced from 17% to 0%. The UV/Vis signal was detected at a wavelength of 280 nm with a reference wavelength of 450 nm. The initial flow rate (1.5 mL/min) was decreased after 25 s to 1.0 mL/min. The retention time of the radioactive peaks was compared to the reference standard of (+)-PHNO in order to identify the respective product peak.All crude, small-scale reaction mixtures were analyzed by analytical HPLC measurements as previously published by Nics et al. . The staμm; Phenomenex Ltd., Aschaffenburg, Germany) as a stationary phase. The mobile phase consisted of 25 mM phosphate-buffered saline (PBS) (pH 7.0)/ACN (60/40 v/v) with a flow rate of 5.8 mL/min. The UV/V is signal was measured at 254 nm. The peaks were collected and correlated with the respective peaks of the analytical HPLC.A part of the partial runs was additionally investigated on a semipreparative HPLC system with a Phenomenex Luna C18 column  in THF (200 μL) for 3 min under He atmosphere. Afterwards, a solution of SOCl2 (5 μL) in THF (400 μL) was added to generate the acyl chloride. After 4 min, the precursor solution of (+)-HNO  in TEA (50 μL) and THF (200 μL) was added to the reaction mixture. The resulting suspension was extracted three times with CH2Cl2 (about 500 μL). The combined organic phases were injected into the semipreparative HPLC system. The collected peaks were analyzed by analytical HPLC and HRMS , 3 (or 4) , and 5 .The intermediate amide Characterization and identification of the compounds separated by HPLC was performed by high-resolution mass spectrometry (HRMS) on a Bruker maXis UHR-TOF device (electrospray ionization (ESI); qQ-TOF, mass accuracy < 5 ppm) in either positive or negative mode depending on the respective molecular structure.11C]CO2 and can be expressed in [%] as radiochemical yield not corrected for decay CO2. Values are given as mean values ± SD as calculated with Microsoft Excel 2010 if not stated otherwise. Statistical analysis was performed by using a t test as implemented in Graph Pad Prism 6 .Study data were obtained from all syntheses performed in our department between 01/2017 and 08/2017 with starting activity of 100–150 GBq at the end of bombardment (EOB) of [11C]PHNO preparation by Rami-Mark et al. involves a heating step to 80°C for the acylation reaction, followed by the addition of LAH at −15°C [The previously described radiosynthetic process for (+)-[at −15°C . Afterwa11C]PHNO synthesis were performed at room temperature. In detail, the acylation with [11C]propionic acid chloride was performed without heating and cooling as well as the addition of LAH. In this study, the improved synthesis (n = 16) was compared with two experimental setups: (1) (+)-[11C]PHNO synthesis was conducted as previously described by our group (n = 9) PHNO synthesis can be successfully performed at room temperature without a forfeit of activity yield. The starting activities and activity yields of the experiments are given in In the improved method, all reaction steps of (+)-. Th. Th21]. 11C]CO2 and [11C]propionic acid. The species at 47 s and 48 s are probably the carbonyl intermediates. The peak at a retention time range of 1 min 55 sec to 2 min 10 sec originate from (+)-[11C]PHNO. However, after purification by semipreparative HPLC, (+)-[11C]PHNO could be obtained in an excellent purity of >95% for all applied methods.The limiting parameter of the reaction at room temperature is the reactor volume as the solvent is not evaporated and an overflow of the reaction mixture must be avoided. If an adequate reactor is available, this facilitated four-step synthesis procedure can be performed in every radiochemistry laboratory that is equipped with a cyclotron. All experiments showed a similar pattern of the crude reaction mixture in semipreparative as well as analytical HPLC chromatograms but a different intensity of the peaks . The pea11C]PHNO synthesis -CO2 and [11C]propionic acid. Further peaks with smaller intensities were observed at longer retention times (4.5 min and 6 min in semipreparative HPLC) showing minor amounts of potential carbonyl byproducts (compound 3 or 4) resulting from an insufficient deprotonation of the precursor, which impedes the acylation process.In route A, semipreparative HPLC showed the main species at a retention time of 2.5–3 min , and in 11C]PHNO synthesis. In conclusion, a chromatogram with intensive signals at early retention times, poor signal at retention times of 5–6.5 min, and no formation of product may originate from nonintact Et3N.Reaction scheme B, which included the reduction step with LAH, shows similar chromatographic pattern. However, no product formation was observed showing that deprotonation of (+)-HNO is pivotal for a successful (+)-CO2 (route F)Hydrolysis of the Grignard reagent after the addition of [11C]CO2 (routes G and H)Adding water to the Grignard reagent before addition of CO2, are supposed to elute at earlier retention times.The route F led to semipreparative chromatographic peaks at retention times of 2-3 min, 3–5 min, and 6.0–6.5 min. In this respect, analytical HPLC displayed a main signal at a retention time of 15–20 s (75 ± 7%). Surprisingly, we also observed peaks at 45 s and 2.5 min, although the expected species, propionic acid and PHNO, whereas the addition of 500 μL of water (H1) completely quenched all further reactions. As a consequence, the presence of peaks with a retention time less than 4 min in semipreparative HPLC may originate from a decomposed Grignard reagent.When the reaction process was stopped after amide formation, an intensive peak at 6 min was visible in semipreparative HPLC for an addition of 5 Besides, the peaks of the semipreparative