MXenes are crucial for electromagnetic interference (EMI) shielding, thermal camouflage, and other essential applications due to their exceptional mechanical, electrical, and thermal properties. However, the currently available synthetic methods for single-layer MXene nanosheets require a two-step etching and exfoliation process, leading to prolonged processing times and low yields, which hinder the applications of MXene nanosheets. We present a one-step process for preparing high-quality MXene nanosheets that combines graphite-assisted ball milling (GABM) and etching at room temperature. The MXene nanosheet yield reached 97?wt %, exceeding previously reported methods. Our innovative GABM method achieved synergistic exfoliation and etching in a single step by incorporating graphite, which enhanced shear forces and reduced the perpendicular impact energy that can cause damage. The MXene nanosheets exhibited a Young’s modulus of 0.37?±?0.03?TPa and an exceptional electrical conductivity. We demonstrated that the resulting high-quality MXene nanosheets could be readily assembled into fibers of >50?m long, as well as films and three-dimensional (3D) bulk assemblies. The resulting large-scale MXene films possessed excellent mechanical and electrical properties, and showed outstanding EMI shielding effectiveness, along with low mid-infrared emissivity. Our developed high-yield synthetic method for MXene nanosheets has established a viable pathway for scalable production, thereby promoting their commercial applications.
Significance Statement
Herein, we demonstrate a synergistic one-step process that integrates exfoliation and etching via graphite-assisted ball milling (GABM) with an etching agent, achieving record-high yields and high-efficiency production of single-layer MXene nanosheets. The synergistic mechanical exfoliation and chemical etching mechanism was thoroughly revealed by comprehensive experiments. The MXene nanosheets not only exerted high quality but also exhibited strong assembly capabilities. The resulting large-scale MXene films demonstrated high tensile strength and conductivity, showing promising potential for applications in electromagnetic shielding and infrared camouflage.
Introduction
MXenes, an emerging family of two-dimensional (2D) transition metal carbides and nitrides,1,2 hold significant promise for applications in flexible electrodes,3 wearable textile devices,1,4 electromagnetic interference (EMI) shielding,5 and infrared (IR) thermal camouflage,6,7 due to their excellent mechanical,8 electrical,9 and electromagnetic properties.5 To realize their full potential, mass production of MXenes nanosheets is essential. Thus far, the wet chemical etching method has been favored for the mass production of MXenes nanosheets, thanks to the scalability and availability of bulk MAX (M?=?early transition metals; A?=?groups IIIA or IVA elements; X?=?carbon or nitrogen) phase and the convenient dispersibility of MXenes nanosheets.1,2,10,11 However, the current process for preparing single-layer MXenes nanosheets via selective etching is a time-consuming two-step procedure that involves both etching and delamination.2,9–13 This results in prolonged processing times and a low yield of single-layer MXenes nanosheets ( Supporting Information Table S1). Therefore, a strong impetus exists to develop an efficient one-step process that allows for high-yield preparation of single-layer MXenes nanosheets at a low cost for industrial-scale applications. Ball milling can generate the necessary shear forces and has been successfully employed for the high-yield exfoliation of 2D nanomaterials such as graphene, h-BN, MoS2, and others.14–18 Attempts have been made to prepare MXenes using ball milling, yet none have proved satisfactory due to the absence of in-situ etching agents or size-matched grinding materials. This contributes to the reliance on multistep processes and a maximum yield of around 73%.19,20 Additionally, conventional ball milling is an aggressive method that produces strong perpendicular impact forces on the surfaces of layered materials, typically resulting in small-sized, thick nanosheets, which limits their applications.18,21
Experimental Methods
The Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene nanosheets were exfoliated using the GABM method. Comprehensive comparative experiments were conducted to reveal the synergistic mechanical exfoliation and chemical etching mechanisms of the GABM method. Detailed experimental parameters and methodologies are described in the Supporting Information in the Materials and Methods section and Supporting Information Movie S1. Various characterization methods were applied to analyze the morphology and structure of the prepared MXene nanosheets, including atomic force microscopy (AFM), scanning electron microscopy (SEM), transmission electron microscopy (TEM), high-angle annular dark-field scanning TEM (HAADF-STEM), energy-dispersive spectroscopy (EDS), X-ray diffraction (XRD), X-ray photoelectron spectroscopy (XPS), Raman spectroscopy, and Ultraviolet–visible (UV–vis) spectroscopy. The mechanical and electrical properties of single-layer MXene nanosheets were conducted using AFM nanoindentation and field-effect transistor (FET) devices, as detailed in the Supporting Information. Subsequently, the MXene nanosheets were assembled into scalable fibers, films, and lamellar scaffold aerogels through wet spinning, layer-by-layer blade coating, and bidirectional ice template assembly techniques. The EMI shielding and IR camouflage performance of the MXene films were tested using a vector network analyzer and an FTIR spectrometer equipped with an IR integrating sphere. All detailed experimental parameters and characterization methods are included in the Supporting Information.
Results and Discussion
Production of MXene nanosheets
We report a high-yield, high-efficiency, scalable production method for synthesizing
Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene nanosheets through GABM in the presence of an etchant that facilitated both
exfoliation and etching in a single step. This preparation process, utilizing the
GABM strategy, is illustrated in Figure?1a. Taking Ti3C2Tx MXene as the first example, Ti3AlC2 MAX phase powder ( Supporting Information Figure S1d), graphite flakes ( Supporting Information Figure S2), zirconia balls ( Supporting Information Figure S3), and an etching agent (LiF/HCl) solution were mixed in an optimized ratio before
conducting vacuum ball milling at room temperature. The GABM process provided both
shear forces and chemical etching. For the Ti3AlC2 MAX phase, the bond energy of Ti–Al is lower than that of Ti–C, which aided in breaking
the Ti–Al bonds under the combined mechanical and chemical forces.22,23 In situ-generated hydrofluoric acid (HF) molecules from the etching agent (LiF/HCl
solution) penetrated the interlayer of the Ti3AlC2 MAX phase from the surfaces and etched the Al atomic layer in a layer-by-layer process,
resulting in MXene layers with reduced interlayer interactions (van der Waals forces
and hydrogen bonds).24,25 Concurrently, the shear forces generated by GABM continually exfoliated the etched
MXene layers into nanosheets as soon as they were formed through the collaborative
process of etching and exfoliation in one step. The introduction of graphite as a
grinding aid was crucial in enhancing the exfoliation of larger MXene nanosheets,
as discussed in detail in the section covering the mechanism and process for controlled
synthesis of MXene nanosheets. This synergistic process of etching and exfoliation
using the GABM strategy not only enhanced the exfoliation yield of MXene nanosheets
but also accelerated the etching process, as the timely exfoliation of the etched
MXene layers promoted the release of etching byproducts (AlF3, H2) and their exchange with HF molecules.20,26 This method effectively produced various MXene nanosheets, including Ti2CTx, Ti3C2Tx, and Ti3CNTx. Additionally, to assess the efficiency and quality of MXene nanosheets prepared
using the GABM strategy, we also synthesized Ti3C2Tx MXene nanosheets using current mainstream methods, including HF etching combined
with LiCl intercalation (HF/LiCl)27 and minimally intensive layer delamination (MILD) methods.5,28 Figure 1 | Preparation of single-layer MXene nanosheets using the GABM strategy. (a)?Schematic
illustration of the one-step exfoliation and etching process of MXene nanosheets.
