The integrated method reduced the RSME for both the hip and the knee joints below 5°, and no statistically significant differences were found between MoCap and InCap kinematics. This was consistent across all the different analyzed movements. The developed method was integrated on an MSK model workflow, and it increased the sensor-to-segment calibration accuracy for an accurate estimate of 3D joint kinematics compared to MoCap, guaranteeing a clinical easy-to-use approach.Inertial capture (InCap) systems combined with musculoskeletal (MSK) models are an attractive option for monitoring 3D joint kinematics in an ecological context. However, the primary limiting factor is the sensor-to-segment calibration, which is crucial to estimate the body segment orientations. Walking, running, and stair ascent and descent trials were measured in eleven healthy subjects with the Xsens InCap system and the Vicon 3D motion capture (MoCap) system at a self-selected speed. A novel integrated method that combines previous sensor-to-segment calibration approaches was developed for use in a MSK model with three degree of freedom (DOF) hip and knee joints. The following were compared: RMSE, range of motion (ROM), peaks, and R The assessment of locomotor function needs to be taken out of the lab and into the patient’s ecological environment to evaluate the true impact of specific clinical treatments such as gait retraining or pain medication on joint function. To do so, joint kinematics and ultimately loading should be accurately measured in real-life conditions. An inertial capture (InCap) system combined with a musculoskeletal (MSK) model that represents the complex kinematics patterns and the mechanical interactions of the joint can be used for instance to evaluate patients’ knee adaptations to instructed gait pattern modifications , aiming to decrease knee loading and eventually contributing to intermediate pain alleviation ,2,3,4,5 Laboratory-based 3D motion capture (MoCap) combined with in-ground embedded force plate systems has long been considered the gold standard method for capturing human movement and musculoskeletal function. Consequently, the use of 3D motion capture data as an input for MSK modelling workflows has become an attractive method to allow the estimation of different movement parameters , however restricting the evaluation of locomotor function to highly controlled laboratory settings. In the context of degenerative joint diseases, the validity of MoCap based methods for the evaluation and the analysis of human movements may not reproduce real-life movement patterns. Indeed, laboratory-based methods that document locomotor function represent a limited view of an individual patient’s condition and the need for more ecological studies of locomotor function are being advocated . Recently, wearable sensor technologies have gained massive development, with increased performance due to reduced weight and size, long battery life, and versatility . In partBased on the current limiting accuracy, especially for non-sagittal plane kinematics, it is crucial to improve the accuracy of InCap 3D joint kinematic estimations to reach a similar level of accuracy as seen in 3D MoCap. Such improvements would provide clear applications in several MSK disorders  and cerebral palsy), where the joints range of motion (ROM) or movement extremes are necessary to properly analyze and monitor joint loading in clinical practice ,15.Kinematics parameters such as gait speed, stride length, or joint angles in patients with osteoarthritis (OA) can discriminate the OA population from healthy subjects . Furthermanual calibration procedure utilizes manual alignment of the sensor reference system axes with the relative body segment axes [static calibration procedures, which use a sensor’s gravity vector aligned with the appropriate body segment axes during static poses such as the N-pose or the T-pose [third calibration method is a functional calibration method that assumes a perfect alignment of the sensor and the joint rotation axes during specific movements . This method has proven to be more appropriate for estimating frontal and transverse angles. For example, methods based on hip adduction or cycling motion show errors below 3.8° for sagittal plane angles and below 6° for non-sagittal plane angles during walking and cycling [The final method is anatomical calibration which uses external devices with integrated sensors to locate the joint axes with respect to the sensor axes system [The main factor limiting kinematic accuracy using InCap systems is the sensor calibration . InCap sent axes . This mee T-pose . This mee T-pose ,28,29,30e T-pose  informede T-pose . However cycling ,35,36,37s system ,39. Thiss system . Althougs system  combinedAdvanced algorithms have been developed to overcome other current calibration accuracy limitations such as soft tissue motion artifacts caused by no-rigid connection between the body and the sensors ,43 and mThis study aims to (1) develop an integrated functional sensor-to-segment calibration open-source workflow applicable in a clinical context and therefore independent of 3D MoCap for computing three-dimensional pelvis, torso, and lower limb joint kinematics  with IMU sensors based on an MSK model in the open-source software OpenSim; and (2) evaluate its accuracy compared to the gold standard MoCap system during walking, stair ascent and descent, and running. Ideally, the accuracy of such a workflow would allow estimation of kinematics within an accuracy of ±5°, which allows discriminating joint kinematic of healthy and OA subjects, and it would allow for the measuring of differences in joint kinematics targeted during gait retraining. Furthermore, performance would be expected to be similar during a variety of different relevant, functional movements.3D MoCap data was collected at the Movement & Posture Analysis Laboratory Leuven (MALL), at the Department of Movement Science, KU Leuven, Leuven, Belgium. This research was in accordance with the ethical guidelines provided by the ethical research committee KU Leuven (reference no. G-2021-3436). Eleven healthy adults with no musculoskeletal or neurological disorders volunteered for the study . All participants provided written informed consent, prior to data collection.Full-body InCap data were collected using the Xsens MVN Awinda system , for which 17 IMUs were positioned based on the manufacturer guidelines . To record the motion at 60 Hz, Xsens MVN Analyze Pro 2021.0 was used. Simultaneously, a 10 infrared camera MoCap system  was used as the gold standard reference system to record the 3D position of the 54 reflective markers attached on anatomical landmarks of the different body segments (Plug-in Gait Reference Guide—Nexus 2.12 Documentation—Vicon Documentation) at 120 Hz . Both sy3D InCap and MoCap data were collected, while each subject performed four barefoot movement tasks at self-selected speed: (a) 10 m of over ground walking, (b) 10 s of treadmill running, (c) ascending, and (d) descending stairs. For calibration, participants started each task with specific postures/movements: (i) N-pose, standing with arms to the sides, feet shoulder width apart, and facing forward for a period of five seconds, (ii) sitting in a wooden chair facing forward, (iii) knee extension followed by knee flexion for each leg (iv) standing in N-pose for three seconds, (v) hip abduction/adduction movement (three repetitions for each leg) with no movements in knee or ankle joints; the latter movement execution enforced similar angular velocities of the three segments—thigh, shank and foot, and (vi) participants performed the tasks (a—b—c—d). Raw data recorded by the InCap and the MoCap system were exported to .mvnx and .c3d files using Xsens MVN Analyzer Pro 2021.0 and Nexus 2.12, respectively, and they were imported into MATLAB . For both systems, all movements were time-normalized  and processed using custom-built MATLAB scripts.3D MoCap kinematics were calculated using the inverse kinematics tool in OpenSim v4.2 using an MSK model with 21 segments and 41 degrees-of-freedom (Hamner 2010 ). The knThis integrated sensor-to-segment calibration method (method 3—M3) can be used independent of 3D MoCap data, and it relies on a combination of previously published methods ,46,47,48Performance of the developed method M3 (combination of M1 and M2) was evaluated against the anatomical calibration that uses an anatomical position as an initial guess for the standard static calibration with the OpenSense toolkit (M0). Specifically, the anterior–posterior axis of the pelvis sensor was aligned with the anterior–posterior axis of the pelvis segment of the MSK model; then, torso, thigh, shank, and foot sensors were assumed to be aligned with that of the pelvis. For the calibration reference system, segment and joint coordinate systems were defined according to ISB recommendations ,51.2, range of motion (ROM) difference, and the maximum peak difference between each calibration method. Significance was set at p ≤ 0.05. A false discovery rate (FDR) statistical approach in multiple assumptions testing was used for multiple comparisons in order to correct for random events that falsely appear significant. A q-value threshold of 0.05 (FDR of 5%) among all significant variables was used. To evaluate the functional relevance of the observed differences between methods for the different joints, the error in kinematics between different InCap methods and MoCap was evaluated against reported differences in kinematics between control and OA subjects [For multiple comparisons, the Wilcoxon rank test was implemented in MATLAB, and it was used to statistically assess significant differences in the root mean squared error (RMSE), the determination coefficient Rsubjects ,53,54,55Complete tables and figures for comparison of all sagittal and non-sagittal plane kinematics are presented in the 2 > 0.7 ± 0.2. ROM errors were 1.1 ± 0.6°, 2.2 ± 1.8°, and 2.9 ± 2.2°, and peak magnitude errors were 1.6 ± 2.5°, 3.2 ± 2.1° and 4.3 ± 2.7° for knee flexion, abduction, and internal rotation angle, respectively (see Knee: vely see . M3 showThe RMSE between kinematics estimated from MoCap and the different methods (M0–M1–M2–M3) during walking are shown in Similar to the knee, the integrated method M3 significantly decreased on average the estimated angle offset error for all the other joints observed—lumbar, pelvic, hip, and ankle joint—compared to the other methods: M0 (10°), M1 (8°), M2 (4°), M3 (2°)— during w2 > 0.7 ± 0.2. ROM errors