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. 2020 Oct 8;20(11):8250–8257. doi: 10.1021/acs.nanolett.0c03358

Thermomechanical Nanostraining of Two-Dimensional Materials

Xia Liu †, Amit Kumar Sachan ‡, Samuel Tobias Howell †, Ana Conde-Rubio †, Armin W Knoll §, Giovanni Boero †, Renato Zenobi ‡, Jürgen Brugger †,*
PMCID: PMC7662931  PMID: 33030906

Abstract

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Local bandgap tuning in two-dimensional (2D) materials is of significant importance for electronic and optoelectronic devices but achieving controllable and reproducible strain engineering at the nanoscale remains a challenge. Here, we report on thermomechanical nanoindentation with a scanning probe to create strain nanopatterns in 2D transition metal dichalcogenides and graphene, enabling arbitrary patterns with a modulated bandgap at a spatial resolution down to 20 nm. The 2D material is in contact via van der Waals interactions with a thin polymer layer underneath that deforms due to the heat and indentation force from the heated probe. Specifically, we demonstrate that the local bandgap of molybdenum disulfide (MoS2) is spatially modulated up to 10% and is tunable up to 180 meV in magnitude at a linear rate of about −70 meV per percent of strain. The technique provides a versatile tool for investigating the localized strain engineering of 2D materials with nanometer-scale resolution.

Keywords: 2D materials, strain nanopattern, molybdenum disulfide, local bandgap, thermal scanning probe lithography, tip-enhanced Raman spectroscopy

Introduction

Straining atomically thin 2D materials (2DMs) is a promising strategy to modify their electronic and optoelectronic properties,1−3 enabling high-performance device applications such as single-photon emitters,4 high-responsivity photodetectors,5,6 and flexible nanoelectronic systems.7 For the case of graphene, it has been computed that by applying a strain of about 12% along C—C bonds a bandgap of about 0.5 eV is obtained,8 which enables graphene to be used as a semiconductor. For the case of molybdenum disulfide (MoS2) and molybdenum ditelluride (MoTe2), their phase change between semiconducting and metallic phases can be modulated by biaxial and uniaxial strain.9,10 Additionally, it has also been demonstrated that the mechanical strain in transition metal dichalcogenides (TMDCs) increases the carrier mobility by almost 2 orders of magnitude11 and influences their local bandgap12,13 as well as their quantum emission properties.14 So far, methods to strain 2DMs include introducing strain during growth,2,15 stretching/bending the supporting substrate,13,16−21 or transferring the 2DM onto a topographically patterned substrate.22−24 These methods create a strain before or during the growth/deposition of the 2DM and generally induce a global strain covering large surface areas.

To further improve strain engineering in 2DMs and to be able to integrate them into devices, a method is required that allows for two assets, namely to achieve spatial control of the strain pattern and to adjust its magnitude. Probe-based indentation techniques offer a convenient solution to these challenges because strain can be locally generated and the ability of indenting strain profiles makes it possible to control the strain distribution. In previous works, an atomic force microscopy (AFM) tip was used to temporarily induce strain in a suspended graphene membrane.25 In another approach, permanent local strain patterns in graphene on a SiO2 substrate were created by means of an AFM tip at room temperature yielding an indentation depth below 2 nm resulting in a strain magnitude of only 0.1%.26

Though various types of strain engineering methods have been proposed for 2DMs, both the strain magnitude and spatial distribution are far from being controlled with the high resolution that is needed to fabricate integrated nanodevices. Thermal scanning probe lithography (t-SPL) is an advanced direct-write method that uses a heated nanotip for 2D and 3D subtractive/additive manufacturing.27,28 Recently, we have shown that t-SPL can be used as a one-step nanolithography process at high resolution for directly nanocutting 2DMs, provided the thermomechanical energy is sufficient to break the interatomic bonds in the 2DM and sublimate the polymer underneath.29 By expanding from there, we demonstrate here that by applying a suitable combination of indentation force and tip temperature below the cutting threshold, we are able to induce a well-controlled and permanent local deformation of the 2DM. This permanent local deformation determines the local strain, which results finally in a local bandgap modification. By raster scanning the tip over the sample, arbitrary patterns, such as lines, squares, and ripples, can be written through consecutive indentations. A key benefit of our technique is that the indentation depth can be controlled by the indentation force of the tip, the strain in the 2DM being locally confined to the writing area due to van der Waals interactions with the underlying polymer. We show by our results that this novel strain nanoengineering technique is applicable to a variety of 2DMs, including MoS2, MoTe2, molybdenum diselenide (MoSe2), and also graphene. The proposed strain nanopatterning technology allows for a deterministic and precise modulation of the local bandgap of 2DMs with nanometer-scale resolution.