HPLC and analytical HPLC were correlated to each other as follows : retentiμL water was added to the precursor solution -HNO)) prior to addition of intact propionic acid chloride. The synthesis was stopped after formation of the amide (I) or after the reduction (J). For partial run I, the ratio of the reaction intermediates was as follows: 15–20 s and 45 s with 37.5% and 34.0%, respectively. Due to the high intensity of the peak at 45 s, it can be assumed that the reaction was not completely quenched and one of the carbonyl species was partly formed. Likewise, semipreparative HPLC revealed signals at 4–7 min indicating carbonyl formation. Performing the entire synthetic route for (+)-[11C]PHNO production (J) leads to a peak at a retention time similar to (+)-[11C]PHNO (9 min) with low intensity. However, analytical HPLC showed a peak at 1 min 26 s with 8.3% conversion but no signal at 2 min for the product. This chromatographic pattern is quite similar to the one observed for the partial run without TEA.Here, 20 11C]PHNO synthesis is [11C]propionic acid, which was identified by comparison of the nonradioactive reference compound [natC]propionic acid. The analytical chromatogram showed a signal at 15–40 s at 280 nm wavelength -PHNO radiosynthesis can be attributed to various reasons, but the chromatogram of the semipreparative HPLC provides information on error sources. Based on our results, we created a decision scheme that helps troubleshooting in case of a failed (+)-[11C]PHNO production -[11C]PHNO.As shown in this study, failure of (+)-[oduction . The pre11C]PHNO is a challenge for every radiochemist in terms of time efficiency, reduction of failed syntheses, and error evaluation. Here, we present a new method allowing a tremendous reduction of the radiosynthetic duration of approximately 5 min by omitting heating and cooling steps, which enhances the activity yield significantly. Moreover, the investigation of side products and intermediate species facilitates the error evaluation after a failed synthesis. Accordingly, we propose a decision tree to support troubleshooting and facilitating a stable and continuous radiotracer production that is a necessity for clinical studies.The complex synthesis of (+)-[


All‐pure and Gd‐doped CeO2 NPs are synthesized using a coprecipitation method and characterized using X‐ray diffraction (XRD), absorption spectroscopy, surface‐enhanced Raman Spectroscopy (SERS), X‐ray photoelectron spectroscopy (XPS), and superconducting quantum interference device (SQUID). The effect of Gd‐doping on properties of CeO2 is discussed along with defects and oxygen vacancies generation. The XRD confirms the incorporation of Gd3+ at the Ce3+/Ce4+ site by keeping the crystal structure same. The average particle size from transmission electron microscopy (TEM) images is in the range of 5–7 nm. The XPS spectra of Ce 3d, O 1s, and Gd 4d exhibits the formation of oxygen vacancies to maintain the charge neutrality when Ce4+ changes to Ce3+. The gradual increase in hydrogen production is observed with increasing Gd concentration. The observed results are in good correlation with the characterization results and a mechanism of water splitting is proposed on the basis of analyses. The absorption spectra reveal optical band gap (2.5–2.7 eV) of samples, showing band gap narrowing leads to desired optical absorbance and photoactivity of NPs.Enhanced visible light photocatalytic activity of Gd‐doped CeO Li et al.3+‐doped CeO2 (10 at% Gd) nanostructure due to Gd3+ ions or formation of oxygen vacancies. In most of these studies, the local ordering of oxygen vacancies on grain boundaries in heavily RE‐doped CeO2 samples has been reported.2 nanoparticles (NPs) with doping concentration 3%, 5%, 7%, 9%, and 11%. On the basis of their results based on X‐ray absorption spectroscopy, extended X‐ray absorption fine structure, Raman, and scanning transmission electron microscope‐electron energy loss spectroscopy measurements they have discussed that below 7% and above 7% distribution of defects strongly depend on the concentration of Sm3+ ions in CeO2 NPs. Since, there is a lack of literature available on the evidence of the distribution of defect study with small doping concentration of RE‐doped CeO2 NPs.Among all RE‐doped CeO2 materials the impurities tend to shift the band position and can tune the band gap because of their effects on electronic transition.2 is used for thermolysis of water that can produce 101.6 H2 (µmol g−1) hydrogen.2; however, the effect of interaction between the rare earth dopant and oxygen vacancy defects on enhanced visible light photocatalytic activity of CeO2 is still not investigated so far.Metal oxide photocatalyst have attracted increasing attention due to their potential applications in the environmental protection and energy utilization, such as water splitting for hydrogen production. In the Zn‐, Mg‐, and Ca‐doped CeO2. To undertake this study we have investigated the structural properties, dopant distribution, and their association with oxygen vacancies in Gd‐doped CeO2 NPs. The CeO2 NPs have been doped with different concentrations of Gd3+ ions  to discuss the presence of defect induced oxygen vacancies (either intrinsic or extrinsic) and their association with doped cation with surface‐enhanced Raman spectroscopy (SERS) measurements. Transmission electron microscopy (TEM) has been used to observe changes in the surface morphology and particle size with increased fluencies of Gd3+ ions in CeO2 lattice. The optical absorption spectra have been measured using ultraviolet–visible–near infrared (UV–Vis–NIR) spectrometer to find out the band gap energy. Finally, we