Mechanical Properties and Piezoresistivity of Tellurium Nanowires

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C: Physical Processes in Nanomaterials and Nanostructures

Mechanical Properties and Piezoresistivity of Tellurium Nanowires Sijia RAN, Tom S Glen, Bei Li, Tianye Zheng, In-Suk Choi, and Steven T. Boles J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.9b05597 • Publication Date (Web): 13 Aug 2019 Downloaded from pubs.acs.org on August 13, 2019

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The Journal of Physical Chemistry

Mechanical Properties and Piezoresistivity of Tellurium Nanowires Sijia Ran1, Tom S. Glen1, †, Bei Li1, Tianye Zheng1, In-Suk Choi2, Steven T. Boles1, * 1 Department

of Electrical Engineering, The Hong Kong Polytechnic University, Kowloon, Hong

Kong SAR, China 2

Department of Materials Science and Engineering, Seoul National University, Seoul 151-747,

Korea *

Corresponding author e-mail: [email protected]

ACS Paragon Plus Environment

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ABSTRACT

Among

elemental

semiconductors,

tellurium

(Te)

exhibits

unique

mechanical

and

electromechanical properties due to its highly anisotropic crystal structure and mixed interatomic bonding modes. A lack of experimental investigations of these properties inhibits its adoption in new applications both in bulk form, as well as at the nanoscale. In this study, uniaxial tensile tests were conducted in a scanning electron microscope (SEM) on [0001] orientated Te nanowires (NWs) with diameters ranging from 15 to 35 nm. An average elastic modulus is estimated to be 38.6 ± 4.7 GPa. Both elastic and elastic-plastic behaviors are observed in tested NWs, with a large fracture strain of up to 18% achieved in the latter case. Regardless of the deformation type, electromechanical tests of Te NWs show a trend of decreasing resistance with increasing strain at low-to-moderate tensile strains (0 – 4%). This piezoresistive effect provides for new opportunities for tellurium to be utilized either in nanoscale devices, or in systems that can utilize the extraordinary properties of single crystal tellurium.


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INTRODUCTION The pronounced interactions between electrical and mechanical states of various semiconducting materials have been broadly adopted to achieve micro-electro-mechanical systems (MEMS) for sensing and actuation1,2, where the piezoresistive effect can be utilized. This effect refers to the change in resistance when the material subjected to mechanical stress, and is related to the straininduced change in electron band structure in semiconductors3. Therefore, in piezoresistive semiconductors, the electron transport properties can be greatly altered by applied stress via modifications of the band gap energy, carrier effective mass, and the scattering probability4. As a narrow bandgap (0.34 eV5) p-type semiconductor, elemental tellurium (Te) holds a unique position among these materials. Under a variety of strain conditions, piezoresistive6-8 and piezoelectric9 effects have been observed in Te during experimental investigations, and a trivial insulator to topological insulator transition in strained Te has been predicted by theoretical simulation10.

Figure 1. A schematic illustration of the lattice structure and the slip systems of hexagonal Te. Te atoms are packed in hexagonal lattice and are bonded by vdW force in the c-plane. Along the helical chain, Te atoms are covalently bonded to two nearest atoms, which are illustrated by the solid grey lines. The slip planes {1010} of Te are indicated in green color.


