**Polysiloxane Side-Chain Azobenzene-Containing Liquid Single Crystal Elastomers for Photo-Active Artificial Muscle-Like Actuators**

Jaume Garcia-Amorós and Dolores Velasco

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/50436

### **1. Introduction**

24 Will-be-set-by-IN-TECH

[36] Luckhurst, G.R. and Stephens, R.A. and Phippen, R.W.(1990) Computer simulation studies of anisotropic systems. XIX. Mesophases formed by the Gay-Berne model

[37] Allen M. P., M. A. Warren, M. R. Wilson, A. Sauraon, and w. Smith(1996) molecular dynamics calculation of elastic constants in Gay-Berne nematic crystals. J. Chem. Phys.

[38] Rappe, AK and Casewit, CJ and Colwell, KS and Goddard Iii, WA and Skiff, WM (1999) UFF, UFF, a full periodic table force field for molecular mechanics and molecular

[39] Allen M. P., M. A. Warren, M. R. Wilson, A. Sauraon, and w. Smith(1985) molecular dynamics calculation of elastic constants in Gay-Berne nematic crystals. Mol. Phys. 55:

[40] Cleaver, D.J. and Care, C.M. and Allen, M.P. and Neal, M.P.(1996) Extension and

[41] Parrinello, M. and Rahman, A.(1981) Polymorphic transitions in single crystals: A new

[42] Nose, S.(1984) A unified formulation of the constant temperature molecular dynamics

[43] Hoover, W.G.(1985) Canonical dynamics: Equilibrium phase-space distributions. Phys.

[44] Beeman, D.(1976) Some multistep methods for use in molecular dynamics calculations.

[45] Li, M.H. and Keller, P. and Li, B. and Wang, X. and Brunet, M.(2003) Light-driven side-on

generalization of the Gay-Berne potential. Phys. Rev. E 54: 559-567.

mesogen. Liq. Cryst. 8: 451.

dynamics simulations. JOSA 114: 10024-10035.

molecular dynamics method. J. Appl. Phys. 52: 7182.

nematic elastomer actuators. Adv. Mater. E 15: 5569-572.

methods. J. Chem. Phys. 81: 511.

Rev. A 31: 1695-1697.

J. Comput. Phys. 20: 130.

105: 2850-2858.

549.

Liquid crystals (LCs) are unique materials with amazing properties and uses. Since their discovery in 1889 by F. Reinitzer1, they have experienced an explosive growth because of their successful application in a wide variety of areas such as information displays2, cosmetics and health care3, thermography4-5, artificial muscle-like actuators6-7 and enantioselective synthesis8, among others. Thus, liquid crystals play an important role in modern technology and they are present in the most common devices used in our daily live. Research into this field is growing day by day and new promising applications for such materials are discovered and developed continuously.

The liquid-crystalline state or mesophase is a different state of condensed matter, which is intermediate between solids and liquids (**Figure 1**).9 Liquid-crystalline materials combine both the molecular order typical of the solid phases and the fluidity characteristic of the isotropic liquid state. Compared with conventional liquids, liquid crystals exhibit longrange orientational order and, in some cases, some partial positional order. Although the degree of order present in a mesophase is lower in comparison to that in conventional solids, it is high enough to induce a great anisotropy in the properties of the system. Hence, liquid crystals are fluids with anisotropic properties because of the mesogens tendency to point to a common direction called director, *n*.

Liquid-crystalline elastomers (LCEs) consist on high-molecular mass liquid crystals where the mesogenic moieties can be connected either head-to-tail, forming the polymer main chain (main-chain LCEs), or attached as side-chain groups to the main polymer backbone (side-chain LCEs), generally *via* flexible spacers, which allow the polymer main chain to accommodate the anisotropic arrangement of the mesogenic side groups (**Figure 2**).10-11

© 2012 Garcia-Amorós and Velasco, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

**Figure 1.** Schematic representation of the molecular distribution in the solid, nematic liquid-crystalline and liquid state, respectively.

Among all liquid-crystalline phases, the nematic one, where the mesogens are approximately oriented parallel to their longest axis, is the less ordered and, therefore, the less viscous. As a consequence, the mesogens alignment can be more easily manipulated than in the more ordered ones. Hence, nematic liquid-crystalline materials are the most commonly used for technical applications. Specifically, this chapter focuses on the photoactuating properties of nematic liquid-crystalline elastomeric materials.

The properties of liquid-crystalline materials can be easily modulated by applying external perturbations such as light, temperature, electro-magnetic fields, changes of solvent or pH, and so on, which induce changes in the molecular and supermolecular organization of such systems. This feature is the basis of all their further applications.12-13

**Figure 2.** Main-chain (left) and side-chain (right) liquid-crystalline elastomeric materials.

Among all the possible external inputs, light enables a rapid and punctual wireless control of the properties of the material and, moreover, it is a clean, cheap and environmentallyfriendly energy source. Several chromophores are known in photochemistry: spiropyranes, diarylethenes, fulgides, stilbenes, viologens, *etc*. However, azobenzenes are doubtlessly the most used ones for designing optically-controlled materials because of their totally clean and reversible isomerisation process between their two isomers of different stability: *trans* and *cis*. The most interesting feature of azobenzenes is that both isomers can be switched back and forward with light of particular wavelengths: UV light (*hν*1), for the *trans*-to-*cis* conversion, and visible light (*hν*2), for the *cis*-to-*trans* isomerisation. Moreover, *cis* isomer is less stable than the *trans* one because it has a bent shape which increases the steric hindrance and somehow breaks the conjugation of the *trans* linear form. Thus, the metastable *cis* isomer will also relax back spontaneously to the thermodynamically stable *trans* form in the dark isothermally (Δ, **Figure 3**).14

62 Advanced Elastomers – Technology, Properties and Applications

and liquid state, respectively.

**Figure 1.** Schematic representation of the molecular distribution in the solid, nematic liquid-crystalline

Among all liquid-crystalline phases, the nematic one, where the mesogens are approximately oriented parallel to their longest axis, is the less ordered and, therefore, the less viscous. As a consequence, the mesogens alignment can be more easily manipulated than in the more ordered ones. Hence, nematic liquid-crystalline materials are the most commonly used for technical applications. Specifically, this chapter focuses on the photo-

The properties of liquid-crystalline materials can be easily modulated by applying external perturbations such as light, temperature, electro-magnetic fields, changes of solvent or pH, and so on, which induce changes in the molecular and supermolecular organization of such

actuating properties of nematic liquid-crystalline elastomeric materials.

systems. This feature is the basis of all their further applications.12-13

**Figure 2.** Main-chain (left) and side-chain (right) liquid-crystalline elastomeric materials.

Among all the possible external inputs, light enables a rapid and punctual wireless control of the properties of the material and, moreover, it is a clean, cheap and environmentallyfriendly energy source. Several chromophores are known in photochemistry: spiropyranes,

**Figure 3.** Photochromism of azobenzene and energetic profile for its *trans*-to-*cis* and *cis*-to-*trans* isomerisation processes.

The rod-like structure of *trans*-azobenzene allows its easy introduction in both low- and high-molecular mass nematic liquid crystals without causing the destruction of the host mesophase. When the system is irradiated with light of the appropriate wavelength (*hν*1), the bent *cis* form is generated. As a consequence, the orientation of all the mesogenic molecules of the sample will also change (domino effect). In this way, the *cis* azo-moiety acts as an impurity lowering the nematic order. This effect can be nicely observed experimentally when azobenzene-doped liquid-crystalline mixtures are irradiated at a constant temperature, *T*, within *T*N-I (trans) and *T*N-I (cis). In this case, the sample changes from the ordered nematic phase to the disordered isotropic one isothermally, that is, a photoinduced nematic-to-isotropic phase transition occurs (**Figure 4 and 5**).15-17 The initial state is restored on turning off the irradiation due to the thermal back *cis*-to-*trans* isomerisation of the azo-dye in the dark (Δ).18 This effect can be exploited for obtaining promising lightcontrolled liquid-crystalline materials such as photo-active artificial muscle-like actuators and light-driven optical switches.

**Figure 4.** Photo-induced nematic-to-isotropic phase transition in azobenzene-doped nematic liquid crystals.

**Figure 5.** Photo-induced nematic-to-isotropic phase transition in an azobenzene-containing nematic liquid-crystalline mixture. Destruction of the host nematic mesophase by irradiation with UV light at a constant temperature *T* between *T*N-I (cis) and *T*N-I (trans) (up). Regeneration of the nematic mesophase when the thermal *cis*-to-*trans* relaxation of the azobenzene occurs in the dark (down).

## **2. Liquid-crystalline elastomers for light-induced artificial muscle-like actuation**

Photo-active artificial muscle-like actuators convert light into mechanical quantities such as displacement, strain, velocity and stress. Moreover, materials used to produce musclelike movements should be soft and deform easily upon irradiation. Specifically, polymers are very interesting materials for this purpose due to their many attractive properties and characteristics; they are lightweight, inexpensive, easily manufacturable and implementable, fracture tolerant, pliable and biocompatible.19 So, light-driven polymerbased actuators are valuable materials to be implemented in a wide range of micro- and macro-scale devices.

Generally, photo-active liquid-crystalline polymers and elastomers are multidomain systems and, therefore, the director changes abruptly from one domain to another. Hence, both LCPs and LCEs deform themselves when non-polarised light falls on them in an isotropic way, that is, there is no preferential direction for the deformation and, as a consequence, the degree of deformation of the material is generally small.20

64 Advanced Elastomers – Technology, Properties and Applications

crystals.

**actuation** 

macro-scale devices.

**Figure 4.** Photo-induced nematic-to-isotropic phase transition in azobenzene-doped nematic liquid

**Figure 5.** Photo-induced nematic-to-isotropic phase transition in an azobenzene-containing nematic liquid-crystalline mixture. Destruction of the host nematic mesophase by irradiation with UV light at a constant temperature *T* between *T*N-I (cis) and *T*N-I (trans) (up). Regeneration of the nematic mesophase

**2. Liquid-crystalline elastomers for light-induced artificial muscle-like** 

Photo-active artificial muscle-like actuators convert light into mechanical quantities such as displacement, strain, velocity and stress. Moreover, materials used to produce musclelike movements should be soft and deform easily upon irradiation. Specifically, polymers are very interesting materials for this purpose due to their many attractive properties and characteristics; they are lightweight, inexpensive, easily manufacturable and implementable, fracture tolerant, pliable and biocompatible.19 So, light-driven polymerbased actuators are valuable materials to be implemented in a wide range of micro- and

Generally, photo-active liquid-crystalline polymers and elastomers are multidomain systems and, therefore, the director changes abruptly from one domain to another. Hence,

when the thermal *cis*-to-*trans* relaxation of the azobenzene occurs in the dark (down).

Highly noteworthy photo-actuating properties have been successfully achieved in polydomain azobenzene-based liquid-crystalline elastomeric materials by Ikeda.21-26 These materials show a great variety of three-dimensional contraction and expansion movements when they are exposed to polarised light of the appropriate wavelength. Within the last few years, it has been reported that it is possible to control the bending direction of the LCE film depending on the role of the azoderivative within the elastomeric network, that is, if the azo-chromophore acts as a cross-linker or as a simply pendant group.27 Moreover, the first prototypes of real light-controlled plastic rotating motors have been fabricated with laminated films of photo-sensitive LCEs.28

However, if non-polarized light is used instead, it is strictly necessary that the director points to a unique direction in the whole sample, that is, a monodomain sample is required.

Liquid single crystal elastomers (LSCEs) are a subclass of liquid-crystalline polymers, which were synthesized for the first time by Küpfer and Finkelmann in the early nineties29, although their possible use for both thermally- and photo-controlled artificial muscles was predicted earlier theoretically by P. G. de Gennes.30 LSCEs consist in weakly cross-linked polymer networks with a macroscopic orientation of the director, *n*. These materials combine the elasticity typical of conventional rubbers with the anisotropic properties characteristic of liquid-crystalline systems. As a result of this coupling, the macroscopic dimensions of the elastomeric sample can be easily modified by the appropriate modification of the LC order through the application of different external stimuli such as temperature variations or, when an azo-dye is introduced in the system, by irradiating the probe with light of the appropriate wavelength.31

Indeed, LSCEs experiment a spontaneous contraction along the director direction when the system is driven from the ordered nematic phase to the disordered isotropic one by heating the sample over the nematic-to-isotropic phase transition temperature, *T*N-I, due to the mesogens misalignment. If the temperature is lowered below *T*N-I, the LSCE expands back thereby recovering its original shape. This effect is the so-called thermo-mechanical effect (**Figure 6a and 7**).32-36

The introduction of photo-sensitive azoderivatives in LSCEs, not only as side-chain groups but as cross-linkers as well, affords a new way to control the macroscopic dimensions of the sample isothermally just by applying light. So, azobenzene-containing LSCEs are very attractive materials for the fabrication of light-controlled artificial muscle-like actuators.37-38 During the last decades, both the preparation and properties modulation of light-driven artificial muscle-like actuators based on azobenzene-containing LSCEs are topics of intensive research. Thus, in this chapter we will describe the last investigations that have been performed in our group towards not only to enhance the mechanical efficiency of such materials but also to optimise the response time of the final artificial muscle.

**Figure 6.** Thermo-mechanical (a) and opto-mechanical (b) effect in nematic LSCEs.

**Figure 7.** Uniaxial thermal expansion of an LSCE sample on lowering the temperature as a consequence of the isotropic-to-nematic phase transition.

Besides the thermo-mechanical effect commented above, LSCEs that contain isomerisable azobenzenes as light-sensitive molecules also undergo macroscopic contractions in a preferential direction when they are exposed to non-polarised light of the appropriate wavelength (opto-mechanical effect, **Figure 6b**). This happens due to the *trans*-to-*cis* isomerisation of the azo-dye, which drops the nematic order parameter of the elastomeric sample.38 On turning off the light, the system recovers its initial dimensions due to the thermal back isomerisation of the azo-chromophore. Photo-mechanical effect in nematic LSCEs was first observed by Finkelmann *et al.*39 in 2001 and it has been deeply investigated during the last decade from both the theoretical and the experimental point of view. A mathematical description of the opto-mechanical effect will be given in Section 4 of this chapter.

### **3. Preparation and characterization of polysiloxane-based liquid single crystal elastomers**

As it has been commented above, both low- and high-molecular mass liquid crystals present different domains and, therefore, there are some places in the sample where the director changes abruptly. However, for many applications like artificial muscle-like actuation, it is strictly necessary that the system should be macroscopically oriented, *i.e.* a monodomain sample used.

66 Advanced Elastomers – Technology, Properties and Applications

of the isotropic-to-nematic phase transition.

**crystal elastomers** 

**Figure 6.** Thermo-mechanical (a) and opto-mechanical (b) effect in nematic LSCEs.

**Figure 7.** Uniaxial thermal expansion of an LSCE sample on lowering the temperature as a consequence

Besides the thermo-mechanical effect commented above, LSCEs that contain isomerisable azobenzenes as light-sensitive molecules also undergo macroscopic contractions in a preferential direction when they are exposed to non-polarised light of the appropriate wavelength (opto-mechanical effect, **Figure 6b**). This happens due to the *trans*-to-*cis* isomerisation of the azo-dye, which drops the nematic order parameter of the elastomeric sample.38 On turning off the light, the system recovers its initial dimensions due to the thermal back isomerisation of the azo-chromophore. Photo-mechanical effect in nematic LSCEs was first observed by Finkelmann *et al.*39 in 2001 and it has been deeply investigated during the last decade from both the theoretical and the experimental point of view. A mathematical

description of the opto-mechanical effect will be given in Section 4 of this chapter.

**3. Preparation and characterization of polysiloxane-based liquid single** 

As it has been commented above, both low- and high-molecular mass liquid crystals present different domains and, therefore, there are some places in the sample where the director Several techniques are used for this purpose with low-molecular mass LCs like rubbing, surface treatments with polymers or the application of external fields in a determined direction. High-molecular mass LCs, like linear LCPs, can be macroscopically oriented by extruding the polymeric mass in order to create very thin oriented fibres. However, this technique is not useful for LCEs. In such systems, the macroscopic orientation of the sample should be carried out during its preparation.

Many smart and functional liquid-crystalline elastomeric materials have been prepared using different polymer backbones. However, silicone is doubtlessly one of the most commonly used polymer backbone for preparing LSCEs due to its great advantages like good thermal stability, constancy of properties over a wide range of temperature which leads to a large operating temperature range – from −100 to 250 ºC −, hydrophobicity, excellent resistance to oxygen, ozone and sunlight, good flexibility, anti-adhesive properties and low toxicity. An additional noteworthy point is that silicone polymers, of which polyhydrogenomethylsiloxane is an example, have very low glass transition temperatures (*ca*. *T*g = −120 °C). This feature allows polysiloxane-based LSCEs to present photo-actuation even at room temperature.

Generally, polysiloxane-based side-chain liquid single crystal elastomers are prepared following the synthetic methodology in three steps, which was first developed by Küpfer and Finkelmann (**Figure 8**).29

In the first step, the different monomers, those are, mesogen/s, cross-linker/s and comonomer/s, react through their terminal olefin with the polysiloxane backbone through a Pt-catalysed hydrosilylation reaction which takes place at 70 °C (**Figure 9**). This reaction is carried out at 5000 rpm by means of the spin-casting technique. The aim of this step is to obtain a stable elastomeric system, which can be handled properly, but partially cross-linked for its further macroscopic orientation. For this purpose, the hydrosilylation reaction is stopped before it is completely finished.