were 3.1 ± 2.6°, 3.6 ± 1.8°, and 3.3 ± 3.1° for hip flexion, hip adduction, and hip rotation, respectively. The errors for the peak were 3.2 ± 1.1°, 3.5 ± 2.3°, 4.0 ± 3.3°. Compared to MoCap, significant difference was found in the peak hip rotation in M0 and between M0 and M3. The RMSE averaged over all joint-planes improved for M3: 3.4 ± 1.2° compared to M1: 5.6 ± 3.5° and M2: 5.2 ± 2.8°.Hip: M3 showed the highest accuracy compared to the other InCap methods, with average RMSE being significantly reduced from 7.1 ± 3.1° to 3.2 ± 1.9°, for hip non-sagittal plane angles  compared2 > 0.8 ± 0.2. ROM errors were 1.4 ± 2.2°, 2.9 ± 3.9° and 2.9 ± 2.1° for lumbar extension, bending, and rotation, respectively. The errors for the peak were 3.4 ± 3.2°, 3.1 ± 2.7° and 2.2 ± 2.4°. Compared to MoCap no significant difference was found between methods. RMSE averaged over all joint-planes improved for M3: 3.9 ± 2.2° compared to M1: 8.8° ± 3.5° and M2: 6.7 ± 3.2°.Lumbar: M3 showed the highest accuracy compared to the other InCap methods, with average RMSE being significantly reduced from 9.1 ± 4.0° to 3.9 ± 2.2° over all joint-planes compared to M0, with average R2 > 0.7 ± 0.3. ROM errors were 1.3 ± 1.5°, 3.2 ± 1.7° and 2.8 ± 1.7° for pelvic tilt, list, and rotation, respectively. The errors for the peak were 2.6 ± 2.0°, 1.6 ± 1.6° and 2.2 ± 2.4°. Compared to MoCap significant difference was found for pelvic rotation in M0, M1, and between M2 and M3. RMSE averaged over all joint planes improved for M3: 4.0 ± 1.3° compared to M1: 8.0 ± 4.0° and M2: 4.2 ± 1.8°.Pelvic: M3 showed the highest accuracy compared to the other InCap methods with average RMSE being significantly reduced from RMSE 10.5 ± 3.7° to 4.0 ± 1.3°, over all joint planes compared to M0, with average R2 > 0.8 ± 0.2°. ROM errors were 3.7 ± 2.9° and 3.1 ± 1.5° for dorsi-plantar flexion and prono-supination. The errors for the peak were 4.1 ± 2.2° and 4.6 ± 2.7°. Compared to MoCap, no significant difference was found. RMSE averaged over all joint planes improved for M3: 4.8 ± 1.9° compared to M1: 5.6 ± 3.5° and M2: 5.0 ± 3.1°.Ankle: M3 showed the highest accuracy compared to the other InCap methods with average RMSE being significantly reduced from RMSE 7.3 ± 3.8° to 4.8 ± 1.9° over all joint planes compared to M0, with RFor running , accuracThe aim of this study was to develop an integrated functional sensor-to-segment calibration applicable in a clinical context for computing three-dimensional pelvis, torso, and lower limb joint kinematics  during common daily life activities based on IMU sensors combined with an OpenSim MSK model. The developed method addresses previously documented challenges within literature, specifically sensor-to-segment calibration to estimate the body segment orientations from the sensor orientations without MoCap data and the integration with a complex MSK modeling workflow to estimate 3D joint kinematics ,32,56,57In contrast to previous studies that only focused on specific activities ,35,58, tIn addition, we evaluated the performance of the integrated method M3 by calculating the differences in ROM and in angle peaks compared to MoCap. The ΔROM and Δpeak error of sagittal and non-sagittal plane angles varied between 2° and 5°. Such ROM and peak differences have been demonstrated to be important parameters for a variety of clinical applications ,63. For Given that the InCap system is an easy-to-use system capable of accurately detecting the difference in ROM of 5°, it is a valid alternative to conventional MoCap systems for the assessment of 3D joint kinematics in clinical and ecological settings. However, its accuracy decreases for frontal and transverse plane angles: in particular, for joints with small ROM such as knee adduction (2–5°), where a RMSE threshold of 5° may represent an error of 100% of ROM. Here, correspondence in terms of correlation of the kinematics may yield a more accurate comparison for evaluating the accuracy of the InCap systems ,53,65. CTherefore, for applications where the specific clinical intervention, e.g., gait retraining in patients with knee OA, needs accurate measurement of knee non-sagittal plane angles the developed functional calibration approaches in combination with an MSK model with multiple degrees-of-freedom could be used if future studies confirm similar levels of accuracy in OA patients ,69. HoweOur study showed that the accuracy of the developed calibration method is higher than previously developed methods—M0, M1, M2—and that the effect is consistent across multiple movement tasks. Nevertheless, there seems to be a negative effect on the calculated error due to the more dynamic movement tasks such as running. This may be due to the limited sample frequency used by the InCap system, which was only 60 Hz. Therefore, future studies should consider a higher sample frequency during such high dynamic movements. Another limitation of the presented method is that it relies on the ability of the subjects to perform specific joint movements—the hip abduction/adduction movement, and the sit-to-stand movement. Therefore, functional calibration may be challenging in subjects with limited joint mobility or impairments. Furthermore, the repeatability of sensor placement was not investigated in this study. Sensor placement repeatability is operator dependent and the introduced misalignment is known to account for up to a 55% error in the sagittal plane angles, while the sensor-to-segment calibration misalignment is known to account for up to 40–50% in non-sagittal plane angles ,70. By iThis study impacts the use of wearable sensor technology for documenting 3D knee joint kinematics . However, to evaluate the impact of specific clinical treatments out-of-the-lab such as gait retraining or pain medication, the knee loading should ideally be accurately measured in real-life conditions, e.g., to estimate the impact of the patient’s gait pattern modifications  on knee loading ,73,74,75In conclusion, in this study we developed a novel integrated sensor-to-segment calibration method for IMU sensors that improves non-sagittal plane kinematic estimates in a range of movements. Moreover, keeping the focus on future clinical applications, this method aimed to calculate the 3D kinematics of the torso, pelvis, and lower limbs, allowing for a not-predefined location of the IMU-sensors thereby guaranteeing a clinical easy-to-use approach.


Optimization of degradation conditions and the degradation levels were studied by high performance liquid chromatography, the surface morphology of the ZVZ and AC changed before and after the reaction were observed by scanning electron microscope. The degradation intermediates were detected by gas chromatography- mass spectrometry and ion chromatography, and inferred the degradation pathway. The degradation effect of [C4mim][BF4] was best with ultrasonic assistance, pH 3 and an AC/ZVZ ratio of 1:1. The degradation of [Cnmim][BF4] in aqueous solution exceeded 91.7% in 120 min, and the mineralization level exceeded 88.9%. The surface of smooth and dense ZVZ particles became loose flocculent and the porous surface of AC became larger and rougher after reaction. The degradation pathway suggested that the imidazolium ring was sulfurized or oxidized, and then the ring was opened to form N-alkyl formamide and N-methyl formamide. ZVZ/AC micro-electrolysis combined with ultrasonic irradiation is an effective method to remove ILs, which provides new insight into IL degradation.Increased attention has been given to the removal of ionic liquids (ILs) from natural water environments. In this work, 5 kinds of 1-alkyl-3-methylimidazoliumtetrafluoroborate ([C An increased number of ILs is used as media and solvents in cutting-edge science and technology. However, because of the low vapor pressure, good water solubility and stability of ILs, waste liquid that is discharged during the use of ILs remain in the soil or water body for a long time, which results in different extents of environmental pollution9. The potential toxicity of ILs has attracted attention, including the toxicity of ILs to higher plants, and the toxic effects on aquatic organisms, microorganisms, cells and skin11. Therefore, it is necessary to find a suitable and effective method to remove ILs from waste liquid.Ionic liquids (ILs) have unique physical and chemical properties, which make them widely applicable in synthetic chemistry, electrochemistry, pharmacy, biology, nanotechnology, separation technology and other fields12, including Fenton oxidation degradation of ILs14, ultraviolet (UV)-H2O2 or UV-TiO2 photo-assisted degradation19, ultrasonic (US)-H2O2/acetic acid system oxidative degradation20, electrochemical oxidation and electro-Fenton system degradation25. Among the degradation systems reported above, some are time-consuming and inefficient, some are complex in subsequent treatment , whereas others have a high energy consumption.Because most ILs are difficult to biodegrade, the chemical degradation of ILs has been discussed27. On this foundation, zero valent zinc (ZVZ) was chosen to replace zero valent iron, because the potential of zinc was lower than that of iron, so it was speculated that the galvanic effect was stronger.Micro-electrolysis technology is considered to require low energy consumption and is a high-efficiency wastewater treatment method. Previous studies have shown that zero-valent iron activated carbon (ZVI/AC) micro-electrolysis can degrade ILs2+ and free radicals. The IL structure is destroyed by redox reaction. In this work, 1-alkyl-3-methylimidazolium tetrafluoroborate ) IL was degraded by a battery reaction without a power supply, which provided a scientific basis for the treatment of IL wastewater.In this study, ZVZ and activated carbon (AC) were combined to form a micro-electrolysis system. The potential difference between low-potential zinc and high-potential carbon is large, and the IL wastewater to be degraded acts as an electrolyte to form numerous microgalvanic cells, in which zinc is the anode and carbon is the cathode and produces nascent hydrogen, ZnIt was worth mentioning that the degradation of ILs in the micro-electrolysis system was related to their electrochemical stability, the results of this paper will provide relevant reference for the design and their electrochemical stability of ILs.4mim][BF4] (1 mmol/L) was used as a model imidazolium IL to optimize the degradation of the imidazolium ILs at 30 °C in the US-ZVZ/AC micro-electrolysis system, for the initial pH, dosages and ratio of ZVZ and AC.