Results and Discussion

The Creation of Strain Nanopatterns in MoS2

The concept of thermomechanical strain nanopatterning is illustrated in Figure 1a. A monolayer of 2DM, such as TMDC or graphene, is transferred onto a 50 nm thick polyphthalaldehyde (PPA) polymer layer that serves as thermosensitive deformable material. With the combination of heat and indentation force from the t-SPL nanotip, the 2DM and the underlying polymer are locally deformed as shown in Figure 1b. The deformation in the polymer is due to a combination of thermoplastic indentation and sublimation of PPA at a temperature of about 150 °C, where PPA decomposes into monomers.30 At this temperature, the 2DM layer stays physically and chemically intact and remains in conformal contact with the entrapped polymer, linked through van der Waals interactions. Forcing the tip out of contact results in fast cooling and resolidification that induces the formation of permanent ripples in a shape corresponding to the written pattern (Supporting Information Section 1). The ripples induce a permanent strain in the 2D layer at a resolution primarily given by the tip apex.

Figure 1.

Figure 1

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t-SPL based strain nanopatterning in 2D materials (2DMs). (a) Conceptual illustration of the thermomechanical nanoindentation process for strain nanopatterning in the 2DM, such as TMDCs or graphene. Drawing is not to scale. (b) Cross-section scheme showing details of the heated nanotip indenting a monolayer TMDC layer on a PPA layer. Drawing is not to scale. (c) AFM topography of the written MoS2 ripple nanostructures. (d) Three-dimensional representation of the area marked in panel c. (e) Depth profile of the selected line in panel c. The nanoindentation depth is around 4 nm and the pileup height is around 1.5 nm.

To demonstrate this approach, we first wrote a pattern consisting of horizontal and vertical lines to create arrays of designed ripples on a monolayer (1L) MoS2 flake (Supporting Information Section 2). Figure 1c shows the topography of the nanoripples having a width of 60 nm. The probe was scanned at a velocity of 0.5 mm/s and the total writing and imaging time for a 1.5 × 1.2 μm2 area (Figure 1c) was 38 s. Surface imaging with a conventional AFM confirms the topography measured with the t-SPL tool (Supporting Information Section 3). The three-dimensional (3D) topography demonstrates how accurately the strain nanopatterns in the MoS2 layer can be written (Figure 1d). Figure 1e shows the cross-section profile of the selected area in Figure 1c with nanoripples consisting of 4 nm deep valleys and 1.5 nm high pileups. We notice by detailed evaluation of the surface topography that the integrated volume of the valleys is larger than that of the pileups. This volume loss can be attributed to the local densification of the PPA polymer, an interpretation that is consistent with a similar nanoscale densification observed in earlier thermal scanning probe studies for nanolithography on PPA31 and data storage using cross-linked polymers.32 Here, in addition, because the 2DM layer is impermeable, the locally created monomers cannot outgas but rather diffuse into the polymer matrix, occupy free-volume sites of the polymer, and thereby cause further densification.

Versatility of the Strain Nanopatterning of 2DMs

To further investigate the capabilities of this approach, we designed additional patterns with the width ranging from 2 to 100 nm. Using monolayer MoS2, this results in measurable indents with a feature size down to 20 nm, which is mainly limited by the size of the tip apex (Figure 2a). We also successfully demonstrate a resolution of 20 nm over larger areas (5 μm × 3.5 μm) in the same material (Supporting Information Section 4). The depth for each indent can be quantitatively tuned by the combination of temperature and force of the tip during the t-SPL writing (Supporting Information Section 5). The temperature of the tip–sample contact increases monotonically with the heater temperature but due to the high and hardly measurable effective thermal resistance of the tip it is difficult to quantify accurately.33−35 The actual temperature of the tip in contact with the sample is approximately 40–70% of the heater temperature as discussed also in previous works.33−35 We performed experiments with a linear increase of the heater temperature from 200 to 1200 °C. A noticeable increase in line depth is observed when the heater temperature reaches 700 °C and above, as shown in Figure 2b. We attribute this relatively sharp transition to the softening transition of the polymer which coincides with the decomposition temperature.30 At temperatures significantly above the softening transition, the force to deform PPA was observed to be strongly reduced in nanolithography experiments.27 Thus, here the force applied by the tip is balanced by the stress induced by the deformed 2D material layer, which leads to a constant line depth at higher temperatures.

