Atomic Defects in Monolayer Titanium Carbide (Ti3C2Tx) MXene

Sep 6, 2016 - (d) Scatter plot of defect concentration from images acquired from samples produced using different HF concentrations. ..... National En...
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Atomic Defects in Monolayer Titanium Carbide (Ti3C2Tx) MXene

Xiahan Sang,*,† Yu Xie,*,† Ming-Wei Lin,† Mohamed Alhabeb,‡ Katherine L. Van Aken,‡ Yury Gogotsi,‡ Paul R. C. Kent,†,§ Kai Xiao,† and Raymond R. Unocic*,† †

Center for Nanophase Materials Sciences and §Computer Science and Mathematics Division, Oak Ridge National Laboratory, Oak Ridge, Tennessee 37831, United States ‡ Department of Materials Science and Engineering, and A. J. Drexel Nanomaterials Institute, Drexel University, Philadelphia, Pennsylvania 19104, United States S Supporting Information *

ABSTRACT: The 2D transition metal carbides or nitrides, or MXenes, are emerging as a group of materials showing great promise in lithium ion batteries and supercapacitors. Until now, characterization and properties of single-layer MXenes have been scarcely reported. Here, using scanning transmission electron microscopy, we determined the atomic structure of freestanding monolayer Ti3C2Tx flakes prepared via the minimally intensive layer delamination method and characterized different point defects that are prevalent in the monolayer flakes. We determine that the Ti vacancy concentration can be controlled by the etchant concentration during preparation. Density function theory-based calculations confirm the defect structures and predict that the defects can influence the surface morphology and termination groups, but do not strongly influence the metallic conductivity. Using devices fabricated from single- and few-layer Ti3C2Tx MXene flakes, the effect of the number of layers in the flake on conductivity has been demonstrated. KEYWORDS: MXene, conductivity, defect, vacancy, minimally intensive layer delamination (MILD) catalysis.15−18 About 20 MXenes including Ti2C, Ta4C3, (Ti0.5,Nb0.5)2C, (V0.5,Cr0.5)3C2,19 Nb2C, V2C,20 and Nb4C321 have been reported. MXenes are either metallic or semiconducting, and their surface functional groups make them hydrophilic. These properties and the large size of the MXene family provide versatility in designing 2D nanoelectronic devices, sensors, and catalysts. Defects have been extensively studied in 2D materials, such as graphene,22,23 BN,24−26 and MoS2,27,28 and have been shown to influence electronic, electrical, and optoelectronic properties.22,23,26,28−32 Therefore, a detailed understanding of the single-layer structure and point defects is crucial for exploration of the physiochemical properties of MXenes. Here, aberrationcorrected atomic-resolution scanning transmission electron microscopy (STEM) imaging is performed on a prototypical MXene phase Ti3C2Tx (where T stands for surface termination, such as −OH) with different off-tilt angles from the c axis, in order to confirm the single-layer crystal structure from different perspectives. STEM imaging also revealed the presence of

T

he past decade has witnessed a rapid increase in research on two-dimensional (2D) materials, inspired by the discovery of attractive properties of freestanding single-layer graphene. 1 The search for other technologically relevant and scientifically important 2D materials has led to the discovery of several classes of 2D materials in recent years.1−6 MXene phases are a group of 2D materials that have already shown great promise in energy storage applications.7−11 The MXene phase is synthesized by etching the bulk layered MAX phase, where the “M” element is a transition metal, “A” is a group A element such as Al, and “X” is either C or N,12 using hydrofluoric acid (HF) or HFcontaining etchants.7,10,11 The layered crystal structure of the prototypical and most widely used for MXene synthesis MAX phase, Ti3AlC2, is shown in Figure 1a.10 The Ti3C2 layers, consisting of three Ti sublayers (Ti1, Ti2, Ti3) with C located at octahedral interstitial sites, are connected by the more reactive Al layers. After the Al atoms are selectively removed by HF etching,13 the Ti3C2 layers can then be exfoliated by sonication. The top and bottom Ti sublayers (Ti1 and Ti3) tend to terminate with functional groups (Tx) such as O2−, OH−, or F− (Figure 1a).14 These functional groups are very important for their applications in energy storage and © 2016 American Chemical Society

Received: August 3, 2016 Accepted: September 6, 2016 Published: September 6, 2016 9193

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functional theory (DFT). The excellent conductivity suggested by DFT is further confirmed by direct electrical measurements on devices made of single- and few-layer flakes of Ti3C2Tx devices.

