Nanoscale

PAPER

Cite this: Nanoscale, 2016, 8, 17598

Received 27th July 2016,Accepted 15th September 2016

DOI: 10.1039/c6nr05937f

www.rsc.org/nanoscale

The influence of interfacial tensile strain on thecharge transport characteristics of MoS2-basedvertical heterojunction devices†

Fu Huang,‡a,b Byungjin Cho,‡c Hee-Suk Chung,a Seung Bae Son,a Jung Han Kim,a,d

Tae-Sung Bae,a Hyung Joong Yun,e Jung Inn Sohn,f Kyu Hwan Oh,d

Myung Gwan Hahm,g Jung Hee Parkb and Woong-Ki Hong*a

We demonstrate the charge transport characteristics of MoS2-based vertical heterojunction devices

through the formation of interfacial strain. Atomically thin MoS2 bilayers were directly synthesized on a

p-type Si substrate by using chemical vapor deposition to introduce an interfacial tensile strain in the ver-

tical heterojunction diode structure, which was confirmed by Raman, X-ray and ultraviolet photoelectron

spectroscopy techniques. The electrical and optoelectronic properties of the heterojunction devices with

the as-grown MoS2 (A-MoS2) on p-Si were compared with those of transferred MoS2 (T-MoS2)/p-Si

devices. To clearly understand the charge transport characteristics induced by the interfacial tensile strain,

the Fowler–Nordheim (FN) analysis of the electrical properties of the diode devices was conducted with

the corresponding energy band diagrams. All of the fabricated MoS2-based vertical diodes exhibited

clearly rectifying behaviors, but the photoresponse properties of the A-MoS2-based and T-MoS2-based

heterojunctions exhibited distinct differences. Interestingly, we found that the tunneling barrier heights of

the A-MoS2-based heterojunction devices were relatively higher than those of the T-MoS2-based devices

and were almost the same before and after illumination due to the interfacial tensile strain, whereas those

of the T-MoS2-based devices were lowered after illumination. Our study will help further understand the

charge transport properties of 2D material-based heterojunction devices in the presence of interfacial

strain, ultimately enabling the design of electronic and optoelectronic devices with novel functionalities.

Introduction

Two-dimensional (2D) atomic layers, such as graphene, boronnitride, and especially transition metal dichalcogenides(TMDs), have garnered considerable interest as the thinnest

building blocks for next-generation electronic and opto-electronic devices.1–6 Unlike conventional three dimensional(3D) materials, 2D TMDs with weak van der Waals interactionsbetween the adjacent layers have prompted a wide range of1D–2D,7,8 2D–2D,6,9–12 2D–3D,3,13–15 and 2D–organic hybrid16–18

heterojunction architectures due to both the wide tunability oftheir physical properties and their easy integration capability. Ap–n heterojunction based on 2D TMDs is practically usefulin implementing electronic and optoelectronic applicationssuch as transistors, photovoltaic cells, and light emittingdevices.1,3,12,14,15,19 For example, Lee et al.1 demonstrated a gate-tunable photovoltaic response under white-light illuminationfor van der Waals-stacked MoS2/WSe2 heterojunction devices.Ajayan and co-workers proposed a new 3D band diagram for theheterojunction formed between n-type monolayer MoS2 andp-type Si, which shows the flow of charge carriers inside thedevice in a 3D manner.14 They also demonstrated an atomicallythin optoelectronic memory array for image sensing withlayered Cu–In–Se and MoS2 atomic layers.19 Li et al.15 alsoreported on electric-field-induced strong electroluminescence inmultilayer MoS2-based vertical heterojunctions.

†Electronic supplementary information (ESI) available. See DOI: 10.1039/c6nr05937f‡These authors contributed equally to this work.

aJeonju Center, Korea Basic Science Institute, Jeonju, Jeollabuk-do 54907,

Republic of Korea. E-mail: wkh27@kbsi.re.krbDivision of Biotechnology, Advanced Institute of Environment and Bioscience,

College of Environmental and Bioscience Sciences, Chonbuk National University,

Iksan 54596, Republic of KoreacDepartment of Advanced Functional Thin Films, Surface Technology Division, Korea

Institute of Materials Science, Changwon, Gyeongnam 51508, Republic of KoreadDepartment of Materials Science and Engineering, Seoul National University,

Seoul 08826, Republic of KoreaeAdvanced Nano Surface Research Group, Korea Basic Science Institute,

Daejeon 34133, Republic of KoreafDepartment of Engineering Science, University of Oxford, Oxford OX1 3PJ, UKgSchool of Materials Science and Engineering, Inha University, Incheon 22212,