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The vertical and horizontal red arrows represent the slip directions ([0001] and ) on the slip planes, respectively. The diverse effects of strain on Te are governed by its highly anisotropic crystal structure. Single crystalline Te exhibits a trigonal structure consisting of two types of atomic interactions11. As shown in Figure 1, covalently bonded Te atoms spiral along the crystalline c-axis ([0001]), which form the helical chains aligned in parallel within a hexagonal lattice. Between these chains, a weaker bonding, i.e., van der Waals (vdW), dominates the interaction between the nearest atoms. The difference between the intrachain and the interchain interactions results in the anisotropy in mechanical properties of Te. Bulk Te has the largest elastic modulus of 43 GPa along the covalently bonded c-axis12. Some of the earliest work exploring the mechanical properties of Te suggested that when the tensile loading direction is misoriented off the c-axis, bulk Te may show plasticity from room temperature to extremely low temperature down to -195 ºC13,14. This is in contrast to more conventional semiconductors where high temperature is needed to enable plastic deformation15. The plasticity of Te is directly related to its bonding characteristics: Dislocation slip tends to occur on planes parallel to the weakly bonded helical chains to prevent the dissociation of the covalent bonds16. As is well established, bulk Te has a slip system consisting of only one set of slip planes (m-planes, {1010}) with slip directions along the a-axis () and the caxis13. One-dimensional (1D)17,18 and two-dimensional (2D)19,20 Te nanostructures have been successfully synthesized with controllable morphologies and excellent electron transport properties. The combination of the ultra-thin thicknesses (down to 4 nm17,20) and interesting electrical properties make Te nanostructures promising candidates for applications in nanoelectronic devices, such as transistors21, nanogenerators22, force sensors23,24, and


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thermoelectric devices25. Particularly, the synthesis of nanoscale Te may benefit strain-related applications, because nanoscale materials can exhibit novel mechanical and electromechanical properties compared to their bulk counterparts (e.g., size-dependent elastic modulus26, high fracture strength27-29, unusual plasticity30, and remarkable piezoresistive and/or piezoelectric properties31,32). Quantitative characterization of these properties at the nanoscale is therefore important for the development of high performance and reliable nanoscale devices. Currently, however, only a few studies of mechanical and electromechanical properties of Te nanostructures have been reported. Simulations and experiments to quantify the elastic modulus focus on different size scales and show clear discrepancies. Ghosh et al.12 simulated the elastic properties of Te nanowires (NWs) with diameters less than 1 nm using first-principles densityfunctional theory calculations. The results suggested Young’s modulus of ultrathin Te wires (28.535.1 GPa) is smaller than the bulk value (43 GPa) and decreases with reducing diameters12. A different trend was observed in an experimental work using contact resonance atomic force microscopy (CR-AFM), which showed a nonlinear elastic modulus increases from 45 GPa to 85 GPa with NW diameters decreasing from 100 nm to 30 nm33. Beyond the elastic limit, a firstprinciples study predicted the possibility of slip between each helical chain before the breaking of covalent bonds in Te NWs34, but no experimental data is currently available to support this finding. Apart from the mechanical properties, the investigations of electromechanical properties were conducted on Te microwires23,24 and nanobelts35 using bending tests but were limited to low strain levels (< 0.4%) and imperfect testing design since the bending configuration generated electrical responses to stresses along the radial and axial directions at the same time. Therefore, uniaxial tension and/or compression tests are needed to decouple the effects of piezoresistive and


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piezoelectric properties and can provide more accurate quantifications of the relationship between electromechanical properties and the longitudinal strain. In this work, uniaxial tensile tests were conducted to investigate the mechanical and electromechanical properties of freestanding Te NWs with the diameter down to 16 nm during in situ scanning electron microscopy (SEM). The mechanical behavior and electrical properties of Te NWs under high strain levels (> 4%) were explored for the first time. The in-situ testing technique provided a real-time observation of elastic-plastic-fracture processes and yielded a quantitative relationship between stress and strain. It was observed that Te NWs underwent either elastic or elastic-plastic deformation before fracture. In both cases, NWs sustained much larger fracture strength than previously demonstrated for bulk Te. Meanwhile, to explore the possible strain related applications of these NWs, the relationship between the electrical properties and the applied strain was studied and a resistance decrease under increasing tensile strain in the elastic regime is seen, supporting the previous work evidencing piezoresistive properties6-8. METHODS Nanowire synthesis. Te NWs were prepared using a vapor transport method in a single-zone horizontal tube furnace (OTF-1200X-S, MTI). 2 mg germanium telluride (GeTe) powder (Alfa Aesar, 99.999%) was placed at the middle of a quartz tube (inner diameter of 1.5 cm) as the precursor. Silicon (Si) substrates were placed approximately 14 cm downstream from the GeTe source. Prior to NW growth, the quartz tube was flushed with argon (Ar) and 5% hydrogen (H2) forming gas and pumped down to vacuum. The forming gas was used as the carrier gas at a flow rate of 120 sccm, and the low pressure in the tube was maintained during the whole process by keeping the pump on. During the NW growth, the GeTe source vaporized at 450 ºC via sublimation at the hot central region of the furnace, and the vapor was transported to the colder region


