carbon nanotube high-range piezoresistive sensor

carbon nanotube high-range piezoresistive sensor

Journal Pre-proofs Enhanced electromechanics of morphology-immobilized co-continuous polymer blend/carbon nanotube high-range piezoresistive sensor K...

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Journal Pre-proofs Enhanced electromechanics of morphology-immobilized co-continuous polymer blend/carbon nanotube high-range piezoresistive sensor K.A. Dubey, R.K. Mondal, Jitendra Kumar, J.S. Melo, Y.K. Bhardwaj PII: DOI: Reference:

S1385-8947(20)30103-0 https://doi.org/10.1016/j.cej.2020.124112 CEJ 124112

To appear in:

Chemical Engineering Journal

Received Date: Revised Date: Accepted Date:

4 September 2019 8 January 2020 11 January 2020

Please cite this article as: K.A. Dubey, R.K. Mondal, J. Kumar, J.S. Melo, Y.K. Bhardwaj, Enhanced electromechanics of morphology-immobilized co-continuous polymer blend/carbon nanotube high-range piezoresistive sensor, Chemical Engineering Journal (2020), doi: https://doi.org/10.1016/j.cej.2020.124112

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1

Enhanced electromechanics of morphology-immobilized co-continuous polymer blend/carbon nanotube high-range piezoresistive sensor K.A. Dubey1,2*, R.K. Mondal1,2, Jitendra Kumar1,3, J. S. Melo1,3, Y.K. Bhardwaj1,2 1Homi

Bhabha National Institute, Mumbai 400094, India

2Radiation 3Nuclear

Technology Development Division

Agriculture and Biotechnology Division

Bhabha Atomic Research Centre, Trombay, Mumbai-400085, India [*Corresponding author E-mail: [email protected]; Fax: 91-022-25505151] Abstract This study reports the development of novel piezoresistor which offer excellent sensitivity for strain measurements, can be used for the detection of high compressive stress and have low creep and fatigue characteristics. Morphology immobilized composites of hard-soft polyolefin blends with co-continuous (co-PEP) and dispersed (d-PEP) morphology were prepared by melt compounding polyolefin blends with Multiple-Walled Carbon Nanotubes (MWCNT) followed by radiation crosslinking. In comparison to d-PEP/MWCNT composites, co-PEP/MWCNT composite exhibited superior electromechanics, in terms of superior strain sensing range and gauge factor, along with low hysteresis. Most importantly, d-PEP did not exhibit compressive stress-dependent piezoresistivity whereas co-PEP showed reversible piezoresistivity in the range 25 kPa to 1 MPa, which is considerably higher than the highest stress-sensing limit typically observed in MWCNT based systems. Fractal dimension analysis of MWCNT network using non-linear rheology revealed that co-PEP/MWCNT composites exhibit soft-aggregates contact breakage and had a much milder dependence on MWCNT volume fraction than binary composites. The cyclic and creep measurements confirmed excellent recovery under static and dynamic conditions, revealing the high potential of such material for standalone stress sensors for robotics, human-machine interfacing, and intelligent transport monitoring. Keywords: Carbon nanotubes, polymer alloys, electrical conductivity, piezoresistors, radiation, hysteresis

2 1

Introduction Conducting polymer nanocomposites (CPCs) are actively explored for piezoresistive applications, owing to

their high sensitivity and versatility [1-4]. Carbon nanotubes (CNT) have intrinsically high electrical conductivity, high aspect ratio and therefore can impart electrical conductivity in polymers at much lower volume fractions than that is needed with conventional fillers such as conducting carbon black and metallic particles [5-6]. Once percolated inside a polymer matrix, the CNT network provides conducting channels that are highly sensitive to external strain/stress, enabling their use as piezoresistors, tensoresistors and other sensors[7-8]. Polymer composites, in general, manifest high mechanical hysteresis, morphological instabilities, uncertain percolation profile and poor up-scalability [9-11]. To minimize mechanical hysteresis and to obtain piezoresistor with a good gauge factor; it is desirable to achieve optimal electrical conduction within the matrix, at low loadings of CNT [12-13]. However, in industrial processes of producing polymer composites such as melt compounding, or in the matrices with low crystallinity (such as elastomers), the percolation threshold is very high [6]. It may be noted that melt compounding does not involve any solvent, and has high industrial fidelity. The functionalization of CNT has been tried to improve the percolation threshold of CNT in a melt-processed polymer matrix. However, both covalent and non-covalent functionalization adversely affects intrinsic electrical conductivity and mechanical properties of CNT, which are desirable for superior electromechanical response [14-15]. Selective percolation, dynamic percolation and use of hybrid fillers are some of the other effective tools to achieve high electrical conductivity in polymer matrices at much lower CNT loading, without affecting any of the CNT intrinsic attributes [16-18]. There is however little understanding, on how such multiphase systems behave under external strain. Yamada et al, in their seminal work, demonstrated that CNT film-based sensors can offer gauge factor (GF) of around 0.82 and can be used for human motion monitoring[19]. Zheng et al. recently developed CNT/CB/PDMS based sensors with a gauge factor of 0.91 by solution casting.

By using laminated fabrication design, in the

CNT/PDMS system, Mortez et al. achieved GF of 2 while marinating a good stretchability[20]. CNT infiltration approaches were used in polyurethane [21] and in PDMS [22] and GF of 1.4 and 0.99 were reported respectively. CNT/PDMS foams were also reported to have a GF of 1.6 [23]. These studies demonstrate high efficacies of CNT based systems for piezo-tensoresistive measurements. However, most of these sensors are based on soft (elastomer such as PDMS) matrix or rely on porous architecture, therefore inherently have very low modulus/stiffness and have challenges as standalone tensoresistors[24]. Furthermore, such sensors being highly soft, are suitable for tactile

3 sensing or for strain monitoring at the surface; however, the stiffness and/or strain at break of such composites is generally too low to allow their use in applications wherein the sensor is expected to withstand a preload and stress. Considering the intricacies and the diversities of different applications wherein strain measurements are needed, the availability of tensoresistors with tailor-able micromechanics and the electromechanical response is expected to provide a lot of flexibility, in terms of sensor design and integration with different objects, structures, and devices. High modulus is directly correlated with the efficacy of a material to withstand load-bearing stress at a particular strain [25]. There have, therefore, been efforts on investigating the electromechanical response of CNT percolated high modulus polymers[26]. In a notable development, Srivastava et al. developed polystyrene /CNT sensors by solution mixing and demonstrated high modulus and GF of 1.5, though the strain range was quite limited[26]. Li et al. have recently fabricated a high modulus CNT based composite with GF of around 2.36 and a strain range of 16% [27]. There is however not much success in developing tensoresistors with high GF, high strain range, and high modulus. Furthermore, there is also not much information available on the morphological, fillerfiller disaggregation dynamics and electromechanical response in CPCs with different morphologies and micromechanics. Polymer alloys/blends can be an effective means to tailor micromechanical attributes of the polymer matrix in CPCs. By varying the composition and morphology of constituents, different physicomechanical property sets can be achieved in polymer alloys. Ethylene-Propylene Terpolymer (EPDM) is an elastomer with an excellent elongation at break and flexibility and is therefore suitable for piezoresistive applications; however, percolation threshold of Multiple-Walled Carbon Nanotubes (MWCNT) in EPDM is very high (this study). On the other hand, polypropylene (PP), a thermoplastic polymer, has a low percolation threshold of MWCNT; however, PP/MWCNT nanocomposites are highly brittle. These facts suggest that binary PP/MWCNT or EPDM/MWCNT composites are of little use for tensoresistive applications. EPDM/PP blends could be a tool to circumvent this limitation and offer tunable physicomechanical properties [28-30]. However, not much is understood about the MWCNT-MWCNT aggregation and disaggregation dynamics under deformation in EPDM/PP blends with different morphologies. Furthermore, the applicability of the morphology immobilization of such nanocomposites for piezoresistive applications is not explored. PP/EPDM system is also interesting on electrical conduction and micromechanics stand-points. Particularly, in addition, to providing a possibility of interfacial or preferential percolation, and,

4 PP/EPDM system offers hard (PP)-soft (EPDM) domains within the matrix, which can be helpful in tailoring different mechanical and electromechanical responses [31-32]. Network density is expected to have a profound effect on the mechanical and electromechanical response of a conducting polymer composite [33]. High energy radiation is an additive-free, room temperature technique by which different crosslinking densities can be imparted to a polymer matrix. EPDM predominately undergoes crosslinking on exposure to high energy radiation [34-35]; therefore, the high energy radiation can be used to immobilize or freeze the morphology hard (PP)-soft (EPDM) –MWCNT based CPCs. This approach may allow the development of optimal crosslinked networks in the EPDM phase without affecting the characteristics of the hard PP phase, rendering a soft phase immobilized thermoplastic conducting composite. This work harnesses the phase-morphology associated variations in the electrical conduction and the micromechanics of an MWCNT percolated hard-soft polyolefin blend system for the development of tensoresistors with improved electromechanics.