Figure 2.

Figure 2

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Versatility of the method and application to other 2D materials. (a) AFM topography of an array of nanostripes in 1L MoS2 with designed width in the range from 2 to 100 nm and corresponding depth profile using the voltage of 7.5 V and the temperature of 950 °C. The writing direction is starting from upper right to bottom left. (b) AFM topography of nanopatterns produced with a heater temperature from 200 to 1200 °C and corresponding depth profile under the voltage of 7.5 V. The designed width is 20 nm. The writing direction is starting from bottom left to upper right. (c) Topography of 2L MoTe2 nanoripples. (d) Topography of 1L graphene nanowells array. (e) Topography of 1L MoSe2 nanopattern with arbitrary strain distribution (the EPFL logo is given as an example). Logo is used with permission. The writing direction in panels c, d, and e is starting from bottom right to upper left.

The applied indentation force Find at which the heated t-SPL tip presses against the surface can be estimated as follows: Find = FV,T – Fk, where FV,T is the applied force due to heating and electrostatic bending, and Fk is the cantilever’s spring force (detailed in Supporting Information Section 6). In the experiments performed here, the indentation force is estimated to be ∼280 nN for a heater temperature of 950 °C and an indentation tip–sample voltage of 7.5 V (see Supporting Information Section 6). To determine the local pressure exerted on the 2DM during the process, one has to consider the contact area. If we assume a tip diameter of 15 ± 5 nm, we obtain a local pressure in the range from 1 to 4 GPa.

Besides MoS2, we demonstrate here the use of nanoprobe based strain engineering also on other single and multilayer 2DMs of interest, such as MoTe2, MoSe2, and graphene. To show the versatility of the method we patterned nanoripples in 2L MoTe2 (Figure 2c), nanowells in 1L graphene (Figure 2d), and arbitrary design patterns, such as the EPFL logo in 1L MoSe2 (Figure 2e). So far, the underlying thermosensitive polymer used was PPA, which under the effect of heat and force breaks into monomers. To elucidate the importance of the decomposing polymer, we also performed nanoindentation experiments with the identical parameters described in Figure 2a on thermoplastic poly(methyl methacrylate) (PMMA) polymer that has a glass transition temperature (Tg) of about 105 °C and a much higher thermal decomposition temperature of 325–500 °C.36 In the case of PMMA, the indent causes only a thermoplastic deformation and no sublimation, thus pushing some material into the site of and around the preceding indent. Similar to the case of PPA, the PMMA as supporting polymer layer deforms, which results in local strain in the 1L MoS2 layer; however, the pattern formation is, as expected, less uniform due to the higher thermal decomposition temperature (Supporting Information Section 7).

Scanning Raman Spectroscopy of Nanopatterned 1L MoS2

The spatially varying strain distribution over the nanopatterned monolayer MoS2 is verified by micro-Raman spectroscopy. Figure 3a shows the topography of a representative deep nanopatterned 1L MoS2 created by the t-SPL at a heater temperature of 900 °C and an indentation voltage of 8.5 V. The valley depth and pileup height of the strained area is about 8 and 5 nm, respectively. There are two columns of strained ripples with a valley width of 60 ± 3 nm. The Raman spectrum of the patterned MoS2 shows significant redshifts of the E12g and A1g peaks as compared to pristine MoS2 (Figure 3b) as a result of the induced strain. The E12g mode is assigned to the in-plane vibration of the two S atoms with respect to the Mo atom, and the A1g mode is assigned to the out-of-plane vibration of the two S atoms (Figure 3b). The typical Raman spectrum of pristine 1L MoS2 presents two dominant peaks at 385.7 cm–1 (E12g) and 405.3 cm–1 (A1g).37 No oxidation of the nanopatterned MoS2 is observed in wide-range Raman spectra (Supporting Information Section 8). Figure 3c,d shows the scanning Raman spectroscopic maps of the E12g and A1g peaks of the strained MoS2, respectively, with the redshifts for the E12g peak position by 2.4 cm–1 and the A1g peak position by 1 cm–1. It should be noted that the micro-Raman spectrum is an optical average of the strained valleys and pileups due to the laser spot diameter (about 0.5–1 μm). From an overall fit to the redshift magnitudes and the height profile, we estimate that an average strain of 1.3% is induced in the nanopatterned MoS2 according to the theoretical calculation discussed in Supporting Information Section 9. To corroborate our hypothesis that the strain is maximized by writing ripples, we performed another test and wrote a single square (without ripples) of the same size as the rippled area (1 μm × 1 μm) with the same indentation depth. As expected, we do not observe the Raman shift in the single square region (Supporting Information Section 10), which indicates that the strain is a result of the nanostructures written by t-SPL as illustrated in Figure 1a. In addition, the Raman spectra of the strained 2L MoTe2 (Figure 2c) and 1L graphene (Figure 2d) also show significant redshifts of the Raman peaks as compared to unstrained MoTe2 and unstrained graphene, respectively (Supporting Information Section 11).