have attempted to investigate how doping concentration affect the oxygen vacancies and cation (Ce3+) defects, present in pure and Gd‐doped CeO2 NPs. These results offer a physical understanding for the available experimental results to explain the enhanced photocatalytic activities of Gd‐doped CeO2 NPs, which can be useful for designing and understanding the novel doped CeO2 photocatalyst.Hence in this study, we systematically explore to develop the correlation between the rare earth dopants, its concentration and oxygen vacancy defects to enhance the photocatalytic activity of doped CeO22.1x1−GdxO2 nanoparticles for x = 0.0, 0.02, 0.04, 0.06, 0.08, and 0.10 at room temperature are shown in Figure2 samples exhibit fundamental Bragg reflections corresponding to the fluorite type face centered cubic structure in the space group of Fm‐3m, in which Ce and Gd atoms are located at 4a position, surrounded by eight O (located at 8b) positions.2O3 or other impurity peaks indicates well incorporation of Gd3+ ions on CeO2 lattice site, which confirms the single phase formation of all the Cex1−GdxO2 nanoparticles. The intensity of XRD diffraction peaks is found to vary with incorporation of Gd3+ ions in CeO2 NPs  diffraction peak (111) is shifted toward higher angle side  peak is again shifted toward lower angle side. This shifting of (111) peak toward lower and higher angle side is attributed to the lattice expansion and reduction, respectively, which is induced by incorporation of Gd3+ ions in CeO2 NPs.3+ ions in CeO2 NPs, indicating that crystal size and crystallinity of the samples are affected with the fluencies of Gd3+ ions. The average nanocrystalline particle size (D) of Gd‐doped CeO2 samples has been calculated with XRD diffraction spectra using the Debye–Scherrer's formulaa” refers to the lattice parameter, d is the crystalline lattice spacing, and h, k, l, are the miller indices of crystal. The calculated value of lattice parameter with fluency of Gd3+ ions in CeO2 NPs is tabulated in TableThe rising intensity signifies an improvement in the crystalline nature while falling intensity signifies low crystallinity of Ce3+ (0.1053 nm) ions are substituted for the smaller ionic radii Ce4+ (0.097 nm) ions and created the larger radii Ce3+ ions (0.114 nm).2 lattice, Gd3+ and Ce3+ ions are collectively creating oxygen vacancies in the CeO2 lattice, which causes further lattice expansion.The variation in the calculated values of lattice parameter can be directly related to the ionic radii of the dopant ion. Since, the larger ionic radii Gd2 does not affect the cubic fluorite structure of the CeO2, as no additional diffraction peaks related to possible impurity phases of Gd and oxides of Gd are observed in these Gd‐doped CeO2 samples. It is further confirmed the formation of a single phase of Ce1‐xGdxO2. The refinement results clearly indicate that Gd ions are well incorporated in the CeO2 matrix and Gd‐doping in CeO2 leads to small enhancement in the unit‐cell volume.Rietveld profile refinements  cations in CeO2 NPs, the maximum value of strain is for x = 0.08 doping concentration. Some theoretical investigation revealed that tensile strain promotes the formation of oxygen vacancies rather than compressive strain.2 samples, increased value of tensile strain can be directly related to the endorsement of oxygen vacancies in doped CeO2 samples, which may be associated to the bonding length and the strength between the surface O and Ce atoms.2 system.2 samples are discussed in the next segment by TEM, high‐resolution transmission electron microscopy (HRTEM), and selected area (electron) diffraction (SEAD) images.Looking to Table 2.22 NPs. The particle size and morphology of pure and Gd‐doped CeO2 NPs are analyzed by TEM as shown in Figure 2 NPs, which are in good agreement with the results obtained from Debye–Scherrer formula  indicates that the obtained particles are nanocrystalline. Furthermore, HRTEM and SAED analysis are also used to decipher the information about the nanocrystallinity and impurity phases, if any present in Gd‐doped CeO2 NPs. HRTEM images  for Gd‐doped CeO2 samples but for 8% Gd‐doped CeO2 sample, the interplanar distance (d = 0.32 nm for (111) plane) is slightly increased, which again promotes the crystal growth and indicates the low crystallinity. Some defects, such as dislocations , (200), (220), and (311) planes. These rings indicate that the particles are crystallized and diffraction rings are very well matched with the XRD measurement results. The fluency of Gd3+ ions in CeO2 NPs also affect the optical band gap energy, which is further discussed in the next segment by UV–Vis–NIR spectroscopy.As shown in Table 2.3Figurex1−GdxO2 samples . These samples exhibit a strong absorption below 400 nm with an absorption peak in UV‐range corresponding to the different doping concentration of Gd3+ ions in CeO2 NPs as tabulated in Table 2− to 4f conduction band of Ce4+ ions.2 have wide band gap semiconductor with a forbidden gap of 5.5 eV.4+ (4f) conduction band with the holes in the O2− (2p) valance band. It can be seen from Figure 2 but after incorporation of Gd3+ ions peak is shifted toward lower wavelength (blue shift) up to optimal doping concentration x = 0.04, while for further fluencies x = 0.06 to 0.10 peak is shifted toward higher wavelength (red shift) side. It has been reported that when metal NPs are forming smaller size particles, the λmax shifts toward shorter wavelength (blue shift) side, whereas, when the smaller particles aggregate to form bigger/larger size particles, the λmax value shifts toward longer wavelength (red shift) side.x = 0.04), while with increasing fluencies of