The binary and ternary composites were developed by melt compounding and

phase-continuity, electrical conductivity, mechanical hysteresis, and electromechanical behavior were investigated in detail. The nature of MWCNT network architecture in dispersed and co-continuous blends was investigated by impedance and linear rheological measurements. Tensorestive measurements and non-linear rheology were employed to understand MWCNT-MWCNT disaggregation under strain. Finally, high energy radiation was used to immobilize the composite morphology by crosslinking of soft phase and electromechanics of the phase immobilized composites was investigated. 2 2.1

Materials and methods Materials Polypropylene

(density

946

kg/m3)

and

Oil-Extended

Ethylene-Propylene

Terpolymer

EPDM

(VISTALONTM ; ethylene content 63.0 ± 2.0%, ethylene norbornene Content: 4.2 ± 0.3; Oil Content-Paraffinic: 75 ± 3 phr, density :0.86) were procured from local suppliers. MWCNT used in this work had an outer diameter of 8-20 nm, length of 1-2 m and density, 2100 kg/m3. CNT was supplied by Otto Chemie Pvt Ltd, Mumbai India. Solvents such as toluene and xylene used in this work were of analytical grade (M/s SD Fine Chemicals, Mumbai). 2.2

Sample preparation

5 PP/EPDM alloys were prepared by melt compounding in an internal mixer (Brabender Plasti-Corder® LabStation) at 220 oC, 30 rpm for 5 minutes. Afterward, different fractions of MWCNT were incorporated in blends and mixed continued for another 25 minutes. The components were carefully weighed to maintain 80% filling factor in the mixing chamber. The compounded mass was passed through a two-roll mill and then was cut to small pieces. A compression molding machine was used to make sheets of size 10x10 cm2 of different thicknesses (150 kg/cm2 and 200 oC). Sample compositions and designations have been defined as EPX, where X denotes the percentage of EPDM weight fraction. The co-continuous blend with EPDM 60 wt% was termed as co-PEP and dispersed blend with EP 40 wt% was termed as d-PEP, to stress upon the morphological attributes. Volume fraction was mentioned as  or   here x is the EPDM or CNT. 2.3

Sample irradiation Composites were subjected to different doses of Co-60 gamma radiation using a gamma chamber (GC-5000,

M/s BRIT, India) under aerated condition. The dose rate of the irradiation chamber was determined using Fricke dosimetry and the duration exposure was maintained accordingly. 2.4

Mechanical Characterization Dumbbell-shaped samples were cut from compression-molded flat sheets, using sharp-edged steel die. The

thickness and length of the samples were measured nearest to 0.1 mm and five measurements were conducted for each type of composite. The tensile strength and elongation at break were determined using a universal testing machine (M/s Hemetek, MUMBAI, India). 2.5

Equilibrium swelling The crosslinking density of the irradiated composites was estimated from the solvent uptake studies. Before

swelling measurements, the samples were extracted in xylene using a Soxhlet apparatus for 12 h to remove the sol. The gel was then dried for 4 hrs under ambient conditions followed by vacuum drying at 40°C for 8 hrs. Thus prepared, samples were weighed and placed in a 200-mesh stainless steel compartment immersed in toluene at 25 oC.

The swelled samples were periodically taken out from the swelling compartment and weighed. The saturation

weight was used blotted free of surface solvent using laboratory tissue paper and weighed on an analytical balance (accuracy 0.00001 g).

6 2.6

Electrical conductivity and AC impedance DC resistance of the samples was measured by a two-probe arrangement. Before measurement, the samples

were polished and conducting paint was applied on both surfaces. The impedance measurements were carried out at room temperature (HIOKI IM3570) in the frequency range of 4 Hz to 5 MHz. All samples were of disc shape (Diameter = 15 mm; thickness = 1 mm). All measurements were performed at ~24°C and ~50% relative humidity. At least three specimens were tested for each composite. 2.7

Electromechanical measurements For electromechanical measurements under tension, the samples were subjected to a tensile strain and the

change in resistance was monitored using a PC interfaced multimeter (data acquisition rate: 0.5 sec-1) [2, 33]. The gauge factor (GF) was calculated using the following relation. 𝐺𝐹 =

∆𝑅 𝑅 𝑜 𝜀

(1)

𝜺 is the applied tensile strain, ∆𝑹 is change in the resistance, 𝑹𝒐 is the initial resistance. To accommodate dimensional changes during tensile measurements, the relationship between GF and strain can be expressed as the following relation. 𝑮𝑭 =

∆𝑹 𝑹 𝒐 𝜺

=

∆𝝆 𝝆 𝒐 𝜺

+ 𝟏 + 𝟐𝝊 (𝟐)

where  is Poisson’s ratio and o is the baseline electrical resistivity. For elastomers, the Poisson ratio is generally close to 0.5 GF values greater than 2, therefore, it reflects contributions from piezoresistivity [33]. In CPCs, compressive stress may result in the decrease in the distance between conducting junctions, and thus an increase in electrical conductance or the current flowing at a fixed voltage is commonly observed when CPCs are subjected to compressive stress [36-38]. In this work, therefore, the compressive stress and the resulting change in the electrical conductance (M=1/R, mho) or a change in current at a fixed voltage (10 V) was monitored [25, 36-37]. Since ‘G’ is already used to denote shear modulus, conductance was represented as ‘M’. Furthermore, as making electrical contact can also involve small load or pressure, change in conductance relative to conductance at a fixed load of 2N (M/M2N) was used to compare piezoresistors. The typical setup used for measurements is shown in supplementary figure S1a.

7 2.8

Atomic force microscopy (AFM) and Scanning electron microscopy (SEM) Phase imaging and topography measurements were conducted using an AFM instrument supplied by Nanotec

Electronica S.L. systems and devices, Madrid, Spain. Osmium tetraoxide (OSO4) was used to stain the EPDM in co-PEP. The dynamic model was used in such measurements and the scanning tips experiments were silicon probe Tap300-G (Innovative Solutions Bulgaria Ltd., Bulgaria). The cross-section morphology of cryofractured composites was investigated using a scanning electron microscope (Model PS-230, Pemtron, S. Korea, or TESCAN VEGA3, Czech Republic). 2.9

Viscoelastic properties Linear and non-linear rheology measurements were performed using an MCR 102 Rheometer (Anton Par,

Austria). Parallel plate fixture was used and the dimensions of the sample were 12.5 mm (radius) and 1 mm (thickness). A nitrogen atmosphere was maintained throughout the experiment. 3 3.1

Results and discussion Electrical percolation of MWCNT, elastic modulus, and mechanical hysteresis

The electrical conductivity and AC impedance of the composites had a profound dependence on MWCNT and blend composition (figure 1 A-B). It can be seen that at the same volume fraction of MWCNT, the conductivity of EPDM/MWCNT composites is considerably lower than the conductivity observed in PP/ MWCNT or in coPEP/MWCNT composites. In comparison to co-PEP/MWCNT, PP/MWCNT composite has higher conductivity at lower MWCNT volume fractions. At higher volume fractions, however, their conductivity values are relatively closer. Indeed, at 7.0 vol%, PP/MWCNT and co-PEP/MWCNT had almost similar values of electrical conductivity. A comparable level of conductivity (blue horizontal line, Figure 1A) is obtained in EPDM/MWCNT at 17 vol% MWCNT which corresponds to more than double loading of MWCNTs than that is needed in co-PEP/MWCNT or in PP/MWCNT for similar conductivity. Furthermore, at 7.0 vol% of MWCNT, PP/MWCNT composites were completely brittle and inflexible whereas co-PEP composites were completely flexible, reflecting amenability of coPEP to piezoresistive applications. To further understand the effect of blend morphology on the electric conduction, the AC frequency dependence of the impedance of composites at different volume fractions of EPDM is shown in figure 1B. It confirms that EPDM fraction plays a critical role in the electrical percolation, even at 7.0 vol% loading of MWCNT the, the frequency dependence of impedance revealed EPDM/MWCNT composite behaves as an insulator; interestingly though, at same volume fraction of MWCNT, all blend based composites (co-continuous or

8 dispersed) and PP/MWCNT composites displayed characteristics of a conductor. The high conductivity of PP/MWCNT CPCs can be attributed to the high crystallinity of PP. The dependence of impedance on EPDM volume fraction and the lower percolation threshold MWCNT in co-PEP/MWCNT suggest a preferential or selective percolation of MWCNT [6, 39-40]. Young’s equation is often employed to assign filler location in a multiphase system in an equilibrium state. The equation relies on the interfacial energy of the components; it can be described by the following relation. 𝜔𝑎 ― 𝑏 =

𝑇𝐶𝑁𝑇 ― 𝑏 ― 𝑇𝐶𝑁𝑇 ― 𝑎 𝑇𝑎 ― 𝑏

(3)

Herein, a-b is the wetting constant, TCNT-b is interfacial tension between MWCNT and polymer b, TCNT-a, is interfacial tension between MWCNT and polymer a and 𝑇𝑎 ― 𝑏 is interfacial tension between polymer a and polymer b, respectively [41]. Wetting constant can be used to understand the distribution of MWCNT in different phases. 𝜔𝑎 ― 𝑏> 1 suggests MWCNT gets preferentially distributed in phase a. If -1< 𝜔𝑎 ― 𝑏< 1, MWCNT is expected to distribute at the interface; whereas, 𝜔𝑎 ― 𝑏< -1 suggests that MWCNT may have preferential distribution in phase b. The interfacial tension between the two phases (𝑇𝑎 ― 𝑏) can be estimated from the surface energies and their polar and dispersive components. The following harmonic mean average (equation 5) and geometric mean equation (equation 6) for interfacial tension can be described by the following relations.