AFM image (b)?and height profiles (c)?of Ti3C2Tx MXene nanosheets. TEM image (d)?and corresponding SAED diffraction pattern (e)?of
Ti3C2Tx MXene nanosheets. The SAED patterns were recorded in the region marked by a red circle
in the TEM image. (f)?Preparation time versus yields of MXene nanosheets created by
different methods, including HF etching, LiF and HCl etching, electrochemical etching,
and the current GABM strategy. (g)?Photographs of over twenty liters of produced MXene
nanosheet dispersions (～ 4?mg mL?1). (h)?Photographs of Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene films (from left to right).
AFM images demonstrated the successful exfoliation of single-layer MXene nanosheets with thicknesses of 1.1, 1.6, and 1.8?nm for Ti2CTx, Ti3C2Tx, and Ti3CNTx, respectively (Figure?1b,c, and Supporting Information Figure S4). The thickness of the prepared MXene nanosheets exceeded the theoretical value due to the presence of introduced surface groups.8 TEM and high-resolution TEM (HRTEM) images indicated that the MXene nanosheets had clean surfaces and precise edges, and exhibited a clear lattice fringe structure (Figure?1d and Supporting Information Figure S5). The corresponding selected area electron diffraction (SAED) demonstrated that the MXene nanosheets maintained an intact specific space group of primitive, six-fold screw axis/mixed metal carbides (P63/MMC) hexagonal stacking lattice structure and exhibit high-quality single-crystal periodicity (Figure?1e).22,29 The d-spacings of exfoliated Ti3C2Tx MXene nanosheets were 0.268?nm for the (100) plane and 0.155?nm for the (110) plane, as previously reported.30 As measured by SEM, the average maximum lateral sizes of the MXene nanosheets are 2.7, 3.1, and 2.5?μm for Ti2CTx, Ti3C2Tx, and Ti3CNTx, respectively ( Supporting Information Figures S6–S8). These findings illustrated the capability of the GABM strategy to overcome the limitations of conventional ball milling, which significantly degraded the sheet area in the preparation of single-layer nanosheets.15,18,21 The average maximum lateral sizes of Ti3C2Tx MXene nanosheets synthesized using HF/LiCl and MILD methods were 7.2 and 6.0?μm, respectively ( Supporting Information Figures S9 and S10). Taking Ti3C2Tx MXene as an example, the thickness distribution of 300 MXene nanosheets obtained by the GABM method showed 92% as the single-layer, which was 0.99 and 1.30 times that of the HF/LiCl and MILD methods ( Supporting Information Figure S11), and higher than previously reported.9,31
By optimizing the mass ratio of the MAX phase and graphite used in the exfoliation process, as well as the ball milling time, the GABM method achieved high yields of 94.5% for Ti2CTx, 95.8% for Ti3C2Tx, and 97.2% for Ti3CNTx with preparation times of ～ 12, 24, and 12?h ( Supporting Information Table S2). For the same preparation time of 24?h, the yields of Ti3C2Tx MXene nanosheets increased by 127.6% and 140.7% compared to the HF/LiCl and MILD methods, respectively ( Supporting Information Table S1). Figure?1f compares the preparation times and yields of the one-step GABM process for producing single-layer MXene nanosheets with those of previously reported methods, which required two separate preparation steps—one for etching and another for exfoliation ( Supporting Information Table S1). The GABM method boasts the highest yield, shortest total preparation time, and good scalability among all methods. Utilizing a process that produced >10?grams of MXene nanosheets per batch, we created the bottles depicted, which contained exfoliated MXene nanosheet dispersions with concentrations of ～ 4?mg mL?1 (Figure?1g). Free-standing films of Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene nanosheets were assembled from these colloidal dispersions by vacuum filtration (Figure?1h), demonstrating that the exfoliated MXene nanosheets prepared using our GABM process have good dispersibility and solution processing properties.
Characterization and performance of MXene nanosheets
The structures and surface states of the prepared MXene nanosheets were characterized using microscopy and spectroscopy. Depending on the MXene composition, the colloidal solutions varied in perceived color, and the corresponding free-standing films displayed complementary colors ( Supporting Information Figure S12a). For instance, Ti2CTx MXene nanosheets exhibited a dark purple color in dilute solution and a green film color, while the Ti3C2Tx MXene solution appeared forest green and the film was dark purple. The Ti3CNTx MXene solution was navy blue, with the film appearing dark gray. UV–vis absorption and reflection spectra were employed to quantify the observable differences between MXene dispersions and free-standing films ( Supporting Information Figure S12b,c). The Ti3C2Tx MXene solution showed an absorption peak at 766?nm. Reducing the number of carbon atoms in the formula by one to create Ti2CTx MXene resulted in a qualitatively similar spectrum. However, the significant absorption peak shifted to higher energy (498?nm), as previously reported.1,32 The Ti3CNTx MXene showed no peaks in the visible region.
The atomic-resolution 2D structure of the prepared MXene nanosheets was observed using
HAADF-STEM (Figure?2a–c). As shown in Figure?2b, the atomic-resolution HAADF-STEM image, taken along the vertical direction of the
Ti3C2Tx MXene atomic plane, revealed a perfectly hexagonal symmetric lattice of Ti atoms.
This image aligned precisely with the projected atomic structure model of the Ti3C2Tx MXene monolayer.9,33 Hexagonal reciprocal lattices of the [100] crystal orientation depicted in the fast
Fourier transform (FFT) images further confirmed their high-quality single-crystal
nature (Figure?2d). The integrated pixel intensity image displayed the dominant (100) crystal plane
of Ti3C2Tx MXene with a lattice spacing of 0.26?nm, consistent with the SAED results (Figure?2e). XRD patterns (Figure?2f) showed that the Al (104) diffraction peak was absent in any of the prepared MXene
films, and the (002) diffraction peak had shifted to smaller angles, indicating successful
exfoliation of the MXene nanosheets.10,28 Furthermore, there was no graphite (002) diffraction peak at 26.6° present in any
of the MXene films, providing strong evidence for the complete separation of graphite.
Figure 2 | Characterization of as-prepared MXene nanosheets. TEM image (a), atomic-resolution
HAADF-STEM image (b), denoised HAADF-STEM image (c), and corresponding FFT patterns
(d)?of a single Ti3C2Tx MXene nanosheet on the hollow carbon film substrates. (e)?Intensity line profile
of the (100) crystal plane of the Ti3C2Tx MXene nanosheet acquired along the white square in (b). (f)?XRD spectra of graphite,
MAX phase, and the as-prepared MXenes. (g)?EDS characterization of a single Ti3C2Tx MXene nanosheet, including HAADF-STEM image and elemental maps of Ti, C, O, F, and
Cl. (h)?XPS spectra of the as-prepared MXenes. (i)?High-resolution XPS spectra of
Ti 2p of Ti3C2Tx MXene. (j)?A plot of the molarities of surface termination groups for the as-prepared
MXenes.