Gd‐ions (for x = 0.06 to 0.10) in CeO2 NPs, these smaller sized particles are agglomerated , which is smaller than the predictable Bohr exciton radius for CeO2 (7–8 nm).2 NPs. Generally, quantum confinement effect results when the Bohr radius of an exciton approaches the grain or particle size, spatially confining the electron–hole pair. When this happens, the energy of the lowest excited state increases and the increased band gap produces a blue shift in the absorption spectra. Now at the higher fluencies (for x = 0.06 to 0.10) of Gd3+ ions, the contribution of blue shifting from Ce4+ to Ce3+ valance state change will become small. Therefore, red shifting is occurred in the absorption spectra of Gd‐doped CeO2  samples. This red shifting may be the outcome of an interfacial polaron effect arising from electron–phonon coupling phenomenon.Eg) of pure CeO2 and Cex1−GdxO2  NPs has been calculated using Tauc's equationn = 1/2 and for n = 2 for indirect transition.αhν)2 as a function of the incident photon energy (hν). Table Eg) of all samples. Pure CeO2 NPs shows band gap energy of 2.60 eV that is smaller than the band gap energy reported for bulk ceria, i.e., 3.15 eV.3+ states on grain boundaries. Moreover, the optical band gap energy is found to increase for low fluency of Gd3+ ions (for x = 0.02 and 0.04) in CeO2 NPs while it decreases subsequently with increasing fluencies of Gd3+ ions (for x = 0.06 to 0.10) in CeO2 NPs (as shown in Table x = 0.02 and 0.04) of Gd3+ ions may be correlated with the decrease of Ce3+ concentration as well as oxygen vacancies during annealing process. This may eliminate some localized defect energy states within the band gap due to the corresponding decrease of vacancies content, which results increase in the band gap energy.mvc∗ is effective mass of electrons, ne is the electron concentration, and h is the Plank constant. Now, according to BM effect, above the Mott critical density, the increased number of free electron concentration is leading to fill 4f level partially, which in turn blocks the lowest states and leads to band gap widening.3+ ions into CeO2 sample, the crystalline size is reduced  is caused with the existence of Ce3+ contents at the grain boundaries, which increases with decreasing particle size.2 NPs has been calculated by using the following formulaHere, 2 NPs are tabulated in Table x = 0.04, whereas it is increased for further fluencies  of Gd3+ ions in CeO2 NPs. The variation in the refractive index and band gap energy has been shown in Figure2 NPs. Therefore, absorption of UV light at low concentration of Gd‐ions (x = 0.02 and 0.04) in CeO2 NPs has been increased due to reduction of particle size as well as refractive index, whereas, due to increasing doping concentration the transparency and UV protection qualities are decreased.The obtained values for the refractive index of pure and Gd‐doped CeO2.42 NPs.Surface enhanced Raman spectroscopy is a powerful vibrational technique, which allows for highly sensitivity structural detection of low concentration analyses through the amplification of electromagnetic fields generated by the excitation of localized surface plasmons.Figure2g) mode for pure CeO2 and Cex1−GdxO2  samples, at 463.3 cm−1 and in the range of 461.4–459.5 cm−1, assigned for first‐order scattering.2O3 (360 cm−1) clearly indicating the incorporation of Gd3+ ions into CeO2 lattice, confirms the absence of any impurity phase in the lattice, in agreement with XRD results.2g mode corresponding to pure CeO2 and Cex1−GdxO2 is slightly shifted toward lower wavenumber side (or lower energy side) and broadening in its FWHM can also be observed with doping fluencies of Gd3+ ions in CeO2 sample. These structural changes in Raman spectra with Gd‐doping are attributed to the inhomogeneous strain and defects caused by substitution at the smaller radii Ce4+ (0.97 Å) site by larger ionic radii Gd3+ (1.08 Å) ions.2g mode, weak intensity second‐order Raman peaks are also obtained at 598.5 and 595.6 cm−1 for pure CeO2 and 2% Gd‐doped CeO2, respectively, generated due to nondegenerated longitudinal optical (LO) mode.2. The three possible defect induced mechanism for oxygen vacancies in pure CeO2 sample can be given asing.2O3 3 cm−1 cle4+ state but also Ce3+ state. For doing so, oxygen (O2−) ions are released from the structure and finally oxygen vacancies are formed in the system. The intensity of this peak is increased with incorporation of Gd‐ions indicates the rise of oxygen vacancies in ceria lattice. With increasing the fluency of Gd‐ions, two weak second‐order Raman modes in the range of 554.3–558 and 598–600.1 cm−1 are also obtained  NPs can be given asCecex and GdCe′ are Ce4+ and Gd3+ ions on the CeO2 lattice site, respectively, OOx is O−2 ions on an oxygen lattice site, and VO•• is neutral oxygen vacancy site. In addition, another vacancy peak can also be observed in the range of 598–600.1 cm−1, which is assigned to the defect space including intrinsic oxygen vacancies due to reduction of Ce4+ to Ce3+, i.e., nonstoichiometry of ceria.Cecex and CeCe′ are Ce4+ and Ce3+ ions on the Ce lattice site, respectively, OOx is O−2 ions on an oxygen lattice site, and VO•• is neutral oxygen vacancy site.x = 0.04. With further increase in fluency, the intensity of this mode decreases and then again increases at x = 0.10 concentration. The variation of intensity of vacancy mode is related to the concentration of oxygen vacancies. The quantitative estimation of oxygen vacancies of pure CeO2 and Gd‐doped CeO2 samples is made from the relative peak area of vacancy modes (intrinsic and extrinsic) with area of F2g mode. For doing so, Lorentzian fitting is done for