[

𝑇𝑎 ― 𝑏 = 𝑇𝑎 +𝑇𝑏 ―4

𝑇𝑑𝑎𝑇𝑑𝑏 𝑇𝑑𝑎

+

𝑇𝑑𝑏

𝑇𝑝𝑎𝑇𝑝𝑏

]

+ 𝑇𝑝 + 𝑇𝑝 𝑎

𝑏

𝑇𝑎 ― 𝑏 = 𝑇𝑎 +𝑇𝑏 ― 2[ 𝑇𝑑𝑎𝑇𝑑𝑏 + 𝑇𝑝𝑎𝑇𝑝𝑏 ]

(4) (5)

𝑇𝑎 and 𝑇𝑏 are the surface tension of pure phases a and b. Superscripts d and p denote the polar and dispersive components of the surface tension respectively. Both geometric mean and harmonic mean equations were used to determine interfacial energy (table 1). The surface energy value for MWCNT was taken as 45.3 with a polar component of 26.9 [41-42]. This value was established by Nureil et al. using electron microscopy and was validated by Petra et al. in PE/PC and PE/PA blends [41]. The value of surface energy of PP and EPDM was taken as 30.1 with polar component 0 and the 21.6 with a polar component of 10.9 respectively [43]. The interfacial energy of 29.7 N/m for PP/MWCNT and 7.03 N/m for EPDM/MWCNT was obtained using the harmonic mean equation and the values were not much different from that obtained using the geometric mean equation. Wetting constant (𝜔𝑎 ― 𝑏) for the present system was found to be -1.85 and -2.12 from harmonic mean surface energy and geometric mean results (table 1). These values suggest that MWCNT is preferentially distributed in EPDM domains. van Gurp-

9 Palmen plots depicting the effect of MWCNT loading on the phase angle in co-PEP are shown in figure 1C. With the addition of MWCNT solid-like behavior emerged, indicating the formation of the MWCNT network. Increasing the MWCNT volume fraction from 0.045 to 0.096 resulted in the shift in the complex modulus. Dependence of  on shear modulus indicates the formation of the 3d network [44-45]. Fractured SEM images of co-PEP shows blend morphology and MWCNTs (Figure 2).

The specimens were cryofractured and cross-section was stained with

OSO4 to attain phase contrast. Some of the MWCNTs, which are pulled out due to cryofracturing, are visible at the cross-section. Co-continuity of morphology is discussed in detail in the next section. The above analysis confirms that electrical conduction in PP/EPDM/MWCNT is morphology dependent. Apart from electrical conduction, matrix elasticity is a major determinant of piezo- or tensoresistivity. Elastic modulus and elongation at break of the composites containing different fractions of EPDM are shown in figure 1D. PP/MWCNT composites were very brittle albeit had high elastic modulus. PP/MWCNT nanocomposites had an elongation at break of about 2%; at 40% EPDM content, elongation at break increased to ~15 % and at 60% it increased to ~50%. A sudden rise in elongation at the break between d-PEP and co-PEP suggests phase inversion or formation of co-continuity, as discussed in detail in the next section. It may be stressed that though PP/MWCNT showed considerably higher conductivity at lower MWCNT volume fraction but the composite was too brittle to be suitable to tensoresistive applications. Conversely EPDM requires considerably higher loading of MWCNTs to achieve conductivity; at such a high loading of MWCNT, processability, flexibility, and elasticity of the EPDM/MWCNT are greatly compromised.

The co-PEP composition (co-continuous blend) mitigates these

disadvantages by offering good flexibility and good electrical conduction. Notably, co-PEP had an elastic modulus of 80 MPa which is considerably higher than that for pristine EPDM [34, 46-47]. As discussed in later sections, due to the high Young’s modulus, co-PEP/MWCNT CPCs composites could withstand high load and show reversible piezo-resistive behavior. Five successive loading-unloading cycles for co-PEP and d-PEP are shown in figure 1E. With the increase in EPDM content, there was a significant reduction in mechanical hysteresis, as is evident from the size of the loop observed during loading-unloading cycles. In co-PEP there was insignificant variation in the first two cycles; therefore, I and V cycles were used for evaluating elastic recovery. Recovery parameter during loading-unloading cycles can be defined in different terms such as toughness, stiffness, modulus, or ratio of dissipated energy. As the main objective of cyclic stress-strain analysis to understand energy dissipation and hysteresis, in the current work,

10 the work done in the first and fifth cycles was used to evaluate hysteresis characteristics of composites. The ratio of work done in the first and fifth stress-strain cycle (𝑅𝑊𝐷) was therefore calculated for co-PEP and d-PEP. 𝑅𝑊𝐷 =

𝑊𝑉 𝑊𝐼

(6)

WV and WI are the work done in the first and the fifth cycle respectively. 𝑅𝑊𝐷 was 89.8 for d-PEP and was 94.7 for co-PEP. Polymer chains initially have a coil structure; under uniaxial strain, chains straighten up and regain the coil conformation once the stress is removed[48]. This phenomenon is expected to be more effective in EPDM chains than for PP chains, as the high crystallinity of PP matrix forbids the complete stretching and de-stretching of PP chains [7, 30, 48]. 3.2

Co-continuity characterization Figure 3A presents the van Gurp-Palmen plots of different PP/EPDM blends. A distinct feature emerged

when the volume fraction of EPDM reaches 0.623. The sudden decrease in the phase angle suggests phase inversion and increased the elasticity of the matrix. Since after this composition the phase angle is well below 45 degrees for the entire range of |G*|, it is expected that these composites will display solid-like behavior. Complex viscosity also exhibited a sudden change in this composition confirming the morphological transitions (Figure 3B). Co continuity index was determined by selective leaching of EPDM in toluene at 25 oC for 15 days [49]. PP did not show any weight loss during these conditions, confirming the insolubility of the PP phase (Figure 3C). In blends, however, extractable EPDM displayed deviation from the actual amount of EPDM present. Up to 0.423 volume fraction, the extractable EPDM (weight loss) was significantly less than the EPDM content of the blends, thus pointing out of the dispersed morphology. At 0.623 volume fraction, EPDM co-continuity was evident from the fact that extractable EPDM matches well with the actual EPDM content. To confirm the phase inversion and co-continuity, toluene uptake studies (saturation) were also conducted after crosslinking the blends at 100 kGy of gamma radiation (Figure 3B, right axis). Insignificant toluene uptake was observed till 0.423 volume fraction followed by an abrupt increase at 0.623 volume fraction of EPDM, establishing phase continuity in this composition range (Figure 3C). To understand the compatibility of EPDM and PP phases, DMTA studies were done (Figure 3D). Both phases showed a distinct glass transition, highlighting immiscibility. SEM images of the cryo-fractured leached blends are shown in figure 4. At 0.216 volume and 0.423 volume fraction, the dispersed morphology of EPDM was evident. At 0.623 volume fraction, a co-continuous network was exhibited. AFM images of the leached samples corroborated SEM

11 findings (Figure 5 A-C). AFM phase imaging studies of un-leached blends (EP20) manifested the spherulites of isotactic polypropylene whereas co-continuous blend morphology was evident at 0.623 volume fraction (Figure 5DE). 3.3

Non-linear rheology Strain dependence of the shear modulus is related to polymer-filler interface and filler-filler network

architecture in a polymer composite. In all co-PEP60/MWCNT composites, with an increase in strain amplitude, the shear storage modulus decreased after a critical strain c (Figure 6A). Noteworthily, c decreased with the increase in MWCNT and increased with an increase in EPDM (Figure 6-7).

Shear stress-induced filler-network structure

breakdown is the prime factor determining the magnitude and nature of this non-linear strain dependence. Under shear strain, the rheological response from a polymer composite matrix provides information about contributions from different components, for example at the low strain, the elastic contribution of filler network is dominant on the other hand at higher strain, hydrodynamic and polymer-filler interactions play a major role. The extent of filerfiller contact breakage and polymer-filler interface breakdown was analyzed by fitting the amplitude sweep data to the phenomenological quantitative Kraus model [50-52]. 𝐺′(𝛾) = 𝐺′∞ +

𝐺′𝑜 ― 𝐺′∞ 1+

𝛾 2𝑚 𝛾𝑐

()

(7)

𝐺′(𝛾), 𝐺′𝑜and 𝐺′∞ are shear storage modulus at shear strain , at very low strain and very high strain respectively. c is the critical strain and m is the structure factor to the fractal dimension and depends on the geometric quality of the filler network [2]. The experimental data fitted well to the Kraus model (in all cases: r2>0.99) and respective the fitting parameters are presented in Figures 6 and 7. This implies the dominance of filler-filler disaggregation over polymer-filler interaction in the shear strain-induced structural breakdown. 𝐺′(𝛾) increased with MWCNT fraction showed a power-law dependence on MWCNT volume fraction (the difference between 𝐺′(𝛾) the composite and of the unfilled polymer is plotted in figure 6B). Conversely, the critical strain was found to decrease with an increase in the volume fraction of MWCNT, exhibiting negative power-law dependence (Figure 6C). The value of m, which is suggestive of strain sensitivity of matrix, is expected to be close to 0.5 in filled systems. Interestingly, however, m was found to have volume fraction dependence and only at higher MWCNT volume fractions, it approaches to 0.5 (Figure 6D).

