Construction of an Interconnected Nanostructured Carbon Black

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Construction of Interconnected Nanostructured Carbon Black Network: Development of Highly Stretchable and Robust Elastomeric Conductors Eshwaran Subramani Bhagavatheswaran, Meenali Parsekar, Amit Das, Hai Hong Le, Sven Wiessner, Klaus Werner Stoeckelhuber, Gerd Schmaucks, and Gert Heinrich J. Phys. Chem. C, Just Accepted Manuscript • DOI: 10.1021/acs.jpcc.5b06629 • Publication Date (Web): 28 Aug 2015 Downloaded from http://pubs.acs.org on September 2, 2015

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The Journal of Physical Chemistry C is published by the American Chemical Society. 1155 Sixteenth Street N.W., Washington, DC 20036 Published by American Chemical Society. Copyright © American Chemical Society. However, no copyright claim is made to original U.S. Government works, or works produced by employees of any Commonwealth realm Crown government in the course of their duties.

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The Journal of Physical Chemistry

Construction of Interconnected Nanostructured Carbon Black Network: Development of Highly Stretchable and Robust Elastomeric Conductors Eshwaran Subramani Bhagavatheswaran,a,b Meenali Parsekar,a Amit Das,a,d* Hai Hong Le,a,c Sven Wiessner,a,b Klaus Werner Stöckelhuber,a Gerd Schmaucks,e Gert Heinricha,b

a

Leibniz - Institut für Polymerforschung Dresden e.V., Hohe Straße 6, D-01069 Dresden,

Germany b

Technische Universität Dresden, Institute für Werkstoffwissenschaft,D- 01062 Dresden,

Germany c

Institut für Polymerwerkstoffe e.V., D-06217 Merseburg, Germany

d

Technical University of Tampere, Korkeakoulunkatu 16, 33101 Tampere, Finland

e

Elkem AS, Silicon Materials , 4675 Kristiansand, Norway

* Corresponding author: Amit Das, E-mail address: [email protected], Telephone +49351 4658 579

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ABSTRACT: In the present work, a strong filler-filler network of conductive carbon black was strategically established in an elastomer matrix, which leads to unique combination of electrical and mechanical properties. The novelty of our composites was the development of a strong percolated morphology of nano-structured conducting carbon black particles by the incorporation of relatively large non-reinforcing spherical silica particles, inside the soft elastomer matrix. This technique allowed us to fabricate solution-styrene butadiene rubber (SSBR) composites with outstanding electrical conductivity of 40 S/m, tensile strength ~ 10 MPa and extensibility up to 200%. Furthermore, the electrical conductivity was strainindependent up to 50% elongation strain. The electrical conductivity was found to be unaltered after 2000 loading - unloading cycles. This is the first ever report of a robust elastomeric system with such high electrical conductivity where all the basic ingredients used was selected from well-known commercially available raw materials of rubber industry. This work directly manifests an industrially viable method for preparing high-performance elastic conductors that can be utilised in robust and flexible applications.


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INTRODUCTION Development of conducting elastomeric materials with desirable mechanical as well as electric properties are always a challenge to the scientists, particularly in the recent times, where world moves faster towards electrics and electronics driven life style. Lots of efforts have been paid to develop flexible electronics or electrical devices, including electronic displays and batteries, which need highly conducting and stretchable conductors. As an outcome, a plenty of reports could be found in the literature1-4. Most of the published articles described several flexible materials with high electrical conductivities. None of them are suitable for very robust and highly stretchable applications. For example, De et al.1 developed a

flexible

thin

film

based

on

polymer

(poly-3,4-ethylenedioxythiophene:poly-

styrenesulfonate) nanotube composites with very high electrical conductivity and the material is fairly transparent. In this study the flexibility of the material was verified by the evaluation of electrical changes over several bending cycles. However, such ionomeric mixture might not be that much robust and flexible like crosslinked rubber. A review article2 is also published where the applications of such novel flexible conductors are demonstrated in different types of electronic devices. Several strategies are followed to develop such materials, which take into account, for example, design concepts (insertion of fine metal wires inside the soft polymer), synthesis of conducting polymers and incorporation of carbon based conducting solid particles inside polymers. But the materials are found to have a number of deficiencies, for example, failure under large strain, very limited lifetime and nonreproducible properties after certain service period. Reports also can be found to develop a number of different soft piezoresistive polymeric composite materials for strain gauging application. For these materials, the changes in electrical resistance with applied strain show both linear and nonlinear dependencies between strain and resistance5-6. In most of the reported elastomeric conductors, the elastic properties have been compromised to improve its