measuring the peak area of the respected peaks. All calculated values are tabulated in Tablex = 0.08  and then slightly decreased at x = 0.10 concentration. The increment in the concentration of oxygen vacancies can be explained by considering, with incorporation and rising fluency of large radii Gd3+ ions (0.105 nm), the dislocation density as well as strain has been increased up to x = 0.08 doping concentration in ceria. Due to this reason every two Gd3+ ions substituted the smaller radii Ce4+ (0.097 nm) ions increases the probability of oxygen ion (O2−) to leave the ceria lattice to maintain electrical neutrality in the lattice and creates more oxygen vacancies , and has small additional contribution from F2g symmetry  is full width at half maximum (FWHM) of Raman active (F2g) mode and D is particle size of pure CeO2 and Cex1−GdxO2  samples.AOv corresponding to 598.5 and 595.6 cm−1 for pure CeO2 and 2% Gd‐doped CeO2. The relative ratios of A(OV)1, A(OV)2 and AF2g bands, which are corresponding to 554.3–558, 598–600.1, and for F2g band is also calculated to estimate the oxygen vacancies concentration for further doping concentrations of Gd for 4% onward. The calculated values are shown in Figure The particle size of all the Ce2g mode is calculated by fitting the Lorentzian function for the corresponding modes. Ratio AOv/AF2g for pure CeO2 and 2% Gd‐doped CeO2, whereas, (A(OV)1 + A(OV)2)/AF2g ratios are calculated for Cex1−GdxO2  samples. From these calculated values one can infer that the relative oxygen vacancy concentration is found to gradually increase with fluencies of Gd3+ ions in CeO2 NPs.This relative peak area ratio of oxygen vacancies and F2.52.5.12 and Gd‐doped CeO2 NPs have been further characterized using XPS measurements of Ce 3d, Gd 4d, and O 1s core levels. Figurex1−GdxO2  samples. All binding energies have been corrected for the charge shift using the C 1s peak (binding energy = 284.6 eV) as reference.4+ and Ce3+ ions, respectively. These series of peaks are labeled as “u” and “v,” which are due to 3d3/2 and 3d5/2 spin–orbit states, respectively.vo, v′, uo, and u′ are characteristic peaks of Ce3+, whereas, the peaks labeled with v, v″, v‴, u, u″, and u‴ are characteristic peaks of Ce4+  samples, which are in good agreement with the reported papers.The chemical composition and the valence state of the pure CeO3+ and Ce4+ for pure CeO2 and Cex1−GdxO2  samples. All peak positions for “u” and “v” of pure CeO2 and Cex1−GdxO2  samples have been tabulated in Table3+ and Ce4+ in the samples has been calculated using the following formulaWe can see from the Ce 3d core level spectra that Ce ions are present in mixed valance state of both CeACe3+=vo + v′+ uo+ u′ and ACe4+=v + v″ + v‴ + u + u″ + u‴ are the sum of the integrated area of all characteristics peaks of Ce3+ and Ce4+, respectively. These calculated values are tabulated in Table 3+/Ce4+ shows that the concentration of Ce3+ ions over Ce4+ ions is gradually increasing for Cex1−GdxO2  samples. While, at 6% Gd‐doping concentration Ce3+/Ce4+ value is deceased, which is again increased and then decreased at 8% and 10% doping.Here, 3+ ions (0.105 nm) in CeO2 NPs, replacing the smaller radii Ce4+ ions (0.97 Å) and for maintaining the charge neutrality, the concentration of Ce3+ ions (0.114 nm) is gradually increased for x = 0.02 and 0.04 doping concentrations. The presence of Ce3+ may be due to either the formation of Ce2O3 or the creation of oxygen vacancies in CeO2 lattice. This can be verified by calculating the stoichiometry ratios x = [O]/[Ce] and x′ = [O1s]/[Ce3d], which can be estimated from their integrated peak area while considering their sensitivity factor. In order to calculate oxygen content in the samples, we assume that the total oxygen content is the sum of the required oxygen to fully oxidize Ce3+ and Ce4+ to form Ce2O3 and CeO2. Then, considering the stoichiometry x = [O]/[Ce], which is equal to 1.5 for Ce2O3 and 2 for CeO2. Now, the stoichiometric ratio of the oxygen to the total Ce ions (Ce3+ + Ce4+) can be determined using the concentration of [Ce3+] and [Ce4+] as given in TableThis shows that due to incorporation of larger radii GdAO and ACe of the O 1s and Ce 3d peaks, respectively, which has been calculated according to the following equationSCe = 7.399 and SO = 0.711 are the sensitivity factors of the Ce and O atoms, respectively.Figurex and x′, which is provided the concentration of Ce3+ and Ce4+ ions in pure CeO2 and Cex1−GdxO2  samples  are higher than that (x) calculated by Equation 2 and Ce0.98Gd0.02O2 sample, which exhibits low concentration of Ce3+ ions for pure CeO2 NPs in comparison with Ce0.98Gd0.02O2 sample. This means that due to incorporation of Gd3+ ions in CeO2 NPs, Gd3+ ions replaces the Ce4+ ions with formation of oxygen vacancies in the Ce0.98Gd0.02O2 sample. On the other hand, the value of (x′) is smaller than (x) for x = 0.04, 0.06, 0.08, and 0.10 doping concentrations, which suggested that the entire Ce3+ ions are consumed in the formation of Ce2O3. Simultaneously, the oxygen deficiency with increasing Ce3+ ions suggests that Ce3+ ions are associated with Ce2O3 as well as oxygen vacancies in CeO2 and both kinds may coexist in Gd‐doped CeO2  samples.Although, the calculated values of actual stoichiometry  samples, while from XRD analysis only CeO2 is identified. This Ce2O3 phase has amorphous character and indicates that this phase is located at the grain surface and at the grain boundaries. Patsalas et al.2O3 and CeO2 are located at grain surface and volume, respectively. A linear correlation can be established between third power of [Ce4+] (grain volume distribution) as well as third power of D  