Republic of Korea

17598 | Nanoscale, 2016, 8, 17598–17607 This journal is © The Royal Society of Chemistry 2016

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Importantly, the presence of interfaces for heterojunctionsplays a critical role in charge transport properties for electronicand optoelectronic devices due to the nature of the interface,including the doping levels, defect states, strain and theenergy band-edge alignments between two semiconductors. Inparticular, one recent prominent issue is the straining of 2DTMD-based materials, which modulate the energy band gap,affecting the electronic and optoelectronic properties.20–28

Most recently, Ajayan and co-workers29 reported the electronicstructure and its strain dependence on stacked MoSe2/WSe2heterostructures directly synthesized by chemical vapor depo-sition. Though there are many studies regarding p–n hetero-junctions, fundamental and systematic studies on strai-n-induced charge transport behaviors across 2D/3D p–nheterojunctions are rare. It would be significant to explorephysical phenomena appearing in the 2D TMD-based hetero-junctions in the presence of interfacial strain, ultimatelyenabling the strategic design of electronic and optoelectronicdevices with novel functionalities.

In this study, we investigate the charge transport character-istics of MoS2-based vertical heterojunction devices throughthe formation of interfacial strain. To this end, atomically thinMoS2 bilayers were directly synthesized on a p-Si substrate forthe formation of an interfacial strain and bilayer MoS2 syn-thesized on a SiO2 substrate was transferred onto p-Si torelease the strain; the results of these processes are hereafterreferred to as A-MoS2 and T-MoS2, respectively. The character-ization results by Raman spectroscopy, X-ray photoelectronspectroscopy (XPS), and ultraviolet photoelectron spectroscopy(UPS) revealed distinct differences between A-MoS2 andT-MoS2 in terms of their strain, stoichiometry, and electronic

band structures. To clearly understand the effect of tensilestrain at the interface between the MoS2 and p-Si on verticaljunction charge transport, Fowler–Nordheim (FN) analysis wasconducted with the corresponding energy band diagrams.Interestingly, we found that the tunneling barrier heights ofthe A-MoS2-based heterojunction devices were relatively higherthan those of the T-MoS2-based devices and were almost thesame before and after illumination, whereas those of theT-MoS2-based device were lowered after illumination. Thisresult indicates that the interfacial strain at the junction canstrongly affect the tunneling current through the ultrathinMoS2 layer in charge transport for vertical TMD-based hetero-junction devices. Note that a complete description of theexperimental details can be found in the Experimental section.

Results and discussion

Fig. 1a shows schematic illustrations of the vertical heterojunc-tion structures for A-MoS2 (top) and T-MoS2 (bottom) on a p-Sisubstrate. Fig. 1b displays typical Raman spectra with twodominant peaks of E12g and A1g for the A-MoS2 and T-MoS2layers (see Fig. S1 in the ESI† for the statistical distribution).The E1

2g and A1g modes of MoS2 are related to the in-planevibration of molybdenum (Mo) and sulfur (S) atoms, and theout-of-plane vibration of S atoms, respectively.30 Raman spec-troscopy has been widely utilized as a useful tool for probingthe intriguing physical properties such as the lattice strain, thenumber of layers, doping levels, and van der Waals interactionat the interface for 2D crystals.2,25,26,31–34 In Fig. 1b and S1,†the typical Raman spectra with two dominant peaks are

Fig. 1 (a) Schematic illustrations of the p–n heterojunction based on the as-grown and transferred MoS2 bilayers (A-MoS2 and T-MoS2) on a p-Sisubstrate. (b) Representative Raman spectra for the A-MoS2 and T-MoS2 layers on a p-type Si substrate under 488 nm laser lines. The spectra reveala strong in-plane vibrational mode for the Mo and S atoms (E12g) and an out-of-plane vibrational mode for the S atoms (A1g). (c, d) Statistical histo-gram analysis of the frequency difference between the E12g and A1g on the (c) A-MoS2 and (d) T-MoS2 films. (e) Experimental Raman frequencies(open symbols) for the indicated phonon modes as a function of tensile strain. Solid lines are well fitted by the linear functions with slopes of−4.48 cm−1 and −1.02 cm−1 per 1% of biaxial tensile strain for the E12g and A1g modes, respectively, which are well consistent with ref. 24. (f )Comparison of the frequency difference as a function of the number of MoS2 layers between our MoS2 films and reference data. (g, h)Cross-sectional scanning transmission electron microscope annular dark field (STEM ADF) images of the (g) A-MoS2 and (h) T-MoS2 films on thep-Si substrate.