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electrical conductivity7. Hence, there are hardly any reports on common commercially produced elastomers with high electrical conducting properties. For example, Lin et. al.8 reported preparation of stretchable conductors based on polyurethane and carbon nanotubes. Unfortunately, the elastic properties of polyurethanes is rather poor, if one compares it with those of cross-linked diene rubbers like styrene butadiene rubber (SBR), natural rubber (NR). Sekitani et. al.9 reported on highly conducting rubber like material based on single walled carbon nanotube filled fluoroelastomer (vinylidene fluoride- hexafluoropropylene copolymer) composite. Despite the reported materials have high conducing nature and rubber like character, their dynamic performance under repeated cyclic stress cannot be compared with those of cross-linked rubbers. As far as the common cross-linked diene rubbers for the development of flexible conductors are concerned, the use of carbon nanotubes and graphene are extensively explored10-12. The major problem of the production of such elastomeric composites is the high cost and also difficulties in processing and handling in large scales13-14. Finally, most of the above conductive elastomeric materials have a strong drawback with their strain dependent electrical conductivities i.e. the materials exhibit piezoresistive characteristic, where the electrical conductivity is sharply affected by strain15. On the other hand, the fabrication of conducting rubbers using commercially available conducting carbon black (highly anisotropic black) is a matured and well-accepted technology. However, owing to the aggregated and agglomerated structure of carbon blacks, one can only achieve limited electrical properties16. Recently a special type of silica (silica fume, commercially available as ‘Sidistar’) is found to be very effective as dispersing aids along with other multifunctional features for rubber compounds17-18. Schmaucks et al.17 reported that microsilica could be used as a plasticizer without any adverse effects such as blooming or surface leaching. Moreover, the use of these large silica spheres helps in the processing of high viscosity polymers and speciality polymers, various combinations of polymers and fillers, which were not possible with conventional processing techniques. ACS Paragon Plus 4 Environment

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In general, the conducting properties of the composites are governed by the intrinsic conducting nature of the filler particles, as well as, a well percolated network structure of the filler particles. So it can be assumed that a very well established strong filler-filler network may improve the electrical behaviour of the composites. A loose and weak percolated network of conducing particles inside flexible elastomers can be permanently destroyed during dynamic operations, resulting in an irreversible change in the electrical conductivity. In this context a permanent and stable conducting network would be more desirable, particularly when the conducing rubbers are used in dynamic application. In this situation, stable conducting properties along with high mechanical performance are very important. Keeping the facts in mind, efforts have been given to establish a strong filler-filler network, which is very stable up to a certain strain limit and can contribute a high conducting nature of the composites. For this reason, we have exploited the micro-sized silica fume to find its effect on the filler-filler networking behaviour of nano-structures carbon black particles in a SSBR (solution styrene butadiene rubber) matrix. This work aims to develop a simple and efficient method for preparing highly conducting cross-linked elastomers with durable elastic properties.

EXPERIMENTAL Materials. The rubber matrix is a solution styrene butadiene rubber (S-SBR) with commercial name Buna VSL5025-0HM, supplied by Lanxess with 25 wt% of styrene. Silica fume (Sidistar R320) procured from Elkem is an amorphous, non-reinforcing silicon-di-oxide with an average primary particle size of 150 nm, DBP absorption value of 81 ml/100 g and BET surface area of 20 m2/gm19. This silica is amorphous in nature and has larger particle size as compared to precipitated silica. It would be important to mention here that this is not ‘fumed silica’ where the primary particle size is in the nanometer range. This silica fume or microsilica particles are rather large and it does not form any aggregation/agglomeration. This ACS Paragon Plus 5 Environment