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downstream with a temperature lower than 385 ºC, where the vapor condensed on the Si substrates forming Te nanostructures. The growth temperature relied on the temperature distribution in the tube furnace and fell in a reasonable range for the growth of single crystalline Te NWs36,37. The growth duration was 1 h. Using vapor solid (VS) growth, the crystallinity of synthesized nanostructures has been found to be insensitive to temperature, but the morphologies evolve greatly with growth temperature38,39. To select appropriate samples, as-made products were characterized before testing. Transmission electron microscope (TEM, JEM-2100F, JEOL) samples were prepared by dispersing the NWs into isopropyl alcohol (IPA) followed by dropcasting onto a TEM grid. Some of NWs were also harvested from the edge of the growth substrate and transferring to a copper (Cu) grid in a dual beam focused ion beam (FIB)-SEM (JIB-4501, JOEL) for TEM characterization. Tensile testing setup. The mechanical and electromechanical measurements of the Te NWs were performed using a prober shuttle (PS4, Kleindiek) equipped with a nanomanipulator, a threeaxis substage and a micro-force sensor (FT-S100, FemtoTools), as shown in Figure 2e. The piezoelectrically driven nanomanipulator and the substage provided high-resolution movement for precise probing and positioning of individual NWs, while the nanomanipulator also worked as the piezo-actuator to produce different strain states in NWs during tests. The prober shuttle operated in the FIB-SEM equipped with a platinum gas injection system (Pt-GIS, Figure 2f). Electron beam induced Pt deposition (EBID-Pt) was used to weld NWs to the tungsten tip for harvesting and tests. Besides, the conductive nature of the deposited C/Pt made it suitable to generate electrical contacts between the NW and electrodes for the electrical measurements. Tensile testing processes. The selected Te wire was harvested from the sample substrate and placed between a fixed electrode and the manipulator tip with both sides gripped using EBID-Pt.


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These EBID-Pt grips have been proven to be strong enough to hold NWs up to 500 nm in diameter40. To further ensure strong gripping of the wire, before testing, the deposition was kept for 90 sec at each side to generate significant thicker grips than the NW diameter. Uniaxial tensile tests were conducted following the previously reported testing method40-42. During tensile tests, the NW was elongated along its growth direction by the nanomanipulator in a step by step fashion. In each step, a fixed displacement (~3 nm) was first conducted during 1 sec. The total displacement of the nanomanipulator was a sum of the displacements of the NW and the force sensor probe. To obtain the true elongation along the NW, the strain state was maintained for approximately 20 sec to acquire a SEM image. The deposited C/Pt grips in SEM images were then used as markers and tracked by digital image correlation (DIC). To avoid any inaccuracy caused by displacements of the sensor probe, the grip on the sensor side was aligned to a fixed position in all images beforehand. As the DIC analysis tracked the movements of two grips and gave the distance changes in between as the elongations of the wire, the deduced engineering strain was clear of the influence of the displacements of the manipulator and the sensor probe. Based on the resolution of the images taken during the DIC measurements and the magnification of the samples under tests, the strain resolution was estimated to range from 0.27% to 0.14%, depending on the length of the wire. Stress measurements. In mechanical tests a commercial-available capacitive-based microforce sensor was used, which served as the second electrode. When force was generated in the NW, the same amount of force was applied to the sensor probe to create a capacitance change in the sensor. This capacitance change was transformed to a voltage signal and collected by a digital acquisition card (DAQ, USB-6002, National Instruments), which transferred its readings to a LabVIEW-based program in real-time. With the pre-calibrated sensor gain between the force and the output voltage, the force data was calculated and continuously recorded by the program at an