[C4mim][BF4] was investigated at four different pHs of 3.0, 4.0, 5.0 and 7.0 in the US-ZVZ/AC system. The degradation of [C4mim][BF4] was carried out by using a solution of 25 ml [C4mim][BF4] (1 mmol/L) 0.15 g AC and 0.15 g zinc powder, US 45 kHz, and the data are presented in Fig. 4mim][BF4] increased with a decrease of solution pH, and the most efficient pH was 3.0. Under acidic condition, zinc powder has high activity in the free H+ environment and is easy to become an ionic state, thus accelerating the reaction. This result is analogous to our previous reports on imidazolium ILs in the US-ZVI/AC system27.The influence of initial solution pH on the degradation of [Cim][BF4]  mmol/L 04mim][BF4] was examined at pH 3.0 and the three initial concentration ratios of AC/ZVZ were chosen as 1:2, 1:1 and 2:1, i.e., 0.15/0.3, 0.15/0.15 and 0.15/0.075 (g/g). The degradation effect of [C4mim][BF4] was best when the AC/ZVZ ratio is 1:1, as shown in Fig. According to the method in “4mim][BF4] aqueous solution systems were prepared, the pH was adjusted to 3, and the degradation of [C4mim][BF4] was determined by using 0.1 g/0.1 g, 0.15 g/0.15 g, 0.2 g/0.2 g and 0.3 g/0.3 g AC/ZVZ, respectively. The results are shown in Fig. According to the above experimental method, four 25 ml 1 mmol/L [C4mim][BF4] with magnetic stirring and ultrasound was discussed. Two 1 mmol/L [C4mim][BF4] systems were prepared, and the degradation was carried out under the above conditions with the assistance of 45 kHz ultrasound and magnetic stirring, respectively. Figure 4mim][BF4] can be increased significantly by using US assistance. The degradation extent of [C4mim][BF4] was 47.89% in 120 min and 93% under US-assisted conditions. This outcome may have resulted because of the cavitation effect of ultrasound. Many small bubbles are produced in the reaction liquid and these bubbles move, grow or burst with the vibration of the surrounding medium, which results in a local instantaneous high temperature and high pressure and a chemical effect. In air, water molecules decompose to produce highly active free radicals, such as H⋅, HO⋅ and O⋅, which can decompose highly stable organic compounds20. The mechanical and thermal effects of sound waves on the particle surface contribute to organic matter removal. In the ZVZ/AC micro-electrolysis system, many microgalvanic cells were generated between C and Zn, Zn (II), nascent hydrogen was released, and new free radicals were generated continuously under the action of ultrasound, which results in the effective decomposition of [C4mim][BF4] IL.Based on the results from the experiments, the degradation of Cmim[BF4] 4mim][BF4] occurred under 45 kHz US radiation, whereas the combination of US-ZVZ degraded ~ 13% of the [C4mim][BF4] at 120 min. The degradation efficiency in the US-ZVZ/AC system improved significantly, so a synergistic effect occurred between the US and ZVZ/AC system, which improves the degradation of [C4mim][BF4].In our experiments, no degradation of [C4mim][BF4] degradation were pH3, AC/ZVZ dosage of 0.15 g/0.15 g (ratio 1:1) and ultrasound frequency 45 kHz ([IL] = 1 mmol/L at 30 °C).As can be seen from Table 27.AC used in the experiments had been saturated by IL in advance, and the concentration falling of IL in the degradation process was no longer affected by AC adsorptionIt is worth noting that under some degradation conditions, the degradation level of ionic liquid has a low point in the early stage of degradation, most of which are at 10 min. It is speculated that the degradation kinetics of different reaction stages are different. The kinetics of this degradation reaction needs to be further discussed.4mim][BF4] in advance if necessary). After US-assisted micro-electrolysis, the ILs that had adsorbed in the pores disappeared, the AC pores were larger and the surface was rough  under optimized conditions according to the method described in section “2mim][BF4], [C4mim][BF4], [C6mim][BF4], [C8mim][BF4] and [C10mim][BF4] was 93.6%, 93.0%, 96.8%, 92.8% and 91.7%, respectively, at 120 min. No significant influence of alkyl side chain length was observed, and this result agrees with that for iron carbon micro-electrolysis IL degradation by our team26.Degradation was carried out on 1 mmol/L [C][BF4], Cmim[BF4],nmim][BF4] in the ZVZ/AC micro-electrolysis system assisted by US irradiation, representative UV spectra for the solution were determined as a function of reaction time. Figure 4mim][BF4] in water was characterized by a main band in the ultraviolet region at 212 nm. This peak was associated with the imidazolium ring of [C4mim][BF4]. With an increase in degradation time, the absorption of the degradation mixture at 212 nm decreased gradually. The absorption peak at 212 nm almost disappeared at 120 min, which indicates that the opening of the imidazolium ring and effective degradation of [C4mim][BF4]. [C2mim][BF4], [C6mim][BF4], [C8mim][BF4] and [C10mim][BF4] have similar trends.To clarify the changes in structural characteristics of [Cnmim][BF4] degradation solution at different reaction times , saturated the sample with sodium chloride, extracted the intermediate products with organic solvent, and analyzed and identified the product by GC–MS. The inorganic anions that were produced by degradation were detected by ion chromatography. The chemical structures and names of the intermediates are shown in Table The identification of intermediates in the degradation process can provide valuable information about the degradation pathway. Therefore, we took 10 ml of [C2mim][BF4] and [C4mim][BF4]. e2 in the table is the urea that was formed by the oxidation of three C-H on the [C2mim][BF4] imidazole ring to three carbonyl groups and the opening of the imidazole ring. g2 and i are products that were formed by bond breakage between the two carbonyl groups of e2. Seven, seven and eight degradation products were detected in [C6mim][BF4], [C8mim][BF4] and [C10mim][BF4], respectively. For [C10mim][BF4] as an example, where a10 is obtained by sulfurizing the H of C2 on the imidazolium ring, b10 is the product of oxidation of C-H on C2 of the imidazole ring to carbonyl, c10 is the product of oxidation of C-H on the other two C of the imidazole ring to carbonyl, d10 is the urea that formed after the opening of the c10 imidazole ring, f10 is the product of d10 removing N-methylformamide, g10 is the product of further decomposition of d10 and f10, and h10 is generated after dehydrogenation.Four degradation products were detected in [C2mim][BF4] in which oxidant radicals that are generated in the US-ZVZ/AC reaction system attack H on the 2, 4, 5 carbon atom of the imidazolium ring to produce 1-ethyl-3-methyl-2, 4, 5-trioxyimidazole, and then oxidize to open the ring to obtain e2. g2 and i are obtained by C–C bond breakage between the two carbonyls.According to the chemical structure of the intermediates, we propose a degradation pathway of [C4mim][BF4] are consistent with those of [C2mim][BF4], so it is speculated that [C4mim][BF4] has a similar degradation pathway.The degradation products of [C6mim][BF4], [C8mim][BF4] and [C10mim][BF4] ILs in US-ZVZ/AC system are as follows:The possible degradation pathways of [BF4] is shown in Fig. The possible degradation pathway of [BF4]  can be evaluated by the TOC content. The TOC before and after [Cnmim][BF4] degradation can be determined, and the TOC removal during degradation can be calculated. We measured the TOC of the ILs degradation mixture at 0 and 120 min, and obtained the TOC removal at 120 min  with different side chain lengths was almost unaffected by the length of the alkyl side chain. During the degradation of 1-alkyl-3-methylimidazolium tetrafluoroborate, H on the imidazolium ring was sulfurized to 1-alkyl-3-methyl-2-thioimidazole, or oxidized to 1-alkyl-3-methyl-2-oxyimidazole, and further oxidized to 1-alkyl-3-methyl-2, 4, 5-trioxyimidazole. Through the ring opening of imidazole and bond breakage between two carbonyl groups, corresponding intermediate products formed and decomposed into N-alkyl formamide and N-methyl formamide, and then alkyl chain fatty acids (such as formic acid) and nitrate ions were generated. More than 88.9% of the TOC was removed during the degradation.This study demonstrated that 1-alkyl-3-methylimidazolium IL can be degraded rapidly and effectively in the US-ZVZ/AC micro-electrolysis system, and the degradation exceeded 91.7% in 120 min. The degradation of imidazolium IL [BF4]), 1-butyl-3-methylimidazolium tetrafluoroborate ([C4mim][BF4]), 1-hexyl-3-methylimidazolium tetrafluoroborate ([C6mim][BF4]), 1-octyl-3-methylimidazolium tetrafluoroborate ([C8mim][BF4]) and 1-decyl-3-methylimidazolium tetrafluoroborate ([C10mim][BF4]). Stock solutions of [Cnmim][BF4]  (25 mmol/L) were prepared in purified water  and stored at 4 °C. The names, molecular masses and chemical structures of the ILs are listed in Table ILs were from Chengjie Chem. Co. Ltd.  and were used without further purification. The ILs that were used in our work included: 1-ethyl-3-methylimidazolium tetrafluoroborate ([BF4]. A mixture of acetonitrile  with 25 mmol/L of phosphate buffer (KH2PO4/H3PO4) in 0.5% triethylamine (pH3.0) was used for the analysis of [C8mim][BF4] and [C10mim][BF4]. Each sample was measured in triplicate at various reaction times. According to the measured concentrations of [Cnmim][BF4] at the 0 degradation time, the degradation degree of [Cnmim][BF4] at moment t can be calculated.HPLC  was used to determine the concentration of the ILs as described in literatureSEM  was used for the analysis of the ZVZ and AC particle surface morphology before and after the degradation reaction at a 20-kV and 5-kV beam potential.24. The structural identification of intermediates was based on the NIST05 library for the mass spectra and combined with an interpretation of the fragmentation pathways.GC–MS  was used to identify the degradation intermediates. The degradation products were extracted with a mixture of benzene-ethyl acetate-ether (2:2:1) and analyzed as described in the literature2CO3—1.0 mmol/L NaHCO3 aqueous solution; column temperature: 30 °C; suppressor current: 24 mA; flow rate: 1.2 mL/min; injection volume: 10 μL; and a Chromeleon7 workstation was used for data processing.IC  was used to analyze the anionic species in solution after degradation. The analytical IC conditions were: column: IonPac AS14 (4.0 × 250 mm) + AG14, (4.0 × 50 mm); mobile phase: 3.5 mmol/L NaThe TOC of the initial and degraded samples was measured by a TC/TN analyzer  according to a thermal catalytic oxidation principle.