Figure 3.

Figure 3

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Micro-Raman characterization of the nanopatterned 1L MoS2. (a) Topography of the nanopatterned 1L MoS2 created by the t-SPL. (b) Raman spectra of the strained and unstrained MoS2. Inset: schematic atomic vibration of the in-plane E12g and out-of-plane A1g modes. (c,d) Scanning Raman spectroscopic maps of the E12g and A1g peaks of the strained MoS2, respectively. The laser scanning step in the x-axis direction is 0.2 μm and that in the y-axis direction is 0.1 μm. The Raman shift corresponding to the maximum of each spectral peak is obtained by Lorentzian line fitting.

AFM-TERS Measurement of the Nanopatterned 1L MoS2

To identify the strain with nanoscale resolution below the optical diffraction limit, we performed an AFM-based tip-enhanced Raman spectroscopy (AFM-TERS) investigation over the nanopatterned 1L MoS2, as illustrated in Figure 4a,b. The AFM-TERS combines Raman spectroscopy with scanning probe microscopy.38 Upon illumination, the electric field strength at the tip apex (Supporting Information Section 12) strongly increases thanks to localized surface plasmon resonance.39,40 Therefore, the Raman signal of the probed material is enhanced in the vicinity of the tip. Figure 4c shows an AFM topography of the nanopatterned 1L MoS2, acquired prior to the AFM-TERS study. A small rectangular area of 30 nm × 200 nm in Figure 4d includes two valleys and one pileup, where AFM-TERS signals are acquired with a resolution of 3 × 20 pixels. Each pixel represents a 10 nm × 10 nm area, which is small enough to resolve the strain differences present within the nanopatterned 1L MoS2. At each pixel, tip-enhanced near-field Raman signals are collected using tapping mode AFM feedback at ambient conditions. During TERS measurements, the laser power and signal acquisition time are optimized to minimize far-field Raman signals.

Figure 4.

Figure 4

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Tapping-mode AFM-TERS mapping of nanostrained 1L MoS2. (a,b) Schematic illustration of AFM-TERS mapping of strained 2D material (2DM). Drawing is not to scale. (c) AFM topography of the strained nanoripples in 1L MoS2. The vertical lines are artifacts caused by high-frequency noise. (d) Height map showing the selected area of the nanopatterned structure in panel c. Each pixel represents an area of 10 nm × 10 nm. The green regions represent the indented MoS2 and the orange region represents the pileup created during the nanoindentation process (in the gray pixel the TERS spectrum was accidently not measured). (e) TERS map of the A1g peak corresponding to the height map in d. (f) TERS A1g peak data of nanopatterned MoS2 in the valleys and the pileup. (g) Single TERS spectrum of unstrained MoS2 and the average TERS spectrum of strained MoS2 in d. The spectra were extracted, smoothed and subsequently analyzed for peak shift.

The two first-order Raman active vibrational modes of MoS2, E12g and A1g modes, are selected for screening the nanoscale vibrational changes in the strained 1L MoS2. All spectra are background corrected for enhancement fluctuations by normalization to the local enhancement. Figure 4e shows that the tip-enhanced Raman spectroscopic map of the A1g peak corresponds to the height map in Figure 4d. The TERS map of the E12g peak is also consistent with the height map as shown in Supporting Information Section 13. The patterned MoS2 shows the most redshifts for both E12g and A1g peaks: 4.7 and 2.1 cm–1 for the valley regions and 2.7 and 1.5 cm–1 for the pileup region (Figure 4f). This indicates that the strain has been successfully created in the entire nanoripples including valleys and pileups. The redshift value of the strained MoS2 in the valleys is about two times of that measured using micro-Raman spectroscopy. The most strain in the valleys is estimated to be 2.6% and the most strain in the pileup is 1.5%. Figure 4g plots the average E12g and A1g peak shift within the strained region, which is consistent with that measured in the micro-Raman spectroscopy.