The elemental composition of the prepared MXene nanosheets was characterized using the HAADF-STEM mapping technique (Figure?2g, Supporting Information Figures S13 and S14). EDS maps confirmed the uniform distribution of the desired elements in Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene nanosheets. In addition to the required compositional elements (Ti, C, N), O, F, and Cl were also detected, indicating that the surface of the MXene nanosheets could be grafted with functional groups.22 Supporting Information Table S3 details the element content obtained from EDS analysis for the sheets fabricated via the GABM process. XPS differentiates the elemental compositions and surface bonding features of our prepared MXene nanosheets from those of the corresponding precursor MAX phases (Figure?2h and Supporting Information Figures S15–S18). High-resolution XPS spectra of Ti 2p, Al 2p, C 1s, O 1s, F 1s, and N 1s for Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene are displayed in Figure?2i and Supporting Information Figures S16–S18. While the precursor MAX phases contained a high concentration of Al, the XPS spectra of all our produced MXene films showed no Al 2p signal, confirming the successful exfoliation of MXene nanosheets.34 The presence of C–Ti–O, C–Ti–OH, and C–Ti–F bonds, confirmed through peak fitting of the Ti 2p, C 1s, O 1s, and F 1s high-resolution XPS spectra, indicates an abundance of hydrophilic groups (=O, –OH, –F) grafted onto the surfaces of the prepared MXene nanosheets, resulting from the removal of Al layer atoms and the exposure of unsaturated transition metal surfaces.22,35 To evaluate and quantify the surface groups of the MXene nanosheets, we combined EDS and XPS characterization to obtain the different element compositions, their proportions, and chemical states ( Supporting Information Figures S16–S18 and Tables S3–S5). This analysis yielded the chemical formulas for the three types of MXene nanosheets: Ti2CO0.36(OH)0.60F0.68, Ti3C2O0.52(OH)0.49F0.47, and Ti3CNO0.58(OH)0.29F0.55.35,36 Figure?2j illustrated the molarity of the surface termination groups (Tx) for the MXene prepared using our GABM method.
To assess the quality of the prepared MXene nanosheets, using Ti3C2Tx MXene as an example, we examined the atomic defects in several MXene nanosheets through
atomic-resolution STEM characterization (Figure?3a–c and Supporting Information Figure S19). Compared to direct HF etching, MXene nanosheets produced by the MILD and GABM methods,
which generated HF molecules from LiF/HCl in situ, showed fewer atomic vacancy defects.
The statistical results of the ratio of atomic vacancies to all atomic sites from
ten STEM images are summarized in Figure?3d. The defect concentrations of MXene nanosheets prepared via the HF/LiCl, MILD, and
GABM methods were 2.57?±?0.21%, 0.82?±?0.12%, and 0.76?±?0.11%, respectively. The
defects in MXene nanosheets primarily arose from the etching and removal of surface
Ti atoms by HF during the etching process,33 as well as potential oxidation that might result from extended etching, the delamination
process, and heating during preparation.37 These findings suggested that the in-situ exfoliation one-step process of the GABM
strategy at room temperature did not result in a significant number of atomic defects
in MXene nanosheets.
Figure 3 | Comparison of quality characterizations and properties of MXene nanosheets exfoliated
using HF/LiCl, MILD, and GABM methods. Atomic-resolution HAADF-STEM images of MXene
nanosheets exfoliated by HF/LiCl (a), MILD (b), and GABM (c). Red circles indicate
atomic vacancies. (d)?Comparison of defect concentration based on STEM images obtained
from MXene nanosheets produced via different synthesis methods. Raman intensity mappings
of various vibration peaks in the lattice skeleton region of a single-layer MXene
nanosheet prepared using HF/LiCl (e), MILD (f), and GABM (g)?methods. (h)?The corresponding
Raman spectra at strong and weak signal points were collected at the white arrow (e–g).
(i)?A typical force-deflection curve of a single-layer MXene nanosheet exfoliated
by GABM during AFM nanoindentation. The top inset features an AFM image of a suspended
single-layer MXene nanosheet before the indentation test, as well as the height profile
(red line) along the white line. The bottom inset offers a detailed view of the same
curves, highlighting the center of origin. (j)?IDS–VDS curves for the FET device with MXene nanosheets exfoliated by GABM at gate voltages
of 40, 0, and ?40?V. The inset presents a schematic of an MXene-based FET device.
(k)?Comparison of Young’s modulus and electrical conductivity of MXene monolayers
exfoliated via HF/LiCl, MILD, and GABM methods.
We utilized UV–vis absorption and Raman spectroscopy to characterize the MXene nanosheets ( Supporting Information Figure S20). The transverse surface plasmon resonance peaks of MXene nanosheets prepared through the HF/LiCl, MILD, and GABM methods were 793, 772, and 766?nm, respectively ( Supporting Information Figure S20a). The differences in peak positions of MXene nanosheets obtained by various methods related to size effects and discrepancies in the surface groups.32,38 Using a 785?nm laser for resonance excitation, corresponding with the plasmonic peak, we compared the Raman vibration peaks of MXene nanosheets prepared by the three methods ( Supporting Information Figure S20b), which included the A2g (Ti, C, O) vibration peak, A1g (Ti, C, O) skeleton vibration for all lattice atoms, Eg (Ti, C, O) modes for surface groups, and A1g (C) phonon modes for lattice carbon.39,40 The vibration peaks of the skeleton and lattice carbon in MXene nanosheets exfoliated by HF/LiCl revealed a blue shift. The A2g (Ti, C, O) vibration peak shifted to a higher wave number (124?cm?1), the A1g (Ti, C, O) peak appeared at 202?cm?1, and the A1g (C) peak moved to 728?cm?1, demonstrating low intensity for all peaks, potentially due to the influence of defects.39,40 In contrast, MXene nanosheets exfoliated by the MILD and GABM processes displayed similar spectral features, with the A2g (Ti, C, O), A1g (Ti, C, O), and A1g (C) peaks located at 122, 200, and 723?cm?1, respectively, as previously reported.40 Moreover, the sharper and more intense A2g (Ti, C, O) and A1g (Ti, C, O) peaks indicated high lattice quality in MXene nanosheets.39 The surface groups of Eg (Ti, C, O) peaks of MXene nanosheets exfoliated by all three methods showed similar peak positions and shapes. By comparing the full Raman spectra of graphite flakes and Si/SiO2 substrates, the MXene nanosheets exfoliated by GABM displayed no residual carbon, with characteristic vibration peaks free from substrate interference ( Supporting Information Figure S20c). We employed Raman mapping technology to visualize the intensity of vibration peaks in the lattice skeleton regions (122?cm?1 in red and 200?cm?1 in blue) of individual MXene nanosheets (Figure?3e–g and Supporting Information Figure S20d–f). The MXene nanosheets exfoliated by HF/LiCl present a heterogeneous signal intensity with weak intensity in in-plane Raman signals (Figure?3e). The consistency of signal intensity and clear edges confirmed the uniformity of the lattice and stress in MXene nanosheets prepared by the MILD and GABM methods, demonstrating high in-plane homogeneity and quality (Figure?3f,g). Figure?3h compares the Raman spectra of the strong and weak signal points of the MXene nanosheets in the mapping images. Compared to the strong signal positions, the weak signal points of the MXene nanosheets exfoliated by HF/LiCl demonstrate low intensity and a blue shift in the A2g (Ti, C, O) and A1g (Ti, C, O) peaks, suggesting that lattice vibration was restricted, potentially due to increased defects in that region of the MXene nanosheets.39,40 However, the MXene nanosheets exfoliated by the GABM process exhibited good intensity and peak position retention, indicating high quality.