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acquisition rate of 200 Hz. The sensor provided a measurement range of up to 100 µN and a resolution of 0.005 µN at a measurement frequency of 10 Hz (0.02µN at 200 Hz), which resulted in a stress resolution ranging from 99.5 to 20.8 MPa for wires with 16 to 35 nm diameters during in situ measurements at 200 Hz. Since all force data at a certain strain level was calculated into an average during analysis, more than 100 data points can be obtained even when the maximum force is only about 1 µN. To determine the stresses, diameters of the nanowire were measured from the SEM images. TEM imaging was also performed on part of the tested wires to determine the thicknesses of the C/Pt surface coatings and to establish more accurate measurements of the wire cross-sectional area. I-V measurements. In electromechanical tests, a stainless-steel (SS) electrode was used to replace the force sensor to work as the second electrode. The SS electrode was electrically connected to the substage, which, together with the manipulator, provided two conductive paths from the vacuum chamber to a source-meter (2601B, Keithley). Two-point probe electrical measurements were conducted when a strain state was maintained. Current (I)-voltage (V) curves were obtained by applying a voltage sweep from -3 to 3 V across the NW, during which the applied voltage and the current responses were recorded. To prevent noise from the electron beam, the wire was not imaged during any electrical measurement.


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RESULTS AND DISCUSSION

Figure 2. (a) SEM image of the synthesized NWs grown at the edge of the Si substrate and a harvested NW attached to a manipulator tip. The NW was welded to the tip via EBID-Pt. (b) TEM image of a representative NW. (c) The corresponding HRTEM image of the wire. (d) The FFT pattern converted from (c). (e) Photograph of the tensile testing setup. (f) SEM micrograph of the force sensor probe and the manipulator tip. The sensor tip is fixed opposite to the manipulator tip to sense force along the probe during fluctuations of the manipulator. Te nanowires used for in situ testing. Figure 2a and 2b show the SEM and TEM images of representative nanostructures that were used in tensile testing. The NW lengths range from 1 to 5 μm with diameters between 10 and 45 nm. The high resolution TEM (HRTEM) image (Figure 2c) and the fast Fourier transformation (FFT) pattern (Figure 2d) reveal that the NWs are typically single crystalline hexagonal Te with [0001] growth direction. Since NWs suspended at the edge of the Si substrate can be easily approached and detached from the substrate without significant


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damage, they were chosen for tensile tests and were harvested and examined in TEM. As shown in the TEM image (Figure S1a) and the selected area electron diffraction (SAED) pattern (Figure S1b), the microstructure of the harvested NWs matches well with those dispersed on the TEM grid, which also confirm the reliability of the manipulation process since no significant damage is observed. As seen in Figure S1a, the harvested NWs have surface coatings, with the thicknesses increasing from 4 to 13 nm with extending imaging time or additional deposition times in FIBSEM. This observation implies that the coatings were formed during TEM sample preparation rather than the NW growth. Organic/carbonaceous contaminants in FIB-SEM or diffusion of deposited carbon/Platinum (C/Pt) may be responsible for the generation of the surface coatings40. Consequently, these coatings may contain Pt nanoparticles embedded in organic matrix43, which could result in the O and Pt peaks in energy dispersive X-ray (EDX) spectrum (Figure S1c). Nevertheless, these C/Pt coatings contain much less Pt content than at the targeted region for deposition and similar coatings were found to be benign in the case of other NW mechanical testing41.

Figure 3. (a) A sequence of SEM images (0.0%, 1.4%, 2.8%, and fracture) showing the mechanical testing processes of a Te NW. The engineering strains were deduced from these


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images through DIC. (b) Stress-strain relations of tested Te NWs. Both elastic-plastic (type-I) and pure elastic (type-II) behaviors were observed. The inset shows the elastic regime of the type-I behavior. The elastic modulus of Te core is determined by the slope at low strain levels. Regardless of different types of behaviors, NWs show similar elastic moduli. Mechanical properties. Figure 3a shows the tensile testing processes of a freestanding Te NW. Stress-strain curves of multiple NWs from the elastic regime to fracture were obtained from the mechanical tests. Two types of mechanical behaviors were observed, with typical stress-strain curves of type-I and type-II behaviors shown in Figure 3b. The NWs that showed the type-I behavior underwent elastic deformation at low strain and yielded at 3.3%-4.6% of strain (Table 1). Followed by this regime, they deformed plastically and reached strains of up to 15.7%-18.1% before fracture. After fracture, a permanent elongation of the total length of the NW was observed. In contrast to this behavior, NWs of type-II behavior showed elastic stress-strain relations before fracture and broke with maximum engineering strain less than 10%. Table 1. Measured Elastic Modulus, Yield Strain and Strength, and Fracture Strain and Strength of Te NWs with Different Types of Mechanical Behaviors