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featured at 382.7 ± 0.6 (E12g) and 406 ± 0.6 cm−1 (A1g) for A-MoS2and ∼384.7 ± 0.6 (E1

2g) and ∼406.1 ± 0.7 cm−1 (A1g) for T-MoS2.It is obvious that the A1g peak of A-MoS2 shows a negligiblered-shift of ∼0.1 cm−1, whereas the E12g peak of A-MoS2 showsa significant red-shift of ∼2.0 cm−1, compared to those ofT-MoS2 (see Fig. S1 in the ESI†). Because the E12g in-planevibration mode is highly sensitive to the built-in tensile strainof 2D MoS2, the red-shift of the E12g peak position for A-MoS2can be attributed to tensile strain effects by the thermal mis-match between MoS2 and the substrate.2,22,23,25,26,33 In thecase of the T-MoS2-based heterojunction device in which MoS2was synthesized on a SiO2/Si substrate and then transferredonto the p-type Si substrate, the tensile strain can be spon-taneously released during the transfer process.22,23 Fig. 1c andd show the statistical distribution of the frequency difference(A1g–E

12g) for the A-MoS2 and T-MoS2 layers, respectively. The

magnitude of the frequency difference between the E12g andA1g peaks reveals the layer number of the MoS2.

37–40 From themagnitude of the red-shift in our results, the biaxial tensilestrain on MoS2 directly synthesized on the p-type Si substratecan be estimated as ∼0.398% (Fig. S1 in the ESI†),34–36 whichis comparable to previously reported strain values.2,22 This isalso consistent with theoretical Raman frequencies as a func-tion of strain as reported in ref. 24. Fig. 1e shows the shifts oftwo main peaks (E1

2g and A1g) as a function of biaxial tensilestrain. Interestingly, the E12g and A1g modes of the Ramanpeaks for A-MoS2 and T-MoS2 are well fitted to linear curveswith slopes of −4.48 and −1.02 cm−1 per 1% biaxial strain,respectively. As shown in Fig. 1f, compared to previousreports,37–39 it is reasonable for our MoS2 layers to be assignedto the bilayer thickness despite the frequency difference of∼2 cm−1 between A-MoS2 and T-MoS2 due to biaxial tensilestrain. To verify that both types of MoS2 are bilayer, TEMcharacterization was carried out for A-MoS2 and T-MoS2, as

shown in Fig. 1g, h and S2.† In the case of A-MoS2, weobserved clearly a thin SiO2 layer at the interface between theMoS2 bilayer and the p-Si substrate, compared to T-MoS2. Theformation of the thin SiO2 layer on the p-Si substrate likelyoriginates from the synthesis of 2D MoS2 films through thesulphurization of the MoO3 film (see Experimental details inthe ESI†). This could result in making the A-MoS2 layer be intensile strain.22,23

To study element compositions for A-MoS2 and T-MoS2 onthe p-type Si substrate, we examined XPS. Fig. 2a and b displaythe XPS spectra of Mo and S binding energies for A-MoS2 andT-MoS2, respectively, where the XPS spectra are deconvolutedinto peaks that correspond to the Mo 3d, S 2s, and S 2p com-ponents of the MoS2. From the XPS spectra (Fig. 2a and b), thepeaks at 233.1 and 229.9 eV (232.4 and 229.2 eV) represent theMo 3d3/2 and Mo 3d5/2 for the A-MoS2 (T-MoS2) sample,respectively, and the peak at 226.8 (226.4) eV corresponds toS 2s for A-MoS2 (T-MoS2). In addition, the peaks at 163.8 and162.7 eV (163.3 and 162.1 eV) for A-MoS2 (T-MoS2) are relatedto the S 2p1/2 and S 2p3/2 orbitals of divalent sulfide ions (S2−),respectively. The deconvoluted XPS peaks (Mo 3d3/2, Mo 3d5/2,S 2p1/2, S 2p3/2, and S 2s) also reveal the difference of theFWHMs for A-MoS2 and T-MoS2 (Fig. 2c). Although the XPSresults of A-MoS2 and T-MoS2 are slightly different, these XPSdata are consistent with the results of previous studies (seeTable S1 in the ESI†). The different peak positions betweenA-MoS2 and T-MoS2 are likely due to nonstoichiometry, struc-tural defects, and impurities or interface charge transferinduced by the tensile strain, which can result in the workfunction (WF) difference.41–43 This is also strongly supportedby a plot of MoSx stoichiometry as a function of (Mo 3d5/2–S2p3/2) binding energy (Fig. 2d). Here, it is possible to deter-mine the MoSx stoichiometry to an accuracy of x ± 0.1 by usingthe Mo and S peak positions.44 Fig. 2e and f show the Si 2p