We applied AFM nanoindentation to measure the mechanical properties of single-layer MXene nanosheets prepared using three different methods ( Supporting Information Figure S21). By fitting the force-deflection curves obtained from nanoindentation, we could determine the 2D elastic modulus (E2D) of the MXene nanosheets. Based on the single-layer MXene nanosheet thickness of 0.98?nm,41,42 the effective Young’s modulus of the single-layer MXene nanosheet prepared by the GABM process was 367.0?±?2.6?GPa (Figure?3i and Supporting Information Figure S21e,f). The MXene nanosheets prepared by the MILD method (345.9?±?15.4?GPa) showed a comparable effective Young’s modulus as previously reported ( Supporting Information Figure S21c,d),8 confirming the reliability of our method. However, due to the high concentration of defects, the MXene nanosheets prepared by the HF/LiCl method exhibited the lowest Young’s modulus (336.5?±?12.0?GPa) ( Supporting Information Figure S21a,b). We also employed a mask-assisted evaporation method to fabricate back-gate FET devices with a single-layer MXene channel on SiO2/Si substrates (Figure?3j inset and Supporting Information Figure S22a–c). The IDS–VDS curves of all MXene devices exhibited a linear trend and nearly coincided at three different gate voltages (?40, 0, and 40?V) (Figure?3j and Supporting Information Figure S22d–f), indicating stable ohmic contact between the MXene channel and gold electrodes.43 The sheet resistances of the devices were calculated based on the IDS–VDS curves measured at VG?=?0 and the dimensions of the MXene nanosheets ( Supporting Information Table S6). The average sheet resistances of MXene devices fabricated by the HF/LiCl, MILD, and GABM methods were 941.23?±?71.93, 507.02?±?21.95, and 509.77?±?19.36?Ω?sq?1, respectively. Consequently, the electrical conductivity of MXene nanosheets exfoliated by HF/LiCl was 10,885.54?±?870.31?S cm?1, comparable to previous reports.43 Due to a more complete crystal structure, the MXene nanosheets prepared by the MILD and GABM processes exhibited higher electrical conductivity of 20,150.19?±?851.34 and 20,036.09?±?753.28?S cm?1, respectively. Compared to the HF/LiCl and MILD methods, the MXene nanosheets exfoliated by the GABM method displayed a favorable Young’s modulus and electrical conductivity (Figure?3k), attributed to their high quality.
Mechanism and process for the controlled synthesis of MXene nanosheets
Under the same experimental conditions, the GABM method was employed for ball-milling graphite or MAX?+?graphite. Supporting Information Figure S23 illustrates that the dispersions produced from ball milling graphite alone yield a clear, transparent solution, in contrast to the dark black–green dispersions obtained from ball milling the Ti3AlC2 MAX phase, which indicated the presence of MXene nanosheets in the dispersions. This suggested that the graphite grinding aid did not exfoliate, thus contaminating the liquid. A comparison of the morphology and maximum lateral size distribution of graphite before and after ball milling showed negligible changes ( Supporting Information Figures S2 and S24). Studies on liquid-phase exfoliation have demonstrated that achieving graphene exfoliation from graphite flakes in an aqueous solution is challenging.44,45 Graphite has the highest in-plane fracture strength (130?GPa for single-layer graphene),46 making it resistant to in-plane fragmentation during the GABM process. Additionally, graphite exhibited excellent chemical stability and corrosion resistance against acids and bases, making it an ideal choice for a long-term stable and recyclable grinding aid.
By substituting HF/HCl for LiF/HCl in the GABM process (HF/HCl-GABM), the MAX phase was ball-milled for 12?h and successfully exfoliated into numerous thin flakes ( Supporting Information Figure S25a,b). AFM characterization confirmed that a single-layer (1.5?nm) and a few layers (2.9?nm) of MXene nanosheets were produced through synergistic HF etching and shear force exfoliation ( Supporting Information Figure S25c–e), eliminating the need for LiCl intercalation steps.11,27 The resulting MXene nanosheets displayed a typical forest green color in dilute solution ( Supporting Information Figure S25f,g). XRD analysis combined with AFM validated that the HF/HCl-GABM directly converted the MAX phase into single-layer MXene nanosheets in one step ( Supporting Information Figure S25h). However, its yield (10.25%) was notably lower than that of the LiF/HCl-GABM (86.8%), likely due to Li+ intercalation, which expanded MXene interlayer spacing and weakened the interlayer van der Waals forces.2,12 We also assessed the use of HCl-GABM (without HF) for MAX phase treatment. The HCl-GABM process still reduced the MAX phase thickness ( Supporting Information Figure S26a) through shear force exfoliation, demonstrating the considerable peeling effect of the GABM process, yet it failed to exfoliate MXene nanosheets without HF etching. Furthermore, we compared the effects of Na+ and K+ as alternatives to Li+ as intercalants in the GABM process ( Supporting Information Figure S26b–d), where the emergence of a new low-angle (002) diffraction peak confirmed the presence of MXene. However, the exfoliation of MXene nanosheets remained unsatisfactory due to the mismatch between the ion radius and the intercalation energy of Na+ or K+ and the MXene interlayer.47 These findings indicated that in-situ HF etching and Li+ intercalation were essential for the efficient exfoliation of single-layer MXene nanosheets by the GABM process.
Ball milling in the presence of graphite during the GABM process was crucial in this
work for ensuring the high exfoliation efficiency of MXene nanosheets, defined as
the yield of single-layer MXene nanosheets. The exfoliation efficiency was calculated
by multiplying the yield by the percentage of single layers, determined through numerous
AFM micrographs that measured the thickness of individual MXene nanosheets. Comparative
experiments conducted under room temperature stirring in an identical etching agent
(LiF/HCl) solution showed that the rapid exfoliation and accelerated etching advantages
of the GABM method could not be achieved by merely stirring. SEM images indicated
that significant etching and exfoliation of the Ti3AlC2 MAX phase did not happen during 12 and 24?h of room-temperature stirring-assisted
etching ( Supporting Information Figure S27), with the yield of collected MXene nanosheets being only 2.0% and 12.4%, respectively.