Nevertheless, the high desorption temperature, sluggish kinetics, and rapid capacity decay hamper its commercial application. Herein, 2D TiO2 nanosheets with abundant oxygen vacancies are used to fabricate a flower-like MgH2/TiO2 heterostructure with enhanced hydrogen storage performances. Particularly, the onset hydrogen desorption temperature of the MgH2/TiO2 heterostructure is lowered down to 180 °C (295 °C for blank MgH2). The initial desorption rate of MgH2/TiO2 reaches 2.116 wt% min−1 at 300 °C, 35 times of the blank MgH2 under the same conditions. Moreover, the capacity retention is as high as 98.5% after 100 cycles at 300 °C, remarkably higher than those of the previously reported MgH2-TiO2 composites. Both in situ HRTEM observations and ex situ XPS analyses confirm that the synergistic effects from multi-valance of Ti species, accelerated electron transportation caused by oxygen vacancies, formation of catalytic Mg-Ti oxides, and stabilized MgH2 NPs confined by TiO2 nanosheets contribute to the high stability and kinetically accelerated hydrogen storage performances of the composite. The strategy of using 2D substrates with abundant defects to support nano-sized energy storage materials to build heterostructure is therefore promising for the design of high-performance energy materials. MgHThe online version contains supplementary material available at 10.1007/s40820-022-00891-9. Particularly, a high temperature of at least 280 °C is required at 1 bar for desorption of coarse-grained MgH2, which far exceeds the working temperature of the proton exchange membrane fuel cell (PEMFC). Consequently, it is essential to improve the hydrogen storage performances of MgH2 for possible applications in the hydrogen energy field.With the ever-increasing global energy demands and environmental concerns, clean and renewable energy is urgently demanded worldwide. Hydrogen has been considered as one of the most promising candidates to alternate fossil fuels, due to its intrinsic clean nature, environmental-friendliness, and high gravimetric energy density 142 MJ kg2 MJ kg−1−1) , 6. Comp−1) . Among vtoxicity , 9. Unfo2 [2. These additives include transition metals (TM) [2, Nb2O5, V2O3) strike more attentions and have been proven to be the most effective additives to ameliorate the hydrogen sorption performances of MgH2, since transition metals with multiple valences favor the electron transfer during the de/re-hydrogenation processes [2 by the synthesis of core (Mg)-shell (TM) structure. The results show that Mg-Ti composites, which can release hydrogen even under 200 °C, exhibit optimal desorption performances, owing to the lower electronegativity of Ti. Notably, it has been proved that TiO2 possesses a superior catalytic effect than metallic Ti [2 has many advantages, such as high natural abundance, low cost, and nontoxicity. Though the addition of catalysts through simple ball milling method can indeed ameliorate the kinetics of MgH2 to some extent, the improvement is restricted by the inevitable self-aggregation and growth of the additives during cycling, as well as the low density of exposed active sites.In the last few decades, extensive researches have been executed toward fabricating novel Mg-based hydrogen storage materials with enhanced performances, such as adding catalysts –13, allo2 –19. Fromals (TM) , 21, traals (TM) , 23, nitals (TM) , 25, halals (TM) , 27, andals (TM) , 29. Amorocesses . The Ferrocesses . Cui et allic Ti . Further2, which owes the shortened diffusion path for hydrogen atoms and increased active sites for re/de-hydrogenation, has been successfully synthesized through various nanotechnologies, giving rise to improved hydrogen storage performances. For instance, the nano-confinement method has been proved to be a facile strategy to simultaneously improve the thermodynamics and kinetics of MgH2 while reducing the capacity loss as little as possible. Nano-sized MgH2 confined into porous carbon has demonstrated a noticeable decrease in its initial desorption temperature (~ 90 °C) compared with its bulk counterpart [2 nanoparticles anchored on graphene sheets that exhibit excellent hydrogen sorption properties. The composites absorb 5.4 wt% H2 within 10 min at 200 °C, and identical capacity can be released at the same temperature within 150 min. It is noted, however, that neither graphene nor porous carbon possesses good catalytic activities, and additional catalysts are needed to further improve hydrogen storage performances of MgH2. Although huge achievements have been achieved, it is still far from enough to enhance the hydrogen storage performances of MgH2 by solely depending on any methods mentioned above. In this respect, it is essential to fabricate MgH2 with extraordinary hydrogen storage performances by taking full advantage of both the intrinsic (nano-size) and the extrinsic  strategy.On the other hand, from an intrinsic point of view, nano-sized MgHnterpart . Yu et anterpart  synthesi2 nanosheets (TiO2 NS) and MgH2 by a simple ball milling method would cause irreversible agglomeration of TiO2 NS, leading to seriously reduced catalytic active sites and deterioration of various spectacular intrinsic properties of 2D materials [2 (0D)/TiO2 (2D) heterostructure using graphene-like 2D TiO2 NS is triggered as it can not only facilitate the formation of nano-sized MgH2 through the confinement effect of TiO2 NS but also possess catalytic effect for MgH2. More importantly, the construction of MgH2/TiO2 heterostructure can effectively inhibit the self-stacking of TiO2 layers, leading to increased surface catalytic sites for MgH2. Additionally, defect engineering has also aroused more attention to adjusting the band structure and providing more catalytic active sites [2 has been proved to remarkably promote electron transportation, enhance electrical conductivity, and provide more active sites for the capture and diffusion of hydrogen [Recently, graphene-like two-dimensional (2D) transition metal oxides have been extensively studied due to their potential in energy-related applications with high surface area and unique electronic/thermal properties –37. Howeaterials . Thus, tve sites . Among thydrogen , 41.2/TiO2 heterostructure with remarkably enhanced hydrogen sorption properties through facile solvothermal approach, where 2D TiO2 NS exhibit both nano-confinement and catalytic effects. Surprisingly, we found that the oxygen vacancies were introduced into TiO2 NS in the process of solvothermal treatment. Owing to the presence of oxygen vacancies and unique heterostructure, the obtained MgH2/TiO2 heterostructure exhibits impressive hydrogen storage performances with high stability and accelerated kinetics. Specifically, the 60MgH2/TiO2 heterostructure presents low onset desorption temperature (180 °C) and superior long-term cycling stability with a capacity retention of 98.5% after 100 cycles at 300 °C, remarkably higher than that of the state-of-the-art TiO2 catalyzed MgH2 systems.In light of the above consideration, herein, taking full advantage of the intrinsic (nano-size) and extrinsic  effects, we report a first attempt to design a novel self-assembled MgH2 nanosheets were synthesized by a surfactant self-assembly method [2 NS. Moreover, to improve the crystallinity and remove the impurities in TiO2 NS, the pristine TiO2 NS were annealed at different temperatures  for 2 h in the air .TiOy method . Firstly2/TiO2 heterostructure was prepared by a bottom-up solvothermal strategy of in situ growth of MgH2 nanoparticles on TiO2 substrates. Di-n-butyl-magnesium  and cyclohexane  were used as received states without further treatment. As provided in Table S1, different loading capacities of MgH2 on TiO2 NS of 40, 50, 60, 70, and 80 wt% were synthesized by changing the ratio between MgBu2 and TiO2 NS, which were denoted as 40MgH2/TiO2, 50MgH2/TiO2, 60MgH2/TiO2, 70MgH2/TiO2, and 80MgH2/TiO2, respectively. In a typical synthesis of 60MgH2/TiO2, 30 mg as-prepared TiO2-300 NS and 1.8 mL MgBu2 were firstly dispersed in 40 mL cyclohexane under sonication. The mixture was then transferred into a 100 mL custom-designed autoclave, filled with 4.5 MPa H2 pressure, and heated at 200 °C for 12 h under vigorous magnetic stirring [2/TiO2 heterostructure) were collected by centrifugation, washed with cyclohexane several times, and dried at 80 °C overnight under dynamic vacuum. Moreover, blank MgH2 without TiO2 NS were also prepared under the same conditions for comparison.MgHstirring , 42. Finα radiation operated at 40 kV and 15 mA. The samples were prepared in an argon-filled glovebox and sealed in a custom-designed holder covered by the Scotch tape to avoid the possible oxidation during the XRD tests. The microstructure and morphological analyses were performed by scanning electron microscope  and transmission electron microscope . For TEM observations, the preparation of samples was carried out in the glovebox. The samples were dispersed in cyclohexane, sonicated, dropped on a copper grid, and rapidly transferred to the equipment. The preparation conditions of samples for SEM tests were similar to those for the TEM observation, except that the samples were dispersed in cyclohexane and dropped onto the silicon wafer. The in situ high-resolution transmission electron microscope (HRTEM) observation was applied to explore the microstructural evolution during the decomposition of the MgH2/TiO2 heterostructure under an irradiation beam current of 3 nA [2/TiO2 composites were analyzed by X-ray photoelectron spectroscopy . For XPS analyses, a special air-proof transfer vessel was used to transfer the samples from glovebox to the equipment. Electron paramagnetic resonance (EPR) measurement was performed at room temperature using a Bruker spectrometer (EMXplus-9.5/12). Solid-state 1H magic-angle spinning nuclear magnetic resonance (MAS NMR) was carried out on a Bruker AVANCE NEO 600 spectrometer. Brunner–Emmet–Teller (BET) surface areas of all samples were measured by the N2 isothermal ad/de-sorption tests at 77 K on Autosorb-IQ3 apparatus. Differential scanning calorimetry  measurements were performed at different heating rates  under a flowing Ar atmosphere to analyze the desorption kinetic properties. The DSC data were further processed by the Kissinger method to figure out the desorption apparent activation energy (Ea) according to the following equation:β represents the heating rate, Tp refers to the peak temperature, R is the universal gas constant (R = 8.314 J mol−1 K−1), and C is a constant.The phase composition was characterized by using X-ray diffraction  with Cu K of 3 nA . The val−1. The initial hydrogen pressures for isothermal ab/de-sorption were 3 and 0.001 MPa, respectively. Besides, the hydrogenated samples were heated from room temperature to 350 °C with a heating rate of 2 °C min−1 to perform TPD measurements. The hydrogen pressure/duration/temperature parameters for cycling tests are set as 3 MPa/15 min/300 °C for re-hydrogenation and 0.001 MPa/15 min/300 °C for dehydrogenation. The hydrogen storage capacity was calculated in the weight percent of the entire composites including the TiO2 scaffold.The isothermal ab/de-sorption tests and temperature program desorption (TPD) measurements of samples were carried out by a Sievert-type pressure-composition-temperature apparatus . For the isothermal ab/de-sorption tests, samples were loaded into the vessel in the glove box and were heated to the preset temperature at a heating rate of 10 °C min2/TiO2 heterostructures is schematically illustrated in Scheme 2 NS and subsequent impregnation of MgBu2 as well as the bottom-up self-assembly of MgH2 nanoparticles anchored on TiO2 NS. In a typical surfactant self-assembly synthesis process of TiO2 NS, P123 and ethylene glycol (EG) acted as co-surfactants which played the role of structure-directing agents to form inverse lamellar micelles. The precursor oligomers (TTIP) were then self-assembled into lamellar structures with surfactant molecules, inducing the formation of layered inorganic oligomer agglomerates. Afterward, solvothermal treatment was carried out to promote the organization and crystallization of 2D TiO2 NS. Subsequently, the TiO2 NS were ultrasonically dispersed into the C6H12 solution followed by the dropwise addition of MgBu2. Finally, the MgH2/TiO2 heterostructure was obtained in the above dispersed suspension by hydrogenation of MgBu2 on TiO2 NS through a facile solvothermal method.The preparation procedure of the MgH2 NS were observed using SEM and TEM. As shown in Fig. 2 NS exhibit a morphology of 2D corrugated ultrathin nanosheets with their edges roll up owing to the surface tension, which is also observed in graphene as a general phenomenon. The sizes of the TiO2 NS vary from 150 to 200 nm as seen from the TEM images. The thickness of the prepared TiO2 NS was further characterized using AFM and TEM . The obtained TiO2 NS have a thickness of around 3.7 nm, corresponding to 5–6 monolayers, which confirms the formation of ultrathin nanosheets. In addition, the sheet-like structure of TiO2 can be clearly distinguished in the HAADF and BF images . It is worth emphasizing that TiO2-300 and TiO2-400 exhibit a similar sheet-like structure as the pristine TiO2 NS, except for the slight granulation after annealing at 400 °C. However, the crystallite size and crystallinity of TiO2-500 increase obviously and the 2D nanostructures can hardly be observed. This phenomenon coincides well with the previous report [2 NS in the hydrogen atmosphere and it can be seen that the sheet-like structure of TiO2 is still preserved after treatment in H2 at 300 °C for 2 h . Figure 2, TiO2-300, TiO2-400, and TiO2-500. According to the XRD patterns, pristine TiO2 NS are mainly composed of the anatase phase (JCPDS No. 21–1272), as well as a small amount of rutile phase (JCPDS No. 21–1276). After calcination treatment, the intensities of the diffraction peaks are much sharper and stronger, manifesting an increase in crystallinity, coinciding well with the SEM and TEM results . Additionally, the diffraction peaks of rutile phase disappear and crystalline diffraction peaks assigned to anatase TiO2 crystal can be identified when the annealing temperature increases up to 400 and 500 °C. The specific surface area and pore size distribution of TiO2 NS were further investigated . Besides, it still possesses a high specific surface area of 268.31 m2 g−1 even after calcination treatment (300 °C). Thus, we selected TiO2-300 as the substrates to fabricate the MgH2/TiO2 heterostructure, and the TiO2 mentioned below all refers to TiO2-300 unless otherwise specified. The TiO2 NS with ultrahigh specific surface area will provide more nucleation sites for nanoparticles, which could be ingeniously employed as the excellent scaffold for the self-assembly of MgH2 NPs.Furthermore, the influence of the annealing temperature on the morphology of TiOs report . Moreove2/TiO2 heterostructure, MgBu2 was used as the precursor and added into TiO2 suspension followed by the solvothermal method under 4.5 MPa of hydrogen pressure. Impressively, the color of TiO2 suspension fortuitously changed from white to dark blue with the addition of MgBu2. The obvious blue color indicates the presence of oxygen vacancies caused by MgBu2 [2/TiO2 heterostructure. It can be seen that TiO2 nanosheets still maintain a graphene-like structure after the self-assembly of MgH2 nanoparticles on them. Besides, 2D TiO2 nanosheets are automatically arranged into flower-like structures to support and encapsulate MgH2 nanoparticles, tremendously restricting the agglomeration of MgH2 NPs and simultaneously hampering the re-stacking of TiO2 NS in turn. According to the measurement of lattice spacing in HRTEM images  and (111) plane of TiO2 (JCPDS No. 21–1276), respectively. In addition, the diffraction rings in SAED pattern of 60MgH2/TiO2 , 70MgH2/TiO2 (233 °C), and 60MgH2/TiO2 (200 °C). Thus, we further explored the hydrogen storage performances of 60MgH2/TiO2, and the MgH2/TiO2 heterostructure mentioned below refers to the 60MgH2/TiO2 unless otherwise specified. In the first non-isothermal desorption process, the 60MgH2/TiO2 exhibits the initial desorption temperature of 200 °C with a total hydrogen release of 4 wt%. Notably, the 60MgH2/TiO2 composite possesses a lower onset desorption temperature of 180 °C along with a hydrogen capacity of 3.4 wt% in the second non-isothermal dehydrogenation process. The lower desorption temperature in the second cycle reveals that the catalytic activity of TiO2 NS varies upon de/ab-sorption behavior. The capacity loss after the first desorption may be attributed to the irreversible reaction between MgH2 and TiO2. It is therefore reasonable to speculate that the chemical environment of MgH2 and TiO2 in the first dehydrogenated sample should be considerably different from the as-synthesized sample, indicating the formation of catalytic species during initial hydrogen cycling, which is similar to what shown in previous reports [2 is 408.6 °C under the heating rate of 3 °C min−1, while it tremendously reduces to 220.2 °C for the 60MgH2/TiO2 heterostructure. According to the fitting results of DSC by Kissinger’s equation, the apparent activation energy of 60MgH2/TiO2 is determined to be 106.7 kJ mol−1 H2, much lower than that of the commercial MgH2 (142.27 kJ mol−1 H2).The hydrogen storage performances of the MgH reports –48, and 2 NS on MgH2, the isotherm de/ab-sorption measurements were carried out for the 60MgH2/TiO2 heterostructure as well as control samples . Figure 2/TiO2. As can been seen, the 60MgH2/TiO2 releases hydrogen of 0.51, 1.56, 2.14, 2.83, and 2.43 wt% at the temperature of 180, 200, 250, 275, and 300 °C within 10 min, respectively. Moreover, even at a lower temperature of 180 °C, the composite can still desorb 1.24 wt% of hydrogen within 60 min, corresponding to 41% of the maximum hydrogen storage capacity. At 250 °C, 2.14 wt% of hydrogen can be desorbed within 10 min, while no hydrogen desorption is detected for the blank MgH2 and commercial MgH2 under identical conditions . In addition, blank MgH2 needs 68 min to release 80% of its theoretical hydrogen capacity at 300 °C. In contrast, it takes only 1.8 min for the 60MgH2/TiO2 composite to reach the identical reaction extent  isothermal de/re-hydrogenation was tested at a relatively high temperature of 300 °C  and a capacity retention as high as 98.5% is achieved , which is associated with the formation of the catalytic oxide mentioned above. Moreover, the MgH2/TiO2 heterostructure still maintains flower-like structures and shows no obvious agglomeration of MgH2 and restacking of TiO2 NS even after 100 cycles . The high stability of the MgH2/TiO2 heterostructure is presumably ascribed to the confinement effect of 2D TiO2 NS on MgH2 and robust interfacial contact between host and support materials.Cycling stability is also a key factor for the practical application of hydrogen storage materials. To further evaluate the stability of the 60MgH2/TiO2 heterostructure exhibits rapid hydrogen ab/de-sorption kinetics and excellent long-term cycling stability than those of blank MgH2, highlighting the crucial role of the 2D TiO2 nanosheets, which simultaneously possess catalytic effect to MgH2 and inhibit the agglomeration of MgH2.Therefore, the MgH2/TiO2 heterostructure, we further characterized the phase components and microstructure of samples at different states. Figure 2 NS and MgBu2-treated TiO2 NS. As shown in Fig. 2-treated TiO2 NS exhibit in different colors ranging from blue to black, in marked contrast with the white pristine TiO2 powders. The color change of TiO2 is ascribed to the presence of oxygen vacancies, which coincides well with the previous work [2 NS by a facile strategy, that is, the impregnation of MgBu2 into TiO2 NS. We discover that the MgBu2 could not only act as the reducing agent to fabricate oxygen vacancy-rich blue TiO2 NS but also serve as the precursor of MgH2 to promote the self-assembly of MgH2 nanoparticles onto TiO2 NS through subsequent solvothermal approaches. Such a "one-stone-two-birds" effect enables the MgH2/TiO2 heterostructure as a promising hydrogen storage material with superior performances. Figure 2 samples, indicating that the impregnation of MgBu2 does not destroy the intrinsic crystal structure of the TiO2 nanosheets. Besides, XPS spectra of elemental Ti , oxygen vacancy (OV), and lattice oxygen (OL) in the MgBu2-treated TiO2 NS, respectively . Based on the peak area of OV to the overall peak area of O 1s, the percentage of OV on the surface of the MgBu2-treated TiO2 is determined to be 32.53%. Additionally, the concentration of oxygen vacancies on the surface of the as-synthesized MgH2/TiO2 heterostructure (33.4%) shows no significant change when compared to that on the TiO2 NS treated with MgBu2 .In order to further understand the mechanism of improved sorption performances of the MgHous work –41, 49. 