type-I

type-II

core diameter (nm)

shell thickness (nm)

length (µm)

elastic modulus (GPa)

yield strain

yield strength (GPa)

fracture strain

fracture strength (GPa)

16

12.5

1.92

57.5

4.6%

2.6

18.1%

4.5

18

10

2.60

54.4

3.3%

1.5

15.7%

2.8

35

12.5

1.50

42.1

4.5%

1.8

17.7%

3.9

26

12.5

2.89

49.4

NA

NA

9.9%

4.9

34

10

3.55

48.5

NA

NA

4.5%

1.8

4.1%

2.0

17.2%

3.7

NA

NA

7.2 %

3.4

average values

type-I type-II

50.4


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To quantify the elastic moduli of the NWs under tests, the wire’s composite nature (Te, C/Pt) must be considered. The diameters of the Te cores were determined by subtracting the thickness of the organic contamination from the total diameter of the wires, as obtained from SEM images. Thicknesses of the C/Pt shells range from 10 to 12.5 nm, which were measured from TEM images of multiple wires after loading. Depending on the total testing time and C/Pt-deposition times, the lower and the upper end of the shell thicknesses were used to estimate the diameters of different Te NWs. The elastic modulus of a single NW was determined by the slope of the linear region of its stress-strain curve at low strain. Differences in measured elastic moduli were not distinguishable regarding different types of mechanical behaviors. Thus, by taking an average from all the tested wires, we obtained an elastic modulus (E) of 50.4 ± 6.5 GPa. Compared to the bulk single-crystal Te modulus (43 GPa12), our measured values are slightly higher. However, by accounting for the nanoscale specimen sizes and the thin C/Pt shell layer41, our measured elastic modulus is reasonable and in line with expectations. For instance, if we consider the Te core and the C/Pt shell as the Voigt composite configuration and assume the modulus of the C/Pt shell layer is 10% of the tested NWs40, the estimated elastic modulus is obtained as 38.6 ± 4.7 GPa, in excellent agreement with the bulk value. For wires showing the type-I behavior, 0.2% offset criterion was used to determine their yield point, which gave an average yield strength of 2.0 ± 0.5 GPa. The fracture strength varied from wire to wire with maximum values of 4.5 GPa and 4.9 GPa for type-I and type-II behaviors, respectively. Strong size dependence of elastic modulus as reported by Stan et al33 was not observed in this work. We attribute this difference to the different testing techniques. The CR-AFM method used in the above-mentioned study produced the largest tensile strain in the near-surface layer of the wire and caused compression in the inner core33. On the contrary, uniaxial tensile tests used in this


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work can provide uniform strain across the cross-section of the tested wire44. Therefore, the proposed explanation, i.e., surface stiffness enhancement33, of the CR-AFM method may not valid or has less effect for our results. Plastic deformation was evident in wires showing type-I behavior as indicated by stress-strain relations. Unfortunately, the deformation mechanics could not be studied in detail by HRTEM since the harvested Te NWs were easily damaged and bent under the high energy electron beam. However, the observed large plastic behavior can be understood by considering the slip systems of single crystalline Te and the existence of a weak vdW type bonding. In bulk Te, slip can occur in a and c directions over m-planes, and the activation of two slip modes depends on the crystal orientation and the loading direction13. In [0001] orientated Te NWs, dislocation slip along c-axis is the only reasonable deformation mechanism for the observed large plasticity. This kind of slip will not happen during an ideal uniaxial tensile test along c-axis, as no stress is acting across the slip planes of Te. Nevertheless, alignments remain a challenge for tensile tests at the nanoscale and offset angles were not completely avoided in height in this work. This misalignment can lead to nonuniform bending stresses along the NW, which reach a maximum near the grips and become negligible in the middle of the wire45. Although the maximum bending stress only takes up about 7% of the axial stress with a 10º offset before yielding45 and result in even smaller shear components on the slip planes, they may exceed the critical resolved shear stress due to the weak vdW interactions. In fact, without the need of breaking the covalent bonds during plastic deformation, the room temperature elastic limit of bulk Te is three orders of magnitude smaller than bulk Si at elevated temperature46. Some preliminary measurements of our wires suggest a best case offset angle of 6.4º during testing. According the calculations of Kang and Saif45, these wires were likely to experience several hundred MPa in bending stress when they were stretched to