SINGLE-LAYER Ti3C2TX MXene CRYSTAL STRUCTURE ANALYSIS High angle annular dark field (HAADF)-STEM imaging is an important tool in 2D materials characterization and is used to unambiguously resolve the crystal structure and defect configurations.28,33 The atomic arrangement of 2D materials can be directly investigated with subangstrom spatial resolution. Moreover, the intensity, I, in HAADF-STEM images is proportional to the atomic number Z (I ∝ Z1.64)34 and the specimen thickness, rendering them directly interpretable. Figure 1b shows a typical low-magnification HAADF-STEM image of Ti3C2Tx flakes. The sample was prepared from Ti3AlC2 powder using the minimally intensive layer delamination (MILD) method, which eliminates the need for sonication and produces large Ti3C2Tx flakes (see details in the Methods section). Two large micrometer-size flakes are observed in Figure 1b that partially overlap. The nonoverlapped flakes exhibit even darker contrast than the supporting lacey carbon film, indicating that they should be no more than nanometers thick. Electron energy loss spectroscopy (EELS) confirms the existence of Ti, C, and O and an absence of F in Ti3C2Tx (Figure S1), which suggests that the termination groups Tx for these flakes should mainly be −OH or −O. An atomic-resolution HAADF-STEM image acquired along the c axis of the Ti3C2Tx flake is shown in Figure 1c and is overlaid with the projected atomic structure model of a Ti3C2Tx single layer. Here, the brighter atomic columns arranged in a hexagonal pattern directly correspond to the stacked Ti atoms from the Ti1, Ti2, and Ti3 sublayers. The C atoms and any of the functional groups are not distinguishable because of their

Figure 1. (a) Crystal structure model of MAX phase Ti3AlC2 (left), MAX phase Ti3AlC2 projected along the a axis (middle), and MXene phase Ti3C2Tx projected along the a axis (right) with −OHterminated functional groups. The three Ti sublayers are labeled Ti1, Ti2, and Ti3. (b) Low-magnification HAADF-STEM image showing the sheet-like morphology of two Ti3C2Tx flakes. The number of layers is indicated for areas exhibiting different contrast; two layers (2) are due to overlapping flakes. (c) Atomic resolution HAADF-STEM image of Ti3C2Tx showing the Ti atom hexagonal lattice aligned along the c axis with the overlaid crystal structure.

different point defects in single-layer Ti3C2Tx layers. The defects were formed during the etching process, and the defect concentration can therefore be tuned by adjusting the HF concentration in the etchant. The effects of point defects on conductivity and surface structure are explored using density

Figure 2. Comparison between experimental high-resolution HAADF-STEM images and simulated HAADF-STEM images for Ti3C2Tx MXene from a single layer (SL), double-layer model AA, and double-layer model AB, at four different off-tilt orientations: (a) α = 0 and β = 0 mrad; (b) α = 0 and β = −189 mrad; (c) α = 0 and β = 50 mrad; (d) α = 0 and β = 356 mrad. The insets show the projected arrangement of Ti atoms for each case. (e) HAADF-STEM image showing different layer thicknesses. (f) Intensity line profile acquired from the HAADF-STEM image along the white arrow from bottom to top. 9194

