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Halide Induced Self-limited Growth of Ultrathin Nonlayered Ge Flakes for High-Performance Phototransistors Xiaozong Hu, Pu Huang, Bao Jin, Xiuwen Zhang, Huiqiao Li, Xing Zhou, and Tianyou Zhai J. Am. Chem. Soc., Just Accepted Manuscript • DOI: 10.1021/jacs.8b07383 • Publication Date (Web): 14 Sep 2018 Downloaded from http://pubs.acs.org on September 14, 2018

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Journal of the American Chemical Society

Halide Induced Self-Limited Growth of Ultrathin Non-Layered Ge Flakes for High-Performance Phototransistors Xiaozong Hu†, Pu Huangζ, Bao Jin†, Xiuwen Zhangζ, Huiqiao Li†*, Xing Zhou†*, and Tianyou Zhai†* †

State Key Laboratory of Material Processing and Die & Mould Technology, School of Materials Science and Engineering, Huazhong University of Science and Technology (HUST), Wuhan 430074, P. R. China Shenzhen Key Laboratory of Flexible Memory Materials and Devices, College of Electronic Science and Technology, Shenzhen University, Nanhai Ave. 3688, Shenzhen, Guangdong 518060, P. R. China

ζ

Supporting Information Placeholder ABSTRACT: 2D nonlayered materials have attracted intensive attention due to the unique surface structure and novel physical properties. However, it is still a great challenge to realize the 2D planar structures of nonlayered materials owing to the naturally intrinsic covalent bonds. Ge is one of them with cubic structure impeding its 2D anisotropic growth. Here, the ultrathin singlecrystalline Ge flakes as thin as 8.5 nm were realized via halideassisted self-limited CVD growth. The growth mechanism has been confirmed by experiments and theoretical calculations, which can be attributed to the preferential growth of (111) plane with the lowest formation energy and the giant interface distortion effect of Cl-Ge motif. Excitingly, a Ge flake-based phototransistor shows excellent performances such as a high hole mobility of ~ 263 cm2V-1s-1, a high responsivity of ~ 200 A/W, and fast response rates (τrise=70 ms, τdecay=6 ms), suggesting its great potential in the applications of electronics and optoelectronics.

INTRODUCTION 2D materials such as graphene,1, 2 transition metal dichalcogenides (TMDs),3-8 group III-VIA materials,9, 10 and group IV-VIA materials11-17 have stimulated extensive attention owing to their rich and varied electrical, optical, mechanical, and magnetic properties caused by the ultrathin 2D nanostructures.18-21 Moreover, their 2D nanostructures rendering them high compaticity with silicon process technology and flexible substrates. Up to now, the research of 2D nanostructures is mainly limited to the 2D layered materials such as graphene and TMDs as mentioned above. Because of the relatively weak van der Waals interaction between contiguous layers, 2D layered materials with ultrathin thickness can be effortlessly realized through mechanical exfoliation or chemical vapor deposition (CVD). However, there are also many nonlayered materials such as group IVA,22 II-VI23 and III-VA24 semiconductors with excellent electronic and optoelectronic properties.25 Ge is famed for the high carrier mobilities (electron: µe = 3900 cm2V-1s-1, hole: µh = 1900 cm2V-1s-1), a narrow bandgap of 0.67 eV in bulk phase at 300 K and large absorption coefficient of 2 × 105 cm-1, rendering it excellent candidate for near-infrared (NIR) photodetection. Furthermore, a large Bohr exciton radius (~ 24.3 nm) for Ge makes the quantum confinement effects observed more easily.26 Thus Ge has been significantly explored for field effect transistors (FETs),27 photodetectors,28-29 and nonlinear optics.30, 31 The photodetector based on Ge can be further integrated into silicon process technology based photonic devices, owing to