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silica fume is a specially refined by-product obtained during the production of high-purity silicon metal. The nano-structured carbon black (Printex XE2) is obtained from Orion Engineered Black with the DBP value of 350-410 ml/100g, iodine adsorption number of 9751175 mg/g, BET surface area 1002 m2/g and mean particle size of 35 nm. Zinc oxide, stearic acid, sulfur, cyclo-hexyl-bezothiazole-sulphenamide (CBS) and diphenyl guanidine (DPG) are purchased from Acros organics Belgium. Preparation of the Composites. Table 1 summarizes the different rubber compounds with their reference nomenclature SSBR_S_CB, indicating the silica fume content (S) and carbon black content (CB), in phr. Other ingredients (in phr) like ZnO (3), stearic acid (2), sulphur (1.4), CBS (1.7) and DPG (1.4) were kept constant throughout the study. All rubber compounds were prepared in a Haake internal mixer PolyLab (Thermo Electron Corporation, Haake, Karlsruhe, Germany) with a mixing chamber volume of 69 cm3. S-SBR along with silica, carbon black, ZnO and stearic acid were mixed in the internal mixer at 70 °C and 70 rpm. The fill factor of the mixing was set to 0.7 with a constant mixing time of 9 minutes. The curative package was further added to the compound in a laboratory-size two roll mill (Polymix 110 L, size 203 x 102 mm, Servitech GmbH, Wustermark, Germany) at 40 °C. The friction ratio of the mill was 1:1.25 with a total 11 min mixing cycle. The viscoelastic behaviour of all samples was studied in a moving die rheometer, Elastograph Göttfert, Germany, to evaluate the optimum cure time (t90). The samples were pressed and cured into sheets with dimensions 12 x 12 x 2 mm3 in a compression moulding machine to their respective curing times at 160 °C.


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Table 1: Formulation of the Rubber Compounds with Their Respective Filler Volume Fractions S-SBR

Microsilica

Carbon Black

φmicrosilica

φCB

φmicrosilica+CB

SSBR_0S

100

0

0

-

-

-

SSBR_50S

100

50

0

0.1678

-

0.1678

SSBR_5CB

100

0

5

-

0.0217

0.0217

SSBR_10CB

100

0

10

-

0.0405

0.0405

SSBR_15CB

100

0

15

-

0.0596

0.0596

SSBR_20CB

100

0

20

-

0.0779

0.0779

SSBR_50S_5CB

100

50

5

0.1649

0.0173

0.1822

SSBR_50S_10CB

100

50

10

0.1621

0.0340

0.1960

SSBR_50S_15CB

100

50

15

0.1594

0.0501

0.2095

SSBR_50S_20CB

100

50

20

0.1568

0.0657

0.2224

Compounds

Characterization Techniques. A scanning electron microscope (LEO 435 VP, Carl Zeiss SMT) was used to study morphological features of the filler particles. Morphological investigations of the composites were performed under a transmission electron microscope (Zeiss Libra 200 TEM) with an acceleration voltage of 200 kV. Ultrathin sections of the rubber composite were cut by ultra-microtome (Leica Ultracut UCT) at -120 °C. Tensile tests were done using a universal testing machine (Zwick 1456, Z010, Ulm, Germany) following ISO 527 standards with DIN S2 dumbbell specimens. Dynamic mechanical analysis of the composites is performed using a dynamic mechanical thermal analyser Eplexor 2000 N, Gabo Qualimeter, Germany). The strain sweep measurements were performed at room temperature with a constant frequency of 10 Hz, 60 % static strain and dynamic amplitude sweep from 0.01% to 30%. Electrical volume resistivity was measured with stripes using a 4-point test fixture (distances between the inner electrodes and outer electrode are 10 mm and 16 mm, respectively) combined with a Keithley electrometer 6517A. The electrical conductivity of the samples during mechanical stretching was measured by a homemade stretching device and the conductivity during cyclic dynamic loading was taken using a dynamic mechanical analyzer ACS Paragon Plus 7 Environment


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(Eplexor 2000 N) coupled with electrical measuring unit (Fig. 1). Swelling experiments were done according to ASTM D 6814-02 standard. Rubber samples with known weight were allowed to swell in toluene at ambient conditions for 72 h. The test pieces were then takenout, weighed and dried to a constant weight.

a)

b)

Fig. 1: a) Picture of the stretching device by which the electrical conductivity and the mechanical properties (strain) were studied. b) Photograph of the dynamic mechanical analyser coupled with electrical measuring unit. This device was used for the evaluation of electrical properties during dynamic deformation of the rubber materials.