In contrast, the MAX phase produced using the GABM method displayed significant exfoliation
and etching (Figure?4a and Supporting Information Figure S28). A notable reduction in the thickness of the Ti3AlC2 MAX phase was observed as the ball milling time increased. AFM characterization demonstrated
that the GABM process directly exfoliated single-layer MXene nanosheets (1.6?nm) before
washing and vibration steps ( Supporting Information Figure S29). This illustrated that the GABM process facilitated efficient in-situ exfoliation
of MXene nanosheets from the MAX phase in a single step through synergistic mechanical
exfoliation and chemical etching. The exfoliation efficiency reached 80.0% and 88.1%
when the ball milling time was 12 and 24?h, respectively ( Supporting Information Figure S11 and Table S7). This sharply contrasted with the reported accordion-like multilayer MXene obtained
by the two-step preparation method.10,12,26,34 As shown in Figure?4b,c and Supporting Information Figure S30, a typical accordion-like multilayer MXene was obtained by stirring-assisted etching
with HF/HCl and LiF/HCl for 24?h at 35?°C, but incomplete etching of multilayer MXene
was noted with LiF/HCl ( Supporting Information Figures S30b and S31). Importantly, thin and single-layer MXene nanosheets were produced using GABM after
12- and 24-h ball milling, bypassing the accordion-like MXene intermediate and post-intercalation
delamination steps, demonstrating GABM’s capability to achieve one-step MXene exfoliation
( Supporting Information Figures S29, S30c, and S31). Compared to GABM, the exfoliation efficiency of HF/LiCl and MILD methods was only
39.2% and 28.3% after 24?h, indicating a low degree of complete exfoliation ( Supporting Information Table S7).
Figure 4 | Measurement results of the MXene nanosheets exfoliated using the GABM technique. (a)?The
morphology of the sample was acquired at different stages during the GABM process.
SEM images showcasing accordion-like MXene were obtained through HF/HCl (b)?and LiF/HCl
(c)?etching for 24?h, respectively. (d)?A schematic illustration depicting the exfoliation
mechanism based on ball-to-wall impact models during milling with and without graphite.
(e)?Yields and maximum lateral sizes of MXene nanosheets exfoliated using various
mass ratios of the MAX phase and graphite, with the ball milling duration fixed at
12?h. (f)?Yields and maximum lateral sizes of MXene nanosheets exfoliated over different
ball milling durations, maintaining a constant MAX to graphite mass ratio of 1:1.5.
(g)?A radial plot comparing the yield, percentage of single-layer nanosheets, efficiency
of single-layer production, Young’s modulus, and electrical conductivity of MXene
nanosheets exfoliated using the HF/LiCl, MILD, and GABM methods.
Supporting Information Table S8 presents the washing times and pH changes of MXene nanosheets prepared using HF/LiCl, MILD, and GABM methods. Our one-step exfoliation and etching process using the GABM method also facilitated the rapid washing of interlayer ions (Li+, H+) and byproducts (AlF3), thereby increasing the preparation efficiency of single-layer MXene by a factor of 2–4 compared to conventional two-step methods ( Supporting Information Tables S1 and S8). Since the accordion-like multilayer MXene required a more extensive washing and exfoliation process, the reported two-step preparation strategies necessitated longer processing times and incurred higher costs ( Supporting Information Tables S1 and S8). Additionally, in the absence of graphite during ball milling, the yield of MXene nanosheets was only 59.3% ( Supporting Information Table S2), and the average maximum lateral size was quite small, ～ 0.9?μm ( Supporting Information Figure S32a). Using the GABM method, the resulting exfoliated MXene nanosheets exhibited an average maximum lateral size of 3.0?μm and a yield of 86.8% ( Supporting Information Figure S32b and Table S2), indicating that graphite played a buffering and shearing role. Compared to the two-step processes using HF/LiCl and MILD, the GABM method achieved synergistic exfoliation and etching in a single step, resulting in a high yield of large-sized MXene nanosheets (Figure?4a–c).
A feasible exfoliation mechanism for the high-yield exfoliation of large-sized MXene nanosheets using the GABM method was proposed (Figure?4d and Supporting Information Figure S33). The Ti–Al bonds broke in the MAX phase (e.g., Ti2AlC, Ti3AlC2, and Ti3AlCN) under mechanical and chemical forces because the bond energy of Ti–Al was weaker than that of Ti–C.22,23 During the GABM process, MXene layers with weak interlayer interactions (van der Waals forces and hydrogen bonds) were produced on the MAX phase surface through in-situ HF molecular etching from the LiF/HCl solution and subsequent Li+ intercalation.24,48 Etching starts from the surface of the MAX phase and proceeds layer by layer.25 Simultaneously, the zirconia balls generated strong impact and shear forces during the ball milling process. Horizontal shear forces facilitated in-situ exfoliation of MXene layers, while vertical impact forces caused material pulverization, reducing platelet area.21 The zirconia balls possessed a Mohs hardness of ～ 8.5, whereas the Mohs hardness of the Ti3AlC2 MAX phase was ～ 5.28.49 Thus, the intense vertical impact on the Ti3AlC2 MAX phase could result in in-plane pulverization, yielding smaller lateral sizes of MXene nanosheets. Figure?4d and Supporting Information Figure S33 illustrate the peeling and crushing mechanisms of MAX via the impact/shear forces from ball-to-wall and ball-to-ball impact models in GABM, both with and without graphite. The Mohs hardness of graphite flakes was only 1 to 2. Introducing graphite between the hard zirconia balls and MAX is highly effective in absorbing and deflecting the vertical impact energy transmitted by the zirconia balls onto the MAX surface during the ball milling process,50 thus reducing the likelihood of in-plane fragmentation of the MAX phase and increasing the size of the exfoliated MXene nanosheets. Additionally, a significant size mismatch existed between the zirconia balls (3?mm) and MAX (19?μm for Ti2AlC, 28?μm for Ti3AlC2, and 23?μm for Ti3AlCN), which diminished the effectiveness of shear force transfer between the zirconia balls and the MAX ( Supporting Information Figures S1 and S3). The closely matched size (444?μm) of graphite flakes to the MAX phase enhanced the transfer of shear forces from the zirconia balls to the MAX ( Supporting Information Figure S2), promoting the continual exfoliation of MXene nanosheets from MAX by shear forces.51 Therefore, the ternary system formed by graphite, zirconia balls, and the etching and intercalation agent (LiF/HCl) creates a synergistic effect of chemical etching and mechanical shear forces generated during the GABM process, achieving high-yield exfoliation of large-sized and high-quality MXene nanosheets.