the capability of detecting weak optical signals with high speed.32 Thus it is highly promising that 2D Ge nanostructures will show novel physical properties due to the confined charge carriers and photon transportation in the 2D structure. For example, the quantum-spin Hall effect and high-Tc superconductivity can be predicted in monolayer Ge.33 However, there are few reports on 2D Ge nanostructures. Most reported 2D Ge nanostructures are obtained by molecular beam epitaxy, or smart-cut process of bulk Ge followed by etching technology, which are complex, highly cost, and incompatible with silicon-based transistor technology.29, 33 CVD has been proved to be a valid way to realize high-quality 2D nanostructures, owing to the advantage over the accurate control on structure, morphology, and size of the synthesized materials.34 However, it is still difficult to realize 2D nonlayered Ge nanostructures due to the intrinsically isotropic covalent bonds in the three-dimensional (3D) directions. When the flake becomes thinner to atomically thin thickness, the 2D anisotropic growth will be tremendously impeded by the unsaturated dangling bonds on the surface.35 Thus the realization of ultrathin Ge nanoflakes still remains significant challenge. Herein, ultrathin Ge flakes are realized via halide-induced self-limited growth and the thickness can be as thin as ~ 8.5 nm. The as-synthesized single-crystalline Ge flakes with cubic structure grow along the (111) plane. The growth mechanism is revealed by both theoretical calculations and experiments, which can be ascribed to the preferential adsorption of Cl on the Ge (111) surface resulting in 2D anisotropic growth. Impressively, the Ge based phototransistor exhibits a typical p-type characteristic with a high hole mobility of ~ 263 cm2V-1s-1 and also exhibits excellent photodetection performances such as a high external quantum efficiency (EQE) (~ 1.0 × 105 %) and high responsivity (~ 200 A/W), as well as a fast decay speed of ~ 6 ms. These excellent performances render 2D Ge flakes promising candidates for next-generation optoelectronics. RESULTS AND DISCUSSION Figure 1a illustrates the improved CVD process (Details see Methods). GeS was used as Ge precursor, and H2 was employed as the reductive agent. Most of all, KCl was introduced to the CVD growth and is the key to realize the ultrathin 2D Ge flakes for the following two reasons: i) Mixing metal precursor and salt will produce a molten solution, leading to increase the vapor pressure of the metal precursor. The introduce of KCl will dramatically promote the reaction;36, 37 ii) The preferential adsorption

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Figure 1. Illustration of the CVD growth and characterizations. a, b) Sketch map of the CVD growth. c) Optical image of Ge flakes on SiO2/Si substrate. d) AFM image and height profile of a thin Ge flake. e, f) Raman spectrum and mapping of a typical Ge flake. g) XRD pattern of the Ge flakes. Inset: crystal structure showing the (111) plane. h) EBSD inverse pole figure (IPF) map of a triangular Ge flake on SiO2/Si substrate. of Cl on the (111) plane of Ge surface will break the thermodynamic equilibrium state and suppress the isotropic growth along 3D directions as shown in Figure 1b.23 Figure 1c demonstrates the large-production of triangular Ge flakes on the SiO2/Si substrate. The thickness can be as thin as ~ 8.5 nm confirmed by atomic force microscopy (AFM) in Figure 1d. The Raman spectrum of a triangular flake shows a single peak at ~ 302 cm-1 (Figure 1e), in accordance with previous reports.29 Figure 1f exhibits the Raman mapping of the peak at 302 cm-1 suggesting the uniformity of the flake. The X-ray diffraction (XRD) pattern (Figure 1g) confirms the cubic phase Ge (JCPDS: 04-0545) without any detectable impurities such as GeS. The only detectable peak of (111) plane means that the (111) plane is the preferred growth orientation of these Ge flakes shown in the inset of Figure 1g. We further employed the electron back-scattered diffraction (EBSD) to probe the crystal orientation of these flakes as shown in the inverse pole figure (IPF) image Figure 1f, indicating that the flake has an ordered (111) in-plane orientation over the whole sample in accord with the above XRD result. Transmission electron microscopy (TEM) characterizations were employed to further explore the crystal structure of the assynthesized Ge samples. Figure 2a shows a low magnification image of a typical triangle. The high-resolution TEM image in Figure 2b shows perfect atomic structure with a lattice spacing of 0.2 nm, corresponding to the (220) planes. Figure 2c shows the selected area electron diffraction (SAED) pattern and the sharp spots also identify the single-crystalline structure and the preferential growth (111) planes. The energy dispersive X-ray microscopy (EDX) mapping of the selected area in Figure 2a suggests a uniform spatial distribution of Ge demonstrated in Figure 2d, and the EDX spectrum shows Ge atom peaks without any S atoms in Figure S1. Based on the above characterizations, the assynthesized Ge flakes shows oriented surface (111) planes. The top view of the crystal structures are illustrated in Figure 2e, indicating the ABC-ABC stacking order of the layers. To further explore the growth mechanism and crystal structures, we performed a series of experiments and theoretical calculations. Figure S2a shows the products with rectangular morphology under Ar flow, and these rectangular products are identified