12 In filled systems, the filler network can be treated as a fractal object and following aggregation disaggregation model-based scaling relations can be defined. 𝛽 (𝑑 ― 𝑑 ) 𝑓

𝐺′ ∝ ∅

(8)

(𝑑 ― 𝛽 ― 1) (𝑑 ― 𝑑 ) 𝑓

𝛾𝑐 ∝ ∅

(9)

Herein, df is the fractal dimension of the network, d is the Euclidean dimension and β is given by the following relation. 𝛽 = (𝑑 ― 2) + (𝑥 + 2)(1 ― 𝛼) (10) x represents the fractal dimension of the backbone and has a value between 1 and df.  is an important parameter related to interflock and intraflock strength. For co-PEP60/MWCNT composites, the value of df and  were found to be 1.99 and 2.24 respectively.  was 0.58<<0.69 for 1
13 and m is presented in figure 7C-D, for unfilled and MWCNT filed systems. It is evident that MWCNT filled systems had lower c at all EPDM fractions than the corresponding unfilled composition. Moreover, the unfilled system does not exhibit a systematic EPDM fraction dependent change in critical strain, whereas MWCNT filled systems showed the EPDM dependent increase in critical strain. The differences in the percolation threshold of MWCNT in PP and EPDM phases and preferential portioning of MWCNT in the EPDM phase are the major factors contributing to this shear strain dependence. The chain relaxation spectrum of the blends also analyzed (Figure 7B). PP exhibited a single relaxation peak at around 4.0 s, whereas EPDM showed solid-like behavior. At compositions EPDM=0.216 and EPDM=0.423, one relaxation peak at a higher time scale and a shoulder at lower time scales were noted. The first can be ascribed to the PP domain and the second peak to the interface. At EPDM=0.815, only one relaxation peak at about 2.0 s was observed. co-PEP i.e EPDM=0.623 showed solid type behavior with a shoulder at 4.0 s, confirming the co-continuity of the phases and their individual identity. Compositions with higher EPDM fractions i.e EPDM>0.423 are analogous to dynamically vulcanized systems hardened by physical crosslinking. For the practical standpoint on the filler-filler network disaggregation dynamics, the above analysis critically established the role of EPDM in enhancing the c and thereby affecting mechanical hysteresis and electromechanical response. 3.4

Linear rheology The complex melt viscosity showed frequency dependence and changed with a change in EPDM or in

MWCNT content (Supplementary figure 1b-d). Almost over the entire frequency range (0.1-100 rad/s), EPDM showed higher complex viscosity than PP; though, in the terminal frequency reason, the difference was considerably higher. With an increase in EPDM content, the viscosity was not significantly changed till 40% and abruptly at coPEP60, confirming the development of co-continuous morphology. It is noteworthy that the terminal frequency plateau was observed only in unfilled PEP0, PEP 20 and PEP40. With the increase in MWCNT loading, a volume fraction dependent increase in the viscosity was noted [56]. Variation of storage modulus and loss modulus with angular frequency for unfilled PEP blends, MWCNT filled blends and co-PEP/MWCNT composites were also investigated (supplementary figure 2-4). Pristine PP showed G’’ values higher than G’ and a crossover was noted at higher frequencies. With the addition of EPDM, the crossover was shifted to a slightly higher frequency and at 60% EPDM, the behavior was reversed i.e. G’ was higher than G’’ over the entire frequency range. At MWCNT=0.07, all blends exhibited G’>G’’, confirming the formation of

14 a continuous network of filler [57-58]. In the terminal frequency region, the frequency dependence for G’ was close to zero in PEP20, and d-PEP nanocomposites and small but non-zero for other systems [59]. 3.5 3.5.1

Tensorestivity Phase morphology, electron tunneling, sensitivity and range The relative change in resistance with the increase in uni-axial tensile strain for co-PEP/MWCNT

composite and d-PEP/MWCNT composite is shown in figure 8A. The resistance increased with the applied strain for both CPCs; though, the respective response profiles were considerably different. d-PEP/MWCNT showed an increase in resistance with the applied strain; however, it could sustain up to 13 % applied strain. On the other hand, in co-PEP/MWCNT, the strain range increased close to 50% i.e. around three times higher than that observed in dPEP nanocomposites. This can be attributed to the co-continuous morphology and the preferential percolation of MWCNT in the EPDM matrix [7].

The strain-induced change in the resistance of CPCs can be attributed to the

changes in the number of MWCNT junctions and the MWCNT-MWCNT distance in the percolated network. These factors affect the tunneling resistance and the current-carrying ability of CPCs. As a result, the relative change in resistance increases with strain, and eventually, after a particular distance, the inter-MWCNT distance exceeds the tunneling distance, leading to the complete disruption of conducting channels. The total resistance of the CPCs can be described by the following relation.

(

𝐿 8𝜋ℎ𝑤

𝑅=𝑁

4𝜋𝑤 2𝑚𝜑 ℎ

)𝑒𝑥𝑝(

2 2

3𝑟𝑎 𝑒

)

(11)

L represents the number of MWCNTs forming a single conducting channel and N denotes the total number of such conducting channels. h is the Planck’s constant, w is the width of tunneling junction, e is the electron charge, m is electron mass, 𝜑 is the tunneling potential barrier and a2 is the effective cross-section area. Under strain , N varies as per the following relation. 𝑁0

𝑁 = exp (𝑀𝜀 + 𝑊𝜀2 + 𝑈𝜀3 + 𝑉𝜀4)

(12)

M, W, U, and V are constant and N0 is the total number of conducting paths at zero strain. The width of the tunneling junction will vary as per w = w0 (1+C). The resistance, therefore, can be defined as per the tunneling theory by using the following relation. 𝑅 = 𝐵(1 + 𝐶𝜀)exp (𝐴 + (2𝑀 + 𝐴𝐶)𝜀 + 2𝑊𝜀2 + 2𝑈𝜀2 + 2𝑉𝜀2) (13)

15

A is

8πhw0 4𝜋𝑤 2𝑚𝜑 and B is . The electromechanical response of co-PEP/MWCNT followed the tunneling theory ℎ 3ra2e2N0

(dotted profile, r2~0.98) [60]. The tensoresistor exhibited a high dynamic strain range (strain~50%) and high GF. As observed by Shi et al., the electromechanical response can be categorized into three distinct strain regions [10]. GF in the lower, middle and terminal strain range were 13.05±0.16, 87.85±1.18 and 434.73±31.07 respectively (Fig 8A, linear fit in the region I, II & III). Even if, the entire range is considered, the GF at the terminal point was ~260. Corroborating our observations, Liu et al. observed that the electromechanical response of PU/graphene-based CPCs follows tunneling theory and the variable gauge factor can be obtained by playing with the morphology/composition[61]. Recently, Hu et al. have also reported different gauge factors in different strain ranges[62]. As the linear range has more practical relevance, for d-PEP and co-PEP, the linear gauge factor was determined for initial strain in which both tensoresistors followed a linear response (strain 10%; from 3-13 %; strain rate 1 mm/min). It was 5.43±0.07 (r2>0.99) for d-PEP and 12.73±0.04 for co-PEP/MWCNT (r2>0.99), clearly indicating that higher EPDM (soft phase) not only improves the strain range but also the sensitivity (gauge factor). Nonetheless, d-PEP offers impressive characteristics namely high stiffness (elastic modulus ~ 700 MPa), about 13% stretchability and GF around 5. It is, therefore, a promising standalone tensoresistor for the applications demanding high stiffness/ low deformation under load conditions. Indeed, the linear gauge factors of d-PEP/MWCNT and coPEP tensoresistors was considerably higher than that reported in several other important studies on MWCNT based composites [19-23, 27, 63]. Zheng et al., recently developed MWCNT/CB/PDMS based sensors with gauge factor of 0.91 and, by using laminated fabrication design, in MWCNT/PDMS system, Mortez et al. achieved GF of 2 [20]. On the other hand, MWCNT infiltration approaches were used in polyurethane[21] and in PDMS [22] and GF of 1.4 and 0.99 were reported respectively. MWCNT/PDMS foams were also reported to have a GF of 1.6 [23].