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around 5% of strain. Therefore, the observed plasticity in Te NWs can be a result of misalignment during tensile tests, which initiates dislocation slip when subjected to high bending stresses. Even though the regions close to the grips may carry the largest bending stresses, most of tested NWs fractured close to the middle. One possible reason could be that the regions near the grips were covered by a thicker layer of C/Pt due to the diffusion of deposited EBID-Pt and may result in smaller axial stresses being carried by the Te core close to grips, preventing them from breaking. Furthermore, the thickness of the Te core varies slightly (~4 nm) along the NW length, which could cause higher core axial stresses at the regions with smaller diameters. Depending on the average diameter of the NW, the axial stress of the Te core can be 10% to 20% larger than the average value at the regions with the smallest diameter, making these regions the preferential locations for fracture. Meanwhile, the fracture of these NWs may also be affected by the failure of the quasi-static condition during the short period when it was elongated to a new strain state, which could result in additional shear stresses in the wires. Despite the difficulty in avoiding misalignment, pure elastic behavior was achieved up to 9.9% of strain before brittle fracture with better alignment. Using the corrected elastic modulus to exclude the contribution of the C/Pt shell, the fracture strength of the Te core of this wire is estimated to be 3.8 GPa. This experimental value largely exceeds the fracture strength of bulk Te (140 MPa47) under c-axis tensile loading and is located in a reasonable range of experimental-totheoretical ratio, which represents approximately 61% of predicted theoretical strength (E/2π48) of defect-free Te. For the type-I behavior, the high fracture strain of up to 18% approaches the prediction from first principle simulation34, where the results showed interchain slip could be achieved and reach a maximum strain of 24% in c-axis before the broken of the covalent bonds. The large fracture strength or strain were also reported for other metallic and semiconducting NWs


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(e.g., gold27, Si28, and Ge29), which are believed to be associated with low defect densities in nanostructures. The soft nature of Te may further enlarge the differences between NWs and bulk and even between individual specimens, as mild stress can lead to a dramatic increase in dislocation densities from 103 to 106 cm-3 49. For the above reasons, the achieved high fracture strength and strain confirmed the high quality of synthesized NWs and the testing processes.

Figure 4. (a) I-V curves of a Te NW at a variety of strain levels in the voltage range from -3 to 3 V. (b) The relative resistance change (∆𝑅/𝑅0) of the Te NW under applied tensile strains. The inset shows the current change at a bias of 3 V under tension. Electromechanical properties. To study the electromechanical properties of Te NWs, twopoint probe electrical measurements were conducted on five Te NWs under different strains. Almost symmetric nonlinear I-V curves were obtained from all tested NWs. Figure 4a shows the results of a tested NW of 36 nm in diameter and 2.27 µm in length, which achieved the highest fracture strain of 9.5% in electromechanical testing. The nonlinearity of the curves indicates the formation of back-to-back Schottky barriers at two metal-to-semiconductor contacts50,51. The formation of Schottky barriers was possibly due to the p-type nature of Te NWs and their larger work function compared to the metals (i.e., stainless steel and W) of the two electrodes52. Besides, since the resistivity of the C/Pt shell on the NW surface is at least three orders of magnitude higher than Te, it may also affect the contact type.