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Figure 3. Atomic-scale HAADF-STEM images of defects in single-layer Ti3C2Tx. Comparison between experimental HAADF-STEM image, defect crystal structure determined from DFT, and simulated HAADF-STEM image of VTi (a) and two adjacent VTi within the same sublayer (b). (c−e) Experimental HAADF-STEM images of vacancy clusters: two adjacent VTi forming within two different sublayers (c), three VTi within the same sublayer (d), and 17 VTi within the same sublayer (e). Scale bars are 0.5 nm. (f) VTi formation energy on bare Ti3C2 and terminated single-layer Ti3C2Tx. (g) Formation energy of VCTi clusters as a function of number of VTi. The dashed trend line shows the linear relationship of the formation energy.

the single-layer structure is the only structure that matches the experimental images acquired for the other two off-tilt orientations shown in Figure 2: α = 0, β = −50 (Figure 2c) and α = 0, β = 356 (Figure 2d). A single-layer Ti3C2Tx flake is therefore confirmed since all the experimental images match the single-layer structure. Since the thinnest Ti3C2Tx flakes have been shown to be single-layer, the number of layers composing other Ti3C2Tx flakes can be determined by comparing HAADF-STEM image intensity with the intensity observed for a single layer. Figure 2e shows a low-magnification STEM image of overlapping Ti3C2Tx flakes that exhibit different thicknesses and, therefore, intensities in the image. A line profile across different flakes in Figure 2f shows that the intensity increases from roughly 0.135 to 0.242 to 0.351, which is approximately 1:2:3. Therefore, the image intensity can roughly be used to determine the layer number for different Ti3C2Tx flakes. In our experiments, the multilayer Ti3C2Tx flakes are always stacks of single-layer flakes having different rotation orientations between the layers (Figure S3). This suggests that the current sample preparation method has a very high yield of single-layer flakes, which could lead to industrial level application of single-layer Ti3C2Tx flakes.

low scattering intensity and their overlap with the Ti atoms when projected along or close to the c axis. Figure 2a compares the experimental HAADF-STEM image of a Ti3C2Tx flake oriented along the c axis with similarly oriented simulated HAADF-STEM images that were calculated using a multislice algorithm35 for single-layer Ti3C2Tx and two different double-layer structural configurations (model AA and model AB; see Figure S2). These two model double-layer structures were proposed from DFT geometry optimization15 and differ in their interlayer arrangement (Figure S2). All three simulated images agree well with the experimental image despite slight contrast differences, suggesting that the layer number (single-layer versus double-layer) cannot be unambiguously determined from HAADF-STEM images acquired with the Ti3C2Tx flake directly oriented along the c axis. This is because the projected Ti atom arrangements (overlaid on all the simulated images in Figure 2) are independent of the number of layers. To resolve this ambiguity and accurately determine the number of layers comprising the Ti3C2Tx flakes, the experimental versus simulated image comparison was extended to include HAADF-STEM images acquired with the flake tilted slightly off the c axis (Figure 2c,d), which was achieved either by tilting the flake off-axis in the microscope or by using sample areas that are naturally off-tilt on the rumpled flakes. The off-tilt angles, α and β, are defined as the tilt angles around the a axis and the axis that is perpendicular to both the a and c axes (c × a), respectively. As α and β deviate from zero, the projected Ti atom arrangements for a single layer will be different from the two double-layer configurations. For example, as shown in Figure 2b for α = 0 and β = −189 mrad, the Ti atoms in simulated HAADF-STEM images form a 3 × 3 atom hollow-square-like pattern for single-layer Ti3C2Tx, while simulated images of both double-layer models AA and AB show overlapping Ti atoms and blurred contrast. The experimental HAADF-STEM image exhibits well-separated Ti atoms that match the simulated image and projected crystal structure from the single-layer Ti3C2Tx (Figure 2b). Similarly,