In the second step, the orientation of the different directors of the entire sample is performed. This process is induced by the application of a uniaxial force to the sample along its longest axis. After the suitable time, a monodomain LCE is obtained, that is, a liquid single crystal elastomer (LSCE). The progress of the orientation process can be followed qualitatively *in situ* by checking the optical transparency of the probe. When the orientation of the director of the different domains is successful, one gets an LSCE, which is perfectly transparent (**Figure 10a**). Contrarily, if a macroscopically disordered sample is obtained, it will show opacity due to the scattering of the incoming light through the sample (**Figure 10c**).

As it has been aforementioned, macroscopic changes in shape are coupled to changes at the molecular scale. Hence, the deformation of an elastomer can lead to reorientation of the

**Figure 8.** Synthetic methodology in three steps used for preparing photo-active side-chain polysiloxane-based nematic LSCEs.

**Figure 9.** Attachment of an alkene-terminated monomer to the polysiloxane backbone through the Ptcatalysed hydrosilylation reaction. (Pt(COD)Cl2 = dichloro(1,5-cyclooctadiene)platinum(II)).

**Figure 10.** Qualitative evaluation of the macroscopic orientation of the different directors of the elastomeric sample by means of their optical transparency: monodomain (a), partial monodomain (b) and polydomain (c).

director or vice versa, reorientation of the mesogens can change the shape of the elastomer (see thermo- and opto-mechanical effects in Section 2 of this chapter). In this way, since an LSCE is prepared, its macroscopic dimensions are directly determined by the degree of order created during the second stage of its fabrication. So, the macroscopic orientation of the different directors of the system becomes the key step of the whole synthetic procedure since it will determine the proper actuation of the final material.

68 Advanced Elastomers – Technology, Properties and Applications

polysiloxane-based nematic LSCEs.

and polydomain (c).

**Figure 8.** Synthetic methodology in three steps used for preparing photo-active side-chain

**Figure 9.** Attachment of an alkene-terminated monomer to the polysiloxane backbone through the Pt-

catalysed hydrosilylation reaction. (Pt(COD)Cl2 = dichloro(1,5-cyclooctadiene)platinum(II)).

**Figure 10.** Qualitative evaluation of the macroscopic orientation of the different directors of the elastomeric sample by means of their optical transparency: monodomain (a), partial monodomain (b) In the third stage of the synthesis, the created anisotropy should be fixed by a second crosslinking reaction without removing the applied force. This second hydrosilylation reaction is carried out completely in an oven at 70 °C for 48 h. Afterwards, the LSCE is purified in order to remove both the catalyst and the non-reacted monomers (*ca.* less than 1%) from the network. This cleaning procedure is carried out by a swelling-deswelling process using acetone and hexanes, respectively. Finally, the complete characterisation of the prepared LSCE should be done. The conventional techniques used for characterising nematic LSCEs are briefly described below.

Not only the proper macroscopic orientation but also both the thermal and the mechanical properties of liquid-crystalline elastomers can be studied by means of a wide variety of techniques. Generally, polarised optical microscopy (POM) and X-Ray diffraction (XRD) are used together to get insight into the structure of the mesophase unambiguously but also to evaluate its macroscopic ordering. Moreover, differential scanning calorimetry (DSC) is needed to determine the temperature range of stability of the LC phase. Besides these techniques, additional experiments, such as swelling experiments, are performed on LSCEs to be used further for muscle-like actuation purposes.

The density of cross-linking units present in an LSCE, which is directly related with the mechanical properties of the elastomeric material, can be easily evaluated by means of swelling experiments. Liquid-crystalline elastomers, as a main difference with linear liquidcrystalline polymers and other solids, cannot be dissolved due to the presence of crosslinking points. When an LSCE is immersed in a suitable solvent, it absorbs a large amount of it without dissolving and, as a consequence, it experiments a large deformation producing a small internal stress in the network. Hence, the free energy change that takes place in the elastomer during its swelling process can be separated in two additive contributions, those are, the free energy of mixing, Δ*G*mix, and the free energy related with the elastic deformation of the network, Δ*G*el. 40-42

Traditional swelling experiments are carried out by placing the sample in a vessel containing a suitable solvent (*e.g.* toluene).43 Then, the system is allowed to reach the thermodynamically equilibrium conditions (Δ*G*swelling = 0; Δ*G*mix = -Δ*G*el). The swelling parameter, *q*, is defined as the ratio of volumes of the swollen and original elastomer both measured under thermodynamically equilibrium conditions. Because of the anisotropy of the liquid-crystalline elastomer in its *z*- direction (**Figure 11**), both *x*- and *y*-isotropic dimensions change in the same way but differently from the *z*- one. Thus, the swelling parameter can be expressed as it is indicated in eq. 1.

$$q = \frac{V}{V\_0} = \frac{\mathbf{x} \cdot \mathbf{y} \cdot \mathbf{z}}{\mathbf{x}\_0 \cdot \mathbf{y}\_0 \cdot \mathbf{z}\_0} = \frac{\mathbf{z}}{\mathbf{z}\_0} \cdot \left(\frac{\mathbf{x}}{\mathbf{x}\_0}\right)^2 \tag{1}$$

$$\mathbf{z}\_0 \left\| \underbrace{\mathbf{z}}\_{\text{toluene}} \xrightarrow[\text{t. 24 h}]}\_{\mathbf{t} \text{-. 24 h}} \mathbf{z} \right\|\_{\mathbf{y}} \mathbf{y}\_0$$

**Figure 11.** Change in the elastomer volume on reaching thermodynamically equilibrium after swelling.

According to that mentioned above, those samples that contain a high density of crosslinking units will yield a low value of the swelling parameter. However, it is highly remarkable that the magnitude of the swelling parameter depends also greatly on the chemical nature of the cross-linker. In this way, elastomers with more rigid cross-linkers afford lower values of the swelling parameter than those with more flexible ones, although they contain the same density of cross-linking points.44

### **4. Opto-mechanical effect in liquid single crystal elastomers**

The macroscopic dimensions of an LSCE are directly related with the degree of order induced during its fabrication. So, any reorientation of the mesogens will lead to changes in the shape of the elastomer. When a photo-active LSCE is illuminated with light of the appropriate wavelength, the azo-chromophore changes its geometry from linear to bent due to its *trans*-to-*cis* photo-isomerisation. This fact produces a microscopic disorganisation of those mesogen molecules that are close to the azo-dye ones producing a decrease in the local order parameter. As a consequence of this molecular disorganisation, a shortening of the LSCE in the director direction is observed (see **Figure 6b**). If the network is fixed by both ends, the system cannot shrink and, as a consequence, the appearance of a retractive force in the elastomer is observed.38-39

Opto-mechanical experiments consist in the measurement of the evolution of the internal stress generated inside the LSCE with the time. A typical opto-mechanical experiment for a photo-active liquid single crystal elastomer is shown in **Figure 12**. On turning on the light, the internal stress created in the elastomer, *σ*, grows until the photo-stationary state is reached. The curve describes a *plateau* which corresponds to the maximum opto-mechanical response produced by the artificial muscle-like actuator, Δ*σ*max. When the irradiation is ceased, the thermal back *cis*-to-*trans* isomerisation of the azo-dye occurs and the stress starts to diminish with the time until the initial stress value is recovered.

they contain the same density of cross-linking points.44

the elastomer is observed.38-39

**4. Opto-mechanical effect in liquid single crystal elastomers** 

to diminish with the time until the initial stress value is recovered.

2

(1)

0 000 0 0 · *<sup>V</sup> xyz z x <sup>q</sup> V xyz z x* 

**Figure 11.** Change in the elastomer volume on reaching thermodynamically equilibrium after swelling.

According to that mentioned above, those samples that contain a high density of crosslinking units will yield a low value of the swelling parameter. However, it is highly remarkable that the magnitude of the swelling parameter depends also greatly on the chemical nature of the cross-linker. In this way, elastomers with more rigid cross-linkers afford lower values of the swelling parameter than those with more flexible ones, although

The macroscopic dimensions of an LSCE are directly related with the degree of order induced during its fabrication. So, any reorientation of the mesogens will lead to changes in the shape of the elastomer. When a photo-active LSCE is illuminated with light of the appropriate wavelength, the azo-chromophore changes its geometry from linear to bent due to its *trans*-to-*cis* photo-isomerisation. This fact produces a microscopic disorganisation of those mesogen molecules that are close to the azo-dye ones producing a decrease in the local order parameter. As a consequence of this molecular disorganisation, a shortening of the LSCE in the director direction is observed (see **Figure 6b**). If the network is fixed by both ends, the system cannot shrink and, as a consequence, the appearance of a retractive force in

Opto-mechanical experiments consist in the measurement of the evolution of the internal stress generated inside the LSCE with the time. A typical opto-mechanical experiment for a photo-active liquid single crystal elastomer is shown in **Figure 12**. On turning on the light, the internal stress created in the elastomer, *σ*, grows until the photo-stationary state is reached. The curve describes a *plateau* which corresponds to the maximum opto-mechanical response produced by the artificial muscle-like actuator, Δ*σ*max. When the irradiation is ceased, the thermal back *cis*-to-*trans* isomerisation of the azo-dye occurs and the stress starts

**Figure 12.** Opto-mechanical experiments: increase of the internal stress generated in the network upon irradiation with UV light (left) and decrease of the stress with the time in the dark at a constant temperature, *T* (right).

Besides the maximum stress that the system is able to generate when UV light falls on it, Δ*σ*max, opto-mechanical experiments allow also determining the time required by the network to produce its maximum mechanical response, *τ*irrad, and that necessary to recover the initial state thermally, *τ*th. These three parameters are crucial in the overall performance of LSCEs for artificial muscle-like photo-actuation.

The UV-irradiation process of an LSCE is a bidirectional step, since the *trans*-to-*cis* photoisomerisation reaction competes with the thermal *cis*-to-*trans* isomerisation simultaneously.

$$
tau \xrightarrow[k^{\text{th}}]{} \underset{k^{\text{th}}}{\longleftrightarrow} \text{cis} \tag{2}
$$

Both individual isomerisation processes are well known to follow a first order profile for low-molecular weight azoderivatives in isotropic and nematic solution as well as in dense polymer matrixes.45 Assuming a similar kinetic behaviour for LSCEs, the rate of this process is given by both the disappearance of the *trans* isomer with the time by the photo-induced reaction and its formation through the thermally-activated one (eq. 3).

$$\upsilon = -\frac{d\lbrack T\rbrack}{dt} = k^{\textnormal{ph}} \cdot \lbrack T\rbrack - k^{\textnormal{th}} \cdot \lbrack \mathbb{C}\rbrack = k^{\textnormal{ph}} \cdot \lbrack T\rbrack - k^{\textnormal{th}} \cdot \left( \lbrack T\rbrack\_{0} - \lbrack T\rbrack \right) \tag{3}$$

*k*ph and *k*th correspond to the first-order rate constants for the photo-induced *trans*-to-*cis* and the thermal *cis*-to-*trans* isomerisation processes of the elastomeric network, respectively. [*T*] and [*C*] stand for the concentration of *trans* and *cis* isomer at each moment of the reaction, respectively, and [*T*]0 is the total concentration of azo-dye present in the elastomeric system, that is, the total number of *trans* plus *cis* isomers. [*T*]0 is a constant value and it is given by the initial composition of the material.

Once integrated the differential equation 3, one gets the time-dependence of the concentration of both isomers *trans* (eq. 4) and *cis* (eq. 5). The curve describing this process is

an exponential growth with an apparent rate constant, *k*irrad, which is related with the firstorder kinetic constants of both processes (*k*irrad = *k*ph + *k*th).

$$\mathbb{E}\left[T\right] = \left[T\right]\_0 \frac{k^{\text{th}} + k^{\text{ph}} \cdot \exp\left[-\left(k^{\text{ph}} + k^{\text{th}}\right) \cdot t\right]}{k^{\text{th}} + k^{\text{ph}}} = \left[T\right]\_0 \frac{k^{\text{th}} + k^{\text{ph}} \cdot \exp\left[-k^{\text{irrad}} \cdot t\right]}{k^{\text{irrad}}}\tag{4}$$

$$\mathbb{E}[C] = [T]\_0 \frac{k^{\text{ph}}}{k^{\text{th}} + k^{\text{ph}}} \left( 1 - \exp\left[ - (k^{\text{ph}} + k^{\text{th}}) \cdot t \right] \right) = [T]\_0 \frac{k^{\text{ph}}}{k^{\text{irrad}}} \left( 1 - \exp\left[ - k^{\text{irrad}} \cdot t \right] \right) \tag{5}$$

As we mentioned above, the generation of the *cis*-isomer in the LSCE is the responsible of the mechanical stress observed when the sample is UV-irradiated. In fact, the bent *cis* isomer should be considered as an impurity which shifts the critical temperature of the nematic-toisotropic phase transition of the LSCE (see Section 1 of this Chapter).38 Hence, equation 5 can be transformed into equation 6 which stands for the variation of the measured internal stress produced by the elastomer shrinking with the time when it is illuminated with UV light.

$$
\sigma\_t - \sigma\_0 = \Lambda \sigma\_{\text{max}} \cdot \left(1 - \exp\left[-k^{\text{irrad}} \cdot t\right]\right) \tag{6}
$$

On the other hand, the thermal relaxation process of the LSCE is a unidirectional step.

$$\mathbf{c} \mathbf{i} \xrightarrow{\mathbf{k}^{\mu}} \mathbf{t} \mathbf{r} \mathbf{s} \tag{7}$$

In this case, the rate of the process is given uniquely by the disappearance of the *cis* isomer with the time. The rate equation for this unimolecular first-order reaction is described by equation 8.

$$
\sigma = -\frac{d\left[\overline{\mathcal{C}}\right]}{dt} = k^{\text{th}} \cdot \left[\overline{\mathcal{C}}\right] \tag{8}
$$

The integration of the differential equation 8 yields the integrated rate equation 9 which describes the evolution of the *cis* isomer concentration with the time during the thermal relaxation of the network. The curve describing this process is an exponential decrease with a first-order rate constant, *k*th.

$$\mathbf{I}[\mathbf{C}] = [\mathbf{C}]\_0 \cdot \exp\left(-k^{\text{th}} \cdot t\right) \tag{9}$$

In the same manner than for eq. 6, equation 9 can be transformed into eq. 10 which fits for the variation of the measured internal stress in the elastomer with the time when it is kept in the dark at a constant temperature, *T*:

$$
\sigma\_\text{t} - \sigma\_0 = \Delta \sigma\_{\text{max}} \cdot \exp\left(-k^{\text{th}} \cdot t\right) \tag{10}
$$

The maximum stress that the system is able to generate upon irradiation, Δ*σ*max, is directly obtained by subtracting the initial stress, *σ*0, from the steady-state saturation value reached under irradiation (see **Figure 12**). Both *k*irrad and *k*th are determined by fitting equation 6 and 10 to the experimental data, respectively. The characteristic times for both processes, *τ*irrad and *τ*th, are calculated from *k*irrad and *k*th, respectively, as *τ =* 1/*k.*

72 Advanced Elastomers – Technology, Properties and Applications

order kinetic constants of both processes (*k*irrad = *k*ph + *k*th).

[] [] [ ]

t 0 max

 

On the other hand, the thermal relaxation process of the LSCE is a unidirectional step.

*th*

In this case, the rate of the process is given uniquely by the disappearance of the *cis* isomer with the time. The rate equation for this unimolecular first-order reaction is described by

> th · *d C v kC dt*

The integration of the differential equation 8 yields the integrated rate equation 9 which describes the evolution of the *cis* isomer concentration with the time during the thermal relaxation of the network. The curve describing this process is an exponential decrease with

In the same manner than for eq. 6, equation 9 can be transformed into eq. 10 which fits for the variation of the measured internal stress in the elastomer with the time when it is kept in

The maximum stress that the system is able to generate upon irradiation, Δ*σ*max, is directly obtained by subtracting the initial stress, *σ*0, from the steady-state saturation value reached

t 0 max

 

th

th

equation 8.

a first-order rate constant, *k*th.

the dark at a constant temperature, *T*:

*T T T*

an exponential growth with an apparent rate constant, *k*irrad, which is related with the first-

0 0 th ph irrad

0 0 th ph irrad [] [] 1 exp ( ) [ ] 1 exp *k k C T k kt T k t k k <sup>k</sup>*

As we mentioned above, the generation of the *cis*-isomer in the LSCE is the responsible of the mechanical stress observed when the sample is UV-irradiated. In fact, the bent *cis* isomer should be considered as an impurity which shifts the critical temperature of the nematic-toisotropic phase transition of the LSCE (see Section 1 of this Chapter).38 Hence, equation 5 can be transformed into equation 6 which stands for the variation of the measured internal stress produced by the elastomer shrinking with the time when it is illuminated with UV light.

th ph ph th th ph irrad

*k k k* 

ph ph ph th irrad

irrad

1 exp *k t* (6)

*<sup>k</sup> cis trans* (7)

(8)

<sup>0</sup> [ ] [ ] exp *C C kt* (9)

exp *k t* (10)

exp ( ) exp

*kk k kt kk k t*

(4)

(5)

The thermal activation parameters, the enthalpy (Δ*H* ≠) and entropy (Δ*S* ≠) of activation, can be easily determined by studying the evolution of both isomerisation rate constants as a function of the temperature. This behaviour is well described by means of Eyring's equation (eq. 11).46-47

$$
\ln \frac{k}{T} = \frac{-\Delta H^{\ast}}{R} \cdot \frac{1}{T} + \ln \frac{k\_{\text{B}}}{h} + \frac{\Delta S^{\ast}}{R} \tag{11}
$$

*R* is the universal gas constant (*R* = 8.314 J·K-1·mol-1), *k*B is the Boltzmann's constant (*k*B = 1.381 × 10-23 J·K-1) and *h* is the Planck's constant (*h* = 6.626 × 10-34 J·s), respectively.