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As shown in the inset of Figure 4b, current at a bias of 3 V shows a linear increase at strains smaller than 5%. Beyond that strain level, the current remains at about 0.26 µA with small fluctuations. To reduce the influence of double Schottky barriers, resistances of the NW were calculated using the high bias region from 2.5 to 3 V51. The relative change in resistance (∆𝑅/𝑅0) is plotted as a function of engineering strain (see Figure 4b), where 𝑅0 is defined as the initial resistance of the unstrained NW. The resistance change also clearly shows two regions. ∆𝑅/𝑅0 decreases by approximately 25% when engineering strain reaches 4% and then fluctuates around a constant value until the wire fractured at approximately 9.5% of strain. To be noted, though the voltage increase is small across the forward and the reverse-biased Schottky barriers of the contact at high bias51, their contributions are still included in the measured resistance and lead to an overestimation of the NW resistance. This means the determined ∆𝑅/𝑅0 represents a minimum value of the relative resistance change under strain In the elastic regime, strain-induced resistance change can be related to piezoelectricity and/or piezoresistivity in Te. The uniaxial strain along the a-axis can break the symmetry of Te chains in the c-plane ((0001)) and bring about displacements of positive Te ions and negative charge centers relative to the Te core22. Thus, electric dipoles arise along the a-axis, and the polarization charges accumulate on the NW surface exhibiting the piezoelectric effect. This effect was considered as the main factor in previous research on electromechanical tests of Te microwires23 and nanobelts35, as the employed bending test methods produced stresses both along and perpendicular to the caxis. In this work, however, the piezoelectric effect may be negligible. This is because shear components on the c-plane should be small under tensile loading along the c-axis even with possible slight misalignments. Also, the metal contacts that surrounded the surface of the NWs could rule out the influence of polarization charges.


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With awareness of the origin of the observed electromechanical interactions, the longitudinal piezoresistive coefficient (𝜋𝑙) of Te NWs can be determined. 𝜋𝑙 relates the relative resistivity change (∆𝜌/𝜌0) to the applied longitudinal stress (𝜎𝑙), i.e.,

∆𝜌 𝜌0

= 𝜋𝑙𝜎𝑙 . ∆𝜌/𝜌0 is given by

∆𝜌 ∆𝑅 ∆𝑅 = ― (1 + 2𝜐)𝜀 ≈ 𝜌 0 𝑅0 𝑅0 where 𝜐 and 𝜀 are the Poisson ratio and the engineering strain of the wire, respectively. This equation is valid at strain levels close to zero strain and is normally suited for the elastic regime. Regardless of different types of mechanical behaviors, tested Te NWs deformed elastically under tension smaller than an average strain of 4.1%. Thus, the first six data points of the relations between ∆𝑅/𝑅0 and 𝜀 should be within the elastic regime and were fitted linearly to find 𝜋𝑙. By employing the stress-strain relations in this regime, i.e. 𝜎 = 𝐸𝜀, and introducing the elastic modulus of 38.6 GPa, the average 𝜋𝑙 for the elastic regime is determined to be -15.4×10-11 Pa-1. Considering the underestimation of ∆𝑅/𝑅0, the exact 𝜋𝑙 could be slightly larger. At the bulk scale, change in band gap energy was proved to be the main effect of the piezoresistivity in Te7. In a recent work, ab initio electronic structure calculations indicated c-axis uniaxial tension (or hydrostatic pressure) lowers conduction band minima of Te and result in a band gap narrowing, which can contribute to an increase in carrier concentration and thus a decrease in resistivity10. Experimentally, Takaz6 observed a resistivity increase in single crystalline bulk Te under uniaxial compressive stress parallel to the c-axis, which gave a longitudinal piezoresistive coefficient of -400 ×10-11 Pa-1 at room temperature. At the microscale, however, Liang et al.24 obtained a positive gauge factor from Te wires, where they found the I-V characteristics can be largely affected by the change in Schottky barrier heights associated with the change in band gap under strain. In our work, the trend of strain-induced resistivity change is