ATOMIC-SCALED DEFECTS IN SINGLE-LAYER Ti3C2TX Similar to other 2D materials, the electronic and surface properties of single-layer Ti3C2Tx can be modified by the presence of intrinsic defects. Although Ti vacancies have previously been observed in MXene flakes,36 a detailed knowledge of the structure of Ti vacancies and their impact on the electrical and chemical properties have not been investigated. Figure 3 shows HAADF-STEM images of the different Ti point defects observed in single-layer Ti3C2Tx, including single Ti vacancies, VTi (Figure 3a) and Ti vacancy clusters, VCTi, where C equals the number of clustered vacancies (Figure 3b−e). The vacancies likely form on the top (Ti1) and bottom (Ti3) sublayers due to their direct contact with the etchant solution, but it is also possible to have vacancies in the 9195

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ACS Nano middle layer that are inherited from the original MAX phase. The HAADF-STEM image in Figure 3b shows two adjacent V2Ti from the same sublayer that exhibit similar contrast levels, while Figure 3c shows two adjacent V2Ti from two different sublayers that clearly exhibit varying contrast levels. It is frequently observed that vacancies form as clusters within the same sublayer. Figure 3d and e show VCTi clusters with 3 and 17 vacancies from the same sublayer, respectively. A honeycomb structure is formed when a large number of vacancies cluster in the same sublayer. DFT simulations were performed to gain insight regarding the energetic costs associated with Ti defect formation. We start by examining the case for a single vacancy VTi on a bare single-layer Ti3C2. Considering the underlying energy barrier for a Ti atom to penetrate the layer of functional groups, a VTi is likely to form during the first step of exfoliation (eq 1 in ref 10), when the surfaces of the Ti3C2 flakes are not terminated. The calculated VTi formation energies are 2.842 eV for the outer sublayers Ti1 and Ti3 and 6.485 eV for the inner Ti2 sublayer. Therefore, it is much easier to form a VTi in the outer Ti sublayers than within the inner Ti sublayer. Assuming all the VTi are within the Ti1 sublayer, different defect structures of −OH-terminated Ti3 C 2T x were optimized using DFT. Experimental HAADF-STEM images, optimized crystal structures, and the corresponding simulated HAADF-STEM images are shown in Figure 3a for one VTi and Figure 3b for two adjacent VTi in the Ti1 sublayer. More examples and detailed line profile comparisons can be found in Figure S4. The simulated HAADF-STEM images agree well with the experimental images. The influence of the different termination functional groups on the formation energy of VTi (Figure 3f) will also determine the stability of the vacancies for each case. The formation energy of VTi on −O-terminated Ti3C2Tx is much higher than that for bare Ti3C2, suggesting a strong interaction between Ti and the −O functional groups. On the contrary, −F- and −OHterminated Ti3C2Tx show a comparable VTi formation energy to that of bare Ti3C2. To simplify the DFT calculations, we focus on the formation of a VCTi on single-layer bare Ti3C2. The calculated formation energies of VCTi clusters and up to six VTi are summarized in Figure 3g and demonstrate a linear relationship with the number of VTi. The VCTi cluster has a similar formation energy to the sum of separated VTi when C is less than 3. The energy of the VCTi cluster becomes larger than the sum of separated VTi when C is larger than 3, which may be induced by the interaction between VCTi clusters due to the limited simulation cell size. Nevertheless, the per atom formation energy is 3.017 and 2.808 eV/atom for a V6Ti cluster and a single VTi, respectively. The small energy difference (∼0.2 eV/atom) suggests that the formation of large VCTi clusters might be feasible. The ability to tune the surface defect concentration is crucial toward controlling the surface properties for catalytic, energy storage, and other applications. As discussed before, the defects are formed when the surface Ti atoms are etched/removed by the HF in the etchant solution. Thus, changing the HF concentration in the etchant should result in changing the VTi concentration. Figure 4 shows HAADF-STEM images acquired from single-layer Ti3C2Tx flakes prepared using etchants with different HF concentrations (more HAADF-STEM images can be found in Figure S5). Vacancy clusters are rarely observed after etching with a 2.7 wt % HF concentration, but are relatively common after etching with a 7 wt % HF