It

may be noted that GF is only one of the several criteria for strain sensing applications, and strain sensors with GF as low as 0.82 can be used human motion detection as demonstrated by Yamada et al.[19]. The integration of tensoresistors with objects in different applications warrants different physicomechanical and electromechanical characteristics. For example, tactile sensing demands low Young’s modulus sensors whereas structural health monitoring sensors will benefit from higher values of Young’s modulus [3, 64]. It is therefore important to have tensoresistive materials with different physicomechanical attributes. In that respect, d-PEP/MWCNT and co-PEP both offer distinct advantages.

16 In the recent past, there has been a substantial effort to develop advanced strain sensors for different application requirements. Boland et al. fabricated PDMS/graphene (6.8 vol % of graphene) CPCs which offer a high gauge factor ( 65 and 450 up to 1% strain) [2]. In polystyrene/MWCNT based CPCs, Shrivastava et al. reported a gauge factor close to 4 within a micro-strain range[26]. All these and other studies on the electromechanics of piezoresistors/tensoresistors highlight the complex interplay of filler-filler network disaggregation-aggregation dynamics under strain and the stress transfer within the matrix[65]. In MWCNT/PMMA composites, a gauge factor of 4.4 was observed along with high modulus, though the strain range was less than 5%[64]. Amjadi et al. reported a gauge factor ranging from 2 to 14 in silver nanowire/ elastomer-based CPCs[11]. Georgousis et al. have reported strain sensing behavior of PVDF/MWCNT nanocomposites and claimed that the resistance increases slowly with the applied strain (lower slope) but a rapid increase near breakpoint (6% strain) [25]. These results point out the advantages of co-PEP60/MWCNT based in terms of high modulus (80 fold higher than that of EPDM) yet high dynamic strain sensing range. In another important study wherein conductive thermoplastic vulcanizates based CPC were developed, high strain sensitivity was noted; depending on the CPC synthesis procedure, the composites showed low GF at lower strains and high GF at the higher strains[66]. Such deviations are expected to stem from the fact that resistance measurements were conducted much beyond the linear elastic region. Indeed, phenomenally high GF at the terminal region is a characteristic of the conductor to insulator transition in CPC and/or of the mechanical failure of the matrix [4, 12, 32, 40]. In pursuits to obtain super-sensitive sensing, considerable emphasis has been made to develop CPCs with high gauge factors. PDMS has been an attractive choice in this regard and markedly high gauge factors have been demonstrated in several systems [67-68]. However, in most of the cases, GFs are generally reported in the extreme of the stress-strain curve where conductor to insulator transition takes place or such CPCs offer low strain range or/and extremely low elastic modulus. For example, in recently reported advanced systems such as PEDOT-SS[69] and PDMS/graphene[68], the strain range and the modulus are very low. Boland et al. have reported GF of 450 and 65, only at less than 1% strain [2]. Such sensors are suitable for tactile sensing and similar applications; however, the stiffness and/or strain at break of such composites are too low to allow their use as standalone sensors, particularly in applications wherein the sensor is expected to withstand a preload. This situation is common in structural health monitoring applications [3, 64]. High modulus along with acceptable gauge factor, as is offered by co-PEP CPCs, therefore, makes possible the tensoresistive measurements in intricate objects/structures.

17 3.5.2

Improvement in GF and strain dependence by high energy radiation Network density affects the elasticity of the matrix as well as the strain sensing behavior. Figure 8B

represents the relative change in resistance with an increase in strain for the samples crosslinked by employing gamma radiation. Supplementary figure 5 depicts the stress-strain profile of irradiated composites. The increase in the elastic modulus after irradiation confirms the formation of crosslinked networks and an increase in network density with an increase in radiation dose [70-72]. It can be seen that radiation dose also significantly affects the strain sensing response (Fig 8B). The sample irradiated to 50 kGy had almost as high strain sensing range as was observed for an un-irradiated sample but its slope in linear range was significantly higher. The un-crosslinked sample showed an abrupt rise in resistance at 25% strain whereas the crosslinked one had a linear increase up to 35% (r2 ~0.99). Interestingly, the sample irradiated to 100 kGy showed an almost linear rise in resistance before it broke at 30% elongation. Gauge factor for all three systems was determined in 3-13% strain range and was found to be 12.5±0.04, 15.36±0.08, 16.02±0.05 at 0, 50 and 100 kGy respectively (strain rate 1 mm/min). The irradiated samples were melting processed again and there was an insignificant change in the EPDM domain size, swelling, and dielectric permittivity, confirming crosslinking and immobilization of EPDM domains. Radiation-induced enhancement in the gauge factor can be attributed to the immobilization of the EPDM phase. Due to immobilization, more force is required to attain a specific strain, as the elastic modulus is directly correlated to network density (Inset Fig S5). Since the extent of disaggregation of filler network depends on the applied strain as well as the magnitude of the force applied to attain such strain (internal stress), the change in resistance is expected to be more in the matrices that have higher network density [13, 39, 73]. Furthermore, the stress transfer, due to physical crosslinking as well as grafting of polymer chains onto the nanotube surface can also contribute to the enhanced sensitivity of irradiated piezoresistors [74]. Dependence of strain rate on the electromechanical response of coPEP60/MWCNT irradiated to 100 kGy is shown in figure 8C. It can be seen that the tensoresistors had a little dependency on the strain rate. Notably, at all strain rates, the phase immobilized PEP tensoresistors offer excellent linear gauge factor (>15), much higher than Wheatstone bridge-based conventional tensoresistors, comparable or higher than recently reported GF (linear) for MWCNT based advanced tensoresistors[62, 69, 75-76]. 3.6

Piezorestivity

co-PEP/MWCNT and d-PEP/CPCs were subjected to the different magnitude of cyclic compressive loads. Unlike under tensile (stretching) load, under compressive load, the neighboring MWCNTs in CPCs are expected to comes

18 closer, leading to an increase in electrical conductance (M=1/R). Electrical conductance or the relative change in the current at constant voltage is therefore generally used for evaluating the electromechanical response of piezoresistors under compressive load[77]. A systematic increase in the conductance was observed with an increase in compressive stress for co-PEP/MWCNT (Figure 9A), and the conductance attained baseline values on the removal of load (Inset Figure 9A). The baseline value of conductance was determined at the load of 2N. dPEP/MWCNT CPCs did not show any load-dependent increase in the conductance (Figure 9B); though there was an increase in the conductance with the application of compressive stress, the increase does not vary in proportion to the applied stress. This observation reflects the critical role of co-continuous morphology is attaining reversible electromechanical response. Recently, Wang et al. developed a self-segregated structure-based PDMS/MWCNT nanocomposites and demonstrated significant improvement in the reversible electromechanical behavior under compression[78]. This enhanced piezoresistivity was attributed to the presence of segregated domains of PDMS within the continuous matrix of PDMS. In co-PEP/MWCNT, a similar effect is expected to contribute to the electromechanical response, except for the fact that herein segregation is maintained by two phases having considerably different values of compressive modulus (PP~ 1.5 GPa; EPDM~ 5MPa). When PP and EPDM phases have co-continuity, the elastic properties of EPDM can be extended throughout the matrix; furthermore, since the EPDM phase has been crosslinked by radiation there is a further improvement in matrix elasticity. On the other hand, when the EPDM phase is a dispersed phase, the mechanics of the system are largely governed by the thermoplastic PP phase which has limited elasticity. Restoration of the original conducting network on the removal of stress is an essential criterion to get a reversible electromechanical response from a CPC. In a seminal work, Amjadi et al demonstrated the role of polymer encapsulation layer in mitigating irreversible disconnections in conducting particle networks [11]. They demonstrated that in silver nanowires (AgNW)/PDMS based strain sensors if AgNWs are encapsulated with PDMS, the AgNW follow their path back without buckling after removal of load/strain. These studies point out that, under cyclic loads, the elasticity of the host matrix is critical in assuring reversible changes in the microstructure of CPCs. Significant changes in the properties on the development of cocontinuous morphology were also evident in the mechanical and rheological properties discussed above. In coPEP/MWCNT composites, the non-linear rheological analysis revealed breakage of soft fractals in co-PEP/MWCNT composites and much milder dependence of critical strain on MWCNT volume fraction. The critical strain reflects the extent of deformation at which irreversible changes in percolated MWCNT network start to emerge. These

19 observations indicate better stability of conducting particle networks within co-PEP/MWCNT to irreversible disconnections under cyclic loads. These results also suggest a superior range, fatigue and creep stability of coPEP/MWCNT piezoresistors. These properties are discussed in the following paragraphs. Recently, electrical properties of CPCs under compressive load have received considerable attention for pressure and tactile sensors for biomedical, structural health monitoring and robotics applications. Most of the reported piezoresistors perform considerably well in the low-pressure regime but the response tends to saturate after 50 kPa[37, 77]. In this regard, the range exhibited by co-PEP/MWCNT CPCs (Figure 9A) is noteworthy and such piezoresistors may find applications in conditions wherein 50 kPa or more pressure is experienced; namely, machine-human interface, intelligent transportation system, sports injury and foot ulceration monitoring [11, 23, 77, 79]. Figure S6 shows the effect of two MWCNT volume fractions on the resistance of the CPC, it can be seen that at lower volume fraction there was no detectable signal till 400 kPa; furthermore, the change in relative change resistance with increase in compressive stress was also lower.