2 NS treated with MgBu2. As shown in Fig. 2 NS treated with MgBu2 exhibit typical EPR signals at about g = 1.972 and g = 2.002, which could be attributed to Ti3+ and oxygen vacancies in the lattice, respectively [2 nanosheets by the "one-stone-two-birds" strategy. Due to the higher electrical conductivity, numerous active sites, and enlarged lattice space derived from oxygen vacancies, the formation of oxygen vacancy-rich blue TiO2 could not only accelerate the migration of electrons and diffusion of hydrogen atoms but also provide more active sites for the capture of hydrogen and nucleation of MgH2/Mg. Thus, the MgH2/TiO2 heterostructure displays superior hydrogen ab/de-sorption kinetic properties.Moreover, EPR measurement was carried out to further explore the existence of oxygen vacancies in TiOectively , 50. All2/TiO2 composites during the hydrogen de/ab-sorption process was further investigated. Figure 2/TiO2 heterostructure. Surprisingly, there are no significant differences in morphology among the as-synthesized  show no obvious change. Additionally, as shown in Fig. S8, the phase composition of the MgH2/TiO2 heterostructure remains unchanged after 100 de/re-hydrogenation cycles by showing very similar XRD patterns. Such a Mg-Ti oxide phase may possess a catalytic effect for the de/re-hydrogenation of MgH2 [2; hence, higher temperature is required for the release of hydrogen [2/TiO2 composites perform excellent hydrogen desorption properties. To further prove that these two peaks belong to the highly catalytic Mg-Ti oxide instead of less-catalytic MgO, 1H NMR test was conducted  compared with that of blank MgH2 (~ 4.50 s) indicates a faster H motion in MgH2/TiO2. In particular, the MgH2/TiO2 heterostructure exhibits a much shorter spin–lattice relaxation time of 1.49 s after the 1st re-hydrogenation than that of the as-synthesized MgH2/TiO2, indicating that new species with catalytic activity have formed during the first dehydrogenation process. Thus, it is demonstrated that two characteristic peaks centered at 42.7° and 62.0° are assigned to the ternary Mg-Ti oxide phase (MgTi2O4), having much higher catalytic activity than MgO. It can be deduced that the Mg-Ti oxide is formed by the chemical reaction at the interface between MgH2 and TiO2 NS during the first hydrogen desorption process. The similar phenomenon, for which Ti/Nb oxide could react with MgH2 and result in the formation of a new ternary oxide, has also been reported in the previous work [2, the broader diffraction peaks of the ternary Mg-Ti oxide phase manifest its smaller crystallite size than that of Mg/MgH2.We further explored the phase transformation and chemical valence state change of MgHgH2 Fig. a. No pea of MgH2 , 47. As hydrogen . Howeverous work . Compare2/TiO2 during the hydrogen desorption process  and TiO2 nanosheets could be identified according to the lattice fringe spacings. After irradiation under a beam current of 3 nA for 1.5 min, a portion of MgH2 react with TiO2, generating a thin Mg-Ti oxide layer. After 2.5 min of irradiation, the MgH2 near Mg-Ti oxide begins to release hydrogen to form Mg. The observed area releases hydrogen completely after 3 min. It can be seen that, unlike the neat MgO layer, the in-built ternary Mg-Ti oxide catalyst could act as a pathway to facilitate hydrogen diffusion, resulting in improved reaction kinetics. Through the pathways provided by Mg-Ti oxide, the hydrogen can diffuse into/out of Mg/MgH2, possibly by a formation of metastable titanium hydride (TiHx) species [2/TiO2 interfaces could also strongly anchor MgH2 nanoparticles, which significantly inhibits the growth and agglomeration of MgH2 particles during cycling. Moreover, the high ductility and flexibility of the Mg-Ti oxide [2 particles during de/re-hydrogenation cycles, preventing the detachment of MgH2 from the TiO2 substrate and thereby ensuring the excellent cyclic stability.To further clarify the role of the ternary Mg-Ti oxide in the dehydrogenation process, in situ HRTEM observations were carried out by real-time recording of the structure and composition evolution of hydrogenated MgHess Fig. c. Additi species , 52, whiTi oxide  can signp XPS spectra, the signals of Ti4+ species (458.8 eV for 2p3/2 and 464.5 eV for 2p1/2) can be detected in all samples at different states, indicating that TiO2 nanosheets still exist in the MgH2/TiO2 heterostructure during hydrogen ab/de-sorption process. Additionally, the typical doublets located at 456.8 and 462.5 eV of Ti3+ for as-synthesized MgH2/TiO2 composites further confirm the successful introduction of oxygen vacancies. Moreover, it can be seen that the Ti 2p signals of Ti2+ appear in the dehydrogenated samples and Ti2+ peaks shift to higher binding energies in the re-hydrogenated samples. The presence of Ti2+ peaks testifies the formation of low-valence Ti-H active species. The slight change of the peak position of Ti2+ demonstrates that Ti-H active species could act as the intermediates to promote the transportation of electrons and hydrogen [2+ and H−, providing a favorable condition for weakening the Mg-H bond. In addition, numerous multi-phase interfaces composed of MgH2 and multivalence Ti promote the diffusion of hydrogen and electron and provide more nucleation sites for MgH2/Mg. Thus, the existence of multivalence Ti-based catalysts accounts for the significantly enhanced hydrogen storage performances of the MgH2/TiO2 heterostructure.Moreover, XPS analyses were employed to explore the evolution of the bonding state of Ti in the process of de/re-hydrogenation and the results are shown in Fig. hydrogen , furtherhydrogen , 54. Var2/TiO2 heterostructure can be attributed to the intrinsic and extrinsic synergistic effects, as illustrated in Fig. 2 nanosheets with a high specific surface area provide the unimpeded channels for the rapid diffusion of hydrogen and more nucleation sites for MgH2/Mg, thereby improving the kinetics of de/re-hydrogenation. (2) Nano-sized MgH2 are uniformly confined on TiO2 nanosheets to form a flower-like heterostructure, which significantly suppresses the growth and agglomeration of MgH2 during the hydrogen ab/de-sorption processes, ensuring the excellent cycling stability. Meanwhile, the re-stacking of TiO2 nanosheets can be hampered in turn. (3) Abundant oxygen vacancies introduced by the "one-stone-two-birds" strategy significantly enhance the electrical conductivity of TiO2 NS and provide additional active sites for transportation of electrons and hydrogen, leading to greatly enhanced hydrogen sorption kinetics. (4) The ternary Mg-Ti oxide catalyst located at the interface of MgH2 and TiO2 acts as the pathway to facilitate hydrogen diffusion, resulting in improved hydrogen storage properties. In addition, Mg-Ti oxide possesses better ductility and flexibility than pure MgO. Thus, due to the stable interfacial bonding strength between oxide layers and the MgH2 nanoparticles, Mg-Ti oxide may exhibit a long-term confinement effect on MgH2 nanoparticles and significantly buffer successive expansion and contraction of MgH2 particles during de/re-hydrogenation cycles, improving reversible cycling performance. (5) Multi-valence Ti-based catalysts favor the electron transfer between Mg2+ and H−, weakening the Mg-H bond. Additionally, numerous multi-phase interfaces composed of Mg2+ and multi-valence Ti containing species could provide more diffusion pathways for hydrogen and more nucleation sites for MgH2/Mg.Based on the experimental results and analyses, it is concluded that the superior hydrogen storage performances of the MgH2/TiO2 heterostructure has been successfully synthesized using oxygen vacancy-rich 2D TiO2 NS as the functional scaffold. Benefitting from the intrinsic (nano-sized MgH2) and extrinsic  effects as well as the incorporation of oxygen vacancies, the prepared MgH2/TiO2 heterostructure exhibits superior hydrogen storage performances. Specifically, the composite shows excellent desorption performances with a low onset desorption temperature down to 180 °C and a rapid initial dehydrogenation rate at 300 °C (2.116 wt% min−1), surpassing those of the state-of-the-art MgH2 systems catalyzed by TiO2. Meanwhile, the MgH2/TiO2 heterostructure displays an excellent cycling stability with a capacity retention of 98.5% after 100 cycles at 300 °C. This work offers a promising approach for applying graphene analogues to the construction of advanced hydrogen storage materials with 0D/2D heterostructures.In the present work, a flower-like MgHSupplementary file1 (DOCX 4837 KB)Below is the link to the electronic supplementary material.