Using Ti3C2Tx MXene as an example, the yield of exfoliated MXene nanosheets was significantly improved by adding graphite (Figure?4e and Supporting Information Table S2). When the mass ratio of MAX to graphite was 1.5, the overall process yield of MXene nanosheets was highest at 86.8%, a substantial increase compared to the yield before the introduction of graphite. As the mass ratio of graphite was further increased, the yield decreased because excess graphite hindered the movement of zirconia balls and reduced the shear energy, which had been previously observed during the ball milling of boron nitride nanosheets using other milling agents.21 The maximum lateral size of the exfoliated MXene nanosheets significantly increased with the addition of graphite, reaching 3.0?μm (Figure?4e and Supporting Information Figure S32), due to graphite’s dual role in reducing the direct impact of milling balls on the MAX surface and providing buffering-shearing effects during GABM. The ball milling time was another key factor in improving the yield of MXene nanosheets. Figure?4f illustrates the evolution of the yield and size of MXene nanosheets at different ball milling times. The yield of exfoliated MXene nanosheets increased steadily with prolonged ball milling time. After 24?h of ball milling, the yield of thoroughly exfoliated MXene nanosheets reached 95.8%. The size of the MXene nanosheets was 2.9?μm after 12?h of ball milling ( Supporting Information Figure S34 and Table S2). However, the size of the MXene nanosheets decreased with longer ball milling times. For Ti2CTx and Ti3CNTx MXene nanosheets, the yields reached 94.5% and 97.2%, respectively, when the ball milling time was 12?h ( Supporting Information Figure S35). Compared to methods like HF/LiCl and MILD, GABM demonstrated comprehensive advantages in the preparation of MXene nanosheets, including high yield, single-layer percentage, exfoliation efficiency, Young’s modulus, and electrical conductivity (Figure?4g).
The GABM strategy for controlling ball milling time ( Supporting Information Figure S36), combined with liquid cascade centrifugation for size selection, allowed for the controlled preparation of MXene nanosheets of varying sizes.18,52 Using cascade centrifugation, MXene nanosheets ranging from 54?nm to 3.5?μm could be separated from the collected dispersions ( Supporting Information Figure S37), fulfilling the application needed for exfoliated MXene in various fields. The UV–vis absorption spectra of MXene nanosheets with differing sizes showed size-dependent behavior. As the lateral size decreased from 3.5?μm to 54?nm, the plasmon resonance peak shifted from 766 to 742?nm, indicating a size effect on the optical properties of the MXene nanosheets ( Supporting Information Figure S38). Supporting Information Figure S39a and Table S9 depict the relationship between yield and lateral platelet sizes and ball milling time. Extending the milling time significantly increased the yield of exfoliated MXene nanosheets of different sizes. However, the yield of large-sized MXene nanosheets (3.5?μm) reached a maximum after 12?h, which aligned with the results shown in Figure?4f. Further milling beyond this point decreased the yield of large-sized (1.1–3.5?μm) nanosheets, while the yield of small-sized (54–170?nm) nanosheets continued to rise, although larger MXene nanosheets still prevailed. After 28?h of ball milling, the yield of 54?nm nanosheets peaked at 21.4%, showing substantial application potential in electrochemistry and other areas.53,54 Additionally, the yield distributions of MXene nanosheets of various sizes were examined before and after the introduction of graphite ( Supporting Information Figure S39b and Table S9). The introduction of graphite significantly enhanced the yield of large-sized (1.1–3.5?μm) MXene nanosheets, while simultaneously decreasing the yield of small-size (54–353?nm) MXene nanosheets, further demonstrating that graphite played a crucial role in utilizing the GABM method for the high-yield preparation of large-sized exfoliated MXene nanosheets.
Scalable macroscopic assembled materials from MXene nanosheets
The remarkable dispersibility of MXene nanosheets produced using the GABM strategy provides them with excellent solution processability, attributed to the abundant hydrophilic groups, which are essential for assembling macroscopic materials.55–57 Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene nanosheets can be arranged into scalable macroscopic structures with various stacking configurations, such as one-dimensional (1D) fibers, 2D films, and 3D bulk structures, through simple solution assembly techniques (e.g., wet spinning, layer-by-layer blade coating, and the use of a bidirectional ice template) ( Supporting Information Figures S40–S44). The cross-section of high-density MXene fibers displayed a wrinkled surface structure resulting from the deposition of planar MXene platelets on a highly curved fiber surface ( Supporting Information Figure S41). The cross-section of MXene films showed a dense, well-stacked, and planar structure ( Supporting Information Figure S42). Additionally, the MXene lamellar scaffold aerogels featured an oriented porous lamellar structure ( Supporting Information Figure S44). The tensile strengths of pure Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene fibers were 158.2?±?2.3, 120.9?±?5.3, and 122.4?±?10.4?MPa, respectively ( Supporting Information Figure S45 and Table S10), surpassing the tensile strengths previously reported for pure MXene fibers.58,59 The tensile strengths of the pure MXene films were 96.0?±?1.2?MPa for Ti2CTx, 100.6?±?2.1?MPa for Ti3C2Tx, and 127.0?±?3.7?MPa for Ti3CNTx, respectively ( Supporting Information Figure S46 and Table S10), which were greater than those reported for MXene films prepared via vacuum filtration ( Supporting Information Table S11). The MXene lamellar scaffold aerogels demonstrated high compressive elasticity and regained their original shape after being compressed by over 60% ( Supporting Information Figure S47). Thanks to the highly oriented structure resulting from the layer-by-layer blade coating method, the electrical conductivities of Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene films were 4,306?±?461, 10,756?±?497, and 3,494?±?40?S cm?1, respectively ( Supporting Information Figure S48), which were higher than the electrical conductivities previously reported for pure MXene films ( Supporting Information Table S11). This improvement was attributed to the high quality of the MXene nanosheets produced using the GABM method, confirming their potential as ideal building blocks for assembling high-performance MXene composites.
Using Ti3C2Tx MXene as an example, we demonstrated the step-by-step construction of large-area,
high-performance sequentially bridged MXene (SBM) films through layer-by-layer blade
coating and ionic solution immersion processes (Figure?5a). Supporting Information Figure S49 illustrates the preparation process for creating SBM films. For film preparation,
the glutaraldehyde (GA) bridging agent solution was thoroughly mixed and reacted with
the MXene solution to achieve a uniform dispersion of MXene-GA hybrid building blocks
( Supporting Information Figure S50). A covalently bridged MXene (CBM) film was formed using blade coating. In an additional
step, further ionic bonding was introduced by Zn2+ infiltration to obtain dense, highly oriented SBM films. The thickness of the SBM
film could be easily adjusted by controlling the number of layers deposited using
the layer-by-layer blade coating method, and all films exhibited a dense, uniformly
stacked structure as the thickness increased ( Supporting Information Figure S51). The structural evolution from MXene-GA sol to SBM films is demonstrated in Supporting Information Figure S49b. The same layer-by-layer blade coating technique was employed to assemble the MXene
sol into scalable pure MXene films. Ionic bonding bridged MXene (IBM) films were created
by treating the MXene films through the same immersion, washing, and drying process.
Figure 5 | Applications of MXene nanosheets produced using the GABM process. (a)?Photograph of
a roll of SBM film extended, measuring 25?cm in width. (b)?Cross-sectional SEM image
of the SBM film. (c)?Structural model of the SBM film interface bridged by Zn2+ and GA. (d)?WAXS pattern for an incident Cu-Kα X-ray beam parallel to the film plane.