Importantly, small buttons of the dimensions of 3.0

mm x 1.5 mm (Figure 9C) are used in the loading-unloading cycles depicted in Figure 9A, demonstrating the miniaturization feasibility of these piezoresistors. A very simple circuit using these 3 mm diameter piezoresistors, a 10 V power supply, and an LED is shown in figure 9D. The LED glows when the piezoresistor is subjected to 4 MPa compressive stress, depicting the utility of piezoresistors as a stress-switch for robotic grips or for humanmachine interfacing. Depending on input voltage switches or stress sensors can be designed with the range mentioned above. In addition to excellent electromechanical properties, complete flexibility and melt processability are other advantages of co-PEP/MWCNT CPCs. Such a broad range of piezoresistivity is not previously reported in melt compounded polymer-based flexible sensors. To explore the mechanism of piezoresistivity, AC impedance measurements under different compressive stresses were conducted (figure 10 )[80]. The total AC impedance of the matrix percolated with conducting filler is expected to be governed by an R-C circuit shown in the inset figure 10A. RCNT is the resistance of MWCNTs, Rc is the contact resistance between two MWCNTs and CGap is the capacitance of the gap. In the case of heterogeneous systems such as CPCs, the capacitance of the gap i.e Cgap can be modeled as a constant phase element (CPE) with the following relation with the impedance (ZCPE)

𝑍𝐶𝑃𝐸 =

1 𝑄(𝑗2𝜋𝑓)𝑛

(14)

20 where n and Q are constants, and f is frequency. The total impedance of the circuit and its dependence on angular frequency () can be described as 𝑅2𝑐 𝐶𝑔𝑎𝑝

𝑅𝑐

𝑍 = 1 + 𝜔2𝑅2𝐶2 -j 1 + 𝜔2𝑅2𝐶2 𝑐 𝑔𝑎𝑝

𝑐 𝑔𝑎𝑝

(15)

The plot of real impedance versus imaginary impedance was plotted at different compressive stress and was found to be semi-circular in nature. The equivalent R-C circuit fitted well with the experimental data. The Rc was 2.5x106 ohm at 0 kPa but with an increase in compressive stress, it decreased to 1.6x105 ohm at 200kPa and to 2.6x104 ohm at 2000 kPa. These results suggest that the piezoresistivity in co-PEP/MWCNT CPC can be attributed to the decrease in contact resistance between MWCNTs with an increase in the compressive stress. Figure 11 shows a schematic representing hard-soft co-continuous morphology of co-PEP. Recently, Tsui et al. presented a detailed discussion of the behavior of piezoresistors under cyclic loads. They developed graphene-coated SWCNT aerogels which exhibit remarkable reversible piezoresistivity properties up to 80 kPa[36]. Tsui et al. attributed the excellent reversible electromechanical properties of graphene-coated SWCNT to the graphene led reductions in drastic microstructure changes under cyclic deformations. In particular, they suggested that under compression, graphene bends and these bent junctions exert spring-like restoration force. In d-PEP/ MWNCNT CPCs wherein the EPDM phase is a dispersed phase, the elastic effects of EPDM are localized; therefore, spring-like restoration forces due to the crosslinked network are not as effective as they are in co-PEP/ MWNCNT CPCs. This is also evident from the fact that d-PEP/MWCNT CPCs are highly brittle whereas co-PEP/MWCNT composites are completely flexible and can withstand thousands of cyclic deformation. Fatigue and creep behavior of the CPCs are important for long term applications. Due to viscoelasticity, stressdependent changes also depend on the duration for which stress is applied on a CPC. Though creep and fatigue cannot be completely avoided, it is desirable to minimize the influence of such factors during application time. coPEP/MWCNT CPC was subjected to the constant compressive stress ~ 2.0 MPa for a period of 4 hrs followed by a recovery and creep at a lower load of 10 kPa for a 2 hrs period (Figure 12 A). It can be seen that with the application of load the conductance increased sharply; most importantly, on the removal of the load, there was an instantaneous drop in the conductance. The stress-strain profile during these creep measurements is shown in figure 12B. It can be seen that during the constant load period of 4 hrs there is about a 10% increase in the strain; however, upon the removal of the load, the strain accrued during a 4-hours period recovered. Furthermore, as the residual

21 stress removed from the matrix, the strain became almost zero and original dimensions were regained. As pointed out by Tsui et al. creep is an undesirable factor and affects the reliability of piezoresistors[36]. Due to the viscoelasticity inherent in polymers, it is difficult to avoid creep; however, in the present case, the creep extent is mitigated by the formation of co-continuous morphology and radiation-induced crosslinking. Co-continuous morphology can be approximated as an EPDM filled PP sponge, combining elasticity of EPDM with deformational stability of PP, enabling excellent recovery of residual strain. Figure 12C shows stress-strain profiles for increasing cyclic compressive stress from 50 kPa to 6 MPa. For each stress, 10 loading-unloading cycles were used. The elastic recovery was calculated in each cycle and was found to be more than 95% in each compressive loading-unloading cycle. These results demonstrate the capability of these co-PEP/MWCNT CPCs to withstand cyclic compressive stress of different magnitude. Such results are uncommon in CNT based CPCs wherein mechanical hysteresis a serious concern. For example, recently Yu et al. have developed carbon nanotubes in one-dimensional polymer fiber structure and demonstrated superior piezoresistivity to improve elastic recovery [81]. They reported about 90% elastic recovery and related it to the excellent dynamic reliability and durability of CNT based piezoresistors. Considering the hysteresis inherent in percolated polymer composites, low residual strain observed in co-PEP/MWCNT CPCs in dynamic measurements is particularly noteworthy. Dynamic stability of the signal from piezoresistors was also evaluated by subjecting them 10 loadingunloading cycles of 1 MPa at different stress rates from 25 kPa.s-1 to 250 kPa.s-1. It can be seen that the sensor output followed the input stress signal (Figure 12D). Yu et al. recently used the attenuation degree of sensitivity to estimate the dynamic reliability and durability of piezoresistors[81]. In dynamic measurements, this term essentially denotes the change in the sensor response in a successive loading-unloading cycle with respect to the response obtained in the first cycle. In SBS/CNT based one-dimensional superior piezoresistors, Yu et al. observed that the attenuation degree of sensitivity was considerably high when piezoresistors were subjected to 5-cyclic deformations at different rates. Notably, in co-PEP/MWCNT CPCs, which was subjected to 10 cycles of 1 MPa load at different rates, there was only 5.6% change in the sensor response when the compressive stress rate is increased by 1000%, reflecting good dynamic reliability and durability of these piezoresistors.

22 Conclusion Co-continuous polyolefin/MWCNT composites with immobilized phases were found to be effective for developing advanced tensoresistors and piezoresistors with superior elastic modulus, gauge factor, linear range, and mechanical hysteresis. Even at a fixed loading of MWCNT, the gauge factor was higher in the nanocomposites with cocontinuous blend morphology than it was in the dispersed blend morphology and only co-continuous morphology exhibited load-dependent piezoresistivity. Non-linear rheology suggests the formation of soft aggregates and robust networks in co-PEP/MWCNT composites which enable matrix to sustain larger deformation strains. High energy radiation-induced cross-linking and immobilization of morphology of co-PEP/MWCNT composites further improved the gauge factor. The results suggest that different electromechanical responses can be attained by varying hard/soft domains ratio and radiation-induced phase immobilization. The selective crosslinking of the soft phase coupled with a preferential percolation of MWCNT is, therefore, an effective modality to develop advanced tenso/piezoresistors with tunable electromechanics. As these composites are prepared by melt compounded, the process can be easily upscaled. Furthermore, as the elastic modus of these tensoresistors is significantly higher than the conventional elastomer-based tensoresistors, these tensoresistors will provide to match design material properties that can be used in applications that demand preload, particularly in human-machine interface, intelligent transportation systems and robotics. Acknowledgment Authors thank Dr. P.K. Pujari, Associate Director, Radiochemistry & Isotope Group, BARC for his encouragement and a keen interest in this work.