RESULTS AND DISCUSSIONS SEM images of silica fume as delivered material are shown in Fig. 2a and Fig 2b. Agglomerates of different sizes and shapes appear to be in the millimetres range. Some of the broken agglomerates showed hollow sphere like structure. However, at higher magnifications (Fig. 2b) the micrograph demonstrates the perfect spherical nature of primary particles with a smooth surface. These particles also display a large distribution of sizes ranging from 20 nm to 400 nm19. On the other hand, the carbon black particles are found to form also aggregated and agglomerated structures but the size of these structures are very small (Fig. 2c). It can be found from Fig. 2d that the primary particles of CB tightly form a chain like structure with a dimension of primary particles only a tens of nanometer. The average primary particle size ACS Paragon Plus 8 Environment

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seems to be around 20-40 nm. A highly structured morphology of CB particles could also be visualised from the images.

(a )

(c )

b)

d)

Fig. 2: Scanning electron microscopic images of silica fume a) scale bar = 20 µm, b) scale bar = 200 nm and nano-structured carbon black c) scale bar = 200 nm, d) scale bar 100 nm

Fig. 3a and Fig. 3b depict the TEM micrographs of CB and silica, respectively. It can be very clearly observed that the primary particles of CB form an aggregated structure, developing a highly anisotropic characteristic of the particles. Additionally, presence of pores and formation of highly structured morphology of the primary particles are also noticed. Similar like the SEM images in the TEM micrographs, the silica particles are found to be ACS Paragon Plus 9 Environment


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existed as perfect spherical structures without any agglomeration or cluster formation. This is a just contrary to the precipitated silica where the silica particle always forms big aggregated and agglomerated structure. Fig. 3c and 3d show the TEM micrographs of S-SBR composite with 50 phr silica and 20 phr CB (SSBR_50S_20CB). The CB particles are found to form a network structure by the interaction of different fine CB particles, whereas the silica particles are remaining as single particles without forming any silica-silica network (Fig 3c). At higher magnification several fine layers are visible in the carbon black. Probably, the CB particles are dissociated into finer structure forming layered morphology and indicating laminar graphitic nature of the CB particles. The conducting nature of the particles may pertain to this ordered graphitic structure. These TEM images also demonstrate the rigidity of silica particles as they are neither damaged nor broken during the mixing of the compounds. Now it can be envisaged that the rigid silica exerted high shear force to the CB during the mixing process, which broke the secondary structure/aggregates of the CB resulting in more fine CB particles. Due to presence of large silica sphere and space confinement, the CB particles are forced to interact with each other and form a conductive network structure. The tensile properties of the composites are summarised in Table 2 and some of the results are representatively shown in the stress-strain plots (Fig. 4a). It could be noted from Fig. 4a that the stress-strain plot of 50 phr silica fume composite is very similar to that of gum rubber (without any filler) indicating non-reinforcing nature of silica fume. As expected, for CB filled composites, the mechanical properties like 100% modulus and tensile strength are increasing with the increase of CB loading. The 300% modulus and the tensile strength of the composites containing 20 phr CB are found to be improved 10 and 6 times with respect to the gum vulcanizate. A decrease in the elongation at break values at high filler concentration of CB is reflected in the stress-strain experiment indicating the reinforcing nature of nanostructured CB. The hybrid composite with 50 phr of microsilica and 20 phr CB shows considerable improvement in 100% modulus values as compared with only 20 phr CB filled ACS Paragon Plus 10 Environment

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(a )

(b )

(c )

(d )

Fig. 3. Transmission electron microscopic images of a) pure carbon black (scale bar = 100 nm), b) pure silica fume (scale bar = 100 nm), c) silica fume and carbon black (scale bar = 200 nm) and d) silica fume and carbon black (scale bar = 100 nm ) in the solution-styrene butadiene rubber (S-SBR).