Figure 4. HAADF-STEM images from single-layer Ti3C2Tx MXene flakes prepared using etchants with different HF concentrations: (a) 2.7 wt % HF, (b) 5.3 wt % HF, and (c) 7 wt % HF. Single VTi vacancies are indicated by red circles, while vacancy clusters VCTi are shown by blue circles. (d) Scatter plot of defect concentration from images acquired from samples produced using different HF concentrations. The red line shows the error plot with the average and standard deviation for different HF concentrations.

concentration, with otherwise identical synthesis conditions, demonstrating the important role of HF in defect formation. For each case, the ratio between number of VTi and all the Ti sites was calculated for tens of images, and the results are summarized in Figure 4d. The average VTi concentration is clearly correlated with the HF concentration. Surface defects would provide possible sites for intercalation of ions, potentially increasing the capacity for MXene, but also possibly decreasing the achievable voltage window by catalyzing the electrolyte decomposition. The defects are also important to modify the surface structure and potentially influence the conductivity of single-layer Ti3C2Tx flakes.

SURFACE STRUCTURE AND RELATED ELECTRONIC PROPERTIES To theoretically study the influence of the defects on the surface chemistry of Ti3C2Tx, we used ab initio molecular dynamics (AIMD) simulations. We show the predicted structures of functionalized Ti3C2Tx with Ti vacancy clusters up to V6Ti, by immersing the bare Ti3C2 monolayer with pristine surface on one side and defective surface on the other side in 10% HF solution, as shown in Figure S6. In all cases, the Ti3C2 surfaces are mainly terminated by −OH groups. However, we do observe differences between the pristine and defective surfaces. First, while water molecules are fully dissociated on the pristine surface, there are water molecules attached to the defective surface around Ti vacancies. This is probably due to the change of the charge distribution around the vacancies. Second, with V2Ti or larger vacancy clusters, protons are found in the defective site forming C−H bonds with carbon atoms. This may explain the observed C−H vibrations in NMR spectra that remain even after high-temperature annealing.37 With 9196

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Figure 5. (a−c) Side views of snapshots at 0, 2, and 5 ps of AIMD simulations of a Ti3C2 monolayer V6Ti cluster heated at 350 K. (d) Surface structure of a Ti3C2 monolayer by removing all the water molecules. The blue circles show the two partially function group terminated Ti atoms. Calculated DOS of OH-terminated Ti3C2 monolayer with pristine (e), V2Ti (f), V4Ti (g), and V6Ti (h) cluster surface. No significant difference between pristine and defective surfaces is observed.

Figure 6. (a) Optical image of single- and few-layer Ti3C2Tx flakes on a Si/SiO2 substrate. (b) Optical image showing top-down view of Ti3C2Tx devices with patterned electrodes. (c) AFM image demonstrating that the Ti3C2Tx flake in the device is single layer, with line scan height profile inset along the white arrow. (d) Output characteristics of devices made from single-layer Ti3C2Tx exhibiting linear I−V characteristics. Inset is the transfer characteristics of a single-layer Ti3C2Tx flake device exhibiting absolutely no gate dependence. (e) I−V curves for devices with different Ti3C2Tx thicknesses. (f) Resistivity as a function of Ti3C2Tx thickness for different devices. The scale bars in (a), (b), and (c) are 20, 20, and 4 μm, respectively.

compared to pristine surfaces, Ti atoms around the defective sites are transformed from three-O-coordinated into twocoordinated after removing the water molecules. This generates partially free Ti sites on the MXene surfaces that may enhance the catalytic activity. The effect of the Ti vacancies on the electronic properties of Ti3C2Tx MXene has also been studied. We assume that Ti3C2 surfaces are fully −OH terminated, except for the defects. Protons were added according to AIMD results. The optimized structures are plotted in Figure S7. From the calculated density

water molecules present on the surfaces, it is straightforward to assume they can be removed by heat treatment.38 Therefore, we examine the surface chemistry of a predicted Ti3C2 monolayer with a V6Ti cluster at high temperature via AIMD simulations as an example. Snapshots at 0, 2, and 5 ps of AIMD simulations of a Ti3C2 monolayer with a V6Ti cluster heated at 350 K are displayed in Figure 5a−c, respectively. Intriguingly, water starts to evaporate at 350 K, as shown more clearly in Figure 5d from a top view, suggesting a weak interaction between Ti3C2 surfaces and water molecules. Moreover, 9197