with square of [Ce3+] (surface distribution).This means that core level Ce 3d spectra prove the existence of CeFigure3+]2 and [Ce4+]3 that confirms the distribution of Ce2O3 and CeO2 at the grain surface and volume. While, the experimental points of D3 versus [Ce3+]2 are more scattered around straight line of the dimensional analysis, which is attributed to the strain in the grain that affects the broadening of the XRD peaks with Gd‐doping in CeO2 samples. From Table x and x′ increases and decreases with increase and decrease of Ce3+ ions for Cex1−GdxO2  samples, which show an up and down in the formation of oxygen vacancies in these samples.]2 and Ce+3 that c2.5.22 and Cex1−GdxO2  samples are shown in Figure2− (denoted as OL) in pure CeO2 and Cex1−GdxO2  samples, while peak at higher binding energy side ≈530.3–533.3 eV and ≈533.5–536.4 eV are possibly assigned to oxygen vacancies (denoted as OV) corresponds to Ce3+ species originated from Ce2O32O speciesα and Oβ) on the surface of the samples, respectively  but also oxygen vacancies peak (OV) are found to increase for 2% Gd‐doped CeO2 NPs. The quantitative estimation of OL and OV peaks shows that due to incorporation of Gd3+ ions the oxygen vacancies are formed on the surface of the Gd‐doped CeO2 samples. These vacancies are found to show variation with change in the concentration of Gd‐ions in the CeO2 NPs. Furthermore, as 1s electron of oxygen atom attached more tightly bound to Ce3+ rather than Ce4+ oxidation state. Thus, change in the oxidation state of Ce‐ions (+4 to +3) due to incorporation of Gd3+ ions in the CeO2 NPs, may also be responsible for the change in the formation of oxygen vacancies. The quantitative percentage of Ce3+ oxidation state from core level spectra of Ce 3d for pure CeO2 and Cex1−GdxO2  samples, if compared with quantitative percentage of OV from O 1s core level spectra, one can infer that the increasing Ce3+ concentration is also helpful in increasing the oxygen vacancies on the surface of samples  due to spin–orbit coupling for Cex1−GdxO2  samples, as shown in Figure5/2 and Gd 4d3/2 states, respectively, which are indicating the presence of Gd3+ ions in Cex1−GdxO2  doped lattice.The deconvoluted Gd 4d core level XPS spectra are split into doublet  versus magnetic field (H) curves for pure CeO2 and Cex1−GdxO2  NPs. It is observed that pure CeO2 nanoparticles are found to exhibit weak ferromagnetic (FM) behavior at room temperature with saturation magnetization Ms = 0.049 emu g−1. Although, it has been reported that bulk CeO2 exhibit diamagnetic behavior where it is reported that at nanoregime the undoped CeO2 NPs exhibit weak ferromagnetism with small value of saturation magnetization by few reports.Hc = 77.95 Oe has been observed for pure CeO2 NPs, which ensures the ferromagnetic nature in our pure CeO2 sample. The weak ferromagnetic behavior in pure CeO2 NPs at room temperature is associated with oxygen vacancies that have been originated by the conversion of Ce4+ to Ce3+ oxidation state of cerium.3+ ions in CeO2 NPs, Ce0.98Gd0.02O2 sample still exhibit weak ferromagnetic (FM) behavior with increasing Ms = 0.140 emu g−1 while Hc has been decreased 22.48 Oe as compared to pure CeO2 NPs. While, further increase in Gd3+ ions concentration are not able to maintain this FM behavior, that can be clearly seen from the hysteresis curves  samples. The magnetization of the Gd‐doped CeO2 samples is increased with increasing dopant concentration  μB, where Jg is the Lande g‐factor and for Gd3+ ion ground state 8S7/2, S = 7/2, L = 0, J = 7/2, Jg = 2).3+ is anticipated to get weaker. These noninteracting and localized magnetic spins of Gd3+ ions have induced the paramagnetism with increase in magnetization.3+ ions incorporated into the CeO2 lattice increases with increasing the dopant concentration, which results in reduction of ferromagnetic ordering in Gd‐doped samples. Therefore, 4%, 6%, 8%, and 10% Gd‐doped samples have small ferromagnetic behavior in addition to linear paramagnetic signals, which is gradually increasing with the fluency of Gd3+ ions in CeO2 NPs. Though Raman and XPS analyses are showing an increase in the oxygen vacancies but this increase in oxygen vacancy concentration may not enhance the ferromagnetic ordering in Gd‐doped samples.2 NPs but incorporation of Gd‐ions enhanced the paramagnetic nature, which has been reported due to oxygen defects.B) and due to presence of secondary phases of Gd2O3.Although, after incorporation of Gd2 and Gd‐doped CeO2 NPs, respectively. The origin of FM behavior has been discussed in this paper accounting the F‐center exchange (FCE) mechanism as a subcategory of bound magnetic polaron (BMP) model.VO). These oxygen vacancies and magnetic ions constitute a BMP that produces the ferromagnetism in these systems. In pure CeO2 NPs, the origin of ferromagnetism is supposed to the reduction of the oxidation state of Ce ions, i.e., Ce4+ to Ce3+. The formations of oxygen vacancies give rise to the reduction of Ce4+ to Ce3+ state. The formation of oxygen vacancy left two electrons which may be transferred to a Ce4+ ion converting Ce4+ into Ce3+. Due to this process, mixed Ce3+ and Ce4+ states yield in the pure CeO2 NPs, which has already been confirmed by Ce 3d core level spectra analysis. The ferromagnetism in pure CeO2 NPs may be arise from the nearest‐neighbor interaction, i.e., either double exchange (Ce3+–VO–Ce4+) or superexchange (Ce3+–VO–Ce3+), which is mediated by oxygen ions.