Furthermore, a flexible polyimide (PI) composite film with superior EMI shielding effectiveness is prepared by in situ polymerization. The 40% content of Gr-PANI10:1 (the mass ratio of Gr to PANI is 10 : 1) shows a superior electrical conductivity (σ) as high as 2.1 ± 0.1 S cm−1, 1.45 times higher than that of Gr@PI film at the same loading. Moreover, the total shielding effectiveness (SET) of EMI of the Gr-PANI10:1@PI reaches ∼21.3 dB and an extremely high specific shielding effectiveness value (SSE) of 4096.2 dB cm2 g−1 is achieved. Such a “non-covalent welding” approach provides a facile strategy to prepare high-performance PI-based materials for efficient EMI shielding.With the increasing demands of the electronics industry, electromagnetic interference (EMI) shielding has become a critical issue that severely restricts the application of devices. In this work, we have proposed a “non-covalent welding” method to fabricate graphene-polyaniline (Gr-PANI) composite fillers. The Gr sheets are welded with PANI  With the increasing demands of the electronics industry, electromagnetic interference (EMI) shielding has become a critical issue that severely restricts the application of devices. Recently, carbon materials, such as carbon nanotubes (CNT),6,7 electrical carbon black,8 carbon fibers,9 and graphene10 have been explored and mixed with polymers as conductive fillers to fabricate lightweight and high shielding effectiveness polymer-based materials.PI, as a kind of high-performance engineering polymer, has been widely used in electronic industries as a promising polymer matrix for EMI shielding because of its superior flexibility, excellent thermal stability as well as high wear resistance.11 Feng and coworkers fabricated rGO@PANI/PI films via in situ polymerization. And the dielectric properties and thermal stability increased significantly.12 Ma et al. prepared rGO–poly(ethylene glycol) (PEG)/PI with chemically reduced method.13 However, it is very difficult to uniformly disperse rGO sheets in organic solvent owing to van der Waals force and the π–π stacking interaction between rGO sheets. Meanwhile, the interaction between the rGO sheets and PI matrix is relatively weak, leading to the inferior mechanical properties of PI/rGO film. In this regard, the thermal reduction of GO has advantages in mild reaction conditions and less deterioration of GO compared to chemical reduction of GO. A lightweight porous PI/rGO film was obtained using non-solvent induced phase separation (NIPS) process.14 They found that the PI/rGO film exhibited electrical conductivity of 1.5 × 10−4 S cm−1, corresponding to a SSE value of 693 dB cm2 g−1. The GO reduced through thermal reduction still maintains some oxygen functional groups, contributing to construct strong interfacial interaction with PI matrix. Therefore, these rGO@PI films exhibit the better mechanical properties. Nevertheless, these composite films reduced at high temperature have poor electrical conductivity limiting the improvement of electrical property and EMI shielding performance. An effective way to improve the dispersibility of Gr is needed to enhance the performance for EMI shielding and electrical conductivity.Among these potential carbon-based fillers, graphene, due to its outstanding electrical conductivity, has become a promising candidate as electrical filler in polymer.via π–π non-covalent bonding. The Gr-PANI@PI film prepared through in situ polymerization, contributing to disperse Gr sheets homogenously in the PI matrix. The composite film displays superior mechanical properties (∼56.8 MPa), electrical conductivity (2.1 ± 0.1 S cm−1), EMI shielding performance (∼21.3 dB) and an extremely high SSE value of 4096.2 dB cm2 g−1 with a thickness of only 0.04 mm. The reported “non-covalent welding” conception is also an effective technique to prepare polymer materials with Gr sheets in other fields.In this work, we propose a “non-covalent welding” method to fabricate Gr-PANI filler to improve the electrical conductivity, mechanical and EMI shielding properties of the PI composite film. PANI, as a solder, not only successfully connects adjacent Gr sheets to provide the pathway for electron transportation, but also effectively prevents the stacking of Gr sheets 2.2.1.N-methylpyrrolidone (NMP), ferric chloride hexahydrate (FeCl3·6H2O) and ethanol were all purchased from National Pharmaceutical Group Chemical Co., Ltd.Exfoliated graphene power, pyromellitic dianhydride (PMDA) and 4,4′-diaminodiphenyl ether (ODA) are purchased from Aladdin . Hydrochloric acid (HCl), aniline (AN), 2.2.−1 HCl under ultrasonication. Then, the solution of FeCl3·6H2O  was rapidly transferred to the above solution containing AN and Gr. Next, the mixture was magnetically stirred for 4 h in an ice-water bath. After that, the Gr-PANI composite filler was obtained by filtration, washed with deionized water and methanol until the washing solution was completely colorless. In this process, free PANI particles and unreacted substance were separated and removed. Finally, the mixture was freeze-dried for 24 h. The as-obtained was named as Gr-PANI10:1(mass ratio of Gr to AN was 10 : 1). The Gr-PANI15:1 and Gr-PANI5:1 composite filler were prepared in the same manner.The preparation of Gr-PANI@PI composite films was illustrated in 2.3.10:1@PI composite film involved three steps. First, 1.25 g Gr-PANI10:1 was dispersed in anhydrous NMP (15.33 mL) by ultrasonication (1 h) at ambient temperature in order to homogeneously disperse Gr-PANI10:1 in the solution. Then, ODA  and PMDA  were added into the solution under a nitrogen atmosphere. After continuous stirring at 25 °C for 2 h, the homogeneous and viscous Gr-PANI10:1@PAA  solution was obtained. The mixture was bar-coated on a clean glass dish and treated at 140 °C for 1 h, and 350 °C for 2 h. The as-obtained composite film was coded as x%Gr-PANI10:1@PI, where x was the weight percentage of the Gr-PANI filler to the matrix. The 30%Gr-PANI10:1@PI obtained had a thickness of 40 μm. The 35%Gr-PANI10:1@PI and 40%Gr-PANI10:1@PI films were prepared in the same manner. In the same manner, we fabricated the x%Gr@PI (without PANI) composite film, where x represented the weight percentage of Gr filler to the matrix.The fabrication process of Gr-PANI2.4.−1 at a resolution of 2 cm−1 and for an accumulation of 8 times. Raman spectra  were collected from 200 to 3000 cm−1 with a laser length at 532 nm. X-ray photoelectron spectra  were performed to acquire the chemical bonding and the valence states of Gr and Gr-PANI with a monochromatized Al Kα X-ray source (1486.71 eV). X-ray diffraction  patterns were conducted using nickel-filtered Cu Kα (k = 0.154 nm) radiation with a generator voltage of 40 kV and a current of 40 mA. The scanning speed was 5° min−1 and the step size was 0.02° with the range from 10° to 60°. The surface and cross-section morphology of samples were observed by scanning electron microscopy  with an accelerating voltage of 20 kV. The elemental distribution mapping of Gr-PANI10:1 was measured by energy dispersive spectrometer . The electrical conductivity of films was measured by in-line four-point probe method  at room temperature, and three specimens with the thicknesses of 0.04 mm were used for each sample. The tensile strength and the elongation at break of the Gr-PANI@PI and Gr@PI films were measured by a universal stretching machine . Specimens were tested at a speed of 2 mm min−1 and a preload of 0.1 N. The size of each sample was 10 mm in length and 4 mm in width. Five specimens were used for each sample in the tensile test. The EMI shielding properties of the composite films were performed in the frequency range of 8.2–12.4 GHz (X band) using the waveguide method via the vector network analyzer (Agilent N5222B). All samples were cut into rectangle plates with a size of 22.8 × 10.1 mm2 to fit the waveguide sample holder. The scattering parameters (S11 and S21) of the Gr-PANI@PI composite films were collected to calculate the EMI shielding effectiveness. Thermogravimetric analysis  was performed to determine the thermal stability of films. Thermal gravimetric analysis (TGA) was conducted under the N2 atmosphere from 50 to 900 °C at a heating rate of 10 °C min−1.Fourier transform infrared spectroscopy  was collected on a spectrum from 400 to 4000 cm3.3.1.10:1@PI film with different loading of 30%, 35% and 40% are displayed in 10:1@PI films become darker compared to the transparent yellow PI film. Furthermore, the 40%Gr-PANI10:1@PI still remains excellent flexibility. It can be twisted and folded into a paper plane  in ESI.15 In our experiment, the PANI is in situ synthesized on the surface of Gr sheets to connect adjacent Gr sheets through π–π non-covalent bonding. When the Gr-PANI composites are added into the NMP, the PANI can effectively prevent Gr sheets from agglomerating.16,17 As displayed in Fig. S1(b) in ESI,10:1 composite filler shows excellent dispersibility due to the π–π non-covalent interactions between Gr and PANI, indicating that the PANI can avoid the agglomeration of Gr sheets in the organic solvent.18The optical photos of PI and Gr-PANIer plane . The siz10:1, PI and 40%Gr-PANI10:1@PI are shown in 19,20 The size of the PANI absorbed on the Gr sheets is smaller than that of pure PANI  in ESI21 Additionally, the EDS mapping of Gr-PANI10:1, as demonstrated in Fig. S2,10:1 filler. The atomic percentage of N reaches 22.19 at%, indicating the successful attachment of PANI in the filler  in ESI10:1@PI are further investigated, as shown in 10:1@PI exhibits a typical stake-up morphology, demonstrating that Gr sheets are uniformly distributed in the PI matrix and form coarser fractured surfaces compared with pure PI. It can be attributed to the excellent compatibility and interfacial interaction between the Gr-PANI10:1 composites and the PI matrix.22 The PANI acts as the solder to connect adjacent Gr sheets into an integrated conductive structure with less gaps, thereby improving the electrical performance.The SEM morphologies of Gr, Gr-PANI Table S1. While i10:1 filler. The typical adsorption peaks of PANI are located at 1572 cm−1, 1483 cm−1, 1307 cm−1, 1243 cm−1 and 1140 cm−1, which represent C PBM data was replaced with SVG by xgml2pxml:<glyph-data id="z.dbd" format="PBM" resolution="300" x-size="8" y-size="10" xml:space="preserve">00000000000000000000000000000000111111110000000011111111000000000000000000000000</glyph-data><svg xmlns="http://www.w3.org/2000/svg" version="1.0" width="13.200000pt" height="16.000000pt" viewBox="0 0 13.200000 16.000000" preserveAspectRatio="xMidYMid meet"><metadata>Created by potrace 1.16, written by Peter Selinger 2001-2019</metadata><g transform="translate scale" fill="currentColor" stroke="none"><path d="M0 440 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z M0 280 l0 -40 320 0 320 0 0 40 0 40 -320 0 -320 0 0 -40z"/></g></svg>C vibration in the quinoid ring, CC stretching in the benzene ring, C–N stretching in the secondary aromatic amines, and C–H bonding in the benzenoid ring and the quinoid ring,23 respectively. In the Gr-PANI10:1, these characteristic peaks are shifted to 1566 cm−1, 1476 cm−1, 1296 cm−1, 1232 cm−1, and 1086 cm−1, respectively, indicating that PANI is distributed intimately on the surface of the Gr sheets,24 due to the electrostatic interactions of π–π non-covalent bonding between the Gr and PANI.25 The peak shift can be also observed in Gr-PANI15:1 and Gr-PANI5:1  (ESIS4(a) ESI).10:1 is conducted, as displayed in −1) of the first-order scattering of E2g mode of sp2-hybridized carbon atoms and the D-band (∼1345 cm−1) of the disordered amorphous carbon C–C bonds at the defects or edge boundaries,26 respectively. After the in situ polymerization of PANI, the G-band and the D-band of the Gr-PANI10:1 are shifted to 1561.3 cm−1 (vs. 1574.1 cm−1 for Gr), and 1343.7 cm−1 (vs. 1345.0 cm−1 for Gr), respectively, suggesting the π–π interaction between PANI and Gr sheets.27 Moreover, the ID/IG values of Gr and Gr-PANI10:1 are 0.194 and 0.088, respectively. With the decreasing mass ratio of Gr to PANI, the decrease in ID/IG values can be also observed in Gr-PANI15:1 and Gr-PANI5:1   ESI). The in10:1@PI. The XRD pattern of pure PI shows a broad peak centered at 2θ = 18.5°.30,31 The sharp peaks centered at ∼26.50° in 40%Gr@PI and 40%Gr-PANI10:1@PI are attributed to the (002) of Gr. The 40%Gr@PI shows an obvious diffraction peak at 2θ = 26.54° corresponding to the d(002) of 0.3362 nm, revealing the incorporation of PANI enlarges the interlayer space of graphene.3010:1 composites under nitrogen atmosphere. The 15 wt% weight loss temperature of Gr-PANI10:1 shifts to higher temperature compared with pure PANI and Gr. The thermal stability of Gr-PANI10:1 is remarkably enhanced due to the presence of PANI and the π–π non-covalent between the Gr sheets and PANI.12 Meanwhile, TGA results demonstrate that the as-fabricated Gr-PANI10:1 composites contain 12.5 wt% of PANI by calculation. 