Fig. 2 XPS spectra showing Mo 3d, S 2s, and S 2p core level peaks for the (a) A-MoS2 and (b) T-MoS2 films on the p-Si substrate. (c) FWHMs ofMo 3d3/2, Mo 3d5/2, S 2p1/2, S 2p3/2, and S 2s for A-MoS2 and T-MoS2. (d) MoSx stoichiometry as a function of the (Mo 3d5/2–S 2p3/2) binding energyfor A-MoS2 and T-MoS2. The data show a linear relationship, which is consistent with the open symbol data based on ref. 44. (e, f ) Si 2p XPS corelevel peaks for (e) A-MoS2 and (f ) T-MoS2 films on the p-Si substrate.

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core level XPS spectra of A-MoS2 and T-MoS2, respectively. It isclear that the Si 2p core level peak in SiO2 of the A-MoS2sample is stronger than that of the T-MoS2 sample, indicatingthat the A-MoS2 sample undergoes a stronger oxidation ofthe Si surface. These observations are in excellent agreementwith the results of the TEM characterization shown inFig. 1g and h.

UPS analysis was performed to study the electronic struc-tures determined by the tensile strain and composition vari-ation for vertical heterojunctions based on A-MoS2 andT-MoS2. The WF is defined as the energy difference betweenthe vacuum level (Evac) and the Fermi level (Ef ). Accordingly,from the UPS measurements, the WF can typically be calcu-lated from the difference between the secondary electroncut-off of the highest binding energy and the exciting photonenergy of 21.22 eV (Fig. 3a) (see also Table S2 in the ESI†). Thevalence band maximum (VBM) can also be determined fromthe cut-off of the lowest binding energy (Fig. 3b) (see Table S2in the ESI†). As shown in Fig. 3 and Table S2,† the WF is 4.25and 4.2 eV for A-MoS2 and T-MoS2, respectively, which is com-parable to the values presented in previous reports.45–50 TheWF difference of 50 meV between A-MoS2 and T-MoS2 orig-inates from the strain effect,51 which can modify the Ef ofMoS2 by changing the positions of energy band edges. Anincrease in the WF implies a shift of the Ef below the conduc-

tion band, which can be closely related to the shift of the core-level binding energy in XPS (Fig. 2). In Fig. 3b, the position ofthe VBM below the Fermi level (Ef ) implies that our samplesare n-type semiconductors.43,50 The energy difference betweenthe Ef and VBM is ∼1.3 and ∼1.43 eV for A-MoS2 and T-MoS2(Fig. 3b), respectively, indicating tensile strain-inducedreduction of the energy bandgap from A-MoS2.

23–25,28 The WFvalues of MoS2 in previous reports and our work are summar-ized in Table S2.† Based on the UPS results, the energy bandalignment of the MoS2-based heterojunctions under an equili-brium can be constructed, as shown in Fig. 3c and d. Whenthe electron affinity of MoS2 is ∼4.1 eV,52 the calculated bandgap is 1.45 and 1.53 eV for the A-MoS2 and T-MoS2 bilayersamples, respectively. Assuming that T-MoS2 is tensi-le-strain-free, the energy gap of 1.53 eV is comparable to thevalues calculated using the DFT-PBE calculation in ref. 35 (alsosee the calculation of the decreased rate of the energy gapunder biaxial tensile strain in the ESI† for details). Here, wecan estimate that the decreased rate of the energy gap underbiaxial tensile strain is approximately −0.2 eV per % strain.

Next, the electrical characteristics of the A-MoS2- andT-MoS2-based heterojunctions were examined to investigatethe charge transport behavior under dark and light illumina-tion conditions, as shown in Fig. 4a. Fig. 4b and c show therepresentative semi-logarithmic current–voltage (I–V) curvesunder dark and light illumination with white-light power den-sities of 0.51 and 16.3 mW cm−2 for the A-MoS2-based andT-MoS2-based heterojunction devices, respectively. Under darkconditions, both devices presented typical rectification charac-teristics in which the current at reverse bias increased with thelight power density, whereas the current change at forwardbias was negligible. In particular, the current of theA-MoS2-based heterojunction device increased up to ∼7 nA atreverse bias, which is likely due to the scattering and trappingof the carriers in the SiO2 layer, whereas that of theT-MoS2-based heterojunction device was in the range of ∼1 nA.When the light power density increased by 16.3 mW cm−2, thereverse photocurrent further increased by approximately ∼3 μAat a bias of −5 V. From the I–V characteristics of the p–ndevices under dark conditions, we extracted ideality factors (IF,n) and turn-on voltages (see Fig. S4 in the ESI†). The IF of theA-MoS2-based heterojunction and the T-MoS2-based hetero-junction devices was calculated to be 2.9 ± 0.7 and 4.8 ± 1.7,respectively (see Fig. S4a and c in the ESI†). The high IF values(over 2) may arise from defects or impurities at the MoS2/Siheterointerface and/or multiple junctions device structure.53,54