Figure 2. TEM characterizations and DFT calculations. a) Lowmagnification TEM image of a Ge flake. b) Corresponding HRTEM image of the flake. c) SAED pattern from the same flake. d) Ge elemental mapping of the selected area in (a). e) Structure model of surface (111) plane. f) Formation energy of (100) and (111) planes. g) System energy comparison for ClGe defect distributed from inner bulk phase to surface region. h, i) Illustration of the reaction pathway, from which the surface distortion and unfavorable bonding caused by adsorbed Cl can be observed unambiguously. to be GeS by Raman spectrum in Figure S2b. When H2 was introduced to the CVD system, small and thick islands were obtained in Figure S2c, and these islands are confirmed to be Ge in Figure S2d. While adding KCl, the large-production of triangular Ge flakes on the SiO2/Si substrate can be realized in Figure 1c. Obviously, the addition of KCl plays a key role of promoting the 2D anisotropic growth of these ultrathin Ge flakes. Then, we employed density functional theory (DFT) calculations to further explore the role of the introduced salt and the detailed growth mechanism. Generally, (111) plane possesses the smallest formation energy in cubic phase structure and keeps the most densely stacking mode. It is identified by the DFT calculations in Figure 2f, the formation energy (Ef) of (111) plane is 0.48 eV obviously lower than that of (100) plane (0.72 eV), indicating (111) plane the preferential growth surface. Considering the decreased average energy per atom (~ 28 meV) for the (111) surface plane with Cl adsorption and the gradually reduced system energy for ClGe distributed from inner bulk phase to outermost layer, we believe Cl-Ge bonds are more inclined to form at the surface region (Figure 2g, h). Actually, the crystal structure of the flake along [111] direction is composed of Ge tetrahedron, which indicates every Ge atom deposited at the (111) surface only contains one unsaturated dangling bond. The remaining isolated unpaired electron can be easily captured by Cl to form stable octet electron structure. Thus, the Cl-Ge motif leads to a giant lattice distortion in the aspect of bond angle/length and suppresses the continued growth along [111] direction (Figure 2i). Consequently, the surface Cl-Ge motif results in a distortion structure and 2D anisotropic growth of ultrathin Ge flakes.23 Meanwhile, the weak signal of Cl has been identified by X-ray photoelectron spectroscopy (XPS) in Figure S3, in accordance with the above inference. As well known, second-harmonic generation (SHG)


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Journal of the American Chemical Society ly, resulting in a symmetric tilt boundary. Thus, it is reasonable that the typical angles between neighboring grains are 120° (Figure 3i), which is also in accordance with the six-fold rotational symmetry of Ge (111) plane.40 Then the Ge flake-based phototransistor was constructed as illustrated in Figure 4a with a back-gated configuration (the thickness of Ge flake is ~ 20 nm). Since high-k dielectrics such as Al2O3, HfO2 have been proposed to be valid to screen carriers scattering from charge impurity in devices,43, 44 ~ 20 nm HfO2 as the top passivation layer was employed to promise the optimal performance. Figure 4b shows the Ids-Vds characteristics with gate bias varying from -40 - 10 V, and the linear curves suggest the ohmic contact between metal electrodes and Ge flake. The transfer characteristics are demonstrated in Figure 4c and indicate obviously p-type semiconducting character. Figure S4 shows the leakage currents of the bottom gate with ultralow values of ~ 10-12 A suggesting the high-quality gate of SiO2. Thus the hole mobility can be determined by dIds L , (W and L: width and µ=