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CONCLUSION In summary, we have performed a comprehensive study of the atomic scaled structure and defects in free-standing single-layer Ti3C2Tx produced through MILD etching using aberrationcorrected STEM imaging. Ti vacancies mainly form on the two surface sublayers, and we demonstrated how the density can be tuned by varying the HF etchant concentration. The influence of the defects on the structure and electronic properties have been explored using DFT and suggest high conductivity even for defective samples. We obtain direct measurements of the electrical conductivity as a function of number of layers in flakes in fabricated devices. The conductivity of single-layer Ti3C2Tx is comparable to graphene and much higher than 1T MoS2. Our combination of atomic resolution characterization, theory, and electrical property measurements thus demonstrates that the single layer Ti3C2Tx is a promising 2D material for future nanoelectronic devices and catalytic applications.

of states (DOS) (Figure 5e−g and Figure S8), it is clear that the defective −OH-terminated single-layer Ti3C2 is still metallic. Comparing to pristine MXene, the Fermi level of defective MXene shifts downward, resulting in a decrease of the DOS at the Fermi level, which should reduce the conductivity of the MXene monolayer. Interestingly, the DOS of −OHterminated Ti3C2 with different concentrations of Ti vacancies are very similar. This suggests that for moderate concentrations of Ti defects −OH-terminated Ti3C2 should exhibit high conductivity. The insensitivity of conductivity to defect concentration for Ti3C2Tx is similar to the cases of 2D phosphorus and halide perovskites, where the defects were also reported to be electrically inactive based on DFT calculations.30,31

CONDUCTIVITY OF Ti3C2TX In order to experimentally study the conductivity of 2D Ti3C2Tx flakes, we made devices on Ti3C2Tx flakes. The flakes were first dispersed in ethanol and then drop-cast on a Si wafer with thermally grown SiO2. Micrometer-size flakes (Figure 6a) were selected using an optical microscope to fabricate the devices. Standard electron beam lithography was conducted on a thin flake (indicated by the red circle in Figure 6a) followed by sequential electron-beam evaporation of a 5 nm Ti layer and a 30 nm Au layer. As shown in Figure 6b, the Ti/Au current collectors were patterned on the Ti3C2Tx flakes on a Si/SiO2 substrate. The flake thickness was measured using atomic force microscopy (AFM), and the larger flake in Figure 6b has a thickness around 1 nm and should be the single-layer Ti3C2Tx (Figure 6c). Figure 6d shows the output and transfer characteristics of the device. We can see from the output characteristics that the Schottky barrier is almost suppressed as the characteristic is perfectly linear. There is absolutely no effect in the drain−source current with gate modulation as shown in Figure 5d, which is a clear indication of the metallic nature of the single-layer Ti3C2Tx. The I−V curve was measured for different layer thicknesses, and the slope increases as a function of the layer thickness (Figure 6e). The resistivity as a function of thickness is plotted in Figure 6f. It is interesting to note the resistivity increases as the thickness increases. The conductivity of single-layer Ti3C2Tx is measured to be around 6.76 × 105 S/ m (see Supporting Information for more details), comparable to chemical vapor deposited (CVD) graphene with a conductivity of 6 × 106 S/m39 and 2 orders of magnitude higher than 1T MoS2 with a conductivity of 8 × 103 S/m.40 The conductivity measurement also agrees with previous measurement from two different groups.41,42 As we know, the contact issue significantly limits the device performance of 2D electronics and requires a 2D material as electrodes to achieve ohmic contacts. The highly conductive Ti3C2Tx can be used as electrical contacts and reduce the contact resistance, to benefit the performance of short-channel field-effect transistors for 2D nanoelectronics, as well as many other devices. This opens up opportunities in 2D nanoelectronics to consider a single-layer MXene phase, which currently includes more than 20 different phases, and the number is still increasing. We hope this can also inspire other synthesis methods, such as CVD, to grow single layers of similar transition metal carbides.