The relatively higher IF of the T-MoS2-based heterojunctiondevices might be attributed to polymer and water residues atthe T-MoS2/Si interface which were caused by a transferprocess.12 The average turn-on voltages of the A-MoS2-basedand T-MoS2-based heterojunction devices were estimated to be3.2 ± 0.5 and 1.8 ± 0.5 V, respectively (Fig. S4b and d in theESI†), indicating that the barrier height for charge transport inthe A-MoS2-based heterojunction is relatively much higher.

From the statistical data on the photocurrent as a functionof light intensity (Fig. 4d and e), the forward photocurrent was

Fig. 3 UPS spectra including (a) the secondary electron cut-off of thehighest binding energy and (b) the cut-off of the lowest binding energyfor the A-MoS2 and T-MoS2 films on the p-Si substrate. (c, d) Energyband diagrams before and after contact for the p–n heterojunctiondevices based on the (c) A-MoS2 and (d) T-MoS2 films. EF is the Fermilevel energy, Φ is the work function, χ is the electron affinity, Ev is thevalence band energy, Ec is the conduction band energy, and D is thedifference between EF and Ec.

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almost the same regardless of the light power (Fig. 4d),whereas the photoresponsive behavior at the reverse voltagewas strongly power-dependent (i.e., the photocurrent increaseswith light power) (Fig. 4e). Interestingly, compared to theA-MoS2-based devices, the T-MoS2-based devices showed rela-tively higher current levels at forward bias (Fig. 4d) and higherphotocurrent at reverse and forward biases (Fig. 4d and e). Inaddition, the T-MoS2-based devices exhibited relatively higherphotoresponsivity at the reverse bias than that of theA-MoS2-based devices (Fig. S5, S6 and Tables S3 and S4 in theESI†). We also studied the transient photocurrent (Fig. 4f) andthe photovoltaic effect of the p–n heterojunction diodes(Fig. S7 and Table S5 in the ESI†). In Fig. 4f, the photocurrentratio (Ilight/Idark) of the photodiodes based on A-MoS2 andT-MoS2 at −5 V at the light intensity of 16.3 mW cm−2 was∼150 and ∼1000, respectively (see also Fig. S6a in the ESI†).Intriguingly, we observed the photocurrent decay phenomenonof the A-MoS2-based heterojunction device under light illumi-nation while it was negligible on the T-MoS2-based heterojunc-tion device. This result can be attributed to the difference ofcharge trap states or charge recombination at the interfacebetween MoS2 and p-Si for both heterojunctions, which maybe mainly induced by an interfacial tensile strain through theformation of the SiO2 thin layer at the A-MoS2/p-Si inter-face.14,19,22,23 According to ref. 55, the trap density can bequalitatively calculated from the dynamic curve of photo-switching although the calculation in our work is rough (seeFig. S8 and Table S6 in the ESI†). The calculated trap densityof the A-MoS2-based and T-MoS2-based devices is approxi-mately 16.2 × 1015 and 12.2 × 1015 cm−2, respectively. Thetrapped charge density of T-MoS2 is lower than that of A-MoS2.

This can be attributed to the increase in positively chargedtrap sites due to thicker silicon oxide at the MoS2–Si interfacefor the A-MoS2 based device.56,57 Consequently, the amount ofthe decreasing current of the T-MoS2-based device during thelight illumination is lower than that of the A-MoS2-baseddevice. For the A-MoS2-based heterojunction, tensile strainarising from thermal expansion mismatch and the SiO2 layerformation during the direct CVD-synthesis of MoS2 on the p-Sisubstrate can create charge trap sites such as disorder, defects,and S vacancies during the direct synthesis of MoS2 on the p-Sisubstrate.55,58–60 Additionally, the photovoltaic effect forA-MoS2 is smaller than that for T-MoS2, as shown in Fig. S7.†The photovoltaic parameters of the 10 fabricated devices arelisted in Table S5.† From Fig. S7 and Table S5,† theA-MoS2-based heterojunction devices showed larger open-circuit voltages (VOC) and smaller short-circuit currents (ISC)and fill factors (FF) compared to those of the T-MoS2-basedheterojunction devices. This can be attributed to the carrierloss due to the funnel effect20,24,28 during the charge transportfor A-MoS2 under light illumination, which results from theinterfacial tensile strain across the A-MoS2-based heterojunc-tion. This result implies that the interfacial strain can stronglyaffect photonic applications such as photodiodes and photo-voltaic cells.