system. The microsilica when combined with CB offers synergistic effect in terms of 50% and 100% moduli. The non-linear behavior of filled rubber’s dynamic modulus (E’) at relatively low strain is often referred to as the Payne-effect20-22. This non-linear behavior of filled



vulcanizates can be explained by the breakdown of carbon black network structure during dynamic stretching of the samples. Fig.4b shows the dependence of dynamic modulus on the dynamic strain amplitude. The gum compound and the compound filled with 50 phr of microsilica hardly show any strain dependency, whereas composites containing only CB (indicated by dash lines) show higher values of moduli change, particularly, at low strain amplitude. Furthermore, at higher loading of CB, very prominent strain dependency is noticeable. The composite with 50 phr of silica fume has a very less strain dependency indicating the lack of filler-filler networks inside the composite. This fact corroborates with findings in the TEM analysis. The magnitude of the ‘Payne Effect’ is found to be remarkably higher for the composites containing 50 phr silica fume and 20 phr CB. For hybrid filler system as the microsilica does not contribute to any strain dependency, the CB particles are hence thought to form a stronger percolated filler-filler network as compared to the composite containing only same amount of CB.

a) SSBR Gum SSBR_50S SSBR_20CB SSBR_50S_20CB SSBR_80S_20CB

0.8 0.4 20 0.0 0

2

4

6

8 10

10

0

Storage Modulus (E') (MPa)

b)

30

Stress (MPa)

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60

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5CB 10CB 15CB 20CB Gum

30

50 S_5CB 50S_10CB 50S_15CB 50S_20CB 50S

20

10

0

200

400 Strain (%)

600

0

10-1

100

101 102 Dynamic Strain (%)

103

Fig. 4: a) Stress-strain plot for various composites (inset: stress-strain plot magnified in the low strain regime), b) strain sweep data of the composites fitted and extrapolated by the Kraus equation.


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Table 2: Mechanical Properties of S-SBR (gum) and S-SBR Filled with Microsilica and CB 50% Modulus (MPa) 0.52

100% Modulus (MPa) 0.73

300% Modulus (MPa) 1.50

Tensile strength (MPa) 2.61

Elongation at break (%) 412

SSBR_50S

0.69

0.97

1.95

6.12

547

SSBR_5CB

0.73

1.07

3.23

6.68

456

SSBR_10CB

1.00

1.49

5.19

11.15

493

SSBR_15CB

1.51

2.60

8.87

12.51

355

SSBR_20CB

2.19

3.97

14.12

15.48

317

SSBR_50S_5CB

0.97

1.42

3.91

7.60

466

SSBR_50S_10CB

1.34

2.07

6.36

9.56

418

SSBR_50S_15CB

1.78

2.88

9.46

10.21

339

SSBR_50S_20CB

2.81

4.81

--

12.96

272

SSBR_60S_20CB

3.13

5.83

--

9.66

196

SSBR_70S_20CB

4.62

9.15

--

9.95

110

SSBR_80S_20CB

5.01

9.65

--

10.12

101

Sample* SSBR

SSBR_90S_20CB 6.55 --8.75 *The mechanical data is the average value obtained from five different specimens.

70

For more insight on the filler networking process, a phenomenological model (Eq.1) is considered, which is based on the concept of agglomeration-deagglomeration mechanism 23. According to this mechanism the filler network is developed by the agglomeration and these agglomerates are broken down at larger dynamic strain.

E ' (γ ) − E '∞ = E '0 − E '∞

1 γ  1 +   γc 

2m

(1)

where, E ' (γ ) is the modulus of the composite at a particular strain γ , E ' (γ ) = E ' ∞ at very large strain; E ' (γ ) = E '0 at very low strain γ 0 ; E '0 − E '∞ is the quantitative measure of the Payne effect determined from the amplitude sweep plot; γ c is the critical strain, at which the magnitude of E '0 − E '∞ becomes half. m is a constant which is related to specific fractal dimensions of the fractal agglomerate structures, which governs the shape of the curve. It is ACS Paragon Plus 13 Environment