METHODS Materials. Lithium fluoride (LiF, Alfa Aesar, 98.5%), hydrochloric acid (HCl, Fisher Scientific, 37.2%), hydrofluoric acid (HF, Acros Organics, 49.5 wt %), and tetrabutylammonium hydroxide (TBAOH, Acros Organics, 40 wt % solution in water) were used as received. Synthesis of Ti3AlC2 (MAX). Ti3AlC2 was synthesized as described elsewhere,10 and the powders were crushed and sieved through 400 mesh size (≤37 μm particle size) and collected for etching. MILD Synthesis of Ti3C2Tx (MXene). Ti3C2Tx was synthesized using an improved etching route from the previously reported clay method.7 Using the MILD method, we have eliminated excessive processing that was previously needed for Ti3C2Tx delamination.41 Briefly, the etchant solution used in the MILD method was prepared by dissolving 1 g of LiF in 20 mL of 6 M HCl in a 100 mL polypropylene plastic vial after which 1 g of Ti3AlC2 was gradually added to it, and the reaction was allowed to proceed for 24 h at 35 °C. The acidic product was washed copiously with DI H2O via centrifuging at 3500 rpm until pH ≥ 6, at which point a dark green supernatant solution of large Ti3C2Tx flakes was collected after 1 h of centrifuging at the same rpm. Up to 1.5 mg/mL of Ti3C2Tx colloidal solution was collected. STEM Characterization and Simulation. Samples for STEM of the Ti3C2Tx flakes were prepared by drop-casting onto lacey carbon films supported on a Cu grid from the ethanol solution. STEM images were acquired using a Nion UltraSTEM operated at 60 or 100 kV and equipped with aberration correction of the probe forming lens. All the images in Figure 2 were acquired at 100 kV to enhance spatial resolution. The HAADF-STEM images of the defects shown in Figures 3 and 4 were acquired at 60 kV to minimize beam damage. A convergence angle of 31 mrad was used, with HAADF detector inner and outer collection angles of 86 and 200 mrad, respectively. The beam current was between 50 and 70 pm depending on different sessions. To enhance the signal-to-noise ratio, atomic resolution HAADF-STEM images were acquired using a sequential imaging code that acquires 10−20 fast frames sequentially, which were then aligned using cross correlation. One fast frame typically was acquired using a dwell time of 4−8 s with a frame size of either 512 × 512 or 1024 × 1024. HAADF-STEM image simulation was performed using the code from Kirkland.35 An orthogonal supercell with a1 = a, b1 = a + 2b, and c1 = c was created to be compatible with the simulation code. The projected potentials were sampled at 1.7 pm/pixel and were used to simulate the HAADF-STEM images with 19 × 33 probe positions. The simulated data were enlarged using Fourier interpolation and blurred with a Gaussian function of 0.08 nm to approximately account for the finite size of the electron probe for 100 kV and 0.1 nm for 60 kV. Device Fabrication and Electrical Property Measurement. Ti3C2Tx flakes were first drop-casted onto 290-nm-thick, heavily doped Si/SiO2 substrates. Samples were characterized by both a Nikon 9198