The charge transport across the vertical heterojunctions(A-MoS2-based or T-MoS2-based heterojunction devices) underdark and light illumination can be understood by the energyband diagrams illustrated in Fig. 5 (see also Fig. S6 in theESI†). Fig. 5a and c display the energy band diagrams atforward and reverse biases under dark conditions for theA-MoS2-based and T-MoS2-based heterojunctions, respectively.

Fig. 4 (a) Schematic illustration of the p–n heterojunction devices based on the A-MoS2 (top) and T-MoS2 (bottom) films (TE: top electrode, BE:bottom electrode). (b, c) Representative photoresponsive characteristics as a function of light intensity for the heterojunction devices based on the(b) A-MoS2 and (c) T-MoS2. (d, e) Photocurrents of the A-MoS2-based and T-MoS2-based heterojunction devices at a (d) forward voltage of 5 V and(e) reverse voltage of −5 V. (f ) Photoswitching characteristics of the A-MoS2-based and T-MoS2-based heterojunction devices at −5 V under thelight power of 16.3 mW cm−2.

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Fig. 5b and d illustrate the energy band diagrams under lightillumination, depicting photo-excited charge carriers, chargeseparation, and charge transport. For the A-MoS2-based hetero-junction (Fig. 5a, b and S9a†), the conduction band (Ec) andvalence band (Ev) edges at the interface between p-Si and

A-MoS2 are unlikely to be modulated at both forward andreverse biases due to the interfacial tensile strain effectthrough the formation of a SiO2 buffer layer (Fig. 5a(i), (ii) andS9†), which results in the reduction of the band gap andcharge carrier trapping.20,28,53,54 In contrast, the band edges

Fig. 5 Schematic illustrations of energy band diagrams for the vertical MoS2-based heterojunctions under dark and light illumination conditions.The energy band structures of (a, b) A-MoS2-based and (c, d) T-MoS2-based heterojunctions under dark and light illumination. The dashed and solidlines present band structures at the equilibrium state and specific conditions (biasing, illumination), respectively. In particular, for the T-MoS2-basedheterojunctions, the tunneling barrier heights are changed at forward and reverse biases under illumination in (d). The yellow- and red-solid lines in(d) present the energy bands in T-MoS2 before and after light illumination, respectively. The arrows at the forward bias indicate that the tunnelingcurrent through the MoS2 bilayer is the dominant component in charge transport for the two types of vertical MoS2-based heterojunction devices.

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with a rather weak interaction (physisorption bonding) at theinterface between MoS2 and Al 61,62 can be effectively modu-lated by using forward and reverse biases (Fig. 5a(i) and (ii)).For the T-MoS2-based heterojunction (Fig. 5c, d and S9b†), theband edges at the interfaces between Al/T-MoS2 and T-MoS2/p-Si can be strongly modulated by forward and reverse biases.This finding implies that during charge transport, carrier lossor interface charge transfer across the A-MoS2-based hetero-junction device is much more dominant than that for theT-MoS2-based heterojunction device, as shown in Fig. 4. Theformation of the SiO2 layer can decrease the effectiveness ofthe carrier tunneling due to the scattering and trapping of thecarriers. Under illumination, because there is band-bendingnear the p-Si/MoS2 interface for both devices, the separation ofphoto-excited carriers within the depletion layer of p-Si couldoccur and then photo-excited holes move to the p-Si/MoS2interface (Fig. 5b(i), d(i) and S9†). The photo-excited electronsgenerated in the MoS2 layer move to the p-Si/MoS2 interfaceand finally cause charge recombination with the photo-excitedholes in p-Si (Fig. 5b(i), d(i) and S9†). Therefore, at the forwardbias under illumination, the contribution of the photo-excitedcarriers to the electrical current was negligible, as shown inFig. 4. In particular, the photoelectrons in the T-MoS2 layercan modulate the band edges at the interfaces by increasingthe charge density within the MoS2 layer, which can result inSchottky barrier reduction for the carrier injection from theelectrodes at the forward bias and, thus, an increase of thetunneling current through the MoS2 layer (Fig. 5b(i), d(i) and

S9†). In contrast, at the reverse bias under illumination(Fig. 5b(ii) and d(ii)), the photo-excited electrons and holes inboth Si and MoS2 are well separated and moved to the respect-ive electrode, resulting in a large photocurrent (Fig. 5c(ii) andd(ii)) and a photovoltaic effect (Fig. S7 in the ESI†).Consequently, the charge transport through the ultrathinMoS2 layer for both vertical MoS2-based heterojunction devicesis governed by the tunneling current, which will be deeply dis-cussed in Fig. 6.