The characteristic parameters (Δ*σ*max, *τ*irrad and *τ*th) for all the artificial muscle-like actuators presented herein have been determined according to the kinetic model described above. The relaxation time registered for the different photo-active LSCEs reported in this chapter ranges from several hours to a few seconds. This fact demonstrates that not only fast-responding artificial muscles but also bistable systems can be obtained easily from polysiloxane-based LSCEs just through the proper substitution of the light-sensitive azo-chromophore.

## **5. Mechanical efficiency of photo-active liquid single crystal elastomerbased actuators through the variation of the azo-cross-linker flexible spacer**

Light-controlled artificial muscle-like actuators based on photo-active polysiloxane LSCEs have been growing interest during the last decade and deeply investigated from both the theoretical and the experimental point of view. As it has been aforementioned, two key parameters are needed to characterise properly the actuation ability of such materials, those are, the maximum mechanical response that they are able to generate by irradiation with light of the appropriate wavelength as well as the time required to produce it and that to recover the initial state further. At the present moment, many efforts are being put forward to improve and optimise both of them.

At this point, we will turn our attention towards the enhancement of the mechanical response produced by those photo-actuators based on polysiloxane azobenzene-containing LSCEs. One of the factors that has been probed to play a clear role on the mechanical response produced by such light-driven actuators is the variation of the photo-active azocross-linker spacer length, in other words, the flexible alkyl chain that links the azobenzene core with the main polysiloxane backbone (**Figure 13**).48

In order to describe this effect, five different photo-active elastomers containing the nematic mesogen 4-methoxyphenyl-4-(3-butenyloxy)benzoate (**M4OMe**, 90 % mol), the isotropic cross-linker 1,4-di-(10-undecenyloxy)benzene (**V1**, 5 % mol) and one of the light-sensitive

cross-linkers shown in **Figure 13** (**AZOX**, 5 % mol) will be considered (**Figure 14**). All elastomers were prepared and characterised according to that described previously in Section 3 and 4 of this chapter. The photo-active azo-cross-linkers present lateral alkoxyl chains of different lengths, bearing 3, 4, 6, 8 and 11 carbon atoms, *n*, and thereby producing a total number of methylene units in the flexible spacer, *n*total, of 6, 8, 12, 16 and 22, respectively. All the systems exhibited a broad enantiotropic nematic phase between their glass transition temperature at *T*g = 261-276 K and their nematic-to-isotropic phase transition temperature at *T*N-I = 336-342 K (Δ*H*N-I = 1.6-2.1 J·g-1). Table 1 displays the temperature range of stability of the nematic liquid-crystalline phase for each LSCE. Moreover, all of them showed a clear macroscopic orientation of the director as it reveals their orientational order parameter which falls between 0.71 and 0.73. The swelling ratio, *q*, of the different elastomers ranges from 2.5 to 4.3. According to that discussed in Section 3 of this chapter, the slight differences observed in the swelling parameter for the different systems can be related with the different length of their cross-linker spacer. In this way, elastomer **EAZO3**, which has the most rigid cross-linker, shows the lowest *q* value (*q* = 2.5) whereas **EAZO11**, with the most flexible one, exhibits the highest swelling parameter (*q* = 4.3, Table 1) since it can create bigger empty spaces inside the elastomeric system.

**Figure 13.** Concept of flexible spacer and general chemical structure of the different photo-active azo cross-linkers used in this study.


**Table 1.** Nematic order parameter, *S*; glass transition and nematic-to-isotropic phase transition temperatures, *T*g and *T*N-I; nematic-to-isotropic phase transition enthalpy, **∆***H*N-I; and swelling parameter, *q*.

#### Polysiloxane Side-Chain Azobenzene-Containing Liquid Single Crystal Elastomers for Photo-Active Artificial Muscle-Like Actuators 75

**Figure 14.** Chemical composition of the different photo-active nematic LSCEs **EAZOX**.

74 Advanced Elastomers – Technology, Properties and Applications

can create bigger empty spaces inside the elastomeric system.

cross-linkers used in this study.

parameter, *q*.

**Elastomer** *S T***<sup>g</sup>**

cross-linkers shown in **Figure 13** (**AZOX**, 5 % mol) will be considered (**Figure 14**). All elastomers were prepared and characterised according to that described previously in Section 3 and 4 of this chapter. The photo-active azo-cross-linkers present lateral alkoxyl chains of different lengths, bearing 3, 4, 6, 8 and 11 carbon atoms, *n*, and thereby producing a total number of methylene units in the flexible spacer, *n*total, of 6, 8, 12, 16 and 22, respectively. All the systems exhibited a broad enantiotropic nematic phase between their glass transition temperature at *T*g = 261-276 K and their nematic-to-isotropic phase transition temperature at *T*N-I = 336-342 K (Δ*H*N-I = 1.6-2.1 J·g-1). Table 1 displays the temperature range of stability of the nematic liquid-crystalline phase for each LSCE. Moreover, all of them showed a clear macroscopic orientation of the director as it reveals their orientational order parameter which falls between 0.71 and 0.73. The swelling ratio, *q*, of the different elastomers ranges from 2.5 to 4.3. According to that discussed in Section 3 of this chapter, the slight differences observed in the swelling parameter for the different systems can be related with the different length of their cross-linker spacer. In this way, elastomer **EAZO3**, which has the most rigid cross-linker, shows the lowest *q* value (*q* = 2.5) whereas **EAZO11**, with the most flexible one, exhibits the highest swelling parameter (*q* = 4.3, Table 1) since it

**Figure 13.** Concept of flexible spacer and general chemical structure of the different photo-active azo

**(K)** 

**EAZO3** 0.71 265 336 1.9 2.5 **EAZO4** 0.73 261 340 2.1 3.0 **EAZO6** 0.73 264 342 1.6 3.3 **EAZO8** 0.73 284 348 1.6 3.0 **EAZO11** 0.72 276 342 1.8 4.3 **Table 1.** Nematic order parameter, *S*; glass transition and nematic-to-isotropic phase transition temperatures, *T*g and *T*N-I; nematic-to-isotropic phase transition enthalpy, **∆***H*N-I; and swelling

*T***N-I (K)**  **∆***H***N-I** 

**(J·g-1)** *<sup>q</sup>*

The maximum opto-mechanical response produced by the different LSCEs **EAZOX** under UV-irradiation, Δ*σ*max, depends clearly not only on the number of methylene units of the flexible spacer (**Figure 15**, left) but also on the temperature (**Figure 15**, right). It is highly remarkable that very much efficient light-controlled artificial muscle-like actuators are obtained on using photo-active cross-linkers with long alkoxyl chains in their flexible spacers. Hence, a maximum opto-mechanical response of Δ*σ*max= 19.0 kPa is registered for the LSCE **EAZO11** at 333 K, which bears the longest spacer. Otherwise, a three-fold lower value of Δ*σ*max= 6.4 kPa is obtained for the nematic elastomer **EAZO3** at the same temperature, which contains the shortest one.

**Figure 15.** Opto-mechanical experiments for the different nematic liquid single crystal elastomers under irradiation with UV light (*λ*irrad = 380 nm) at 333 K (left). Evolution of the maximum optomechanical response, Δ*σ*max, generated by the different elastomers with the temperature (right).

According to the molecular model presented in **Figure 16**, azo-cross-linkers with long spacers can connect two quite separated points of the main polysiloxane backbone in comparison with their shorter counterparts. As the azo-cross-linker is clamped by both ends, the constraint of the elastomer after the photo-isomerisation to the *cis*-isomer should be greater for those systems which connect two points of the polymeric backbone which are far away separated. Indeed, the cross-linker with the longest spacer experiments the highest opto-mechanical response when the azo moiety adopts its bent geometry.

As a representative example, **Figure 16** shows the calculated distance for the all *trans* conformation of both **AZO3** and **AZO11**, which bear a total number of methylene units in their flexible spacer, *n*total, of 6 and 22, respectively. For *trans*-**AZO3** and *trans*-**AZO11**, the calculated distance between both terminal carbon atoms are 17.6 and 35.5 Å, while for their corresponding *cis* form, which has a bent shape, it drops until 9.8 and 18.8 Å, respectively. This fact is reflected in a distance change between both terminal carbon atoms that are attached to the main polymeric chain of the network of 7.8 and 16.7 Å, respectively, when the azoderivative is isomerised upon UV-irradiation. Hence, from the calculated distances, it should be expected a greater effect on the macroscopic dimensions of the elastomer when azo-cross-linkers with larger flexible spacers are used. Although the presented model is a qualitative estimation, since the flexible spacer of the azo-crosslinker undergoes a conformational equilibrium, it stands for the different experimental observations.

**Figure 16.** Schematic model of the bonding points of the photo-active azo-cross-linker with the elastomeric network (left). Molecular model for the all *trans* conformation of azo-dyes **AZO3** and **AZO11** and calculated distances between both terminal carbon atoms (right) (Δ*d* = *d*trans – *d*cis).

On the other hand, a marked decrease of the opto-mechanical response of the artificial muscle-like actuator on rising the temperature is detected. This phenomenon is well understood since the thermal back isomerisation process, which competes with the *trans*-to*cis* photo-isomerisation reaction, takes place faster on increasing the temperature and, consequently, the *cis* isomer concentration at the photo-stationary state, which is the responsible of the observed mechanical response, is lower at higher temperatures.

Moreover, the maximum opto-mechanical response at a determined temperature, *T*, shows a clear dependence on the total methylene units of the cross-linker spacer, *n*total (*n*total = 2*n* + 2, where *n* is the number of methylene units of the alkoxyl chain). A growing exponential behaviour tending to a *plateau* for a long spacer length is observed at each temperature. It can be nicely seen from **Figure 17**, that there is a threshold length value for the flexible spacer from which the mechanical response of the system is not enhanced anymore. The threshold spacer length value for the selected architecture of the photo-active azoderivative corresponds to *ca. n*total ≈ 30 methylene units, that is, 15 carbon atoms in each alkoxyl chain of the flexible spacer.

76 Advanced Elastomers – Technology, Properties and Applications

observations.

According to the molecular model presented in **Figure 16**, azo-cross-linkers with long spacers can connect two quite separated points of the main polysiloxane backbone in comparison with their shorter counterparts. As the azo-cross-linker is clamped by both ends, the constraint of the elastomer after the photo-isomerisation to the *cis*-isomer should be greater for those systems which connect two points of the polymeric backbone which are far away separated. Indeed, the cross-linker with the longest spacer experiments the highest

As a representative example, **Figure 16** shows the calculated distance for the all *trans* conformation of both **AZO3** and **AZO11**, which bear a total number of methylene units in their flexible spacer, *n*total, of 6 and 22, respectively. For *trans*-**AZO3** and *trans*-**AZO11**, the calculated distance between both terminal carbon atoms are 17.6 and 35.5 Å, while for their corresponding *cis* form, which has a bent shape, it drops until 9.8 and 18.8 Å, respectively. This fact is reflected in a distance change between both terminal carbon atoms that are attached to the main polymeric chain of the network of 7.8 and 16.7 Å, respectively, when the azoderivative is isomerised upon UV-irradiation. Hence, from the calculated distances, it should be expected a greater effect on the macroscopic dimensions of the elastomer when azo-cross-linkers with larger flexible spacers are used. Although the presented model is a qualitative estimation, since the flexible spacer of the azo-crosslinker undergoes a conformational equilibrium, it stands for the different experimental

**Figure 16.** Schematic model of the bonding points of the photo-active azo-cross-linker with the elastomeric network (left). Molecular model for the all *trans* conformation of azo-dyes **AZO3** and **AZO11** and calculated distances between both terminal carbon atoms (right) (Δ*d* = *d*trans – *d*cis).

responsible of the observed mechanical response, is lower at higher temperatures.

On the other hand, a marked decrease of the opto-mechanical response of the artificial muscle-like actuator on rising the temperature is detected. This phenomenon is well understood since the thermal back isomerisation process, which competes with the *trans*-to*cis* photo-isomerisation reaction, takes place faster on increasing the temperature and, consequently, the *cis* isomer concentration at the photo-stationary state, which is the

opto-mechanical response when the azo moiety adopts its bent geometry.

**Figure 17.** Evolution of the maximum opto-mechanical response, Δ*σ*max, with the total number of methylene units in the spacer, *n*total, at different temperatures.

The result displayed in **Figure 17** evidences that there is a maximum opto-mechanical response for each type of azo-cross-linker used. In this case, responses up to 30 kPa can be obtained using 4,4'-dialkoxysubstituted azo-dyes as photo-active cross-linkers at 323 K. Considering a linear relationship between Δ*σ*max and *T* (see **Figure 15**, right), it can be concluded that this type of azo-cross-linkers can afford responses up to 60 kPa at room temperature.

The analysis of the characteristic times found for both the whole UV-irradiation (*τ*irrad) and the thermal relaxation (*τ*th) process when the azo-dye is incorporated as a cross-linker in the nematic LSCE yields that the rate of both processes is independent of the length of the azocross-linker spacer. The relaxation time for the irradiation process ranges from 15 to 21 minutes at 323 K. For the reverse process, the thermal isomerisation in the dark, relaxation times ranging from 29 to 35 minutes were registered for all the LSCEs at 323 K, except for **EAZO3** which showed a slightly higher value of 45 minutes, probably related with the lower flexibility of the azo-cross-linker (**Table 2**). The thermal *cis-*to*-trans* isomerisation of

the azo-dye takes place faster inside the nematic LSCE than in isotropic media, but it is similar to that obtained in low molar mass liquid crystals. This fact evidences that the kinetic acceleration of the thermal *cis*-to-*trans* isomerisation of the azo-dye in the LSCE is mainly due to the presence of the nematic mean field in the network.49-51


**Table 2.** Relaxation time for the *trans* isomer under irradiation with UV light, *τ*irrad, and for the *cis* isomer on heating in the dark, *τ*th at different temperatures. All values are given in minutes.

Although all the photo-active LSCEs presented till this moment are mechanically efficient, they need several hours to reach their maximum mechanical response and also to relax back to the initial state.37,38,48 Hence, the time required by the network to recover its initial dimensions is also another crucial parameter to consider for obtaining functional artificial muscle-like actuators. Likewise, our discussion will be focused in the next section on presenting different strategies to get photo-active artificial muscle-like actuators with a low thermal relaxation time.

## **6. Response time of artificial muscle-like actuators by using fast thermally-isomerising azoderivatives**

While the *trans*-to-*cis* photo-isomerisation can be easily accelerated either by using more powerful light sources or the appropriate modification of the azobenzene core, the thermal *cis*-to-*trans* back reaction depends mainly not only on the chemical functionalization of the azo-dye but also on the environment where the azo-chromophore is located. Hence, it is essential for getting fast azobenzene-based artificial muscle-like actuators that the return to the thermodynamically stable *trans* form of the azo-chromophore in the dark elapses as fast as possible.

Only two examples of dye-doped polysiloxane-based photo-active nematic LSCEs with a fast isomerisation rate have been published so far. These systems use the well-known push-pull azo-dyes 4-amino- and 4-*N*,*N*-dimethylamino-4'-nitroazobenzene as photo-active molecules, which are doped into host elastomeric networks but not covalently bonded to the polymeric structure52,53; this fact decreases the stability of the final photo-actuator. Besides those azoderivatives which bear a push-pull configuration, azophenols are also endowed with a very rapid thermal *cis*-to-*trans* isomerisation process since they are capable to establish azohydrazone tautomeric equilibria which allows their thermal back reaction to proceed through the rotational isomerisation mechanism. Hence, these two type of azoderivatives are valuable photo-active molecules to be introduced as covalently bonded comonomers in LSCE to obtain stable and fast responding photo-sensitive elastomeric materials.

### **6.1. Artificial muscle-like actuators using push-pull azoderivatives as chromophores**

78 Advanced Elastomers – Technology, Properties and Applications

**Elastomer** 

thermal relaxation time.

as possible.