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reported in general that the constant γ c is largely dependent on the types of fillers, polymers and the state of the dispersion of the filler in the rubber matrix. The γ c and m values are evaluated and depicted in Fig. 5a. It is found that the value of γ c is decreasing with the increase of CB content for both hybrid and single filler systems. This indicates that with higher amount of CB loading, the filler-filler networks are destroyed at much lower strains. However, the presence of microsilica decreases the critical strain further, indicating stronger CB-CB filler networks in the hybrid filler systems. It was reported that for CB systems, the value of m is nearly constant ~ 0.6 which is independent of the nature and type of CB24. In present case as well the composites with higher amount of CB, particularly, when they show a strong strain dependent modulus, the value of m is found to be ~ 0.6. The composites with lower strain dependencies (lower ‘Payne effect’) show larger values of the m. The hydrodynamic reinforcement in the rubber matrix is evaluated using the modified Guth equation

25-27

taking into account the shape factor ‘f’. Consideration of shape factor is

necessary as CB with high structure usually has a shape factor > 1 and the corresponding expression is shown in Eq 2.

= 1 + 0.67φ + 1.62 φ

(2)

where, M is the low amplitude dynamic storage modulus of the filled composite and Mo is the low amplitude dynamic storage modulus of the gum vulcanizate and φ is the volume fraction of the filler. These values are extracted from Fig. 4b in the low strain regime ( γ 0 ). The GuthGold plot is shown in Fig. 5b. The evaluated shape factor of CB is f=22 and such a high shape factor is very reasonable due to the high structure of the particular grade of CB. To understand the enhancement in modulus in hybrid filler systems, a simple normalisation is performed by subtracting the modulus of only silica filled composite from the hybrid filler systems.Eq.3.

′ = ′ + − ′

(3)


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where, M’(CB) is the estimated low amplitude dynamic storage modulus of the hybrid composites. M’(S) is the low amplitude dynamic storage modulus of the composites containing only sidistar and M’(CB+S) is the low amplitude dynamic storage modulus of the hybrid composites. The value of M’(S) and M’(CB+S) are directly taken from strain sweep experiment. The normalised modulus values as estimated from the above equation is plotted against the volume fraction of CB in the hybrid composites (φCB = VCB/Vhybrid-composite). The plot for the hybrid filler also yields a very good fit similar to only CB filled composites. The shape factor of CB (after the normalization) in hybrid filler system is found to be f ~ 29 whereas the composite with only CB the value of f is ~ 22. The increase in the shape factor can be explained if the CB particles exhibit more anisotropic characteristics. It can be mentioned here, since silica fume is perfectly spherical in shape its aspect ratio could be considered as unity and may not contribute to the f values. So, it can be said that in the presence of microsilica, the aggregated CB particles are raptured into finer particles with higher aspect ratio. These fine CB particles are distributed over the large spherical silica particles and connected to form a conductive network.

a)

b) 15

30 0.6

20 0.3

10 0

f = 28.6

10 f = 22

5

0

0.00

BR

_5

0S

0S

_2

0C

B

B

B

5C _1

Normalized Carbon Black values Carbon Black only

0.04 0.08 Volume fraction of CB (φCB)

SS

SS

BR

_5

_5

0S

_1

CB

0C

B SS

BR

_5

0S

_2 BR SS

BR

BR

_5

0C

B

B

5C

0C

_1 SS

BR

_1 SS

_5 SS

BR SS

BR

_5

0S

CB

0.0

Exponent (m)

Critical Strain (γ c )

0.9

M / Mo

40

SS

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16 17 18 19 20 21 22 23 24 25 26 27 28 29 30 31 32 33 34 35 36 37 38 39 40 41 42 43 44 45 46 47 48 49 50 51 52 53 54 55 56 57 58 59 60


Fig. 5: a) Critical strain and exponent values for the rubber composites obtained from Kraus equation, b) Guth - Gold plot for the rubber composites.



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Kraus equation

28-29

for swelling is a method to explore the rubber-filler and filler-filler

interaction as shown in Eq 4.