Bandgap Modulation of Nanostrained 1L MoS2

The bandgap modulation of the strained MoS2 crystal is examined by scanning photoluminescence (PL) spectroscopy. Figure 5a,b shows the A-exciton and B-exciton peak maps of the PL spectra of the same sample studied in micro-Raman characterization (Figure 3), respectively. The PL measurement is also optically averaged over the laser spot diameter (about 0.5–1 μm). The PL intensity peak of the strained area is consistently observed at lower PL energy (blue color). The unstrained MoS2 shows the typical PL spectrum of monolayer MoS2 with a principal peak at 1.9 eV (A exciton). However, the strained MoS2 clearly exhibits redshift of the A-exciton peak, indicating a strain-induced bandgap reduction. Figure 5c compares the representative PL spectra of the strained MoS2 and the unstrained MoS2, where the strong PL peaks arise from the direct bandgap emissions in monolayer MoS2. The A-exciton peak intensity is more than doubled in the strained MoS2. The enhanced PL intensity has been observed in strained MoS2 nanocones22 and semiconductor nanowires.41 As suggested by other works, this might be due to 2D exciton funnel effect,42 exciton drifting, and concentrating.41,43 By varying the indentation depth and hence the average strain, the bandgap decreases at a rate of about −70 meV per percent of strain (Figure 5d). The local bandgap modulation is estimated to be as high as about 180 meV. Therefore, a significant adjustment of the electronic band structure is generated accordingly due to the lattice deformation of the monolayer MoS2 crystal under strain. Compared with previous work (Supporting Information Section 14, Table S1), the strain tuning range is comparable and exhibits advantages in the nanoscale resolution and controllability of strain distribution.

Figure 5.

Figure 5

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PL characterization and analysis of bandgap modulation. (a,b) Scanning photoluminescence (PL) spectroscopic maps plotting (a) A-exciton energy peaks and (b) B-exciton energy peaks of the strained MoS2 sample studied in micro-Raman characterization (Figure 3). The laser scanning step in the x-axis direction is 0.2 μm and that in the y-axis direction is 0.1 μm. (c) PL spectra of the strained and unstrained MoS2. (d) Bandgap modulation as a function of the average strain. The different strains are obtained by varying the indentation depth with patterns designed to go from 1.8 to 10 nm deep. The designed width of the ripple is 50 nm and the pattern pitch is 120 nm. The average strain is calculated using the model in Supporting Information Section 9.

Conclusions

In this work, we demonstrate strain nanopatterning of 2DMs in ambient conditions on a thermosensitive polymer by using thermal scanning probe lithography. A strain pattern resolution down to 20 nm is achieved on 1–2L TMDCs and 1L graphene. By adjusting the indentation force and heater temperature, the bandgap of the strained 2DMs can be precisely modulated locally at the nanoscale. The bandgap is linearly proportional to the indentation depth, making the proposed technique suitable to precisely tune the bandgap. The versatility and repeatability of the approach proposed here will enable many investigations of strain-related phenomena in 2DMs and will help to elucidate the fundamental mechanical and electronic properties in these materials.

Experimental Methods

Material Preparation

PPA (polyphthalaldehyde, Allresist) solution (3 wt % in anisole) was spin coated on SiO2 (200 nm thick)/Si (500 μm thick) substrate (conditions, 100 rpm for 5 s and then 6000 rpm for 60 s). MoS2 flakes were exfoliated from a MoS2 bulk crystal onto polydimethylsiloxane (PDMS) stamps and then were transferred on the PPA substrate. The MoTe2, MoSe2, and graphene flakes were obtained and processed with the same procedure. The layer number of the flakes is identified by Raman spectroscopy (Supporting Information Section 2). MoTe2, MoS2, MoSe2, and graphene crystals were purchased from HQ Graphene.

Nanostrain Engineering

The strain nanopatterning of 2D materials was performed using a commercial t-SPL (NanoFrazor, Heidelberg Instruments, Switzerland). The cantilever is made of n-doped silicon. The spring constant of the cantilever is around 0.9 N/m.35 The apex diameter is 15 ± 5 nm. During the writing process, the tip was heated at a temperature in the range of 200–1200 °C. This operation was performed under N2 atmosphere to minimize the likelihood of oxidation. The indentation tip–sample voltage was in the range of 4.5–9.5 V. One layer MoS2, 2L MoTe2, 1L MoSe2, and 1L graphene samples were patterned at nanoscale.