**thermally-isomerising azoderivatives** 

due to the presence of the nematic mean field in the network.49-51

the azo-dye takes place faster inside the nematic LSCE than in isotropic media, but it is similar to that obtained in low molar mass liquid crystals. This fact evidences that the kinetic acceleration of the thermal *cis*-to-*trans* isomerisation of the azo-dye in the LSCE is mainly

**EAZO3** 19.8 44.1 15.7 30.0 12.4 20.0 **EAZO4** 18.3 34.5 14.1 22.5 10.8 15.9 **EAZO6** 15.3 29.0 12.8 20.0 10.7 14.7 **EAZO8** 18.5 30.3 15.1 20.2 12.8 14.4 **EAZO11** 20.5 33.5 17.3 23.8 14.4 16.4 **Table 2.** Relaxation time for the *trans* isomer under irradiation with UV light, *τ*irrad, and for the *cis* isomer on heating in the dark, *τ*th at different temperatures. All values are given in minutes.

Although all the photo-active LSCEs presented till this moment are mechanically efficient, they need several hours to reach their maximum mechanical response and also to relax back to the initial state.37,38,48 Hence, the time required by the network to recover its initial dimensions is also another crucial parameter to consider for obtaining functional artificial muscle-like actuators. Likewise, our discussion will be focused in the next section on presenting different strategies to get photo-active artificial muscle-like actuators with a low

While the *trans*-to-*cis* photo-isomerisation can be easily accelerated either by using more powerful light sources or the appropriate modification of the azobenzene core, the thermal *cis*-to-*trans* back reaction depends mainly not only on the chemical functionalization of the azo-dye but also on the environment where the azo-chromophore is located. Hence, it is essential for getting fast azobenzene-based artificial muscle-like actuators that the return to the thermodynamically stable *trans* form of the azo-chromophore in the dark elapses as fast

Only two examples of dye-doped polysiloxane-based photo-active nematic LSCEs with a fast isomerisation rate have been published so far. These systems use the well-known push-pull azo-dyes 4-amino- and 4-*N*,*N*-dimethylamino-4'-nitroazobenzene as photo-active molecules, which are doped into host elastomeric networks but not covalently bonded to the polymeric structure52,53; this fact decreases the stability of the final photo-actuator. Besides those azoderivatives which bear a push-pull configuration, azophenols are also endowed with a very rapid thermal *cis*-to-*trans* isomerisation process since they are capable to establish azohydrazone tautomeric equilibria which allows their thermal back reaction to proceed through

**6. Response time of artificial muscle-like actuators by using fast** 

323 K 328 K 333 K *τ***irrad** *τ***th** *τ***irrad** *τ***th** *τ***irrad** *τ***th**

Since an azo-dye is introduced into a polymer, the polymer backbone undergoes different motions when the thermal relaxation of the azoderivative occurs. These motions may modify the kinetic parameters of the *cis*-to-*trans* isomerisation of the azo chromophore. In this way, structural factors such as chain flexibility and chain conformation play an important role on the rate of the thermal isomerisation process of the chromophore.45,54,55 Moreover, very different kinetic behaviour should be expected when the azo-dye is bonded to the polymer backbone than when it is just dissolved in the host polymer matrix as a doped guest. This effect has been already detected in LSCEs that contain the push-pull azodyes 4-(5-hexenyloxy)-4'-nitroazobenzene and 4-(5-hexenyloxy)-4'-methoxy-2'-nitroazobenzene introduced as side-chain covalently-bonded photo-active moieties (**Figure 18**).56

**Figure 18.** Chemical composition of the photo-active nematic LSCEs **E***p***NO2** and **E***o***NO2**.

For describing this effect, two nematic LSCEs **E***p***NO2** and **E***o***NO2** will be considered being composed by the nematogen **M4OMe** (85 % mol), the photoactive co-monomer (**AZO***p***NO2** or

**AZO***o***NO2**, respectively, 5 % mol) and the isotropic cross-linking agent **V1** (10 % mol) as it is shown in **Figure 18**. Both elastomers, **E***p***NO2** and **E***o***NO2**, which differ only in the placement of the nitro- electron-withdrawing group, showed a broad enantiotropic nematic phase between their glass transition temperature at *T*<sup>g</sup> = 276-277 K and their nematic-to-isotropic phase transition temperature at *T*N-I = 331-332 K (Δ*H*N-I = 1.3-1.7 J·g-1), respectively. Completely transparent monodomain nematic samples were obtained for both elastomers exhibiting orientational order parameters of 0.75 and 0.76 for **E***o***NO2** and **E***p***NO2**, respectively.

The determination of the thermal *cis*-to-*trans* isomerisation rate of both azo-dyes acting as covalently-bonded co-monomers in the nematic LSCE by means of opto-mechanical experiments yielded relaxation times for *cis-***AZO***p***NO2** and *cis*-**AZOoNO2** of 4 and 3 seconds, respectively, at 298 K. The impressive acceleration of the kinetics of the thermal *cis*to-*trans* isomerisation of both nitro-substituted azoderivatives in the LSCE is highly remarkable. The process was found to be more than 103 times faster in the LSCE, in which the azo-dye is chemically bonded to the polymer backbone, than that observed in isotropic solvents (between 16 and 92 min.), and about 102 times quicker than the one shown as a guest chromophore doped in a low molar mass nematic liquid crystal (6-7 min.). This acceleration is associated to the covalent attachment of the push-pull azoderivative to a nematic polymer backbone which possesses an anisotropic chain conformation.56

The great acceleration exhibited by these nitro-substituted azo-dyes when they are chemically bonded into nematic LSCEs has been already exploited for the preparation of fast photo-active artificial muscle-like actuators. The maximum opto-mechanical response generated by the elastomer **E***p***NO2** was of Δ*σ*max = 1.4 kPa, while a lower value of 0.7 kPa was obtained for the LSCE **E***o***NO2** (**Figure 19**). The fast thermal relaxation of these systems is the responsible of the low internal stress generated in both elastomers due to the low presence of *cis* isomer in the photo-equilibrium state. The slightly faster thermal isomerisation for **E***o***NO2** accounts for the lower mechanical response produced by this network. Both LSCEs exhibit a reversible behaviour in the time scale of seconds at room temperature, that is, relaxation times of 2-4 s for the whole irradiation process with blue light and of 3-4 s for the *cis*-to-*trans* thermal relaxation in the dark for **E***o***NO2** and **E***p***NO2**, respectively.

**Figure 19.** Opto-mechanical experiments for both nematic elastomers **E***p***NO2** (left) and **E***o***NO2** (right) at 298 K (irradiation of the networks was performed with blue light, *λ*irrad = 450 nm).

### **6.2. Artificial muscle-like actuators based on azophenol derivatives exhibiting azo-hydrazone taumeric equilibrium**

80 Advanced Elastomers – Technology, Properties and Applications

in the dark for **E***o***NO2** and **E***p***NO2**, respectively.

**AZO***o***NO2**, respectively, 5 % mol) and the isotropic cross-linking agent **V1** (10 % mol) as it is shown in **Figure 18**. Both elastomers, **E***p***NO2** and **E***o***NO2**, which differ only in the placement of the nitro- electron-withdrawing group, showed a broad enantiotropic nematic phase between their glass transition temperature at *T*<sup>g</sup> = 276-277 K and their nematic-to-isotropic phase transition temperature at *T*N-I = 331-332 K (Δ*H*N-I = 1.3-1.7 J·g-1), respectively. Completely transparent monodomain nematic samples were obtained for both elastomers exhibiting

The determination of the thermal *cis*-to-*trans* isomerisation rate of both azo-dyes acting as covalently-bonded co-monomers in the nematic LSCE by means of opto-mechanical experiments yielded relaxation times for *cis-***AZO***p***NO2** and *cis*-**AZOoNO2** of 4 and 3 seconds, respectively, at 298 K. The impressive acceleration of the kinetics of the thermal *cis*to-*trans* isomerisation of both nitro-substituted azoderivatives in the LSCE is highly remarkable. The process was found to be more than 103 times faster in the LSCE, in which the azo-dye is chemically bonded to the polymer backbone, than that observed in isotropic solvents (between 16 and 92 min.), and about 102 times quicker than the one shown as a guest chromophore doped in a low molar mass nematic liquid crystal (6-7 min.). This acceleration is associated to the covalent attachment of the push-pull azoderivative to a

orientational order parameters of 0.75 and 0.76 for **E***o***NO2** and **E***p***NO2**, respectively.

nematic polymer backbone which possesses an anisotropic chain conformation.56

The great acceleration exhibited by these nitro-substituted azo-dyes when they are chemically bonded into nematic LSCEs has been already exploited for the preparation of fast photo-active artificial muscle-like actuators. The maximum opto-mechanical response generated by the elastomer **E***p***NO2** was of Δ*σ*max = 1.4 kPa, while a lower value of 0.7 kPa was obtained for the LSCE **E***o***NO2** (**Figure 19**). The fast thermal relaxation of these systems is the responsible of the low internal stress generated in both elastomers due to the low presence of *cis* isomer in the photo-equilibrium state. The slightly faster thermal isomerisation for **E***o***NO2** accounts for the lower mechanical response produced by this network. Both LSCEs exhibit a reversible behaviour in the time scale of seconds at room temperature, that is, relaxation times of 2-4 s for the whole irradiation process with blue light and of 3-4 s for the *cis*-to-*trans* thermal relaxation

**Figure 19.** Opto-mechanical experiments for both nematic elastomers **E***p***NO2** (left) and **E***o***NO2** (right) at

298 K (irradiation of the networks was performed with blue light, *λ*irrad = 450 nm).

A second strategy to decrease the response time of the artificial muscle-like actuators implies the use of azophenols as photo-active moieties. Azophenols are promising chromophores for designing fast-responding artificial muscle-like actuators since they are endowed with a rapid thermal isomerisation process at room temperature, with relaxation times ranging from 6 ms to 300 ms in polar protic solvents depending on the position of the phenol groups.57 But, the main drawback is that hydroxyazobenzenes show a fast thermal isomerisation rate only when they are dissolved in polar protic solvents.

A successful strategy to transfer this fast thermal isomerisation to solid elastomeric materials consists in the preparation of a co-elastomer which contains a nematic mesogenic monomer (**M4OMe**) and the photo-active hydroxyazoderivative (**AZO-OH**). The liquid-crystalline coelastomer obtained contains only a small proportion of the azo moiety (5% mol) in order to not overly disrupt the nematic order of the elastomer. However, in that system, the azo-dye concentration is high enough for hydrogen bonding being established between the hydroxyazobenzene monomers without niether losing the liquid-crystalline properties of the nematic system nor diminishing the temperature range where the nematic phase exists. As a result of this interaction, the resulting LSCE exhibits an isomerisation rate as fast as that of the azo-monomer dissolved in a protic isotropic solvent.58

Two different liquid single crystal elastomers will be presented, **EAZO-Ac** and **EAZO-OH**, in order to describe properly the influence of the phenolic group in the opto-mechanical properties of the final photo-actuator (**Figure 20**). The uniaxially oriented acetylated elastomer **EAZO-Ac** was prepared following the well-known spin-casting technique commented above. Elastomer **EAZO-OH** was prepared from **EAZO-Ac** in order to assure an identical composition in both systems. For preparing elastomer **EAZO-OH**, the acetylated elastomer **EAZO-Ac** is swollen in a mixture of CH2Cl2 and MeOH (2:1 v/v) with one drop of acetyl chloride at room temperature for 2 days. Acetyl chloride enters the network by diffusion producing a mild and chemoselective cleavage of the ester group of the acetylated side-chain group, **AZO-Ac**, without hydrolysing the other ester group placed in the mesogen.

The characterization of both LSCEs was carried out by using the standard techniques (see Section 3 of this chapter). DSC experiments showed that both elastomers, **EAZO-Ac** and **EAZO-OH**, exhibited a broad enantiotropic nematic phase between their glass transition temperature at *T*g = 277-278 K (Δ*C*p = 0.3 J·g-1·K-1) and their nematic-to-isotropic phase transition temperature at *T*N-I = 335-336 K (Δ*H*N-I = 1.6-1.8 J·g-1), respectively. X-Ray scattering experiments clearly indicate that no change in the degree of order of the LSCE occurred during the cleavage reaction since very close orientational order parameters were obtained for both **EAZO-Ac** and **EAZO-OH** (0.71 and 0.69, respectively). Interestingly, the swelling ratio, *q*, was determined to be 5.0 and 3.3 for **EAZO-Ac** and **EAZO-OH**, respectively. The lower *q* value found for **EAZO-OH** evidenced that hydrogen bonding is established between the phenol groups of the azodyes in the LSCE acting as additional cross-linking units.

**Figure 20.** Chemical composition of **EAZO-Ac** and **EAZO-OH** and synthetic concept for the photoactive nematic LSCE **EAZO-OH**.

The analysis of the mechanical response generated by both elastomers as well as their relaxation time by means of opto-mechanical experiments evidenced that both LSCEs, **EAZO-Ac** and **EAZO-OH**, can act as light-controlled actuators when they are irradiated with UV light (*λ*irrad = 380 nm). **Figure 21** shows the evolution of the internal stress generated inside both LSCEs, **EAZO-Ac** and **EAZO-OH**, with the time at 323 and 298 K, respectively. **EAZO-Ac** produced a maximum opto-mechanical response of ∆*σ*max = 15 kPa at 323 K (**Figure 21a**). The opto-mechanical response of **EAZO-Ac** was enhanced to *ca.* ∆*σ*max = 30 kPa on dropping the temperature down to at 313 K (**Figure 22**). When the same experiment is carried out at 298 K for the elastomer **EAZO-OH**, a lower maximum opto-mechanical response of ∆*σ*max = 0.6 kPa is obtained (**Figure 21b**). Once more, similar results than those aforementioned are obtained. On comparing the non push-pull systems shown in Section 5 with those containing push-pull azoderivatives, which have been presented in Section 6.1, it is clearly seen that the mechanical efficiency, expressed in terms of ∆*σ*max, is greater in the former (up to 60 kPa at 298 K), that is, for those azo-dyes that undergo their thermal *cis*-to*trans* isomerisation through the inversional mechanism. Otherwise, those azocompounds that isomerises by means of the rotational pathway, those are, push-pull azo-dyes and azophenols, show lower mechanical efficiencies around 0.6 - 1.4 kPa at room temperature.

82 Advanced Elastomers – Technology, Properties and Applications

**Figure 20.** Chemical composition of **EAZO-Ac** and **EAZO-OH** and synthetic concept for the photo-

The analysis of the mechanical response generated by both elastomers as well as their relaxation time by means of opto-mechanical experiments evidenced that both LSCEs, **EAZO-Ac** and **EAZO-OH**, can act as light-controlled actuators when they are irradiated with UV light (*λ*irrad = 380 nm). **Figure 21** shows the evolution of the internal stress generated inside both LSCEs, **EAZO-Ac** and **EAZO-OH**, with the time at 323 and 298 K, respectively. **EAZO-Ac** produced a maximum opto-mechanical response of ∆*σ*max = 15 kPa at 323 K (**Figure 21a**). The opto-mechanical response of **EAZO-Ac** was enhanced to *ca.* ∆*σ*max = 30 kPa on dropping the temperature down to at 313 K (**Figure 22**). When the same experiment is carried out at 298 K for the elastomer **EAZO-OH**, a lower maximum opto-mechanical response of ∆*σ*max = 0.6 kPa is obtained (**Figure 21b**). Once more, similar results than those aforementioned are obtained. On comparing the non push-pull systems shown in Section 5

active nematic LSCE **EAZO-OH**.

Besides, the nematic liquid-crystalline elastomer **EAZO-Ac**, which contains the azochromophore as a side-chain monomer, is less effective (∆*σ*max = 15 kPa at 323 K) than those elastomers **EAZOX**, where the azo-dye is introduced as a photo-active cross-linker in the elastomeric network (∆*σ*max = 30 kPa at 323 K).

**Figure 21.** (a) Opto-mechanical experiment for the nematic LSCE **EAZO-Ac** at 323 K.(b) Optomechanical experiment for the nematic LSCE **EAZO-OH** at 298 K (*λ*irrad = 380 nm).

**EAZO-Ac** exhibits a relaxation time for its thermal back reaction of *τ*th = 21.3 h at room temperature, which was determined by extrapollation of the corresponding Eyring's plot (**Figure 23**). Opto-mechanical experiments also reveal a completely different kinetic behaviour between both photo-actuators. While **EAZO-Ac** presents a relaxation time of almost 1 day, **EAZO-OH** gives a value of only 1 s at room temperature. The fast thermal relaxation exhibited by **EAZO-OH** evidences clearly the establishment of hydrogen bonding between the different phenol groups due to their spatial proximity within the elastomeric network. This interaction indicates that the azo-dye isomerisation takes place through the rotational mechanism as it has been described for other azophenol derivatives.57

**Figure 22.** Opto-mechanical experiment for the nematic LSCE **EAZO-Ac** at 313 K (*λ*irrad = 380 nm).

**Figure 23.** Eyring's plot for the thermal *cis*-to-*trans* isomerisation of **EAZO-Ac**.

The low mechanical response registered for the free-phenol-containing elastomer **EAZO-OH** in comparison with that of **EAZO-Ac** is associated to the low proportion of *cis* isomer present when **EAZO-OH** reaches its photo-equilibrium state at 298 K, as it has been commented above for the nitro-substituted azobenzene-containing elastomers **E***p***NO2** and **E***o***NO2** (see above). Similarly, the slower thermal relaxation of **EAZO-Ac** at 313 K than at 323 K makes the photo-equilibrium state richer in *cis* isomer in the former. This fact accounts for the higher opto-mechanical response exhibited by **EAZO-Ac** at 313 K.