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ACS Nano LV150 optical microscope and an atomic force microscope (Bruker Dimension Icon). Electron beam lithography (FEI Nanolab 600 dual beam with Raith pattern writing software) was used for Ti3C2Tx fieldeffect transistor (FET) device fabrication. First, a layer of PMMA 495A4 was spin-coated on top of the Ti3C2Tx flakes, followed by 180 °C baking on a hot plate. After pattern writing and developing, a 5 nm layer of Ti and a 30 nm layer of Au were deposited using an electron beam evaporator. Finally, well-defined source and drain electrodes of Ti3C2Tx FET devices were fabricated after a lift-off process using acetone/isopropyl alcohol. The electrical properties of the devices were then measured in a cryogenic vacuum chamber (Desert Cryogenics, ∼10−6 Torr) with a Keithley 4200 semiconductor analyzer in a back-gating configuration. The gate voltages were swept from −20 V to 20 V at Vds = 0.1 and 0.5 V for transport curve measurements, showing typically metallic behavior without gate dependence. The output I−V curves were swept from −1 to 1 V at different gate voltages, demonstrating linear behavior to further verify good metallic properties. DFT Calculation. First-principles calculations were performed using DFT43 and the all-electron frozen-core projected augmented wave (PAW) method44 as implemented in the Vienna ab initio simulation package (VASP).45 For the exchange−correlation energy, we used the Perdew−Burke−Ernzerhof (PBE) version of the generalized gradient approximation (GGA).46 Dispersion interactions were included using the scheme of Grimme.47 Ab initio molecular dynamics simulations were conducted to investigate the surface structure of MXenes in solution. More computational details are described in the Supporting Information. Estimation of Formation Energy. The formation energy of defects was defined as Ef = Edefect − Eperfect + ΣiniμTi, where Edefect and Eperfect are the total energies of the Ti3C2 monolayer with specific defects and pristine surface, respectively, ni is the number of Ti vacancies, and μTi is the chemical potential of the Ti atom. Under thermodynamic equilibrium, μTi and μC satisfy the equation μTi3C2 = bulk bulk bulk 3μTi + 2μC = 3μbulk Ti + 2 μC − ΔHTi3C2, where μTi and μC are the chemical potentials of Ti and C in bulk form, respectively, and ΔHTi3C2 is the formation enthalpy of the Ti3C2 monolayer. Although it is difficult to obtain the exact value of μTi, the range of μTi can be 1 deduced as μTibulk + 3 ΔHTi3C2 ≤ μTi ≤ μTibulk . Since we are interested only in Ti defects, μTi = μbulk Ti was used.

Research Center funded by the U.S. Department of Energy, Office of Science, Office of Basic Energy Sciences. Aberrationcorrected STEM imaging and device fabrication and measurement were conducted at Oak Ridge National Laboratory’s Center for Nanophase Materials Sciences (CNMS), a U.S. Department of Energy Office of Science User Facility. This research used resources of the National Energy Research Scientific Computing Center, a DOE Office of Science User Facility supported by the Office of Science of the U.S. Department of Energy under Contract No. DE-AC0205CH11231.

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ASSOCIATED CONTENT

S Supporting Information *

The Supporting Information is available free of charge on the ACS Publications website at DOI: 10.1021/acsnano.6b05240. Elemental analysis by electron energy loss spectroscopy, crystal structures of models AA and AB, orientation relationship between different flakes, DFT computational details, Ti vacancy clusters, imaging of defects for different HF concentrations, predicted surface structures of a defected Ti3C2 monolayer in HF solution, density of states of a defective Ti3C2 monolayer, and I−V curve measurement and resistivity (PDF)

AUTHOR INFORMATION Corresponding Authors

*E-mail (X. Sang): [email protected]. *E-mail (Y. Xie): [email protected]. *E-mail (R. R. Unocic): [email protected]. Notes

The authors declare no competing financial interest.

ACKNOWLEDGMENTS Research was supported as part of the Fluid Interface Reactions, Structures and Transport (FIRST) Center, an Energy Frontier 9199

DOI: 10.1021/acsnano.6b05240 ACS Nano 2016, 10, 9193−9200

Article

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NOTE ADDED AFTER ASAP PUBLICATION This paper published ASAP on 9/12/16. Figure 4 was replaced and the revised version was reposted on 9/13/16.

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DOI: 10.1021/acsnano.6b05240 ACS Nano 2016, 10, 9193−9200