The azimuthal scan profiles for the (002) peak of the SBM film were recorded. (e)?Frequency
variation function of EMI SE values for MXene and SBM films before and after storage
for 10?days in humid air at ～ 100% relative humidity. (f)?Comparison of SSE t?1 and thickness between the SBM film and other solid materials (CB, carbon black; MWNT,
multiwalled nanotubes). (g)?Mid-IR emissivity spectra for MXene and SBM films before
and after storage in humid air at 100% relative humidity for 10?days. (h)?IR photographs
of the model airplane before and after coating with SBM film after heating to 60?°C.
The inset displays the corresponding photographs of the model airplane, noting that
the cockpit area of the model aircraft was not coated with SBM film.
The SEM image of the cross-section of the SBM film showed highly oriented, densely stacked lamellar structures (Figure?5b). The EDS maps of the fracture surface of the SBM films indicated that Zn2+ uniformly penetrated the MXene nanosheets, resulting in a weight increase of 2.42?wt % ( Supporting Information Figure S52). XPS spectra ( Supporting Information Figure S53a) revealed a new Zn2+ 2p peak in the IBM and SBM films compared to the MXene and CBM films, confirming the modification by Zn2+. The Ti2+ (I, II, IV) 2p3/2 and Ti3+ (I, II, IV) 2p3/2 peaks of the IBM and SBM films increased to 456.2 and 457.4?eV from 455.8 and 457.1?eV for the MXene and CBM films, respectively ( Supporting Information Figure S53b–e). This increase is attributed to the bonded Zn2+ having higher electronegativity than Ti, which reduced the electron cloud density of Ti.60 Additionally, the atomic percentage in C–Ti–O–X rose to 34.5% and 40.5% for the CBM and SBM films, respectively, compared to 16.2% for the pure MXene films in O 1s spectra ( Supporting Information Figure S54 and Table S12), suggesting the formation of Ti–O–C bonds between the MXene nanosheets and GA. Fourier transform infrared spectroscopy (FTIR) displayed a new peak at approximately 839?cm?1 ( Supporting Information Figure S55), indicating the formation of a Ti–O–C covalent bond between the MXene nanosheets and GA molecules.56 The –OH peak (3458?cm?1) of the MXene films was redshifted to 3455?cm?1 in the IBM films and to 3438?cm?1 in the SBM films, indicating the formation of H–O→Zn2+ coordination.61 Thus, the SBM films were prepared through interfacial reactions, where the MXene nanosheets were initially covalently cross-linked with GA via nucleophilic substitution to form Ti–O–C bonds and subsequently formed ionic bonds with Zn2+ (Figure?5c and Supporting Information Figure S56).
Due to strong interfacial interactions, the XRD spectrum indicated that the (002) diffraction peak of the SBM films shifted to a higher angle compared to the MXene films, which corresponded to a reduction in interlayer spacing from 1.45 to 1.42?nm ( Supporting Information Figure S57 and Table S13). The narrower full width at half-maximum (FWHM) suggested that covalent and ionic bonding could effectively enhance the alignment of MXene nanosheets. The high orientation degree of SBM films (0.951, Figure?5d), described by Herman’s orientation factor (f),55 was either higher than or comparable to that of pure MXene films (0.940, Supporting Information Figure S58), implying that the sequential bridging process improved the orientation of MXene nanosheets. This improvement was attributed to the stabilization of the aligned structure by ionic and covalent bonding, thereby suppressing the wrinkling and misalignment of the sheets caused by capillary forces during the drying process.60,62 Due to the high alignment of MXene nanosheets, the density of SBM films significantly increased from 3.25?±?0.03 to 3.41?±?0.06?g cm?3 ( Supporting Information Table S14).
Typical stress–strain curves for MXene, IBM, CBM, and SBM films are shown in Supporting Information Figures S59 and S60. The corresponding mechanical properties are provided in Supporting Information Table S15. Due to the enhancement of interlayer interfacial interactions from covalent and ionic bonding, as well as induced orientation and densification, the SBM films exhibited higher mechanical properties than the MXene, IBM, and CBM films. This included a tensile strength of 483.2?±?8.1?MPa, Young’s modulus of 51.3?±?4.7?GPa, and toughness of 2.70?±?0.26?MJ m?3 ( Supporting Information Figure S61 and Table S15). These values were 1.27, 1.56, and 0.93 times higher than those of the CBM films, 1.94, 1.89, and 2.01 times higher than those of the IBM films, and 4.80, 3.49, and 5.19 times higher than those of the MXene films, respectively.
The tensile strength and modulus of the SBM films surpassed those of most previously reported MXene composite films ( Supporting Information Table S17). The bridged SBM films displayed curled crack edges, while the unbridged MXene films exhibited flat crack edges ( Supporting Information Figures S62 and S63), confirming the enhanced interlayer interaction in bridged MXene films.55,60 Furthermore, the resistance of different film types to ultrasonic damage correlated positively with their mechanical strength, with bridged SBM films demonstrating greater resistance to ultrasonic damage compared to unbridged MXene films ( Supporting Information Figure S64). Even with insulating binder GA molecules positioned between the MXene layers, the SBM films maintained a high electrical conductivity of 9,443?±?698?S cm?1 due to improved orientation and reduced interlayer spacing ( Supporting Information Table S16). The conductivity of SBM films was on par with some pure MXene films and exceeded that of previously reported MXene composite films ( Supporting Information Table S17).