### **7. Conclusion**

84 Advanced Elastomers – Technology, Properties and Applications

derivatives.57

(**Figure 23**). Opto-mechanical experiments also reveal a completely different kinetic behaviour between both photo-actuators. While **EAZO-Ac** presents a relaxation time of almost 1 day, **EAZO-OH** gives a value of only 1 s at room temperature. The fast thermal relaxation exhibited by **EAZO-OH** evidences clearly the establishment of hydrogen bonding between the different phenol groups due to their spatial proximity within the elastomeric network. This interaction indicates that the azo-dye isomerisation takes place through the rotational mechanism as it has been described for other azophenol

**Figure 22.** Opto-mechanical experiment for the nematic LSCE **EAZO-Ac** at 313 K (*λ*irrad = 380 nm).

**Figure 23.** Eyring's plot for the thermal *cis*-to-*trans* isomerisation of **EAZO-Ac**.

Photo-active liquid single crystal elastomers (LSCEs) are valuable materials for artificial muscle-like applications since their macroscopic dimensions can be easily changed by applying light, an environmentally-friendy energy which can be wireless applied to the material. Those actuators based on the azobenzene chromophore are the most commonly used since it presents a clean and totally reversible isomerisation process.

Two different key parameters are crucial for characterising properly the artificial musclelike actuation of such materials, those are, the maximum mechanical response that they are able to generate by irradiation with light of the appropriate wavelength as well as the time required to produce it and that to recover the initial state further.

The use of non-push-pull azoderivatives, which thermally-isomerise through the inversional mechanism, both as light-sensitive cross-linkers and as side-chain pendant groups in polysiloxane nematic LSCEs produces high mechanically-efficient artificial muscle-like actuators. However, those actuators that bear the photo-active azo-dye as a cross-linker exhibit the highest mechanical efficiency. In such materials, the opto-mechanical response produced by the artificial muscle increases greatly by using larger flexible spacers on the photo-active cross-linker. For 4,4'-dialkoxysubstituted azobenzenes, there is a maximum spacer length, *nth* ~30, above which the opto-mechanical response of the network reaches a constant value and it is not enhanced anymore. Nevertheless, the relaxation time of the photo-actuator remains unaltered on varying the cross-linker spacer length. Hence, such materials can be used as bistable artificial muscles.

On the other hand, the introduction of both alkoxynitro-substituted azobenzenes and azophenols, which isomerise by means of the rotational pathway, yield fast-responding artificial muscle-like actuators. Indeed these two different types of fast-isomerising azoderivatives have been successfully used as covalently-bonded side-chain chromophores in nematic LSCEs. Alkoxynitro-substituted azobenzenes accelerate up to 103 times their thermal *cis*-to-*trans* isomerisation kinetics in isotropic media when they are attached in LSCEs. This rapid thermal isomerisation allows their use as chromophores to get stable and fast artificial muscle-like actuators working in time scale of 1-3 seconds.

Besides, the thermal *cis*-to-*trans* isomerisation process of azophenols depends strongly on the proticity of the environment where they are placed. In this way, their thermal isomerisation rate is accelerated up to 104 times if ethanol is used instead of toluene. The fast

isomerisation rate exhibited by these azo-dyes in ethanol has been successfully transferred to polymeric and elastomeric materials, where no solvent is present, due to the hydrogen bonding established between the side-chain azophenol monomers. Similar fast and stable artificial muscle-like actuators than those based on alkoxy-nitro-substituted azobenzenes have been described.

### **Author details**

Jaume Garcia-Amorós and Dolores Velasco *Grup de Materials Orgànics, Institut de Nanociència i Nanotecnologia (IN2UB), Departament de Química Orgànica, Universitat de Barcelona, Barcelona, Spain* 

### **Acknowledgement**

Financial support from the European project: "*Functional Liquid-Crystalline Elastomers*" (FULCE-HPRN-CT-2002-00169) and from the *Ministerio de Ciencia e Innovación* (CTQ-2009- 13797) is gratefully acknowledged. The authors thank Prof. Dr. Heino Finkelmann for its continuous support and helpful discussions.

### **8. References**


Jaume Garcia-Amorós and Dolores Velasco

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**Author details** 

**Acknowledgement** 

**8. References** 

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isomerisation rate exhibited by these azo-dyes in ethanol has been successfully transferred to polymeric and elastomeric materials, where no solvent is present, due to the hydrogen bonding established between the side-chain azophenol monomers. Similar fast and stable artificial muscle-like actuators than those based on alkoxy-nitro-substituted azobenzenes

Financial support from the European project: "*Functional Liquid-Crystalline Elastomers*" (FULCE-HPRN-CT-2002-00169) and from the *Ministerio de Ciencia e Innovación* (CTQ-2009- 13797) is gratefully acknowledged. The authors thank Prof. Dr. Heino Finkelmann for its

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*Grup de Materials Orgànics, Institut de Nanociència i Nanotecnologia (IN2UB), Departament de Química Orgànica, Universitat de Barcelona, Barcelona, Spain* 

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**Section 3** 

**Elastomer Nanocomposites** 

88 Advanced Elastomers – Technology, Properties and Applications

[54] C. D. Eisenbach, *Polymer*, 1980, 21, 1175.

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[57] J. Garcia-Amorós, A. Sánchez-Ferrer, W. A. Massad, S. Nonell and D. Velasco, *Phys.* 

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### **Chapter 4**

## **Rubber Clay Nanocomposites**

### Maurizio Galimberti

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/51410

### **1. Introduction**

RCN Nomenclature and classification of clays are first summarized, highlighting the most important clay features that affect their behaviour as fillers for rubbers. The modification of clays with organophilic compensating cations, to promote their compatibilization with the polymer matrix, is then presented. The processing methods for the preparation of RCN is then discussed and the main aspects of RCN are reviewed, namely: rheology, vulcanization, barrier and mechanical properties. Finally, commercial applications of RCN are presented.

A large scientific and patent literature is available on RCN and it was taken into consideration for the preparation of this Chapter. As the RCN applications are treated as well, press releases and news available on web sites were considered and commercial products were examined. Patents, when cited, are regarded just as publications, without considering if they were already granted or if, viceversa, they are still patent applications. Literature reporting on polymer clay nanocomposites (PCN) is available [1-13] and it has to be considered, to fully understand structure and properties of RCN. Reviews are also available, dealing specifically with rubbers [14-21].

### **2. Clays**

#### **2.1. Nomenclature and classification**

The Joint Nomenclature Committee of the AIPEA (Association Internationale pour l'Etude des Argiles) and the CMS (Clay Minerals Society) say that a clay is a naturally occurring material composed primarily of fine-grained minerals, which is generally plastic at appropriate water contents and hardens when dried or fired. Hence, the definition of clay refers essentially to the macroscopic clay properties. Precisely determined crystallographic structures should be named as clay minerals [22].

© 2012 Galimberti, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

Clays are layered: this is their most important feature. They are inorganic compounds made by stacked layers, whose atoms are joined together by iono-covalent bonds, that are bound to each other in the perpendicular direction through weaker forces. This implies that layers can be separated from each other by applying a minor amount of energy, whereas remarkable energy is required to break the layers.

Clays are a sub-family belonging to the larger family of layered oxides (or oxyhydroxides). They are not only silicates, as some of them do not contain any silicon atom, though those applied for the preparation of polymer nanocomposites are indeed essentially silicates.

Clays can be classified according to the electrical charge of the layer, as summarized in Table 1. Clays have either (i) neutral layers or (ii) negatively charged layers or (iii) positively charged layers. In the last two cases, the layer charge is exactly compensated by an equal amount of opposite charges located in the interlayer space. Clays with negatively charged layers or positively charged layers are respectively called cationic clays and anionic clays.


**Table 1.** Classification of clays as a function of the electrical charge of the layer

#### **2.2. Organization and structure of clays**

Clays give rise to a multiscale organization. Examining the upper level of said organization, clays particles of micrometric size form millimetric-size agglomerates. In a polymer matrix, the dispersion of these agglomerates is one of most important feature of the polymer nanocomposite. Figure 1 reports the structure of a montmorillonite (Mt) (Figure 1a), the most applied clay for the preparation of polymer nanocomposites, thanks to its large availability, low cost and high surface area. Moreover, purified Mt, with less than 1% of crystalline silica, is considered safe and is handled as a standard powder, as its platy nanoparticles, with only one dimension at the nanoscale, appear to have little chance to cross biological barriers [23].

remarkable energy is required to break the layers.

Type of layers Type of clay Main features

kaolinite

**2.2. Organization and structure of clays** 

cross biological barriers [23].

pyrophyllite, talc,

phyllosilicates: e.g. bentonites (main component: montmorillonite)

hydrotalcite (HT). layered double hydroxides (HT-like family)

**Table 1.** Classification of clays as a function of the electrical charge of the layer

essentially silicates.

negatively charged

positively charged

neutral layers

layers

layers

Clays are layered: this is their most important feature. They are inorganic compounds made by stacked layers, whose atoms are joined together by iono-covalent bonds, that are bound to each other in the perpendicular direction through weaker forces. This implies that layers can be separated from each other by applying a minor amount of energy, whereas

Clays are a sub-family belonging to the larger family of layered oxides (or oxyhydroxides). They are not only silicates, as some of them do not contain any silicon atom, though those applied for the preparation of polymer nanocomposites are indeed

Clays can be classified according to the electrical charge of the layer, as summarized in Table 1. Clays have either (i) neutral layers or (ii) negatively charged layers or (iii) positively charged layers. In the last two cases, the layer charge is exactly compensated by an equal amount of opposite charges located in the interlayer space. Clays with negatively charged layers or positively charged layers are respectively called cationic clays and anionic clays.

*neutral clays* 

*cationic clays*

compensated

*anionic clays* 

Clays give rise to a multiscale organization. Examining the upper level of said organization, clays particles of micrometric size form millimetric-size agglomerates. In a polymer matrix, the dispersion of these agglomerates is one of most important feature of the polymer nanocomposite. Figure 1 reports the structure of a montmorillonite (Mt) (Figure 1a), the most applied clay for the preparation of polymer nanocomposites, thanks to its large availability, low cost and high surface area. Moreover, purified Mt, with less than 1% of crystalline silica, is considered safe and is handled as a standard powder, as its platy nanoparticles, with only one dimension at the nanoscale, appear to have little chance to

in the interlayer space.

in the interlayer space

layers joined together by van der Waals interactions and/or hydrogen bonds

the negative layer charge is exactly

by compensating cations are located

by compensating anions located

the positive layer charge is exactly compensated

**Figure 1.** Structure of montmorillonite (Figure 1a) and a scheme (Figure 1b) to highlight the significant difference in length with respect to thickness of the layers

By examining a clay such as Mt at the lower level of its organization, it can be seen that a single clay layer is characterized by lateral dimensions from 100 to 1000 nm and by a thickness of about 1 nanometer. Mt is a TOT type clay mineral, with two tetrahedral (T) sheets linked to both sides to a central octahedral (O) sheet. These sandwiches are held together by weak interforces and alkaline and alkaline-earth cations are located in the interlayer space. Each silicate layer is terminated on its faces by oxygen atoms and on its periphery by oxygen atoms and hydroxide groups. The hydrophilic nature of Mt implies that a compatibilizer has to be used in order to disperse Mt in a lipophilic matrix, such as the one of most diffused rubbers. To modify a pristine Mt with a compatibilizer, the cations located in the interlayer space are exchanged with organophilic cations, preferentially with long chain quaternary alkyl ammonium cations, an organoclays (OC) are formed [24]. OC are discussed in the next paragraph. The high cationic exchange capacity (CEC) of Mt further pushes its use in polymer nanocomposites.

The term nanocomposites was many times used in previous lines. In fact, nanocomposites are defined as composite materials characterized by the presence of dispersed particles whose size is in the "nanoscale", defined as "having one or more dimensions of the order of 100 nm or less" [25]. Mt can be considered as a nanomaterial and, as it will be discussed later on, as a nanofiller and its polymer composites are thus nanocomposites.

Figure 1 shows as well a scheme (Figure 1 b) that highlights the peculiar shape of the layers, that have a significant difference in length with respect to thickness. Considering Mt as a reinforcing filler, this means that Mt has a high aspect ratio. As it will be discussed in Paragraph 10, the mechanical reinforcement of a polymer matrix strongly depends on the aspect ratio of the filler, i.e. on the ratio between its longest and shortest dimensions. This ratio is often indicated as *shape factor f*. The higher is the shape factor f and the higher are the values obtained for the dynamic-mechanical moduli. It is worth commenting here that to have a high shape factor, the Mt layers have to be brought apart, that means Mt has to be exfoliated. The highest aspect ratio for a clay would be given by the ratio between the longest lateral dimension and the thickness of a single layer.

#### **2.3. X-ray diffraction of clays**

Clays are crystalline materials and X-Ray Diffraction (XRD) analysis is a technique largely used to assess clay features. In the Bragg model of X-rays diffraction, that indicates the plane spacing with d index and the plane orientation with three Miller indexes (h, k, ℓ), a constructive interference is obtained for X-rays scattered from adjacent planes when the angle θ, between the plane and both the incident and reflected beams, and the plane spacing is related with the X-ray wavelength λ. through the equation known as the Bragg law [26]:

$$\text{m}\,\text{d}\,\text{sin}\,\theta = \text{n}\,\text{ }\,\text{\textasciicize}\,\text{\textasciicize}\,\text{\textasciicize}\,\text{\textasciicize}\,\text{\textasciicize}\,\text{\textasciicize}\,\text{\textasciticize}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spaceleftharpoons}\,\text{\spacelefth{\space}}\,\text{\spacelefth{\space}}\,\text{\spacelefth{\space}}\,\text{\spacelefth{\space}}\,\text{\spacelefth{\space}}\,\text{\spacelefth{\space}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{\space}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{}}\,\text{\space"{\text{\spacesum}}}\,\text{\space"{}}\,\text{\space"{}}\,\text{\space"{}}\,\text{\space"{}}$$

From the values of θ angles detected in the XRD pattern, it is possible to determine clay features such as: dhkℓ interlayer distance, Dhkℓ correlation length in a crystallographic direction, for example in the plane of the layers and in the direction orthogonal to the layers, and hence the number of stacked layers giving rise to a crystalline unit.

Figure 2 shows the XRD patterns of two pristine Mt available in the market: Cloisite® Na from Southern Clay and Dellite® HPS from Laviosa Chimica Mineraria.

**Figure 2.** X-Ray diffraction patterns in the 2theta range 0-80° for pristine Mt samples: Cloisite® Na and Dellite® HPS

The d001 interlayer distance was determined to be about 1.20 and 1.46 nm for Cloisite® Na and Dellite® HPS, respectively. As mentioned above, from XRD data it is also possible to calculate the Dhk<sup>ℓ</sup> correlation length of the crystalline domain in the direction orthogonal to the structural layer, through the Scherrer equation:

$$\mathbf{D}\_{\mathrm{hk}\ell} = \begin{array}{c} 0.9 \ \mathcal{X} \ \text{ } \end{array} \Big/ \ \mathcal{J}\_{\mathrm{hk}\ell} \cos \theta\_{\mathrm{hk}\ell} \tag{2}$$

where: λ is the wavelength of the irradiating beam (1.5419 Å, CuK), βhkℓ is the width at half height, and θhkℓ is the diffraction angle [27]. With reference to the (001) reflection and by introducing the correction factor, to be used in case βhkℓ is lower than 1°, a D001 value of about 9 nm and about 4 nm was calculated for Cloisite® Na and Dellite® HPS, respectively. Taking into account that the d001 interlayer distances reported above, a number of about 8 and about 4 stacked layers was calculated for Cloisite® Na and Dellite® HPS, respectively. This elaboration demonstrates that the available pristine Mt are characterized by a pretty typical interlayer distance and by a relatively low number of stacked layers.

### **3. Organoclays**

94 Advanced Elastomers – Technology, Properties and Applications

Clays are crystalline materials and X-Ray Diffraction (XRD) analysis is a technique largely used to assess clay features. In the Bragg model of X-rays diffraction, that indicates the plane spacing with d index and the plane orientation with three Miller indexes (h, k, ℓ), a constructive interference is obtained for X-rays scattered from adjacent planes when the angle θ, between the plane and both the incident and reflected beams, and the plane spacing is related with the X-ray wavelength λ. through the equation known as the Bragg law [26]:

> 2 d sin n

and hence the number of stacked layers giving rise to a crystalline unit.

from Southern Clay and Dellite® HPS from Laviosa Chimica Mineraria.