Figure 3. SHG microscopy of the as-synthesized Ge flakes. a) Schematic diagram of SHG measurement. b, c) Power-dependent SHG spectra and the spectra in logarithmic coordinates. d) SHG mapping of a typical Ge flake. e) Polarization-resolved SHG spectra. f) Optical image and Raman mapping of a flake composed of two grains. h) SHG mapping of the flake with boundary. Inset: illustration of the two edges growth at the boundary. i) Illustration of the armchair directions of the two grains. microscopy has been proposed to be particularly sensitive to structural configurations such as the lattice symmetry, defect, and grain boundary.38-41 Thus the SHG signals on Ge flakes were probed via back-reflecting geometry as illustrated in Figure 3a, and a femotsecond optical parametric oscillation (OPO) laser was employed as the excitation source with a tunable wavelength at a repitition rate of 80 MHz. The SHG signals become stronger with increasing power intensity as demonstrated in Figure 3b. Furthermore, the SHG intensity is proportional to the square of the power intensity as shown in Figure 3c (the fitted exponent is 1.92, which is very close to the theoretical value of 2), in consistent with the nonlinear optical principal. Since the polarization-resolved SHG is crystal-symmetry dependent, we performed the angledependent SHG intensity with excitation field parallel to emission field polarization. Before that, a SHG mapping was collected to ensure the high-uniformity of our tested sample in Figure 3d. The uniform intensity distribution confirms the pure phase and uniform thickness of our sample. The polarization-dependent SHG intensity in Figure 3e exhibits a sixfold rotational symmetry, and SHG intensity can be fitted with I SHG = I 0 cos2 (3θ ) , where θ was defined as the azimuthal angle between a mirror plane of crystal and the polarization of incident laser as illustrated, I0 is the maximum SHG intensity.42 It is observed that SHG intensity reaches a maximum value when the polarization locates along the armchair direction. The character of sixfold rotational symmetry can further confirm the (111) plane surface of the as-synthesized ultrathin Ge flakes, which is in accordance with the XRD, EBSD, and TEM characterizations. Furthermore, SHG microscopy has been proposed to be efficient to identify the grain boundary and the formation mechanism. Figure 3f and g exhibit the uniform optical image and Raman mapping of a flake composed of two grains, of which the boundary cannot be observed. Nevertheless, it is clearly visible in SHG mapping as prominent dark lines in Figure 3h when signals from the adjacent grain interfere destructively.40 The armchair directions of these two triangular grains can be represented by the yellow and blue arrow sets in Figure 3i. When two grains touch and continue to grow, the boundary orientation depends on the relative growth rates of the neighboring edges (inset of Figure 3h). In our case, the like edges merge with the same spacing, and the consequent boundary halves the two edges equal-