+ center because the two electrons left by VO are trapped on Ce4+ ion and VO , while superexchange interaction forms an F2+ center due to the both electrons are trapped on Ce4+ ions.2, it has suppressed the ferromagnetism of CeO2 NPs , which is dominated in this sample. When the size of this BMP is large enough to percolate through the lattice, long‐range (weak) room temperature ferromagnetism can be realized with higher saturation magnetization. However, it is clearly observed that ferromagnetism has been suppressed with the increase in Gd‐doping concentration up to x = 0.10. Due to increase in Gd‐ion doping concentration, the number of Gd‐ions in the interior of CeO2 lattice is less than that on its surface or on the grain boundaries. Only those Gd‐ions are allowed to enter the lattice that is permitted by the host ions and rest is expelled. Due to higher doping concentration the separation among Gd3+ ions is decreased. These largely separated Gd3+ ions suppress the ferromagnetism and undergo superexchange interaction with each other via O2− ions and results in antiparallel alignment of the magnetic spins of Gd 4f shell present in the nearest‐neighbor ions that do not negotiate in ferromagnetic ordering. Thus higher doping concentration of Gd3+ ions tends to destroy the observed ferromagnetism in Gd‐doped CeO2 NPs. Hence in our case the increase in paramagnetic signals may be attributed to increase in oxygen vacancy concentrations without enhancing the ferromagnetic ordering of the samples. This ferromagnetic ordering is further suppressed due to the increased concentrations of Gd3+‐cation as the separation between these Gd3+ ‐ions is decreased results in anti‐parallel alignment of the spins of Gd 4f state due to superexchange interaction.Now, the main issue herein is to understand the possible origin of ferromagnetic dominated paramagnetic behavior in pure CeO2.72 evolved from the samples has been hourly monitored  gets energized after receiving that energy of light and jumped from VB to conduction band (CB), which generates a pair of photohole (at VB) and photoelectron (at CB) at Gd/CeO2 surface. Nascent photoelectrons of CB are arrived at the junction of Pt/electrolyte interface by passing through the electron‐pool of the metallic Pt (that can segregate the photoelectrons from photoholes). These photoelectrons interact with H+ ions2 gas. Hole amassed at VB of the doped semiconductor is responsible for the breakdown of CH3OH in formaldehyde or formic acid or both as mentioned in Equations The amount of photocatalytic Hy Figure a. Accord3+ ion is taking position of Ce3+ or Ce4+ ion), oxygen vacancies, band gap and band positions, and many more. All of the above factors collectively responsible for increase in hydrogen generation activity on increasing the dopant concentration in ceria. Usually the steady decrease in particle size increases water splitting activity with increase in Gd proportion due to the large. Introduction of dopant Gd into the CeO2 lattice, also gradually increases the oxygen vacancy in the lattice arrangement of CeO2 because Gd3+(radius of Gd3+ = 0.105 nm and charge density = 91) replaced the high charged but small Ce4+ ion(radius of Ce4+ = 0.097 nm and charge density = 148) in 2%, 4%, and 8% doped samples but also replaced low charged but bigger sized Ce3+ (radius of Ce3+ cation = 0.114 nm and charge density = 75) in 6% and 10% Gd samples.There are many factors which can dominate the water splitting activity such as particle size of photocatalyst, binding energy, dopant concentration and position (either Gd4+/Ce3+) and low charge M(Gd3+) bonds through bridging O and O as shown in FigureAll of the above changes due to Gd‐doping in ceria lattice maintained the phase purity (checked with XRD) with minor but favorable changes in lattice parameters and suggested the lattice arrangement of atoms with expanded electron clouds between high charge M replaced the high charged but small Ce4+ ion (radius of Ce4+ = 0.097 nm and charge density = 148) in 2%, 4%, and 8% doped samples but also replaced low charged but bigger sized Ce3+ (radius of Ce3+ cation = 0.114 nm and charge density = 75) in 6% and 10% Gd samples.Usually the steady decrease in particle size increases water splitting activity with increase in Gd proportion due to the large. Introduction of dopant Gd into the CeO4+/Ce3+and Gd3+) and anions (O22−) generates a strong local electrostatic field inside the tetrahedra, as confirmed by the XRD, XPS, and Raman results. Residual water molecules are captured by the strong local electrostatic field of the molecular device Cex1−GdxO2 . These water molecules attract the bridging oxygen through the protonic side and the metallic cation, i.e., Ce through the hydroxyl side  doped samples. However, the Gd 4d core level spectra cannot be recorded with good statistics due to very small concentrations but reports the presence of Gd ion in our NPs. In our magnetic measurements we can see that all the samples pure CeO2 and Cex1−GdxO2  are showing the weak ferromagnetism dominated paramagnetic behavior, which is assumed to be triggered due to change of oxidation state of Ce+4 to Ce+3 ions, incorporation of the Gd+3 ions in the lattice and formation of oxygen vacancies. From Raman and XPS analyses, the presence of oxygen vacancy related defects is apparent. From Ce 3d XPS spectra one can reveal that a significant increase in Ce3+ ions is not only able to increase the oxygen vacancies due to formation of Ce2O3 phase in Cex1−GdxO2 , which is not traceable with the XRD analysis but also may be due to the substitution of Ce4+ ions by Gd3+ ions and this substitution actually creates the oxygen vacancies and owe the presence of ferromagnetic ordering in Ce0.98Gd0.02O2 sample.Further for electronic structure of these NPs the core level Ce 3d, O1s, and Gd 4d XPS spectra have been recorded and analyzed in