From the values of θ angles detected in the XRD pattern, it is possible to determine clay features such as: dhkℓ interlayer distance, Dhkℓ correlation length in a crystallographic direction, for example in the plane of the layers and in the direction orthogonal to the layers,

Figure 2 shows the XRD patterns of two pristine Mt available in the market: Cloisite® Na

**Figure 2.** X-Ray diffraction patterns in the 2theta range 0-80° for pristine Mt samples: Cloisite® Na and

The d001 interlayer distance was determined to be about 1.20 and 1.46 nm for Cloisite® Na and Dellite® HPS, respectively. As mentioned above, from XRD data it is also possible to calculate the Dhk<sup>ℓ</sup> correlation length of the crystalline domain in the direction orthogonal to

> D 0.9 / cos hk hk hk

where: λ is the wavelength of the irradiating beam (1.5419 Å, CuK), βhkℓ is the width at half height, and θhkℓ is the diffraction angle [27]. With reference to the (001) reflection and by introducing the correction factor, to be used in case βhkℓ is lower than 1°, a D001 value of about 9 nm and about 4 nm was calculated for Cloisite® Na and Dellite® HPS, respectively. Taking into account that the d001 interlayer distances reported above, a

 

(2)

 

(1)

**2.3. X-ray diffraction of clays** 

Dellite® HPS

the structural layer, through the Scherrer equation:

As mentioned in the previous paragraph, OC are prepared through the exchange reaction of a pristine clay with an organophilic ion. In the case of cationic clays, such as Mt, organophilic ammonium cations are mostly used for preparing OC for polymer nanocomposites. The intercalation of ammonium cations leads to an expansion of the interlayer distance. In the XRD spectrum, (00ℓ) indexes refer to a crystalline order in the direction perpendicular to the structural layers: the (001) reflection is due to regularly stacked layers and higher order reflections, e.g.. (002) and (003), originate from the regular arrangement of intercalants. An expansion of d interlayer spacing corresponds to a shift of (001) reflection towards lower 2 θ angle values.

Figure 3 shows two types of ammonium ions largely used as Mt compensating cations and the expansion of interlayer distance that occurs as a consequence of their intercalation. The ammonium cations are: dimethyl-talloyl-dihydroxyethyl and dimethylditalloyl and the talloyl (T) group can be optionally hydrogenated (HT).

**Figure 3.** Expansion of Mt interlayer distance as a consequence of ammonium cations intercalation. 2(H)T and (H)Tdihydroxyethyl as the ammonium cations

It is worth underlying that the same ammonium cation can give rise to different interlayer distances, as a function of its weight percent and of its bond with the inorganic layer. In fact the ammonium can be ionically bonded or can be only absorbed on the layer surface. Table 2 reports the values of interlayer distances detected for commercially available OC with 2HT as the compensating ammonium cation.


**Table 2.** Organoclays with 2HT as the compensating ammonium cation

## **4. Clays and organoclays used for RCN**

Literature available on RCN essentially refers to cationic clays. Those used for RCN preparation were: Na-Mt, Na-bentonite, Na-fluorohectorite, rectorite, vermiculite and a fibrillar silicate such as attapulgite. Clays of the bentonite family and, among them, Mt in particular, were the most applied ones. This chapter refers thus to data based on cationic clays. As discussed in paragraphs 2 and 3, a lipophilic compensating cation promotes the compatibilization of the cationic clay with the polymer matrix. Those used for the preparation of RCN are listed as follows, indicating for the ammonium cation the substituents of the nitrogen atom: (i) cations from primary alkenylamines, (ii) ammonium cations with three methyls and one long chain alkenyl, (iii) ammonium cations with two methyls, an hydrogenated tallow and a benzyl group, (iv) ammonium cations with two methyls and two ethylhexyl or short chain alkenyl groups, (iv) ammonium cations with a methyl, a tallow and polar groups such as 2-hydroxyethyl, (v) ammonium cations with two methyls and two (hydrogenated) tallow groups.

### **5. Processing methods for the preparation of RCN**

In this paragraph, distribution and dispersion of clays in a rubber matrix are discussed. They will be shown as depending on the chemical nature of the clay (pristine or organically modified) and on the selection of the appropriate processing technology.

Different methods have been developed for the preparation of RCN. A classification can be attempted, based on the type of clay used for the preparation of the nanocomposite: pristine or organically modified, as it is summarized in Table 3.

A pristine clay can be blended with rubber latexes (NR and E-SBR as the rubbers) through what is known as the emulsion blending. It can also be blended with the rubber in the melt state, performing the exchange reaction with an ammonium salt and using the rubber as the reaction medium.

For an organically modified clay, melt and solution blending can be applied. In the melt blending, the OC is mixed directly with the rubber in the melt state. In the solution blending, the rubber is first dissolved in a good solvent, adding the OC swollen in the same solvent.

Methods that use pristine clay and the melt blending of an OC are suitable for an industrial development, whereas the solution blending can be adopted in a laboratory. In this paragraph, the impact of mixing methods on clay distribution and dispersion is examined. An even clay dispersion is a crucial property for a PCN, as it allows to best exploit the properties of the layered filler.


**Table 3.** Processing methods of clays and organoclays with rubber

#### **5.1. Emulsion blending from a pristine clay**

96 Advanced Elastomers – Technology, Properties and Applications

Dellite® 72T 36 – 38 3.42 001 2.59

Closite® 20A 38 3.57 001 2.48

Closite® 15A 43 2.98 001 2.97

**Table 2.** Organoclays with 2HT as the compensating ammonium cation

**4. Clays and organoclays used for RCN** 

methyls and two (hydrogenated) tallow groups.

**5. Processing methods for the preparation of RCN** 

or organically modified, as it is summarized in Table 3.

reaction medium.

modified) and on the selection of the appropriate processing technology.

organic 2θ (deg) Hkl d (nm) d001 (nm)

7.12 002 1.24

7.1 002 1.25

7.15 00? 1.24

2.50 001 3.53

4.66 002 1.89 7.23 003 1.22 9.8 004 0.90

Literature available on RCN essentially refers to cationic clays. Those used for RCN preparation were: Na-Mt, Na-bentonite, Na-fluorohectorite, rectorite, vermiculite and a fibrillar silicate such as attapulgite. Clays of the bentonite family and, among them, Mt in particular, were the most applied ones. This chapter refers thus to data based on cationic clays. As discussed in paragraphs 2 and 3, a lipophilic compensating cation promotes the compatibilization of the cationic clay with the polymer matrix. Those used for the preparation of RCN are listed as follows, indicating for the ammonium cation the substituents of the nitrogen atom: (i) cations from primary alkenylamines, (ii) ammonium cations with three methyls and one long chain alkenyl, (iii) ammonium cations with two methyls, an hydrogenated tallow and a benzyl group, (iv) ammonium cations with two methyls and two ethylhexyl or short chain alkenyl groups, (iv) ammonium cations with a methyl, a tallow and polar groups such as 2-hydroxyethyl, (v) ammonium cations with two

In this paragraph, distribution and dispersion of clays in a rubber matrix are discussed. They will be shown as depending on the chemical nature of the clay (pristine or organically

Different methods have been developed for the preparation of RCN. A classification can be attempted, based on the type of clay used for the preparation of the nanocomposite: pristine

A pristine clay can be blended with rubber latexes (NR and E-SBR as the rubbers) through what is known as the emulsion blending. It can also be blended with the rubber in the melt state, performing the exchange reaction with an ammonium salt and using the rubber as the

~ 2.5

~ 2.5

~ 3.0

~ 3.6

Organoclay % wt

Dellite® 67G 45

Zhang first introduced the emulsion compounding: a mixture obtained from an aqueous clay suspension and a rubber latex was coagulated in an electrolite solution. This approach was applied for blending Mt in rubbers such as NR [28, 29], SBR [28-32], NBR [28, 29, 33], XNBR [29]. Moreover, bentonite was blended with SBR [34-37] and SVBR (V = vinylpiridine) [34, 35] and rectorite was blended in SBR [38, 39]. Different electrolytes were used: triethylenetetrammonium chloride (2% wt solution) [29, 38, 39], diluted sulphuric acid solution [28, 30-32, 36], 1% calcium chloride aqueous solution [32], dilute hydrochloric acid solution [34, 35, 37], dilute dichloroacetic acid solution [33]. A scheme for blending a pristine clay with a rubber latex, as reported by Zhang, is shown in Figure 4.

**Figure 4.** Emulsion blending of a pristine clay with a rubber latex

It was reported that the emulsion compounding led to nanocomposites, whereas conventional microcomposites were obtained when a pristine clay was melt blended with a rubber [28]. The Mt dispersion was commented to be excellent, up to 20 phr as concentration, and best results were obtained with diluted sulphuric acid.

Modifications of this procedure, leading to satisfactory silicate dispersion, were adopted by:

i. adding a slurry containing further ingredients prior to coagulation, with NBR as the rubber and a sodium salt of methylene-bis-naphthalene sulphonic acid as the electrolyte.


Bentonite and fluorohectorite (10 phr) dispersions in NR were observed to be to some extent worse when further ingredients were added prior to coagulation and coagulating agents were not used [42, 43].

When a prevulcanized NR latex was used, a minor amount of fluorohectorite layers and stacks with an increased distance between opposite layers were observed [44]. Fluorohectorite revealed a better swelling ability than bentonite and thus gave a more exfoliated structure [45].

### **5.2. Melt blending from clays and organoclays**

Melt blending was performed with the help of the typical processing technologies of the rubber industry, from internal mixers such as brabender® and banbury® to open mills to twin screw extruders.

The melt blending of a pristine Mt with a rubber leads to composites with undispersed agglomerates, as demonstrated for NR [46-49], IR [50], BR [51-53], E-SBR [48, 54], E-SBR/NBR blend [55], EPDM [56] as the rubbers. Analogous results were presented with fluorhectorite and attapulgite as the clays.

A satisfactory dispersion was achieved when the inorganic clays contained a lipophilic modifier [57]. With primary alkenylamines, exfoliated platelets and aggregates composed of different number of platelets were observed in the final composites, commenting that the C18 alkenyl group promoted a better dispersion with respect to the shorter ones.

An even dispersion of the organoclay in the rubber matrix was achieved in any the composites, though a fully exfoliated morphology was not achieved, when an ammonium cation with either three methyls and one long chain alkenyl substituent or two methyls and two long chain substituents was used. The polar hydroxyl group as ammonium cation substituent led to some layers agglomeration. These results were independent of the type of the rubber matrix.

### **5.3. Solution blending from clays and organoclays**

Microcomposites were prepared when a bentonite or a Mt were mixed with the help of an organic solvent with many different types of rubbers: NR [58], BR [55], SBR [55, 59-61], (H)NBR [55, 62], BIMS [63]. Big lumps of clay aggregates and agglomerates were formed. Viceversa, an even dispersion of organoclays was achieved, when a clay containing a lipophilic ammonium as compensating cation was used. As reported in paragraph 4, different types of ammonium cations were used, in many types of rubber. A rationalization is proposed in ref. [57]. To summarize, the mentioned even dispersion was achieved in NR, BR, SBR, NBR, BIMS and EPDM with a primary alkenyl amine as the clay modifier. In NR, IR, BR, SBR, (H)NBR when ammonium cations having either two or three methyl groups and longer chain substituents were adopted as clay compensating cations. Results appeared very similar to those obtained with melt blending.

### **5.4. Formation of organoclays in situ in the rubber matrix**

The preparation of the organoclay *in situ* in the rubber matrix was reported by Galimberti [57, 64-67], by adopting the melt blending approach, with all the processing technologies mentioned in paragraph 5.2. An even clay dispersion was achieved in the rubber matrix and the organoclays revealed a lower number of stacked layers. This approach was adopted for many types of rubbers: IR, NR, SBR, NBR, XIIR, EPDM.

### **6. Structure of RCN**

98 Advanced Elastomers – Technology, Properties and Applications

**5.2. Melt blending from clays and organoclays** 

fluorhectorite and attapulgite as the clays.

with NR as the rubber [40]

the rubber [41].

were not used [42, 43].

exfoliated structure [45].

twin screw extruders.

the rubber matrix.

ii. adding further ingredients by melt mixing and avoiding the use of coagulating agents,

iii. avoiding both coagulating agents and the addition of further ingredients, with NR as

Bentonite and fluorohectorite (10 phr) dispersions in NR were observed to be to some extent worse when further ingredients were added prior to coagulation and coagulating agents

When a prevulcanized NR latex was used, a minor amount of fluorohectorite layers and stacks with an increased distance between opposite layers were observed [44]. Fluorohectorite revealed a better swelling ability than bentonite and thus gave a more

Melt blending was performed with the help of the typical processing technologies of the rubber industry, from internal mixers such as brabender® and banbury® to open mills to

The melt blending of a pristine Mt with a rubber leads to composites with undispersed agglomerates, as demonstrated for NR [46-49], IR [50], BR [51-53], E-SBR [48, 54], E-SBR/NBR blend [55], EPDM [56] as the rubbers. Analogous results were presented with

A satisfactory dispersion was achieved when the inorganic clays contained a lipophilic modifier [57]. With primary alkenylamines, exfoliated platelets and aggregates composed of different number of platelets were observed in the final composites, commenting that the C18

An even dispersion of the organoclay in the rubber matrix was achieved in any the composites, though a fully exfoliated morphology was not achieved, when an ammonium cation with either three methyls and one long chain alkenyl substituent or two methyls and two long chain substituents was used. The polar hydroxyl group as ammonium cation substituent led to some layers agglomeration. These results were independent of the type of

Microcomposites were prepared when a bentonite or a Mt were mixed with the help of an organic solvent with many different types of rubbers: NR [58], BR [55], SBR [55, 59-61], (H)NBR [55, 62], BIMS [63]. Big lumps of clay aggregates and agglomerates were formed. Viceversa, an even dispersion of organoclays was achieved, when a clay containing a lipophilic ammonium as compensating cation was used. As reported in paragraph 4, different types of ammonium cations were used, in many types of rubber. A rationalization is proposed in ref. [57]. To summarize, the mentioned even dispersion was achieved in NR,

alkenyl group promoted a better dispersion with respect to the shorter ones.

**5.3. Solution blending from clays and organoclays** 

In this paragraph, the overall picture of RCN structure is presented, discussing in particular the interlayer distance of clays and the type of intercalants present in the interlayer space, of low molecular mass or polymeric.

This aspect is of the outmost importance for polymer clay nanocomposites. In fact, it is largely reported [1-22] that the intercalation of polymer chains in the interlayer space is at the origin of the remarkable properties of PCN. The intercalation of polymer chains is considered not only to favour the improvement of the polymer matrix but also to promote the clay exfoliation, leading to the ultimate clay dispersion. The overall picture of RCN, that is proposed for PCN is shown in Figure 5.

**Figure 5.** Overall picture of the structure of polymer (rubber) clay nanocomposites

The study of the interlayer distance is performed through XRD and TEM analysis [57]. In particular, as discussed in paragraph 3, the modification of the interlayer distance, determined from XRD patterns, is taken as the evidence of the occurring of modifications in the interlayer space. It is worth summarizing here the two mechanisms proposed in the literature to explain the occurring of clay exfoliation and the modification of the interlayer distance.

#### **6.1. Two mechanisms for clays intercalation and exfoliation**

The two mechanisms can be visualized as reported in Figure 5 and are summarized as follows.

**Figure 6.** Schemes to visualize the two mechanisms proposed for clay intercalation and exfoliation: first mechanism (a), second mechanism (b) (see text for the explanation)

*First mechanism*. Polymer chains are intercalated in the interlayer space, with a consequent expansion of the interlayer distance. In Figure 6a, the position of the (00ℓ) reflection progressively shift towards lower values of the 2θ angle, from black to red to blue curve, as a consequence of the progressive intercalation of the polymer chains. The intercalated polymer chains are not only responsible for the expansion of the interlayer distance but also cause the separation of clay layers, promoting the ultimate exfoliation. This first mechanism, originally proposed for polystyrene as the polymer, is widely accepted in the prior art.

*Second mechanism*. The intercalation of polymer chains is regarded as unlikely and only low molecular mass substances are considered to be present in the interlayer space, when crystalline OC are formed. In particular, it is commented that it is hard to suppose that highly crystalline structures (as shown by the presence of several (00ℓ) reflections in the XRD pattern), can be generated in the presence of a polymer chain. The variation of the interlayer distance is seen as a consequence of: (i) the different arrangements of the substituents of the compensating cations, that indeed can give rise to different d001 values (as shown in Table 2) and/or (ii) of intercalation of guests in the interlayer space and/or (iii) of molecules absorption on clay surfaces. The clay exfoliation occurs through a progressive peeling off of the clay stacks thanks to the shear mixing. In Figure 6b, the (00ℓ) reflection remains at the same 2θ value and its intensity decreases, passing from the black to the red to the black curve, as a consequence of the progressive exfoliation of the clay stacks. Authors presenting this mechanism propose the scheme reported in Figure 7, where the XRD pattern of nanocomposites in the low 2θ region (peaks are due to (00ℓ) reflections are shown as well the mechanism hypothesized for their formation.

**Figure 7.** Schematic presentation for the formation of an organoclay by blending Na-Mt with di(hydrogenated tallow)-dimethylammonium chloride (2HT), in the absence or in the presence of stearic acid (SA) (Scheme reproduced from ref. 66)

### **7. Rheology of RCN**

100 Advanced Elastomers – Technology, Properties and Applications

**6.1. Two mechanisms for clays intercalation and exfoliation** 

mechanism (a), second mechanism (b) (see text for the explanation)

3 4 5 6 7 8 9 10 11

2

distance.

follows.

Intensità (u.a.)

Intensity (a.u.)