= 1−

∅

∅

(4)

where, k is a filler polymer interaction parameter, φ is the volume fraction of the fillers, Vg is the volume fraction of gum rubber in the equilibrium swollen state and Vf is the apparent volume fraction of filled rubber in swollen state. Vf solely refers to the volume fraction of the rubber that can undergo swelling in a filled composite (excluding all insoluble components such as fillers, etc.) and is expressed as

= !

!"α!# ⁄

%$" α!# ⁄$% ' () ⁄$#

(5)

dr and ds are the density of the rubber and solvent, respectively. wd is deswollen/dried weight of the sample, α the weight fraction of the insoluble components, ws initial weight of the specimen, Ao the weight of the absorbed solvent. Upon evaluating the value of k from Eq.4, the characteristic parameter C of the filler is calculated from Eq. 6.

= 3+1 − ,-.. / + , − 1

(6)

Fig. 6 depicts the plot of , ⁄ against 0⁄1 − 0 to estimate the filler characteristics by evaluating the slop of the fitted straight line. A reduction in the , ⁄ values at higher filler concentrations denotes the presence of strong rubber-filler interactions in the composite. The observed k value of 0.732 for CB when compared to 0.205 for silica indicates the higher reinforcing nature of the CB. According to Eq. 6, the characteristic constant C of the fillers is evaluated to be 2.308 for CB and 1.296 for silica. The much higher reinforcing ability of CB could be visualised by comparing its C values with that of silica. As far as the plots of hybrid fillers systems are concerned, only the volume fraction of CB in the hybrid composite is considered. The hybrid filler system shows higher reinforcing capabilities compared to only CB based composite. Hybrid filler registered much higher values of C = 4.104 and k =1.66 compared to single filler, i.e., CB and silica fume. The appreciable increase in C value for the ACS Paragon Plus 16 Environment

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hybrid system indicates the influence of silica fume, which has efficiently increased the number of CB particles by the dissociation of the large CB particles in to smaller fragments. 1.00

k = polymer - filler interaction parameter C = characteristic constant for the filler

0.96 Vg / Vf

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Sidistar k = 0.205 C = 1.296

Carbon Black k = 0.732 C = 2.308

0.92 Carbon Black in Hybrid k = 1.668 C = 4.104

0.88

0.0

0.1

0.2 φ / (1 − φ)

0.3

0.4

Fig. 6: The plot of , ⁄ against 0⁄1 − 0, based on the Kraus Equation for Swelling.

The electrical conductivity of these hybrid mixtures are depicted in Fig. 7a. The composite with only 20 phr of CB displays a very high electrical conductivity of 0.25 S/m (Fig. 7a). After incorporation of microsilica, the value further increases to 40 S/m. This is much higher electrical conductivity than most of the reported values for the CB filled diene rubber based composites6,12,30-33. With higher loading of silica, the CB particles gets closer, generating more electrically conductive pathways as demonstrated in Scheme 1. With the benefit of carbon black aggregates being broken down into smaller particles, the effective number of conducting pathways in the composite is increased as compared to only CB filled composites. The electrical conductivity reaches the maximum value at 60 phr of microsilica loading and drops with further addition. At this critical concentration there is an optimised balance between the concentration of conducting CB and silica fume, wherein a strong and well percolated electrically conducting filler network is formed. With further increase in the insulating silica particles, the overall proportion of conducting fillers in the composite would



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Scheme 1: Development of a strong filler-filler network of the conducting carbon black particles. Incorporation of large silica fume assisted the disintegration process of the agglomerated structure into more number of finer particles of the carbon black during solid state mixing.

be less and hence, a fall in electrical conductivities is observed. The real time change in resistivity with strain has been plotted in Fig. 7b. Maximum strain of up to 300 % is given to the sample and the corresponding resistance values are found to be strongly increasing, particularly at higher strain (>150%). However, in the low strain region 0-50%, the resistance of the samples is unaltered and after which a very slow increase of the values is observed (up to 100% strain). At much higher strain (>150%) a steep increase in the resistance is noticed. It can be assumed that during deformation of the sample, the resulting destruction in filler-filler networks are higher, causing an increase in the resistance values. The non-linear resistance dependency, particularly, at higher strain could be associated with the breaking of the conducting pathways. In this region the intrinsic resistance is different at different domain of strain due to altered arrangement of the conducting filler particles. ACS Paragon Plus 18 Environment