Fig. 6a and b show the Fowler–Nordheim (FN) plots atforward bias for vertical heterojunction devices based onA-MoS2 and T-MoS2, respectively. The FN plot, which showsthe relationship between the FN tunneling (FNT) current (IFNT)and the applied bias voltage (V), is expressed by the followingequations:63–65

IFNT ¼ Aeffe3m0V2

8πhϕBd2m*exp

�8πffiffiffiffiffiffiffiffiffi2m*

pϕB

3=2d3heV

� �ð1Þ

lnIFNTV2

� �¼ ln

Aeffe3m0

8πhϕBd2m*

� �� 8π

ffiffiffiffiffiffiffiffiffi2m*

pϕB

3=2d3heV

ð2Þ

where Aeff is the effective contact area, ϕB is the barrier height,m0 is the free electron mass, e is the electronic charge, h isPlanck’s constant, d is the thickness of the bilayer MoS2, andm* (0.45m0)

66 is the effective electron mass of MoS2. Here, thelinear dependence at the applied high bias indicates FNT,whereas the logarithmic growth at the low bias means directtunnelling. From the FN plots, we found that the FNT barrier

Fig. 6 The Fowler–Nordheim (FN) plots at forward bias for the (a) A-MoS2-based and (b) T-MoS2-based heterojunctions under dark and light illumi-nation and the corresponding energy band diagrams. The blue lines on the FN plots indicate linear fitting curves based on eqn (2). The tunnelingbarrier height (ϕB) was estimated from the slopes of the linear fits in the FN plots. The yellow dashed-line and red solid lines on the bottom panel of(b) present the energy bands in T-MoS2 at forward bias before and after light illumination, respectively. The arrows on the bottom panels indicatethat the tunneling current through the MoS2 bilayer is the dominant component in charge transport for the vertical MoS2-based heterojunctiondevices.

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heights (ϕB) of A-MoS2 and T-MoS2-based heterojunctions havedistinct differences before/after illumination. In the case ofA-MoS2 (Fig. 6a), the ϕB,A-MoS2 value exhibited a negligiblechange before/after illumination, while the ϕB,T-MoS2 value ofthe T-MoS2 device exhibited a considerable decrease from 0.39± 0.1 to 0.35 ± 0.1 eV (Table S7 in the ESI†). As shown in Fig. 6,for A-MoS2 (Fig. 6a, bottom panel), the charge recombinationbetween the photo-excited electrons in MoS2 and photo-excitedholes in Si can be much larger due to the tensile strain effectin the MoS2 layer and the scattering and trapping of the car-riers in the SiO2 layer than that in T-MoS2 (Fig. 6b, bottompanel). Consequently, unlike A-MoS2, the photo-excited elec-trons in the T-MoS2 layer can be preferably accumulated at theT-MoS2/p-Si interface where the accumulated electronsincrease the Fermi level in the MoS2 layer and thereby reducethe tunnelling barrier height.

Conclusions

In summary, we have demonstrated the strain-induced chargetransport characteristics of MoS2-based vertical heterojunctiondevices. An atomically thin MoS2 bilayer was directly syn-thesized on a p-Si substrate (A-MoS2) for the formation of aninterfacial tensile strain, and the bilayer MoS2 synthesized onan SiO2 substrate was transferred onto p-Si (T-MoS2) to releasestrain created by the interaction between the SiO2 substrateand MoS2. The electrical and optoelectronic properties of theMoS2-based heterojunction devices were evaluated usingcurrent–voltage characteristics and photoresponse properties.To understand junction charge transport, FN analysis was con-ducted with the corresponding energy band diagrams. Wefound that the FN tunneling behaviors of A-MoS2 and T-MoS2exhibited a distinct difference. For the A-MoS2-based hetero-junction devices, the tunneling barrier heights at forward biaswere constant regardless of the light illumination conditionsand relatively higher rather than those of T-MoS2-based hetero-junction devices. This is attributed to the carrier loss thatoccurs across the A-MoS2-based heterojunction as a result ofthe interfacial tensile strain. Our study will help further theunderstanding of the charge transport properties in 2D materi-al-based heterojunction devices and might provide insight intothe strategic design of 2D heterojunction devices via interfacialstrain engineering.