The study of the interlayer distance is performed through XRD and TEM analysis [57]. In particular, as discussed in paragraph 3, the modification of the interlayer distance, determined from XRD patterns, is taken as the evidence of the occurring of modifications in the interlayer space. It is worth summarizing here the two mechanisms proposed in the literature to explain the occurring of clay exfoliation and the modification of the interlayer

The two mechanisms can be visualized as reported in Figure 5 and are summarized as

Intensità (u.a.)

Intensity (a.u.)

a b (a) (b)

**Figure 6.** Schemes to visualize the two mechanisms proposed for clay intercalation and exfoliation: first

3 4 5 6 7 8 9 10 11

2

*First mechanism*. Polymer chains are intercalated in the interlayer space, with a consequent expansion of the interlayer distance. In Figure 6a, the position of the (00ℓ) reflection progressively shift towards lower values of the 2θ angle, from black to red to blue curve, as a consequence of the progressive intercalation of the polymer chains. The intercalated polymer chains are not only responsible for the expansion of the interlayer distance but also cause the separation of clay layers, promoting the ultimate exfoliation. This first mechanism, originally proposed for polystyrene as the polymer, is widely accepted in the prior art.

*Second mechanism*. The intercalation of polymer chains is regarded as unlikely and only low molecular mass substances are considered to be present in the interlayer space, when crystalline OC are formed. In particular, it is commented that it is hard to suppose that highly crystalline structures (as shown by the presence of several (00ℓ) reflections in the XRD pattern), can be generated in the presence of a polymer chain. The variation of the interlayer distance is seen as a consequence of: (i) the different arrangements of the substituents of the compensating cations, that indeed can give rise to different d001 values (as shown in Table 2) and/or (ii) of intercalation of guests in the interlayer space and/or (iii) of molecules absorption on clay surfaces. The clay exfoliation occurs through a progressive peeling off of the clay stacks thanks to the shear mixing. In Figure 6b, the (00ℓ) reflection

Clays compatibilized and evenly dispersed in a polymer matrix tend to build networks at low concentration. Rheological measurements, performed on RCN based on various types of rubbers, revealed the pronounced rubber-clay interaction, when measurements were taken at zero shear. The storage modulus in the low frequency region was investigated as a function of clay content and the clay percolation threshold (as wt%) was found to be about 4 (OC was Mt/methyl tallow bis-2-hydroxyethyl ammonium cation), above 5 (OC was Mt/dimethyl-dialkylammonium halide (70% C18, 26% C16 and 4% C14)) and above 7.5 (OC was Mt/ dimethyl dehydrogenated tallow quaternary ammonium chloride) for IR, EPR and EVA, respectively [68-71]. The filler networking phenomenon was observed as well in

matrices based on IR, ENR [72], SBR [73] and EPR [69]. At zero shear, the viscosity of RCN is thus higher than the one of the neat elastomer.

However, organophilic clays were shown to reduce the steady shear viscosity of RCN. Most evident results were a pronounced shear-thinning behaviour, increasing with the clay content, a higher extent of extrudate, a lower swelling and a better surface smoothness, by increasing the shear rate. These findings are of great importance, as they indicate that OC have a positive effect on the processability of rubber compounds. Data were reported for RCN based on various rubbers, such as BR [60], SBR [60], NBR [60], BIMS [74] and fluoroelastomer [75]. A compensating cation with longer alkyl chains led to a reduction of Mooney viscosity, that was not observed with a short chain substituent [76].

This behaviour of OC, that is opposite to that of traditional fillers, is attributed to the orientation, occurring at high shear rates, of clay platelets along flow direction, and to the slippage of platelets on the chains, thanks to the organophilic clay substituents. Data to demonstrate that were provided with an OC (Mt/octadecyltrimethylammonium) in BR [77] and with another OC (Mt/dimethyl hydrogenated-tallow (2-ethylhexyl) quaternary ammonium methylsulfate) in poly(epichlorohydrine) [78].

The improvement of processability brought about by OC was found to be higher when the matrix was apolar rubber such as NBR [79].

### **8. Vulcanization of RCN**

OC promote a fast sulphur based crosslinking of unsaturated polymer chains. Moreover, they increase the delta torque value indicating a higher value of crosslinking density [80].

Many data were reported in the literature, supporting these conclusions, in rubbers such as NR (hexadecyltrimethyl ammonium, octadecyltrimethyl ammonium, tetraoctyl phosphonium, triphenyl vinylbenzyl phosphonium, octadecylamine chloridrate were used as intercalants) [81-86], SBR (Mt/octadecyltrimethylamonium) [87], (NBR (octylamine; dodecylamine and octadecylamine chloridrates were the intercalants) [76]. A reduction of the activation energy for sulphur based NR crosslinking was found by using a bentonite modified with octadecylamine chloridrate [83, 88]. OC could thus present a warning, as they could lead to premature scorching but, at the same time, they could present the chance of performing vulcanization reactions at lower temperatures.

To explain the behaviour of OC in sulphur based crosslinking, the formation of tertiary amines from the thermal degradation of ammonium cations and the enhanced mobility of sulphur accelerating anionic species were proposed [89].

### **9. Barrier properties of RCN**

The reduction of air permeability thanks to clays dispersed in a poly(isobutene) matrix led to the first commercial application of RCN, that will be discussed in Paragraph 11.2.

### **9.1. The tortuous path model**

102 Advanced Elastomers – Technology, Properties and Applications

thus higher than the one of the neat elastomer.

matrices based on IR, ENR [72], SBR [73] and EPR [69]. At zero shear, the viscosity of RCN is

However, organophilic clays were shown to reduce the steady shear viscosity of RCN. Most evident results were a pronounced shear-thinning behaviour, increasing with the clay content, a higher extent of extrudate, a lower swelling and a better surface smoothness, by increasing the shear rate. These findings are of great importance, as they indicate that OC have a positive effect on the processability of rubber compounds. Data were reported for RCN based on various rubbers, such as BR [60], SBR [60], NBR [60], BIMS [74] and fluoroelastomer [75]. A compensating cation with longer alkyl chains led to a reduction of

This behaviour of OC, that is opposite to that of traditional fillers, is attributed to the orientation, occurring at high shear rates, of clay platelets along flow direction, and to the slippage of platelets on the chains, thanks to the organophilic clay substituents. Data to demonstrate that were provided with an OC (Mt/octadecyltrimethylammonium) in BR [77] and with another OC (Mt/dimethyl hydrogenated-tallow (2-ethylhexyl) quaternary

The improvement of processability brought about by OC was found to be higher when the

OC promote a fast sulphur based crosslinking of unsaturated polymer chains. Moreover, they increase the delta torque value indicating a higher value of crosslinking density [80].

Many data were reported in the literature, supporting these conclusions, in rubbers such as NR (hexadecyltrimethyl ammonium, octadecyltrimethyl ammonium, tetraoctyl phosphonium, triphenyl vinylbenzyl phosphonium, octadecylamine chloridrate were used as intercalants) [81-86], SBR (Mt/octadecyltrimethylamonium) [87], (NBR (octylamine; dodecylamine and octadecylamine chloridrates were the intercalants) [76]. A reduction of the activation energy for sulphur based NR crosslinking was found by using a bentonite modified with octadecylamine chloridrate [83, 88]. OC could thus present a warning, as they could lead to premature scorching but, at the same time, they could present the chance of

To explain the behaviour of OC in sulphur based crosslinking, the formation of tertiary amines from the thermal degradation of ammonium cations and the enhanced mobility of

The reduction of air permeability thanks to clays dispersed in a poly(isobutene) matrix led

to the first commercial application of RCN, that will be discussed in Paragraph 11.2.

Mooney viscosity, that was not observed with a short chain substituent [76].

ammonium methylsulfate) in poly(epichlorohydrine) [78].

performing vulcanization reactions at lower temperatures.

sulphur accelerating anionic species were proposed [89].

**9. Barrier properties of RCN** 

matrix was apolar rubber such as NBR [79].

**8. Vulcanization of RCN** 

Platelets with a high aspect ratio substantially reduce the diffusion of penetrating molecules in polymer matrices. The tortuous path model, represented in Figure 8, is generally accepted to explain this result.

**Figure 8.** Tortuous path for molecules in a matrix containing platelets

Various continuum models were developed and applied to polymer clay nanocomposites, assuming in most cases a random dispersion of plates, parallel to each other and perpendicular to the direction of molecules diffusion, as summarized in refs. [90] and [91]. According to these models, reduction of permeability is enhanced by the increase of the platelet aspect ratio. These models are able to interpret the reduction of permeability in a polymer matrix, without necessarily having the presence of nano-platelets. This means that a nano-effect should not be invoked. It was commented [91] that permeating molecules have a sub-nano level and their permeation is slowed down by platelets, whatever is their size.

### **9.2. Improvement of barrier properties of rubber matrices thanks to the use of clays**

Better barrier properties of a rubber matrix, thanks to clay addition, were demonstrated for many different rubbers. In NR as the matrix, the barrier properties were enhanced at low clay (Mt) content, preparing the nanocomposite through the emulsion blending: 1, 2 and 3 phr of clay led respectively to more than 35% and to about 45% and 50% reduction of oxygen permeability [92]. 3 phr of OC (Mt/ didodecyl methyl amine) gave a 50% reduction of the oxygen permeability and a 40% reduction of toluene absorption at 20 oC [93]. 5 and 10 phr of OC led, respectively, to about 10% and 15% reduction of the oxygen permeability, and, at 15% OC, to a 30% reduction of toluene absorption at 30oC [94]. In BR as the matrix, 5 phr OC (Mt/dimethyl ditallow-ammonium) dispersed from solution blending, led to a reduction of about 80% of water vapour permeability, whereas 10 phr OC gave about 20% reduction of toluene uptake at equilibrium [95]. In NBR as the

matrix, 1 phr of OC (Mt/dimethyl ditallow-ammonium) dispersed from solution blending brought to about 80% decrease of water vapour permeability [96]. Better barrier properties were observed by increasing the AN content of NBR: the relative nitrogen permeability was reduced by 11.5, 10.4, and 9.0% for 42NBR, 35NBR, and 26NBR (the figures indicate the wt% of AN), respectively, with 10 phr OC (Mt/dimethyl dialkyl (C14– C18) ammonium) [97].

A clay with a high aspect ratio such as rectorite was found to give better barrier than carbon black N330 [38].

## **10. Mechanical properties of RCN**

### **10.1. The origin of reinforcement: nano-structured and nano-fillers**

Particulate fillers are used to reinforce traditional rubber compounds: they are carbon black and silica and are made by spherical individual particles of few tents of nanometers as the diameter, fused together to form aggregates extending up to few hundreds of nanometers. These fillers are called nano-structured as the aggregates can not be separated into individual particles by thermomechanical mixing.

Theories on the origin of reinforcement [98], developed on carbon black based compounds, led to present the modulus of a filled rubber compound as due to the sum of different contributions, as it shown in Figure 8.

**Figure 9.** Contributions to the modulus of a rubber compound filled with a particulate filler

Contributions that do not depend on the strain amplitude are due to the polymer network (entanglements, physical and chemical crosslinks), hydrodynamic effects (related to the filler volume fraction and implying the strain amplification mechanism and thus the enhancement of the modulus), immobilization of rubber on filler particles (that transform a highly viscous liquid in a solid). A contribution to the modulus that strongly depends on the strain amplitude is due to the so called filler network: filler particles are joined together either directly or through polymer layers.

104 Advanced Elastomers – Technology, Properties and Applications

**10. Mechanical properties of RCN** 

individual particles by thermomechanical mixing.

contributions, as it shown in Figure 8.

C18) ammonium) [97].

black N330 [38].

matrix, 1 phr of OC (Mt/dimethyl ditallow-ammonium) dispersed from solution blending brought to about 80% decrease of water vapour permeability [96]. Better barrier properties were observed by increasing the AN content of NBR: the relative nitrogen permeability was reduced by 11.5, 10.4, and 9.0% for 42NBR, 35NBR, and 26NBR (the figures indicate the wt% of AN), respectively, with 10 phr OC (Mt/dimethyl dialkyl (C14–

A clay with a high aspect ratio such as rectorite was found to give better barrier than carbon

Particulate fillers are used to reinforce traditional rubber compounds: they are carbon black and silica and are made by spherical individual particles of few tents of nanometers as the diameter, fused together to form aggregates extending up to few hundreds of nanometers. These fillers are called nano-structured as the aggregates can not be separated into

Theories on the origin of reinforcement [98], developed on carbon black based compounds, led to present the modulus of a filled rubber compound as due to the sum of different

**Figure 9.** Contributions to the modulus of a rubber compound filled with a particulate filler

**10.1. The origin of reinforcement: nano-structured and nano-fillers** 

The Guth Equation is used to correlate the compound initial modulus with the filler volume fraction (equation 1 as follows).

$$\mathbf{E} = \text{Em}\left(1 + 0.67\,\text{f}\phi + 1.62\,\text{f}^2\,\phi^2\right) \tag{3}$$

This equation holds up to a threshold of filler content, known as the percolation threshold, at which a continuous network is established in the rubber matrix and thus accounts for the sum of the three contributions commented above and shown in Figure 8 that do not depend on the strain amplitude.

In equation 1, appears the *shape factor f,* that was already commented in Paragraph 2. The *shape factor f* was introduced in the Guth model in order to account for the fact that particle aggregation (clustering) has a significant impact on stiffness at high volume fractions (higher than 0.15). These clusters are formed by the filler aggregates and the impact of the *f*  value on the quadratic term of the Guth equation makes them much important for the total value of the modulus.

In the case of a clay, as it was shown in Figure 1b, the *shape factor f* is given by the ratio between the longest lateral side of the layer and the height of clay stacks, the highest *f* value being obtained in the case of a single layer. Therefore, the exfoliation of a clay not only favours a better clay dispersion but also improves the modulus of the compound. To take into account the lower contribution to the modulus of a platelet like filler, a modulus reduction factor of about 0.7 was determined [99] by fitting experimental data.

An important contribution to the reinforcement is given by the so called "in rubber structure": fillers are able to accomodate polymer chains in the voids of their structure. This mechanism is well known in the case of nanostructured fillers such as carbon black and silica. In the case of clays, the immobilization of polymer chains would require their intercalation in the interlayer space, phenomenon that is however, as explained in paragraph 6.1, still a matter of debate.

OC give an important contribution to the polymer network. In fact, as discussed in paragraph 8, they promote a higher crosslinking density, thanks to the cation-thiolate interaction. Clay layers and aggregates could be represented as framed in a cage formed by short sulphur bonds, that prevents their slip on the polymer chains, phenomenon that occurs in uncrosslinked samples, as commented in paragraph 7.

In the same paragraph, it was commented that a low clay concentration is needed for achieving the percolation threshold in a hydrocarbon rubber matrix. This conclusion was

drawn also by determining the dependence of the Young modulus (obtained from stress-strain curves) on the OC concentration, in a SBR matrix (with Nanomer I.42E from Nanocor as OC) [100] or in an isoprene rubber matrix (OC was Mt/2HT), either synthetic [67] or naturally occurring [101]. The excess of the Young modulus was plotted versus the filler fraction [102], determining the percolation threshold, that was: 2.7 vol % in SBR [100], 2.9 vol % in IR [67] and 4 vol % in NR [101]. The excess modulus scales with a power law with an exponent between 1.8 and 2.5 above the percolation threshold, lower with respect to the one typical of carbon black (about 4). The structural difference of the filler is proposed to justify this difference. As mentioned above, the filler network is responsible for the contribution to the modulus that depends on the strain amplitude. Thanks to the easy formation of a filler network, clays promote a remarkable non linearity of the dynamicmechanical behaviour of rubber nanocomposites, a phenomenon known as Payne effect [103]. Data that the show the remarkable Payne effect of RCN, increasing with the clay content, were reported for NR [104] and IR [67] as the rubber matrices. A strong reduction of the dynamic modulus was observed also by increasing the temperature [50]. The 2-D melting of the paraffinic chains substituents of the compensating ion [105] was commented as a possible explanation [66, 101].

### **10.2. The improvement of rubber mechanical properties thanks to the use of clays**

Many data are available in the literature to show that both pristine clays and OC bring about the improvement of the mechanical properties of a rubber matrix, preparing the nanocomposites from latex, melt and solution blending. Some examples are reported as follows.

Pristine Mt was dispersed in a NR latex, observing an increase of modulus for the obtained nanocomposite up to 30 phr of clay, in the presence of worse ultimate properties [41], as well as higher 300% stress, shore A hardness, tensile strength and tear strength, when Mt was at 20 phr level [29]. In a SBR matrix as well, tensile and tear strengths were improved by adding a pristine Mt via emulsion blending, up to 20 phr [29], without any improvement for a further addition [35].

OC were typically used at lower levels. In NR as the matrix, 5 phr of Mt modified by tetraoctyl phosphonium bromide led to a 3-fold increase of tensile strength [81] and best ultimate properties were obtained with 10 phr of OC (Mt/octadecyltrimethyl ammonium) [106]. In BR as the matrix, best mechanical properties were obtained with a Mt having dimethyl dihydrogenated tallow quaternary ammonium as the compensating cation [107], an ammonium cation largely diffused in RCN field. In a SBR matrix, mechanical properties improved up to 10 phr of OC (Mt/octadecyltrimethylamine, Nanomer I.28E) [87] (Mousa and Karger-Kocsis, 2001), decreasing for a further addition. In a nanocomposite based on maleic anhydride modified EPM (EPM-MA, with 0.42 wt% of maleic anhydride) about 5% of OC (Mt/octadecyl amine) gave a modulus three times higher than that of EPM-MA. Modulus increased and elongation at break decreased with clay content [108]. Reinforcement was obtained from OC also in rubber blends. 2 phr of OC (Mt/octadecyltrimethylamine) in an NR/ENR blend (10 phr ENR50) gave an increase of both strain and stress at break values [109] (Teh et al., 2004).