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Highly interesting and encouraging is the strain independent regime (up to 50 %), where the filler-filler conducting networks are unaltered. However, this critical strain should not be compared with the critical strain ( γ c ) obtained from strain sweep analysis. Dynamic strain

a)

b)

1

Sidistar + 20 phr CB 10

10

0

50

60 70 80 Amount of Sidistar (phr)

160

30

120 20 80 10

20 phr CB

-1

200

0

90

40 0

c)

400

800

150 200 Time (s)

250

0 300

12

600

8

400

4

200

0

100

200 300 Time (s)

400

500

Strain (% )

16

10 0 20

R esist.(KΩ )

800

20

Strain (%)

20

0

100

d)

1000

0

50

10 0 0

1200 1210 1220 1600 2000 Number of cycles

Fig. 7: (a) Plot of electrical resistance for hybrid composites with different microsilica content (b) resistance with strain (c) resistance against strain under prolonged cyclic loading conditions.

sweep measurements yield a critical strain value ~5 %, which is very less, and one must consider the complete experimental procedure and the definition of γ c value. By definition γ c is magnitude of the strain at which 50% of the filler-filler network is broken down. So rest 50% filler-filler still exists and this number are conducting the electrons through the matrix. Moreover, in strain sweep experiments, a pre-strain of 60% is applied on the sample, over ACS Paragon Plus 19 Environment

Strain (%)

10

40 Resistance (MΩ)

Electrical conductivity (S/m)

60

Resistance (KΩ)

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which the dynamic strain is additionally applied at high frequency of 10 Hz. So, the final strain experience of the material in a dynamic mechanical test is very different from a simple quasi-static resistance-strain measurement. Fig. 7c displays the resistance of the sample during different magnitude of dynamic strain and Fig. 7d demonstrates the resistance under a large number of cyclic loading and unloading conditions. The resistance remains almost unaltered when the rubber is stretched and released several times, indicating the electric conductivity not being affected by the deformation of the samples. So, this kind of materials can be easily applied where flexibility and a steady current flow are necessary. This kind of mechanically robust stretchable conductors can be utilised in battery panels of an electric car, flexible joints of robotic machines, pressure sensors (Fig. 8), solar cells, flexible electrode and long range strain sensors.

Rel. Resistance [-]

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10

3

10

2

10

1

10

0

10

-1

10

-2

20

30

40 50 Time (s)

60

70

Fig. 8: The developed conducting rubber (sulphur crosslinked S-SBR with 60 per Sidistar and 20 per CB) is fabricated into a pressure sensor to demonstrate a possible application of the material. A considerable change in the relative electrical resistance was observed when the conducting rubber film was manually deformed. The deformation must be more than 50% strain to see the changes in the resistance.


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CONCLUSIONS The relatively large silica microspheres/silica fume could be utilized to develop a percolated conducting carbon black network inside a diene elastomer matrix. The rigid silica particles assisted in the dissociation of relatively large particles of CB into smaller fragments by exerting high shear forces during melt mixing process. Guth-Smallwood model revealed the synergistic effect of the microsilica-carbon black composites by yielding a larger filler aspect ratio with respect to the single filler composite. At optimal concentration of silica, the carbon black particles were able to form a continuous electrically conducting pathways yielding very high electrical conductivity as high as 40 S/m and simultaneously the mechanical properties of these composites were similar to that of highly reinforced elastomer vulcanizate. The electrical conductivity in the case of hybrid filler was far superior to that of single filler composite due to formation of very strong filler-filler network which was also supported by dynamic mechanical analysis of the samples. Development of such elastomer composites can satisfy robust engineering applications such as artificial muscular joints; that demand extreme flexibility, constant dynamic motion, high electrical conductivity and even slight compromise in either of the properties is inacceptable.

ACKNOWLEDGEMENT: Authors are thankful to Mrs. Uta Reuter and Mrs. Regine Boldt of IPF, Dresden for TEM and SEM images.



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Construction of an Interconnected Nanostructured Carbon Black

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