Large-scale SBM films for electromagnetic interference shielding and thermal camouflage applications
Thanks to the scalability and high yield of MXene nanosheets produced by GABM, large-area, high-performance MXene films could be developed for various applications, including EMI shielding and thermal camouflage. The high conductivity of MXene films positions them as excellent candidates for EMI shielding.5 Supporting Information Figure S65 and Figure?5e illustrate the EMI shielding effectiveness (EMI SE) of various MXene and SBM films at frequencies ranging from 8.2 to 12.4?GHz (X-band). The EMI SE values for Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene films, with thicknesses of 2.72, 2.45, and 2.35?μm, respectively, reached 38.0, 51.8, and 36.6?dB at a frequency of 8.2?GHz. Notably, Ti3C2Tx MXene films exhibited superior EMI shielding performance ( Supporting Information Figure S48), attributed to their high electrical conductivity.5 The SBM-1, SBM-2, and SBM-3 films, with thicknesses of 2.36, 6.98, and 20.85?μm respectively, demonstrated remarkable EMI shielding performance due to their excellent conductivity and highly aligned structure, with EMI SE values ranging from 50.5 to 81.6?dB (Figure?5e and Supporting Information Figures S51 and S66), surpassing the commercial standard of 20?dB.57 Structural densification of the SBM film could inhibit oxygen and water penetration, slowing the oxidation of the MXene film and enhancing long-term performance stability. For instance, after being stored in humid air at ～ 100% relative humidity for 10?days, the shielding capability of the SBM film decreased by only 1.6%, which was less than the 9.8% decrease observed for the MXene film (Figure?5e). The thickness ratio SE t?1 of the SBM film, calculated by dividing SE by thickness, was 21,419?dB mm?1, outperforming that of the MXene film (20,955?dB mm?1) ( Supporting Information Figure S61). The surface-specific shielding effectiveness per unit thickness (SSE t?1) of the SBM film, obtained by further dividing SE by thickness and density, reached 62,811?dB cm2 g?1, surpassed most solid shielding materials, including some commercial options such as copper, aluminum foils, and Cu-Ni alloys (Figure?5f and Supporting Information Table S18). The examination of the EMI shielding mechanism of the MXene films is depicted in Supporting Information Figure S67. The power efficiency of reflection (R) significantly exceeds that of absorption (A) and transmission (T) for all MXene films, indicating that the overall shielding performance primarily arose from reflection rather than absorption, due to the substantial impedance mismatch between the conductive MXene nanosheets and free space.63
In addition to excellent EMI performance, MXene exhibited low mid-IR emissivity, suggesting its potential applications in thermal camouflage and photothermal management.6,7 According to Kirchhoff’s law, the IR emissivity of an object in a state of thermal equilibrium was equal to its IR absorption rate.6 The emissivity spectra of various MXene and SBM films were obtained by measuring reflectance over a wavelength range of 2.5–25?μm using an FTIR spectrometer at room temperature (Figure?5g and Supporting Information Figure S68). The emissivity of all MXene films decreased as the wavelength increased. The average IR emissivity (2.5–25?μm) for pure Ti2CTx, Ti3C2Tx, and Ti3CNTx MXene was 0.35, 0.18, and 0.48, respectively (Ti3C2Tx > Ti2CTx > Ti3CNTx, Supporting Information Figure S48). The IR emissivity of the MXene films diminished with increasing electrical conductivity, exhibiting behavior similar to that of thin metal films.64 When viewed with an IR camera, the Ti3C2Tx MXene films demonstrated superior thermal camouflage compared to Ti2CTx and Ti3CNTx MXene films ( Supporting Information Figure S69). Following interfacial cross-linking, the average emissivity of the SBM films was 13.9% in the mid-IR band, with a wavelength range of 7–25?μm, and the lowest IR emissivity was 10.5% at 19.5?μm. In comparison, the average emissivity of the MXene films prior to cross-linking was 14.9%, with the lowest emissivity at 19.1?μm being 12.1%, attributed to the higher nanosheet orientation of the SBM films.7 Furthermore, the mid-IR emissivity of the SBM films stored in humid air was significantly lower than that of the MXene films stored in humid air for 10?days (Figure?5g). Over the wavelength range of 2.5–25?μm, the average IR emissivity of the SBM films increased by only 3.8%, notably lower than the 143.3% increase observed in the MXene films, indicating that the SBM films offer more stable performance for thermal camouflage applications. When viewed by an IR camera, the SBM films displayed a smaller increase in surface radiation temperature after storage in humid air compared to the MXene films ( Supporting Information Figure S70). Due to structural vacancies and atomic defects, MXene was prone to erosion by O2 or H2O, while the interfacially cross-linked and densely structured SBM effectively blocked O2 or H2O, thereby maintaining its superior performance in complex environments.55,65 Our SBM films achieved a lower mid-IR emissivity and significantly higher strength than that of most reported metal composites, polymer composites, carbon composites, MXene composites, and commercial stainless steel ( Supporting Information Figure S71 and Table S19). The exceptional and stable IR emissivity properties of SBM films could be utilized in thermal camouflage coatings for aircraft and the human body. For instance, the areas of a model aircraft and a human arm covered with SBM films exhibited similar radiant temperatures to the environment under an IR camera (Figure?5h and Supporting Information Figure S72). In stark contrast, the uncoated areas displayed high radiant temperatures, highlighting the effective thermal camouflage of SBM films.
Conclusion
We present a high-yield, scalable strategy for producing high-quality single-layer MXene nanosheets at room temperature through a synergistic one-step process that integrates exfoliation and etching via GABM with an etching agent. We incorporated graphite into the ball milling process to create a ternary cooperative system with the zirconia balls used for milling and the etching/intercalation agent. This effectively enhanced shear forces while minimizing impact forces. This approach resulted in high efficiency and yield of thoroughly exfoliated large-area MXene nanosheets. The prepared MXene nanosheets achieved the highest yield and the shortest preparation time reported for generating highly exfoliated MXene nanosheets. These MXene nanosheets exhibited excellent Young’s modulus and electrical conductivity, possessed abundant hydrophilic functional groups, and could be assembled into various structures, including fibers, films, and aerogels. The resulting MXene composite film demonstrated high tensile strength, high electrical conductivity, outstanding EMI shielding capability, and appealing IR thermal camouflage performance. These findings indicated that the GABM strategy offered an efficient and high-yield method for preparing single-layer MXene at a low cost. Additionally, this strategy could be broadly applied to exfoliate various MXenes across a wide range of MAX phases, paving the way for scalable production and industrial applications of MXene.
Supporting Information
Supporting Information is available and includes supplementary materials and methods, supplementary figures, supplementary tables, supplementary movies, and supplementary references.
Conflict of Interest
The authors declare no competing interests.
Author Contributions
Q.C. supervised the project and conceived and designed the experiments. C.L., F.T., and Z.W. performed the experiments and characterizations. C.L. and F.T. performed the data analyses. F.T. and Z.Z. performed the AFM nanoindentation characterization. C.L., F.T., Z.W., and Z.H. performed electrical conductivity and EMI shielding characterizations. C.L. and Z.Z. performed the TEM characterization. C.L., F.T., Z.W., and Q.C. co-wrote the manuscript. All authors discussed the results and commented on the manuscript.
Funding Information
This work was supported by the National Science Fund for Distinguished Young Scholars, China (grant no. 52125302), the National Natural Science Foundation of China (grant no. 52550002), the National Key Research and Development Program of China (grant no. 2021YFA0715700), the Open Research Fund of Suzhou Laboratory through the SZLAB-1108-2024-ZD002, the New Cornerstone Science Foundation through the XPLORER PRIZE, and Suzhou Key Laboratory of Bioinspired Interfacial Science, China (grant no. SZ2024004).
Acknowledgments
The authors thank H. Peng, Q. Xuan, and G. Wang at the School of Chemistry, Beihang University, Beijing, China, for the EMI shielding measurements; J. Yang at Suzhou Institute for Advanced Research, University of Science and Technology of China for the discussion about the exfoliation mechanism. S. Wan at the School of Chemistry, Beihang University, for discussion about method and properties comparison; C. Liu and Y. Song at the School of Chemistry, Beihang University, for their help in the fabrication and testing of the FET device. We thank the High-Performance Computing Platform at Beihang University, the Analysis and Testing Center of Beihang University, the Physical and Chemical Analysis Center at Suzhou Institute for Advanced Research, and the University of Science and Technology of China for their support.
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