⋅ dVds WCiVds

length of the channel, Ci: the specific capacitance with 11.6 nFcm2 for 300 nm SiO2 layer).45 Thus the hole mobility is ~ 263 cm2V1 -1 s at Vds = 0.01 V, which surpasses many other CVD-grown 2D materials based transistors.46-48 Furthermore, the optoelectronic properties were explored. Figure 4d exhibits the incident power density-dependent photoresponse curves under 532 nm illumination and excellent photoresponse is presented. The excellent cycling stability is demonstrated in Figure 4e, illuminated by a 532 nm laser with the power density of 4.63 mW/cm2. Figure 4f shows the enlarged view of one typical photoresponse cycle indicating the fast rising and decay rates of ~ 70 and 6 ms, respectively. Besides, the responsivity Rλ can be represented by the equation: Rλ = I ph / PS , where Iph=Ilight - Idark (Vds = 0.1 V, Vg = 0 V), S is the illuminated effective area (~ 10 µm2), and P is the power density of illuminated light. Thus the Rλ can be calculated to be ~ 200 A/W, which is better than many other 2D materials based photodetectors.14, 49 The Ge-based phototransistor also shows excellent photoresponse under infrared light illumination (Figure S5 and S6). Then the external quantum efficiency can also be estimated to be ~ 1.0 × 105 % via the formula: EQE = hcRλ / eλ (c, λ, and e are the velocity of light, the light wavelength, and the unit electronic charge, respectively). Besides, the ability of a photodetector to probe the weak signals called detectivity (D*) is also a key index. According to the photoresponse in Figure 4e, the noise may come from both dark current and thermal noise. The thermal noise can be calculated by the formula: I nT = 4κ T ∆f / R , and the noise from dark current can be estimated by I nd = 2eI d ∆f , where k is the Boltzmann constant, T is the temperature, △f is the bandwidth, Id is the dark current, R is the effective resistance of the device. Thus the detectivity can be calculated to be ~ 1.06 × 1010 Jones by D * = ( S ∆f )1/ 2 R / I n , and I n = I nT 2 + I nd 2 , where S is the effective area of the device.50 Then we further performed the gate-dependent photoresponse, and the transfer characteristics in dark and under illumination (532 nm laser) are demonstrated in Figure 4g. It is observed that the photocurrent increases for both “OFF” and “ON”-states, suggesting that photocurrent dominates the whole operating range. Such character can be explained by the carrier transport mechanisms: i) large amounts of carriers will generate under light illumination in the “OFF-state”, resulting in photocurrent increasing; ii) in the “ON-state”, the photocurrent continues to increase due to the contact barriers lowered by the gate, leading to a more efficient photocurrent extraction and increased photoresponsivity (Figure 4g, h).34 Thus the responsivity of the phototransistor shows strong dependence of gate in Figure 4h, and can be further improved to 941 A/W at Vg = - 50 V (4.63



Figure 4. Electrical and optoelectronic performances. a) Illustration of a Ge flake-based phototransistor. b) Ids-Vds curves at different Vg in the dark. c) Transfer characteristics at different Vds in the dark. d) Ids-Vds curves under different power densities. e) Time-resolved photoresponse at Vds = 0.1 V, Vg = 0 V. f) One cycle of the time-resolved photoresponse. g) Transfer characteristics in the dark and under 532 nm laser illumination. h) Gatedependent responsivity illuminated by different power densities. i) Power density-dependent photocurrent fitted by the power law. mW/cm2), and 1473 A/W at Vg = - 50 V (0.42 mW/cm2). The comparison of the phototransistors based on 2D structures is summarized in Table S1. It is obvious that the overall performance of the as-synthesized ultrathin Ge flakes can rival or even surpass many other 2D layered and nonlayered materials. Furthermore, the power-dependent photoresponse was fitted by power law ( Iph ∝ Pθ ) to probe the trap states in Ge flake as demonstrated in Figure 4i, and the factor θ can be estimated to be 0.764. This sublinear dependence may be related to the defects such as Ge vacancies in the flake. Here, the factor θ is less than 1, indicating the exist of defects or vacancies in the as-synthesized Ge flake.34 Thus the high performance such as the fast response rates and the high responsivity may be partially explained by the highquality of the as-synthesized Ge flake, as well as the screening effects of the top passivation layer. Furthermore, the stability of the Ge flakes after stored in the ambient air has been investigated in Figure S7, indicating the excellent stability of the Ge flakes. CONCLUSION In conclusion, we have successfully realized the synthesis of ultrathin nonlayered Ge flakes via the KCl-assisted self-limited growth with the thickness down to ~ 8.5 nm. The growth mechanism has been confirmed by theoretical calculations and experiments, which can be attributed to the preferential growth of (111) plane with the lowest formation energy and the giant interface distortion effect of Cl-Ge motif. Excitingly, the Ge flake based phototransistor shows excellent performances such as a high hole mobility of ~ 263 cm2V-1s-1, high responsivity of ~ 200 A/W, and fast response rates (τrise=70 ms, τdecay=6 ms). These satisfactory results render these ultrathin Ge flakes highly promising for applications in electronics and optoelectronics. METHODS Synthesis of 2D Ge flakes: ultrathin Ge flakes were synthesized in a horizontal vacuum quartz tube furnace. 15 mg GeS and 2 mg KCl powders were mixed in a ceramic boat at the centre temperature zone. Silicon subsrates with 300 nm SiO2 were cutted into pieces of 1 × 1 cm and located at the downstream.