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28 Table 1: Values of interfacial tension (N/m) between various components Component A

Component B

a-b (harmonic mean)

a-b (Geometric mean)

PP

EPDM

12.3

11.6

PP

MWCNT

29.7

28.3

EPDM

MWCNT

7.03

3.68

Table 2: Fractal structure parameters for co-PEP60/MWCNT composites derived from non-linear rheology 

df



Composite

dG/dMWCNT

UPR/ MWCNT[54]

1.71

-0.55

1.276

0.351-0.405

0.339

UPR/ SWNT [53]

2.84

-1.36

1.649

0.054-0.222

0.261

Epoxy/ MWCNT [55]

3.74

-1.38

2.153

0.277-0.478

0.316

co-PEP/ MWCNT

2.22

-0.244

1.988

0.584-0.687

0.45

PP/ MWCNT

3.51

-1.397

2.1

0.226-0.427

0.301

d/dMWCNT

Table 3: Effect of strain rate on the gauge factor of co-PEP60/MWCNT tensoresistors Strain rate (mm/min)

Gauge factor

10

23.10.7

5

21.40.5

2.5

19.30.2

29

10

7

10

5

10

3

10

1

EPDM

( B)

0.04

0.08

0.12

0.16

0.20

( co-PEP

10

 (CNT)

Elastic Modulus (MPa)

0.2

10

0.4

0.6

0.8

3

2

4

10 Frequency ( Hz)

60

40

10

10

2

20

( D)

1

0.2

0.4  0.6 EPDM

d-PEP

PP

0.8

Stress (MPa)

( A)

Elongation at break (%)

Z

1E-12

1E-16 0.00

9

)

EPDM

10

o

PP

)

1E-8

co-PEP

(

DC (S/m)

1E-4

12 10 8 6 4 2 0 -2

10

6

( E)

50 45 40 35 30 25 20 15

CNT=0.045 CNT=0.096 CNT=0.000

( C) 10

4

|G*| (Pa)

co-PEP

10

5

d-PEP

co-PEP

5

6

7 8 9 Strain (%)

10

Figure 1: Electrical, mechanical and MWNT percolation characteristics of composites: (A) Variation of electrical conductivity of PP, co-PEP60 and EPDM with reduced MWNT volume fraction [log-log plot] and linear fitting, percolation threshold (c ) is also presented] (B) Variation in impedance with change in EPDM fraction at a fixed MWNT loading (MWNT=0.07) (C ) phase angle as a function of shear modulus for different MWNT content at a fixed EPDM loading (co-PEP) (D) Mechanical properties of composites as a function of EPDM volume fraction at a fixed MWNT loading (MWNT=0.07). (E)Stress strain loops for five cyclic deformations for d-PEP and co-PEP/MWNT composites (MWNT=0.07).

30

(A)

(B)

(C)

Pulled out MWCNT

(D)

(E)

Figure 2: Morphological features of co-PEP composites (MWNT=0.07) (A) OSO4 stained cryofractured cross-section, secondary electron image, 5kx, arrow shows stained phase (B) OSO4 stained cryofractured cross-section, backscattered electron image, 5kx , arrow shows stained phase (C) OSO4 stained cryofractured crosssection secondary electron image, 30 kx, arrow shows MWCNTs (D) cryofractured cross-section , secondary electron image, 30 kx (E) cryofractured cross-section secondary electron image, 150 kx, arrow shows MWCNTs at cross-section

1.0

EPDM=0.000)

1.5

( A)

(EPDM= 0.216 )

( B)

0

10

1.0 Qi

(EPDM= 0.423)

   PP/ Blend

100 90 80 70 60 50 40 30 20 10

-1

10

(EPDM=1.000) (EPDM= 0.623 ) (EPDM= 0.815) 2

80% 3

10

10

4

0.0 0.2 0.4 0.6 0.8 1.0

10

0.0

EPDM

EPDM Continuity

Continuity index Composition

-2

5

10 |G*| (Pa)

0.5

10

0.5

EPDM phase

( D)

(EPDM= 0.815)

0.4

Co-continunity

PP phase

0.3 0.5

EPDM Dispersed phase

tan 

Extrcatable EPDM (Fraction)

o

( )

31

0.1 0.0

( C)

(EPDM= 0.623 )

0.2 (EPDM= 0.216 )

EPDM=0.000)

0.0

0.0 0.2 0.4 0.6 0.8 1.0 EPDM

-80

-60 -40 -20 0 Temperature [°C]

20

Figure 3: Co-continuity characterization; effect of blend composition on (A) phase angle as a function of shear modulus for different blends (B) relative change in complex viscosity and mass uptake (Qi) as a function of EPDM volume fraction (C) phase continuity (D) dynamic mechanical thermal analysis of different blends.

32

(A)

(B)

(C)

Figure 4: Co-continuity characterization; cryofractured scanning electron micrographs of blends after selective leaching of EPDM phase (A)EP20 (B)EP40 [d-PEP] (C) EP60[co-PEP].

33

(A)

(B)

(D)

(C)

(E)

Figure 5: Co-continuity characterization; atomic force microscopy of blends after selective leaching of EPDM phase (A) EP20 (B) EP40 [d-PEP] (C) EP60[co-PEP]. Phase images of unleached blends (D) EP20 (E) EP60 [d-PEP].

34

5

2.0x10

  rad/s

( A)

CNT=0.096)

1.6x10

(CNT= 0.045 )

5

G' (Pa)

EPDM= 0.643

1.2x10

(CNT= 0.033)

4

8.0x10

(CNT= 0.022 )

(CNT=0.070)

(CNT= 0.011)

4

4.0x10

-1

0

10

1

2

10 10 10 Shear strain amplitude (%)

10

3

10

6

10

5

10

4

10

3

10

2

10

1

10

0

10

( B)

'

2

r =0.99   rad/s

-1

0.01

CNT

0.1

( D)

0.5

m

8 7

( C)



6 c  

0.4 0.3

(EPDM=0.643)   rad/s

2

5



'

G comp-G poly 

0.9 0.8 0.7 0.6

  rad/s (EPDM=0.643)

9 c ( %)

G'comp-G'poly (Pa)

5

10

r =0.82

0.01

(=0)

CNT

0.1

0.2 0.00

0.02

0.04

0.06

CNT

0.08

0.10

Figure 6: Non linear rheology; MWNT structure breakdowns in co-PEP containing different volume fractions of MWNTs (A) dependence of the storage modulus on the small amplitude oscillatory shear strain (molten) (B) variation in G’(0) with increase in MWNT loading (C) Variation in γc with increase in MWNT loading (D) variation in m with increase in MWNT loading in [=1 rad/s, T= 200 oC]. Dotted lines represented the best fit of the experimental data to phenomenological quantitative Kraus model (equation 8).

35 4

80

( A)

EPDM=0.100 (EPDM)

70 60

4

2x10

 Pa.s

Storage Modulus [Pa]

3x10

EPDM=0.623 Co-PEP60 4

1x10

( B)

EPDM= 0.623

50

EPDM= 0.423

40

EPDM= 0.216

30

-1

0

10

10

EPDM= 0.815

10 1

2

3

1

0 -1 10

4

10 10 10 Shear Strain [%]

10

2

EPDM=0.000

20 EPDM=0.000 ( PP)

EPDM=1.000

0

10

1

 ( s)

1

( C)

CNT= 0.00

CNT= 0.00

10

0.6

CNT= 0.07

  rad/s

1

0.0

0.2

  rad/s

0.8

m

c ( %)

100

10

0.4

0.6

EPDM CNT

0.8

1.0

0.4

CNT= 0.07

( D) 0.0

0.2

0.4

0.6

EPDM

0.8

1.0

Figure 7: Effect of blend morphology on MWNT network breakdown under deformation (A) dependence of the storage modulus on the small amplitude oscillatory shear strain (molten) of PP, co-PEP60 and EPDM containg fixed MWNT (B) relaxation spectrum of different PP/EPDM blends (C) Variation in γc with increase in EPDM loading in unfilled and MWNT filled blends (D) variation in m with increase in EPDM loading in PP/EPDM blends [=1 rad/s, T= 200 oC]. Dotted lines represented the best fit of the experimental data to phenomenological quantitative Kraus model (equation 8).

36 140

3

130

(A)

p-PEP/CNT

120 d-PEP/CNT

30 1

20

0.05

0.10

0.15

Tunnel

III

0.20 II

10 0

I

0.0

0.1

p-PEP/CNT

0.2

0.3

0.4

(B)

R/R0

R/R0

2

110

10 9 8 7 6 5 4 3 2 1 0

0.5

100 kGy

0

Strain

3.0 2.7

50 kGy

0 kGy

CNT= 0.07

5 10 15 20 25 30 35 40 45 50 Strain (%)

( C) 5mm/min

R/R0

2.4 2.1 2.5 mm/min 10 mm/min

1.8 1.5

5

6

CNT= 0.07 Dose=100 kGy

8 7 Strain (%)

9

10

Figure 8: Tensorestive behavior of co-PEP/MWNT composites (MWNT =0.07) (A) Relative change in resistance for d-PEP/MWNT and co-PEP/MWNT composites, dotted lines represented the best fit of the experimental data to the tunneling model (B)Effect of radiation dose on the electromechanical response of co-PEP/MWNT composites (C) effect of strain rate for the co-PEP/MWNT composites subjected to 100 kGy gamma radiation.

1000 800

100

(A)

III

2N

I

M/M

2N

600 400

M/M

M/M

2N

II

(B)

3.2

0.4 0.8 1.6

2.4

37

4.0

4.8

5.6

50 6.4

200 0 0.0

Time (S)

2.0x10

6

4.0x10

6

6.0x10

6

Compressive Stress (Pa)

(C)

8.0x10

6

0

0

200

400

600

800 1000 1200

Time(s)

Load

(D)

Unload

Figure 9: Piezorestivity of co-PEP/MWNT and d-PEP/MWNT CPCs (MWNT =0.07) (A) Relative change in conductance for co-PEP/MWNT for different compressive stresses [Inset shows representative loading-unloading cycles] (B) Relative change in conductance for d-PEP/MWNT for different loading-unloading cycles of compressive stresses (C) ) A representative image of co-PEP/MWNT CPCs used in experiments (I) Demonstration of co-PEP/MWNT CPCs as a high compressive stress switch (a red LED was connected to a 10V power supply).