detail. From this analysis the change of oxidation state of Ce2 evolved from the samples is monitored, and the respective release of hydrogen is found to increase for pristine Pt/CeO2, Pt/Gd‐doped CeO2 samples with the increased doping concentrations of the Gd‐ion. The particle size of photocatalyst, binding energies, oxygen vacancy concentrations, band gap, and many other factors are collectively responsible for increase in hydrogen generation activity with increasing dopant concentration in ceria. The observed release of hydrogen is found in good correlation with the characterization results and the proposed mechanism of water splitting is reported on the basis of analyses.In the water splitting results, the amount of photocatalytic H4Materials: Ammonium cerium (IV) nitrate (NH4)2Ce(NO3)6 (Alpha Aesar 99.99%), gadolinium (III) nitrate hexahydrate Gd(NO3)3·6H2O (Alpha Aesar 99.9%), and sodium hydroxide (NaOH) were used.Material Preparation: Nanocrystalline pure CeO2 and Cex1−GdxO2  samples were synthesized using coprecipitation method. The appropriate stoichiometric amount of (NH4)2Ce(NO3)6 and Gd(NO3)3·6H2O were used for synthesizing Cex1−GdxO2 NPs. Initially, (NH4)2Ce(NO3)6 and Gd(NO3)3·6H2O precursor solution was prepared in distilled water with magnetic stirring at the rate of 600 rpm. Then NaOH solution was added drop by drop to this solution until the pH level reached about 11. This solution was stirred about 4 h and then the synthesized pale‐yellow precipitate was collected. The precipitate was dried at room temperature and annealed in the furnace about 500 °C for 8 h. A set of samples, i.e., pure CeO2 and Cex1−GdxO2  were prepared. The main chemical reactions during the experimental process are as followsThe complete chemical reaction can be combined as2 lattice are as followsThe final chemical reaction for the growth of various concentrations 2%, 4%, 6%, 8% and 10% of dopant Gd‐ions in CeONanomaterial's Characterization: The structural properties of all the samples were characterized using XRD measurements on a Brucker D8 Advance diffractometer with Cu Kα radiation (λ = 1.5406 Å). The diffraction patterns were recorded at room temperature in the 2θ range from 10° to 90°. The surface morphology, particle size, and crystallinity of the samples were studied using TEM with Technai G2 20 S‐TWIN (FEI Netherlands) instrument operating at an accelerating voltage of 200 kV. Samples for the TEM investigation were prepared by dispersing the nanopowder in ethanol using an ultrasonicator to produce a dilute suspension. Then a standard holey carbon film supported on Cu grid was immersed in the suspension to produce the TEM sample. The particle size distribution was calculated for a total 150 number of particles using imagej software for TEM images. The optical characterizations were carried out by using SERS. For collecting Raman spectra, SERS of make Thermo Scientific DXRxi Raman Imaging Microscope with charge injection device detector using green laser with 532 nm excitation light source with its power kept at 10 mW were used. The UV–vis–NIR absorbance spectra on the samples in the wavelength range of 200–1000 nm with BaSO4 as standard were recorded employing a Shimadzu UV‐3600 Plus spectrophotometer with an integrating sphere. XPS spectra were recorded on a ultrahigh vacuum based Omicron Multiprobe Surface analysis System  operating at a base pressure of 5 × 10−11 Torr. Mg Kα radiation source (with energy of 1253.6 eV) was used for data acquisition. An OMICRON EA125 hemispherical analyzer equipped with a 7 channeltron parallel detection unit was used to collect the XPS spectra. The calibration of binding energy in photoemission spectra was done referring to standard Au 4f7/2 emission line with energy resolution of ≈0.9 eV FWHM on Au 4f7/2 with pass energy of 20 eV during the measurement. The XPS core level data were analyzed after necessary carbon corrections for the Fermi energy referencing. The magnetic properties of the samples were investigated at room temperature using a Quantum Design MPMS‐3 SQUID system. The magnetization measurements were conducted by varying the applied field from −1.5 T to +1.5 T.2 or Pt/GdxCe1−xO2 or CeO2 or GdxCe1−xO2) that was suspended in 120 mL of aqueous hole‐scavenger electrolyte  in a reaction cell, under the irradiation of 1 sun  visible light. The powder of the photocatalyst (0.2 g with and without Pt loading) was suspended in 120 mL of aqueous electrolyte (20% CH3OH pH = 7.0) in a double walled‐Pyrex glass reaction cell  that was sealed with a rubber septum and plastic wire lock.22 photons per hour Xe lamp‐HX1, Model PE300UV, ISS). All the experiments were carried out under ambient conditions. Photocatalytic responses were hourly monitored in terms of the amount of hydrogen generated at 1–4 h time intervals. Hydrogen gas has very small density and not soluble in water. Therefore, the evolved hydrogen was collected into the inverted gas collection graduated bottle by displacement of water from a container. The collected gas was checked with the gas chromatograph  throughout the course of the reaction.Photocatalytic cleavage of the water for hydrogen generation was carried out using the powder of photocatalytic molecular device (0.3 g powder of Pt/CeOThe authors declare no conflict of interest.


In “Disentangling drivers of the abundance of coral reef fishes in the Western Indian Ocean,” which was published in volume 9 issue 7, April 2019, the word Comoros should be changed to Madagascar in the Discussion section 4.2. The correct information is printed below.Average composite values (2010–2015) of SST varied by ~1°C between Madagascar and the other three countries .