Material and Structure Characterizations

AFM was conducted to collect the topographies of the nanostrained 2D materials. AFM was also used to verify the thickness of the nanoflakes. Images were collected using Bruker’s Dimension FastScan AFM system. The FastScan AFM scanner mode (contact mode) was used. Raman spectroscopy was performed to confirm the monolayer nature of all the exfoliated MoTe2 flakes and to characterize the properties of the patterned nanostructures. Raman spectra and PL spectra were collected using a confocal Raman microscope system (inVia Qontor, Renishaw) coupled with an Olympus inverted optical microscope and using a laser source with an excitation wavelength of 532 nm. The laser power (84 μW) was adjusted to avoid sample damage. Raman spectra were acquired in the range from 61 to 1850 cm–1 with a 15 s exposure time and an average of three measurements. Gratings of 3000 and 1800 gr/mm were used for Raman mapping and wide-range Raman measurements, respectively. The PL spectra were acquired in the range from 1.6 to 2.1 eV with a 10 s exposure time and an average of three measurements. A grating of 300 gr/mm was used for PL mapping. The peak at 520.5 cm–1 from the silicon substrate was used as a reference.

Tip-Enhanced Raman Spectroscopy (TERS)

AFM-based TERS measurements were performed using an integrated atomic force microscope, and Raman spectrometer system (NT-MDT, NTEGRA, Russia) in the top-illumination configuration, equipped with a 0.7 NA 100× objective (Mitutoyo, Japan). To prepare probes for AFM-TERS measurements, commercially available monolithic silicon-based ATEC-NC AFM probes (Nanosensors AG, Switzerland) were coated with an adhesion layer of 10 nm Cr followed by 50 nm Ag using an e-beam evaporator (Evatec, model: BAK501 LL). After preparation of TERS probes, the probes were either directly used for TERS measurements or stored under N2 flow in a chamber. A 561 nm Nd:YAG laser was used as excitation source. The incident laser power on the sample was in the range from 0.5 to 1.5 mW, depending on the spectral quality and to avoid damage in samples. The scattered light from the sample was detected by an electron multiplying charge coupled device (EMCCD, Newton 971 UVB, Andor) thermoelectrically cooled to around −85 °C. The acquisition time of the spectral signal at each pixel was 10 s. Before recording the high-resolution AFM-TER spectra and/or maps, the thermal drift was minimized by prerunning the instrument and pseudo scanning the nearby sample areas with laser illumination.

Acknowledgments

The authors thank the Center of Micro/Nanotechnology (CMi) of EPFL for the AFM facility support and the Advanced NEMS lab for the dry-transfer facility support. We also thank T. Walger for help with drawing the cross-section illustration. This work has received funding from the European Research Council (ERC) under the European Union’s Horizon 2020 research and innovation program (ERC-2016-ADG, Project “MEMS 4.0” Grant 742685 and Project “2DNanoSpec” Grant 741431). The authors would like to acknowledge the operations team of the Binnig and Rohrer Nanotechnology Center (BRNC) at the IBM Zurich Lab for their support during the tip preparation and AFM-TERS measurement.

Supporting Information Available

The Supporting Information is available free of charge at https://pubs.acs.org/doi/10.1021/acs.nanolett.0c03358.

  • Additional information, figures, and table (PDF)

Author Contributions

X.L., J.B., and G.B. conceived the idea of nanoprobe-based nanostrain engineering. X.L., S.T.H., A.C.R., J.B., and G.B. designed the experiments. X.L. performed the strain patterning using t-SPL. X.L. and A.K.S. performed the characterizations of AFM, micro-Raman spectroscopy, tip-enhanced Raman mapping, and scanning photoluminescence spectroscopy. R.Z. supervised the TERS measurement and corresponding data analysis. S.T.H. and A.W.K. were responsible for the model of indentation force and the discussion about polymer interaction. X.L., S.T.H., and A.K.S. analyzed the data. X.L., S.T.H., A.C.R., J.B., and G.B. wrote the paper with input from the other authors. All authors read and approved the paper.

The authors declare no competing financial interest.

Supplementary Material

nl0c03358_si_001.pdf (8.8MB, pdf)

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