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in agreement with the simulation and the experiments on bulk Te, and no significant change in Schottky barrier heights is observed. We trust the different results obtained on Te microwires24 are from the loading condition that the axial strains were introduced by bending the flexible substrate in which the wire was embedded. In addition to the possibility that piezoelectric effect may play a role on changing the barrier heights at contacts under bending35, the anisotropy of Te may result in the opposite sign of piezoresistive coefficients under bending comparing to c-axis uniaxial strain considering the fact that uniaxial strain along the a-axis has an opposite influence on band gap energy comparing to uniaxial strain in the c-axis and can change the band gap more effectively6,10. Thus, with better chance to avoid the large uniaxial transverse strain, tensile testing holds advantages in quantifying the piezoresistivity over bending configuration. In our tests, the measured value of piezoresistive coefficient is much smaller than that in bulk Te. Considering the nonlinearity of piezoresistive coefficients close to zero strain7, we should be aware of the difference in investigated strain ranges, which were about 0.002% and 4% of maximum strains for the bulk and NWs, respectively. Moreover, since the resistance response of Te is a combined effect of change in the number of carriers and the carrier mobility, it should not be symmetric with respect to compressive and tensile strains. The change of the effective mass may not be negligible in our testing range and lead to a hole mobility decrease under tension3,8, thereby resulting in a smaller coefficient under tensile strains. As a material that is both semiconducting and ductile, resistance of Te may change irreversibly under strain due to defect evolutions.

With increasing strain, the crystal may experience

dislocation rearrangements and then a significant change in the dislocation density once it starts to deform plastically. Under plastic deformation, the irreversible part should contribute more to the resistance change than in the low strain levels with nearly negligible dislocation density variations.


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This brings about more complicated mechanisms for strain-induced change in resistance under strains larger than 4%. The influence of dislocations on the electron transport properties have been studied in bulk Te53,54, though the extent of this effect here is unclear. It was found that c-screw dislocations can act as scattering centers and result in a decrease in hole mobility53. However, such dislocations also provide local acceptor levels, which in turn generates higher carrier concentrations54. Once dislocations generate and accumulate in the plastic regime, this factor may dominate the resistance change and result in a slower resistance decrease. CONCLUSIONS In summary, we report mechanical and electromechanical measurements of single crystalline [0001] orientated Te NWs by in-situ tensile tests in SEM and explored the high strain levels (> 4%) for the first time. Both elastic and elastic-plasticity behaviors were observed in Te NWs. Although the exact deformation mechanisms require further investigations, the large plasticity can be understood by dislocation slip in crystalline c-axis which can be initiated by misalignment during tensile tests. Tested NWs sustained high fracture strain of up to 18% with plastic deformation and achieved a pure elastic strain of up to 9.9% with better alignment. The piezoresistive effect in Te NWs was evidenced in electromechnical measurements. The combination of high fracture strain and interesting electromechanical interactions make Te NWs promising candidates for bendable and flexible electronics. However, particular care will be needed for strain related applications since change in electron transport properties may follow a different trend under high strains. By comparing the results from mechanical and electromechnical testing, the operative strain range of Te NWs was determined to be 0 – 4% for the application of piezoresistive effect.


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ASSOCIATED CONTENT Supporting Information. TEM characterization of harvested NWs AUTHOR INFORMATION Corresponding Author * E-mail: [email protected] Present Addresses † (T.S.G.) School of Physics and Astronomy, University of Edinburgh, James Clerk Maxwell Building, Peter Guthrie Tait Road, Edinburgh EH9 3FD, United Kingdom Notes The authors declare no competing financial interest. ACKNOWLEDGMENT The authors would like to acknowledge the funding support of internal projects “In Situ Mechanical and Electrical Testing of Nanomaterials for Next Generation Electronic Devices” (GYBLN/G-YBPQ), “Novel Materials for Emerging Energy, Electronic and Photonic Devices” (1ZVGH), and “Microscopy, Mechanics and Reliability of Novel, Nano-scale Electrode Materials for Electrochemical Energy Storage System” (1-DD6Z) of the Hong Kong Polytechnic University.


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Mechanical Properties and Piezoresistivity of Tellurium Nanowires

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