38

5

1.4x10

5

1.2x10

0 kPa

RCNT

5

CPE

Rc

>>

CGap

1.0x10

200 kPa

''

-Z ()

4

8.0x10

RCNT

Rc

RCNT

4

6.0x10

4

4.0x10

Resistor-Capacitor Network

Fitted Curves

4

2.0x10

0.0 0.00

2000 kPa 4

4.10x10

MWCNT

(A) 4

8.20x10

5

1.23x10

5

1.64x10

(B)

'

Z ()

Figure 10: Effect of compressive stress on the real (Z’) and imaginary part (Z’’) of AC impedance and fitting of equivalent R-C circuit (A) variation of Z’’ with Z’’ under different compressive stress (dotted lines shows fitting of corresponding equivalent R-C circuit), inset shows R-C circuit with CPE (B) schematic showing typical resistorcapacitor arrangement in CPCs.

39

(B)

(A) Continuous PP domains providing structural support

Crosslinked continuous EPDM domains providing spring like elasticity

PP

Voids after Leaching EPDM domains

Grafted/Crosslinked EPDM chains on CNT surface

CNT preferentially distributed in continous EPDM phase

(C) Continous EPDM phase

PP domains

EPDM-PP-CNT composite

Crosslinked continous EPDM phase

Degraded PP chains

Crosslinked EPDM-PP-CNT composite with degraded PP domains

Figure 11: Hard-soft co-continuous morphology of co-PEP (A) schematic showing co-continuous structure with radiation crosslinked EPDM (B) SEM of co-PEP after selectively leaching of EPDM phase (C) scheme showing preferential distribution of MWCNT and radiation crosslinking.

40

800

2 MPa, 4 Hrs

(A)

(B)

M/M2N

600 400 10 kPa, 2 hrs

200 0 0.0

3

4

4

4

5.0x10 1.0x10 1.5x10 2.0x10 Time(s)

1000

6

6

6x10

(C)

6

4x10

6

3x10

(D)

500

6

5x10

10 Cycles

M/M2N

Compressive Stress (Pa)

7x10

Output

0

Input

-500

6

2x10

25 kPa.s

6

1x10

0 0.0

-1000 0.1 0.2 0.3 0.4 Compressive Strain

0.5

0

250 kPa.s

-1

500

1000

1500

-1

2000

Time (S)

Figure 12: Creep and fatigue characteristics of co-PEP/MWNT CPCs (A) Relative change in conductance under constant compressive stress (B) Residual strain and strain recovery at constant compressive stress (C) Stress-strain response for cyclic loading and unloading for different stresses, at each load 10 cycles were used (D) Effect of stress rate on the relative change in conductance for co-PEP/MWNT (10 loading-unloading cycles at each stress rate, dotted lines provides a visual guide to baseline and peak ).

41 Supplementary Figures

42

EPDM=0.000 Load cell

Complex Viscosity,  (Pa.s)

PC

EPDM= 0.216

CNT= 0.00

5

EPDM= 0.423 EPDM= 0.623



Moving Crosshead

10

Copper Electrodes

Specimen Insulation Holding Grips

888888 ohm

(A)

Stationary Base Multimeter

EPDM= 0.815

10

4

10

3

10

2

EPDM=1.000

(B)  

EPDM=0.000 EPDM= 0.216 EPDM= 0.423 EPDM= 0.623

5



10

EPDM= 0.815 EPDM=1.000

10 10

4

3

(C)

CNT= 0.07 2

 

10 -2 10

10

-1

10

0

10  (rad/S)

1

10

2

10

-1

10

3

0

1

10  (rad/S)

10

10

10

3

 

(CNT=0.070)

5

10

2

EPDM= 0.623

CNT=0.096

6

Complex Viscosity,  (Pa.s)

10

6



Complex Viscosity,  (Pa.s)

10

(CNT= 0.045 ) 4

10

3

10

2

10

(CNT= 0.033) (CNT= 0.022 )

(D)

(CNT= 0.011)

1

10

10

-1

10

0

1

10  (rad/S)

10

2

10

3

Figure S1: (a) Setup for electromechnical measurements under compression, for tensorestive measurements a tensile grip was used (b) Complex viscosity of unfilled blends (c) Complex viscosity of co-PEP/MWNT composites (d) Complex viscosity of PEP blends with fixed MWNT.

Storage Modulus Loss Modulus

CNT= 0.00 5

10

4

10

4

G', G'' (Pa)

10

3

10

3

G'' ~ 

10



2

2

1

G' ~ 

1 -1

1

10

10

EPDM=0.000



0

10

10

2

3

10

10

G', G'' (Pa) G', G'' (Pa)

G'' ~ 

0

10

1

10

10

2

10

3

10

REVERSAL G' , G"

G' ~ 



4

10

G' ~ 



-1

EPDM= 0.423 0

10

G' ~ 

1

10

10

-1

10

2

10

G'' ~ 

3

3

10

10

-1

0

10

EPDM= 0.623



1

10

10

2

10

3

10

5

10



G'' ~ 

3

10

0

10



10

1

-1

EPDM= 0.216

5

2

10

10



10

4

3

G' ~ 

0

10

10



10

10

10

G'' ~ 

0

10



EPDM= 0.815 1

10

 [rad/s]

2

10

G' ~ 

G'' ~ 

3

3

10

10



-1

10

0

10



EPDM= 1.000 1

10

2

10

3

10

 [rad/s]

Figure S2: Storage (G’) and loss modulus (G’’) crossover for different unfilled blends. T=200 oC, 

43

44

CNT= 0.07 6

6

G', G'' (Pa)

G', G'' (Pa)

10

10

5

10 5

10

4

10 EPDM=0.000

4

10 -2 10 6 10

-1

10

0

10

1

10

3

10

10

10 -2 10 6 10

-1

10

0

10

1

2

10

10

3

10

5

10 5

10

4

10 EPDM= 0.423

4

10

0.1

10

1

100

1000

10

6

10

5

10

6

-1

10

0

10

1

10

2

10

3

10

5

10

EPDM= 0.815

4

10

EPDM= 0.623

3

10

G', G'' (Pa)

(EPDM= 0.216 )

3

2

-1

10

0

10

1

10

 [rad/s]

2

10

EPDM= 1.000

4

3

10

10

-1

10

0

10

1

10

2

10

3

10

 [rad/s

Figure S3: Storage (G’) and loss modulus (G’’) crossover for different blends at fixed MWNT loading. T=200 oC, 

45

EPDM= 0.623

Storage Modulus [Pa] Loss Modulus [Pa]

4

G', G'' (Pa)

10

G' ~ 

4

10



3

G'' ~  -1



G'' ~ 

10

0

10

G' ~ 



1

10

10

CNT=0.011 2

10

3

10

CNT=0.022

3

10



-1

0

10

1

10

10

2

10

3

10

5

G', G'' (Pa)

10

5

G' ~ 

10



G' ~ 



4

10

G'' ~  -1

0

10

CNT=0.033



1

10

10

2

10

4

10

3

10

G'' ~  -1

6

0

10



1

10

10

CNT=0.045 2

10

3

10

G', G'' (Pa)

10 5

10

G' ~ 

G' ~ 

G'' ~ 

3

10



-1

10

0

10



CNT=0.07 1

10

2

10

 ( rad/s)

3

10

G'' ~ 

4

10



-1

10

0

10



CNT=0.096 1

10

2

10

3

10

 ( rad/s)

Figure S4: Storage (G’) and loss modulus (G’’) crossover for different co-PEP60/MWNT composites at different MWNT loadings. T=200 oC, 

46

6

EPDM= 0.623

CNT=0.070

4

0 kGy 1

3

CNT=0.070 100 kGy

EPDM= 0.623

Stress (MPa)

Stress (MPa)

5

50 kGy

100 kGy

2 1

50 kGy 0

0 kGy 0.24

0.48

0.72

0.96

Strain (%)

0

5 10 15 20 25 30 35 40 45 50 55 60

Strain (%) Figure S5: Stress strain profiles of unirradiated and irradiated co-PEP/MWNT composites [Inset: magnified profile at low strains].

47

10

10

CNT=0.02

Resistance (Ohm)

9

10

8

10

7

10

6

10

5

10

CNT=0.07

4

10

0

200

400

600

800

1000

1200

Compressive Stress (Pa) Figure S6: Effect of MWCNT volume fraction on the compressive stress dependent change in the resistance of coDeclaration of interests PEP/MWCNT composites. ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.

48

☐The authors declare the following financial interests/personal relationships which may be considered as potential competing interests:

Highlights 

CNT-percolated hard-soft polyolefin composites were melt-compounded.



High energy radiation was used to immobilize phase morphology.



Superior piezoresistivity was observed in co-continuous but not in dispersed morphology.



Low fatigue and the residual strain was observed piezoresistive sensors.



Non-linear and linear rheology suggested blend-morphology-dependent breakdown.