10:1@PI at a heating rate of 10 °C min−1 in nitrogen atmosphere. The pure PI film shows an onset weight loss below 100 °C due to the vaporization of water.32 The main mass loss occurs at approximately 500 °C, which is attributed to thermal decomposition of PI.33 The 15% weight loss temperatures of pure PI, the 40%Gr@PI, and 40%Gr-PANI10:1@PI are 592 °C, 599 °C and 600 °C, respectively. Meanwhile, the residual weight at 700 °C for the 40%Gr-PANI10:1@PI is evidently improved from 66.4% (pure PI) to 72.7%, indicating that the Gr-PANI10:1@PI film shows superior thermal stability. This phenomenon is owing to the excellent thermal stability of Gr-PANI10:1, its introduction reduces the TGA rate of the PI film.34The thermal stability is a key property of PI-based composite. Fig. S510:1. The C 1s spectrum of Gr , 10.0 atom% (vs. 10.2 atom% for Gr), 2.0 atom% (vs. 2.3 atom% for Gr), suggesting the interaction between Gr and PANI is not covalent bonding. Furthermore, an additional peak at 285.9 eV (3.5 atom%) corresponding to the C–N bonding is observed, suggesting that PANI has been successfully in situ synthesized on Gr sheets. The core-level N 1s spectrum of pure PANI presents three primary peaks –N at 399.1 eV, –NH2 at 399.8 eV and –N+ at 401.0 eV in 37 The peak shift is also observed in Gr-PANI10:1 , 33.1 atom% (vs. 37.0 atom% for PANI), 33.4 atom% (vs. 31.0 atom% for PANI).The above results confirm the gap between the Gr sheets could be coupled together by PANI via the π–π non-covalent bonding, which can significantly enhance the dispersibility of Gr sheets.XPS spectroscopy is further characterized to analyze the elemental compositions and the valence states of Gr, PANI and Gr-PANIum of Gr  consistsPANI10:1 , and the3.2.σ) of each filler is measured by four-point probe method . The pure Gr has the highest σ of 618.1 ± 46.2 S cm−1 compared with Gr-PANI composites, as displayed in Fig. S6.σ decreases to 541.0 ± 16.9 S cm−1, 490.3 ± 39.0 S cm−1, 392.4 ± 44.5 S cm−1, corresponding to Gr-PANI15:1, Gr-PANI10:1 and Gr-PANI5:1, respectively. Such a decrease in σ is primary attributed to the factor that the electrical property of PANI is inferior to that of Gr. Furthermore, the σ of the 40%Gr@PI is only 1.4 ± 0.1 S cm−1 due to the inevitable agglomeration of Gr sheets in PI matrix. But with the increasing mass ratio of Gr to PANI from 15 : 1 to 10 : 1 at a filler loading of 40%, the σ increases to 1.8 ± 0.2 S cm−1 and 2.1 ± 0.1 S cm−1, but when the ratio is up to 5 : 1, the σ displays a dramatic decline to 1.3 ± 0.1 S cm−1. It is noticeable that the maximum σ reaches 2.1 ± 0.1 S cm−1 for 40%Gr-PANI10:1@PI, indicating the PANI welds the adjacent Gr sheets into conductive network so that it can increase the efficiency of electron transfer in PI film via the “non-covalent welding” strategy. However, with the further addition of the PANI (Gr-PANI5:1), there is an obvious decrease of σ in 40%Gr-PANI5:1@PI. Taking the result of SEM images into consideration, the reason lies in that the surface of Gr sheets is wholly covered by excessive PANI. The excessive PANI solder would lead to agglomeration between the Gr sheets, impeding the electron transfer.The electrical conductivity .σ improves from 0.4 ± 0.1 S cm−1 to 1.4 ± 0.1 S cm−1, respectively. In contrast, with the introduction of PANI, the σ of PI is improved further. When the filler content is at 40%, the σ of Gr-PANI10:1@PI is up to 2.1 ± 0.1 S cm−1, which is approximately 45% higher than that of Gr@PI film. The results demonstrate that the “non-covalent welding” strategy can successfully weld up the Gr sheets to improve the dispersibility of Gr and form conductive network in PI matrix, which can significantly enhance the σ of the composite film.We also make a comparison between the Gr@PI and the Gr-PANI3.3.10:1@PI are shown in Fig. S7 and S8 , the tensile strength of 40%Gr-PANI10:1@PI increases to 56.8 ± 1.4 MPa. The elongation at break of Gr@PI and Gr-PANI10:1@PI presents a slight decline with increasing filler contents 10:1@PI displays a slight increase in the elongation at break, which is about 1.29 times higher than that of 30%Gr@PI. The improvement in mechanical properties can be primarily ascribed to the homogeneous dispersibility of Gr-PANI10:1 in PI matrix because of the “non-covalent welding” strategy.39The mechanical properties and stress–strain curves of the Gr@PI and Gr-PANInd S8 ESI, respect3.4.T) of the Gr@PI and Gr-PANI10:1@PI is measured with vector network analyzer at the X-band . The toR), absorption (SEA) and multiple reflections (SEM), respectively. The EMI performance is attributed to the interaction between the electromagnetic waves and free electrons on the surface of materials. The material with high electrical conductivity can provide electric dipoles. When the incident electromagnetic waves interact with the free electrons on the surface and some electromagnetic waves are reflected from the surface of the shield. If electric dipoles interact with the electromagnetic fields in the radiation, the remaining electromagnetic waves would be absorbed by the dipoles.41–43 In addition, the thickness also affects the value of SEM and SEA. The SEM can be neglected when the value of SET is larger than 10 dB, therefore, SET can be simplified as the equation:44In general, the electromagnetic wave consists of three aspects: reflection (SES11 and S21) to calculate the EMI SE as shown in following equation:45In a vector network analyzer, we can record the parameters . The uniformly dispersed Gr sheets could attenuate electromagnetic radiation easily due to reflection and scattering microwave in the Gr-PANI10:1@PI film, and incident electromagnetic microwaves are transferred to heat by being absorbed. The Gr-PANI10:1@PI shows higher SET value compared with Gr@PI, which can be ascribed to the structure of conductive network by the “non-covalent welding” approach.The SE10:1@PI film. A portion of electromagnetic wave PI is reflected at the interface of the film, because conductive networks with a large amount of charge carriers are beneficial to EMI shielding effectiveness by reflection. The remaining portion penetrating the surface is multiply reflected or absorbed by the Gr layers.46 Only a small amount of electromagnetic waves pass through the film. The formation of an excellent conductive network between Gr and PANI in PI matrix ensures that the electromagnetic wave is greatly reduced, leading to an outstanding EMI shielding performance. On the other hand, the Gr-PANI@PI film possesses the continuously cross-linked network structure entirely consisted of Gr-PANI layers, which could certainly enhance internal multiple reflections more effectively than the agglomerated Gr layers in PI matrix. With the effect of the above two aspects, the EMI performance of Gr-PANI10:1@PI film has been obviously improved.47,48The high electrical conductivity and the formation of conductive networks within the PI matrix play an important role in the EMI shielding performance of the film. T, the density, and the thickness is an appropriate criterion to estimate the shielding performance of various EMI shielding materials. The SET value of reported EMI shielding films in the X-band region is listed in Table S2.10:1 has more excellent SET (21.3 dB) and higher SSE (4096.2 dB cm2 g−1) with the thickness only of 0.04 mm. This value is higher than that of the PS/graphene (∼17.3 dB) prepared by high-pressure compression molding49 and PU/MWCNT (∼20.0 dB) prepared ball mill method,50 respectively. Moreover, it is also superior to that of the WPU/CNT composites by a facile freeze-drying method by 72.6% (∼2143.0 dB cm2 g−1)51 and PI/rGO (17–21 dB)52 with a thickness of 0.8 mm. The survey suggests that Gr-PANI@PI10:1 film has the potential advantages to meet the requirement of the commercial application (more than 20 dB).Generally, the density and thickness of a material are also two significant factors for evaluating its EMI shielding performance. Hence, the SSE on the basis of the EMI SE3.5.in situ polymerization. Without the PANI, during stirring process, the Gr sheets tend to agglomerate due to Vander Waals force and π–π stacking interaction between Gr sheets resulting in an inhomogeneous and discrete network.53 In in situ synthesized on the surface of Gr sheets to weld up adjacent Gr sheets through π–π non-covalent bonding, which effectively avoids the agglomeration of Gr sheets in organic solvent. After the thermal imidization, Gr-PANI forms a continuous conductive network in polymer matrix, making the electron transfer more effective. In this case, the combination of Gr and PANI only depends on π–π non-covalent bonding, and the structure of Gr is not destroyed, ensuring the excellent conductivity of Gr.The mechanism of “non-covalent welding” between PANI and Gr is illustrated in 4.10:1@PI film with enhanced EMI shielding performance and electrical properties. The PANI, as a solder, can successfully weld up the graphene sheets via π–π non-covalent bonding and form a continuous conductive network in polymer matrix by in situ polymerization. The Gr-PANI10:1@PI exhibited a superior SET of 21.3 dB and SSE of 4096.2 dB cm2 g−1 at 40% content with a thickness of 40 μm compared to that of the Gr@PI (SET 17.1 dB). The σ of Gr-PANI10:1@PI was promoted to 2.1 ± 0.1 S cm−1 higher than Gr@PI (1.4 ± 0.1 S cm−1) at 40% content. Last but not least, we believe that such a “non-covalent welding” strategy shows a great potential to fabricate composites with inorganic fillers.In summary, we have reported a novel “non-covalent welding” approach to prepare Gr-PANIThere are no conflicts to declare.RA-010-C9RA08026K-s001