Experimental sectionFabrication of as-grown MoS2 on p-Si

P-type Si wafers were cleaned by using DI water, acetone, andisopropyl alcohol and then dried by blowing nitrogen gas.MoS2 ultrathin layers were directly synthesized on the preparedp-Si substrate using chemical vapor deposition (CVD) asdescribed in our previous work.67 Specifically, MoO3 films(∼3 nm) were deposited on the p-Si substrates using a thermalevaporator. The pre-deposited MoO3 samples were loaded into

the center of a furnace and then 1 g of S powder was loadednear the inlet of the furnace. The sublimation of S could becontrolled by using an independent flange heater. The furnaceand the flange heater were heated to 850 and 180 °C for1 hour, respectively. The process lasted for an additional1 hour under flowing Ar (85%)/H2 (15%) mixed gas at achamber pressure of ∼800 torr. Eventually, the MoO3 filmcould be sulphurized to 2D MoS2 layered films. Lastly, thefurnace was naturally cooled to room temperature.

Fabrication of transferred MoS2 on p-Si

After synthesizing the MoS2 film on a SiO2/p-Si substrate, theas-grown MoS2 ultrathin layers were transferred onto anotherp-Si substrate. Specifically, the polymethyl methacrylate(PMMA, MicroChem) polymer was spin-coated onto a MoS2film that was synthesized on the SiO2/p-Si substrate. The filmwas then immersed on 5% HF etching solution at room temp-erature in order to etch the SiO2 layer and then thePMMA-supported MoS2 film was transferred onto another p-Sisubstrate after DI cleaning processes. The PMMA polymer filmwas removed with acetone and isopropyl alcohol solvents.Finally, the transferred MoS2 film was baked at 90 °C for10 min in order to remove the solvent residue.

Formation of metal electrodes of the vertical heterojunctiondevices

For the formation of electrical contacts of p–n heterojunctiondevices, Al (100 nm)/Au (50 nm) electrodes were deposited onthe as-grown MoS2/p-Si and transferred MoS2/p-Si samplesusing a shadow mask and a thermal evaporator (see Fig. S3 inthe ESI†).

Device characterization

All electrical transport measurements were performed at roomtemperature under ambient conditions using a parameter ana-lyzer (Keithley 4200-SCS).

Material characterization

Raman spectra were obtained using laser Raman microspectro-scopy (Nanofinder 30, Tokyo Instrument Inc.) at a laser exci-tation wavelength of 488 nm in which a laser power of∼0.15 mW was used to avoid the laser heating effect on theRaman shift of MoS2. X-ray photoelectron spectroscopy (XPS)and ultraviolet photoelectron spectroscopy (UPS) were per-formed using a XPS/UPS system (AXIS-NOVA, Kratos Inc.). ForTEM characterization, cross-sectional TEM specimens ofA-MoS2 and T-MoS2 were prepared by focused ion beam (FIB)milling and lift-out techniques. (Detailed description for FIB(FEI NOVA200) work: first, an e-beam (5 keV, 98 pA)-induced Ptcapping layer was deposited to protect the target area and a Gaion-beam (30 keV, 50 pA)-induced Pt capping layer was sub-sequently deposited. After capping layer deposition, Ga ionmilling was performed until the target area became electro-n-transparent (∼70 nm thick) suitable for TEM imaging andthe prepared specimen was put onto a Cu TEM grid with amicromanipulator inside the FIB). The structural analyses of

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A-MoS2 and T-MoS2 were performed using a TEM (JEOLARM200F) with a hexapole corrector (CEOS GmbH) for theelectron probe. All STEM operations were conducted with 200kV accelerating voltage. (Detailed description for STEM work:ADF (annular dark field) STEM analysis was performed with aprobe current of ∼20 pA, a condenser aperture (30 µm), abright field aperture (3 mm) and a camera length (8 cm) pos-sessing a collection inner angle of ∼70 mrad, respectively. Thescanning rate of the ADF and ABF image was applied with 6 µsper pixel and 512 × 512 pixels.)

Acknowledgements

W.-K. H. acknowledges the financial support from the KBSIgrant (T36417). B. C. acknowledges the support from theFundamental Research Program (PNK4890) of the KoreanInstitute of Materials Science (KIMS) and the Basic ScienceResearch Program of the National Research Foundation ofKorea (NRF) funded by the Ministry of Science, ICT & FuturePlanning (NRF-2014R1A1A1036139). K. H. O. acknowledgesthe support from a grant (MPSS-CG-2016-02) through theDisaster and Safety Management Institute funded by Ministryof Public Safety and Security of Korean government.

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