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Before heating, the furnace was evacuated and flushed with Ar to exhaust the O2 in the system. Then the furnace was heated to 550 °C in 20 minutes and the growth time was maintained for 30 minutes with 50 sccm Ar and 10 sccm H2 flowing through the system. The system was naturally cooled after the growth process. Characterizations: the morphology, crystal structure, composition of Ge flakes were characterized by an optical microscope (BX51, OLMPUS), atomic force microscope (Dimension icon, Bruker), XRD (XRD-7000, Shimadzu), confocal Raman/PL system (Alpha 300RS+, WITec), TEM (Tecnai G2 F30, FEI) equipped with an X-ray energy dispersive spectrometer, and XPS (AXIS-ULTRA DLD-600W, Shimadzu). Additionally, EBSD (TEAM Pegasus, EDAX) installed on a scanning electron microscope (Quanta 650 FEG, FEI) was performed to identify the nature of Ge flake. Device Fabrication and Performance: the dry-transfer method51 has been employed to avoid the negative effects of ion solution used in the wet etching method. The phototransistors were fabricated by a standard electron-beam lithography (Quanta 650 SEM, FEI and ELPHY Plus, Raith GmbH). Then the 10 nm Cr/50 nm Au electrodes were deposited by a thermal evaporation system (Nexdap, Angstrom Engineering). The electrical tests were performed in a probe station (CRX-6.5 K, Lake Shore) connected with a semiconductor system (B1500A, Keysight). A laser with fixed wavelength of 532 nm and tunable power intensity was employed as the light source for photodetection test. DFT Calculations: The projector augmented wave method was employed to implement the DFT calculations, as performed in the Vienna Ab Initio Simulation Package (VASP).23, 52 We adopted the generalized gradient approximation (GGA) with the Perdew-Burke-Ernzerhof exchange-correlation functional53 and set a plane-wave energy cutoff to 500 eV. The conjugate gradient (CG) algorithm was employed to completely release all the atoms until the residual forces and the difference of maximal energy were less than 10-3 and 5 × 10-6 eV/Å. We employed 7 × 7 × 1 and 7 × 7 × 7 Monkhorst-Pack k-point meshes for surface structures and bulk counterparts of Ge in Brillouin zone. In order to exclude the interaction between neighbouring supercells’s periodic images, the lateral dimensions were set to be sufficiently large (~15 Å). While the outmost four layers went through a completely atomic relaxation, the innermost five layers were fixed as bulk counterpart during the optimization of flake structure. The surface formation energy was evaluated by the energy difference between Ge flake with N layer and its counterpart with N + 1 layer for (111) and (100) surface. The ClGe defect distribution was depicted by substituting Ge atom with Cl from inner bulk phase to outermost layer. The Ge layer structures were drawn using the VESTA software.54

ASSOCIATED CONTENT Supporting Information Experimental details and supplementary figures (PDF).

AUTHOR INFORMATION Corresponding Authors *H. Q. Li: [email protected] *X. Zhou: [email protected] *T. Y. Zhai: [email protected]

Notes The authors declare no competing financial interest.

ACKNOWLEDGMENT


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Journal of the American Chemical Society This work was supported by the National Natural Science Foundation of China (Grant No. 2182500167, 51727809, 91622117, 11774239, 51722202 and 51472097), the National Key Research and Development Program of “Strategic Advanced Electronic Materials” (Grant No. 2016YFB0401100), the National Basic Research Foundation of China (Grant No. 2015CB932600), and the Fundamental Research Funds for the Central University (Grant No. 2015ZDTD038, 2017KFKJXX007). The authors also thank the technical support from Analytical and Testing Center in Huazhong University of Science and Technology.

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