It was commented in Paragraph 7 the importance of sulphur based vulcanization to prevent the OC slip on the polymer chains. OC (Mt modified with hydrogenated tallow ammonium) was reported to promote a remarkable reinforcement also of a peroxide cured rubber matrix. However, in this case, the rubber was NBR, with acrylonitrile content as high as 50% [110]. It was also shown that higher values of tensile strength and better tear resistance were obtained, with a worse elongation at break, by increasing the AN content, with 10 phr of OC (Mt/dimethyl dialkyl (C14–C18) ammonium) [111].

### **10.3. Clays in a rubber matrix for lighter weight of the compound**

106 Advanced Elastomers – Technology, Properties and Applications

as a possible explanation [66, 101].

**clays** 

follows.

a further addition [35].

drawn also by determining the dependence of the Young modulus (obtained from stress-strain curves) on the OC concentration, in a SBR matrix (with Nanomer I.42E from Nanocor as OC) [100] or in an isoprene rubber matrix (OC was Mt/2HT), either synthetic [67] or naturally occurring [101]. The excess of the Young modulus was plotted versus the filler fraction [102], determining the percolation threshold, that was: 2.7 vol % in SBR [100], 2.9 vol % in IR [67] and 4 vol % in NR [101]. The excess modulus scales with a power law with an exponent between 1.8 and 2.5 above the percolation threshold, lower with respect to the one typical of carbon black (about 4). The structural difference of the filler is proposed to justify this difference. As mentioned above, the filler network is responsible for the contribution to the modulus that depends on the strain amplitude. Thanks to the easy formation of a filler network, clays promote a remarkable non linearity of the dynamicmechanical behaviour of rubber nanocomposites, a phenomenon known as Payne effect [103]. Data that the show the remarkable Payne effect of RCN, increasing with the clay content, were reported for NR [104] and IR [67] as the rubber matrices. A strong reduction of the dynamic modulus was observed also by increasing the temperature [50]. The 2-D melting of the paraffinic chains substituents of the compensating ion [105] was commented

**10.2. The improvement of rubber mechanical properties thanks to the use of** 

Many data are available in the literature to show that both pristine clays and OC bring about the improvement of the mechanical properties of a rubber matrix, preparing the nanocomposites from latex, melt and solution blending. Some examples are reported as

Pristine Mt was dispersed in a NR latex, observing an increase of modulus for the obtained nanocomposite up to 30 phr of clay, in the presence of worse ultimate properties [41], as well as higher 300% stress, shore A hardness, tensile strength and tear strength, when Mt was at 20 phr level [29]. In a SBR matrix as well, tensile and tear strengths were improved by adding a pristine Mt via emulsion blending, up to 20 phr [29], without any improvement for

OC were typically used at lower levels. In NR as the matrix, 5 phr of Mt modified by tetraoctyl phosphonium bromide led to a 3-fold increase of tensile strength [81] and best ultimate properties were obtained with 10 phr of OC (Mt/octadecyltrimethyl ammonium) [106]. In BR as the matrix, best mechanical properties were obtained with a Mt having dimethyl dihydrogenated tallow quaternary ammonium as the compensating cation [107], an ammonium cation largely diffused in RCN field. In a SBR matrix, mechanical properties improved up to 10 phr of OC (Mt/octadecyltrimethylamine, Nanomer I.28E) [87] (Mousa and Karger-Kocsis, 2001), decreasing for a further addition. In a nanocomposite based on maleic anhydride modified EPM (EPM-MA, with 0.42 wt% of maleic anhydride) about 5% of OC (Mt/octadecyl amine) gave a modulus three times higher than that of EPM-MA. Modulus increased and elongation at break decreased with clay content [108]. The most important application of rubber compounds is in tires. Lighter rubber compounds are thus pursued, in the light of their remarkable impact on the environment. Direct comparisons were reported between OC and CB in promoting the mechanical reinforcement of rubber compounds.

In NR as the matrix, a comparable mechanical reinforcement was obtained with 10 phr OC (Mt/octadecylamine) and 40 phr CB, with almost twice elongation at break, a lower compression set and similar abrasion loss with OC [112]. With the same OC, a much lower heat build up was obtained with respect to 50 phr of CB [113].

In BR as the matrix, better mechanical performance was obtained with 3 phr OC (Mt/dimethyl dihydrogenated tallow quaternary ammonium) compared with 10 phr of CB [114]. In NBR as the matrix, a comparable tensile strength was obtained with either 10 phr OC or with 40 phr CB [115]. In EPM-MA, 5% of OC (Mt/octadecyl amine) gave the same reinforcement as 30 wt% CB [108].

### **10.4. Synergistics effects between clays and nanostructured fillers**

An increasing number of studies is becoming available in the literature on RCN based on clays and a nanostructured filler such as silica and, in particular, CB. A large scale application of RCN could reasonably imply the use of hybrid filler systems, with a minor amount of OC added to a major part of a traditional filler. In SBR as the matrix, a hybrid sepiolite-silica filler system was adopted [116]. Hybrid OC/CB filler systems were used in the following matrices: SBR [117-119], brominated poly(isobutylene-co-paramethylstyrene) rubber [117], IR [50, 57, 120], NR [67, 121] NR/SBR blend [122], chlorobutyl rubber [123], ENR [124], EPDM [125] (Malas and Das, 2012). In some of these works [50, 57, 120-122], CB was used at a level typical of commercial applications (higher than 50 phr). A remarkable enhancement of material dynamic-mechanical properties was reported, thanks to the use of the hybrid filler system. The most important finding from these study was the synergism developed by the two fillers [118, 120, 121]. In particular, it was shown [126] that initial modulus values obtained with the hybrid CB-OMt filler system were much higher than

those calculated through the simple addition of the two initial moduli of composites with only CB or only OC. Figure 9, taken from ref. [127], demonstrates the synergistic effect between OC and CB. The dashed line indicates the initial modulus calculated by simply adding the initial moduli of composites with only CB and only OC, whereas points refer to experimental data taken from composites containing the hybrid filler system.

**Figure 10.** Initial modulus values () of IR based composites containing the hybrid OC/CB filler system. The dashed line refers to values predicted through the addition of initial moduli values of composites with only CB and only OC.

### **11. Applications of RCN**

In this paragraph, applications of RCN are examined, moving from products already available on the commercial scale. The most meaningful patents were taken as the reference literature for the considered applications. RCN applications, diffused on the commercial scale, well documented and advertised, are essentially in two fields: tires and sport balls.

#### **11.1. Applications of RCN for tyres**

The most abundant application for rubbers is in tire compounds. NR is the most diffused rubber, with almost 11 million ton in 2010 [128] and SBR is the most produced synthetic rubber, with about 6 million ton. The application of both these rubbers is for about 75% in tire compounds. The world demand for tires is rising, at a pace of about 4.7 per year through 2015, achieving a number of produced tires of about 3.3 billion units. In the same period, the tire market is projected to increase, in terms of value, of about 6.5 percent annually, achieving a total value level of \$220 billion 129]. It is thus evident that the application of RCN in tire compounds can lead to a real commercial diffusion as well as to an important economic impact.

A tire can be defined as a toroidal high performance composite behaving as a flexible membrane, able to contain gases under pressure and with the following capabilities: load carrying, cushioning, road handling. Figure 9 reports the structure of a tire, with the name of constitutive parts.

**Figure 11.** Structure of a tire

108 Advanced Elastomers – Technology, Properties and Applications

with only CB and only OC.

economic impact.

**11. Applications of RCN** 

**11.1. Applications of RCN for tyres** 

those calculated through the simple addition of the two initial moduli of composites with only CB or only OC. Figure 9, taken from ref. [127], demonstrates the synergistic effect between OC and CB. The dashed line indicates the initial modulus calculated by simply adding the initial moduli of composites with only CB and only OC, whereas points refer to

**Figure 10.** Initial modulus values () of IR based composites containing the hybrid OC/CB filler system. The dashed line refers to values predicted through the addition of initial moduli values of composites

In this paragraph, applications of RCN are examined, moving from products already available on the commercial scale. The most meaningful patents were taken as the reference literature for the considered applications. RCN applications, diffused on the commercial scale, well documented and advertised, are essentially in two fields: tires and sport balls.

The most abundant application for rubbers is in tire compounds. NR is the most diffused rubber, with almost 11 million ton in 2010 [128] and SBR is the most produced synthetic rubber, with about 6 million ton. The application of both these rubbers is for about 75% in tire compounds. The world demand for tires is rising, at a pace of about 4.7 per year through 2015, achieving a number of produced tires of about 3.3 billion units. In the same period, the tire market is projected to increase, in terms of value, of about 6.5 percent annually, achieving a total value level of \$220 billion 129]. It is thus evident that the application of RCN in tire compounds can lead to a real commercial diffusion as well as to an important

experimental data taken from composites containing the hybrid filler system.

RCN have been so far essentially applied in three parts of a tire: tread, base (i.e. a compound immediately before the tread in contact with the road), liner. Table 4 explains the position in a tire of such compounds and clarifies their role.

RCN in tire compounds are used to achieve the following properties and performances: reduced weight, reduced energy dissipation, enhanced air retention and extension of the balance of the so called magic triangle performances for a tire tread: rolling resistance, traction, wear.

*RCN in tire treads.* By examining the patent literature [130-132], it appears that clays were used in tire tread compounds aiming at longer life of the compound and lower fuel consumption without negatively affecting the tread grip on the road at low temperature. Clays were used with or without an ammonium cation modifier, in the former case in the presence of a large amount of silica. Compound properties claimed in patents are: better ultimate properties, lower abrasion and an hysteresis that remains analogous at low temperature and is lower at medium – high temperature.

*RCN in base compounds*. OC were used in base compounds of tires commonly referred to as "HP" (High Performance) or "UHP" (Ultra high performance) tires, belonging to classes "V" or "Z", designed to experience extreme driving conditions, as they achieve maximum speeds in a range from 210 to 240 km/h and higher than 240km/h, respectively. In this type

of tires, a low thickness base compound is used to favour performances such as tread block stability, road grip, steering stability, cornering stability and ride comfort. This base compound should have high tensile and dynamic-mechanical properties. OC were reported in the patent literature to give a remarkable improvement of the compound dynamic mechanical properties [133-136] and Pirelli Tire launched in 2007 P Zero tires for the HP and UHP segments, with a base compound containing an OC. In particular, it was reported that OC was able to give a much more isotropic behavior (i.e. an equal performances in longitudinal and lateral directions), with respect to traditional reinforcing fibers, such as the aramide ones.


**Table 4.** Tyre compounds explored for RCN application

*RCN in innerliner.* Most research efforts were dedicated to develop OC based innerliner compounds, with the aim to exploit the barrier property provided by clay platelets, trying in particular to achieve a high clay dispersion. In order to have the best barrier, the following technical solutions were adopted: use of clays with a high aspect ratio, control of the clay organization (stacks of exfoliated layers), use of reactive rubbers to promote the clay exfoliation, use of coating layers with a high impermeability.

A kaolin [137] and a mica [138] with high aspect ratio (at least 50) were used in butyl and in butyl, butadiene and natural rubber, respectively.

Stacks of OC were used in BR [139], in BIMS [140] and in SBR [141]. With an OC content as low as 9 phr, the oxygen transmission was reduced to about one fourth [139].

The clay exfoliation was promoted either by dispersing clays in the emulsion were polymerization of monomers such as isoprene and styrene was performed [142], or by performing the exchange reaction with a cationic polymer latex [143],or by mixing the clay (in a BIMS matrix) in the presence of a tertiary amine. In this latter case, with only 3 phr of clay, a reduction of 20% permeability was detected [144].

An amine terminated oligomer (e.g. a butadiene-acrylonitrile copolymer) was used as the reactive polymer to prepare a 25/75 clay/SBR blend and the oxygen transmission was reduced to less than one half [145]. A clay was mixed with a aqueous dispersion of a dienebased elastomer, with one or more functional groups such as acid or an anhydride [146].

OC was melt blended with poly(isobutylene-co-p-alkylstyrene) and poly(isobutylene-coisoprene) elastomers and at a concentration of 10 phr the permeability was reduced to one half [147]. OC was melt blended with an halogenated copolymer of isobutylene and pmethyl-styrene and, at an OC level of 5 phr, the air retention was improved up to 30% [148]. Poly(isobutylene-co-N,N-dimethylvinylbenzylamine) copolymer was blended with an OC and an improvement of impermeability up to 70% was reported [149].

Barrier coating mixtures were prepared by mixing an elastomer (butyl rubber) an exfoliated clay with a high aspect ratio (preferably a vermiculite) and a surfactant, adopting a polymer/clay ratio from 1:1 to 20:1 and obtaining a permeability reduction of about 25%. This technology found a large commercial success in tennis ball and it will be discussed in the following paragraph [150].

### **11.2. Applications of RCN for sport balls**

As mentioned above, a clay based barrier coating technology was developed [151-154], and applied to sport balls, with the aim to drastically improve the pressure retention without negatively affecting properties such as the bounce, the feel and the reproducibility of performances. A relatively thin coating layer (about 10-30 micron) was deposed, made by an elastomeric nanocomposite prepared by combining aqueous dispersions of a vermiculite (having a aspect ratio up to 1000) and butyl rubber. In a laboratory study [90] (Takahashi et al., 2006) a coating layer containing vermiculite (20 - 30 wt%) in butyl rubber led to a reduction of the diffusion coefficients by two orders of magnitude and of gas permeability by 20–30 fold. This technology was developed by InMat company and Air D-Fense products were commercialized.

### **Author details**

110 Advanced Elastomers – Technology, Properties and Applications

Base (tread cushion) Below the tread.

**Table 4.** Tyre compounds explored for RCN application

butyl, butadiene and natural rubber, respectively.

clay, a reduction of 20% permeability was detected [144].

exfoliation, use of coating layers with a high impermeability.

aramide ones.

Liner or innerliner

of impermeable rubber.

Thin layer

of tires, a low thickness base compound is used to favour performances such as tread block stability, road grip, steering stability, cornering stability and ride comfort. This base compound should have high tensile and dynamic-mechanical properties. OC were reported in the patent literature to give a remarkable improvement of the compound dynamic mechanical properties [133-136] and Pirelli Tire launched in 2007 P Zero tires for the HP and UHP segments, with a base compound containing an OC. In particular, it was reported that OC was able to give a much more isotropic behavior (i.e. an equal performances in longitudinal and lateral directions), with respect to traditional reinforcing fibers, such as the

Tread In contact with road Dry, wet, ice and snow traction.

*RCN in innerliner.* Most research efforts were dedicated to develop OC based innerliner compounds, with the aim to exploit the barrier property provided by clay platelets, trying in particular to achieve a high clay dispersion. In order to have the best barrier, the following technical solutions were adopted: use of clays with a high aspect ratio, control of the clay organization (stacks of exfoliated layers), use of reactive rubbers to promote the clay

A kaolin [137] and a mica [138] with high aspect ratio (at least 50) were used in butyl and in

Stacks of OC were used in BR [139], in BIMS [140] and in SBR [141]. With an OC content as

The clay exfoliation was promoted either by dispersing clays in the emulsion were polymerization of monomers such as isoprene and styrene was performed [142], or by performing the exchange reaction with a cationic polymer latex [143],or by mixing the clay (in a BIMS matrix) in the presence of a tertiary amine. In this latter case, with only 3 phr of

low as 9 phr, the oxygen transmission was reduced to about one fourth [139].

Under the carcass Impermeability to air and

Stability at high speed. Abrasion

Low hysteresis. Good adhesion.

resistance. Protection of compounds below.

Good fatigue, tear and durability. Compatibility

moisture. Prevention of degradation of tire structure (due to air and moisture). Good flex fatigue, crack, long

term aging resistance.

Tyre compound Position in the tyre Performances

Between tread and nylon 0°

Maurizio Galimberti *Politecnico di Milano, Dipartimento di Chimica, Materiali e Ingegneria Chimica G. Natta, Milano, Italia* 

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112 Advanced Elastomers – Technology, Properties and Applications

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	- [154] Feeney C.A., Farrell M., Tannert K., Goldberg H.A., Lu M., Grah M.D. (2000) US Patent 6,087,016 to InMat and Michelin Recherche et technique

**Chapter 5** 

## **Pigment and Dye Modified Fillers as Elastomeric Additives**

Anna Marzec and Marian Zaborski

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/50735

### **1. Introduction**

120 Advanced Elastomers – Technology, Properties and Applications

elastomer nanocomposites. US 5,807,629 to Exxon Mobil

[141] Ishida, K., Fujiki, K., 2003. WO 03/087214 A1 to Bridgestone

nanocomposites. US 2005/0027058 to Exxonmobil.

nanocomposite. WO 04/005387 to Exxonmobil.

Michelin and Herberts Gmbh.

7,078453 to InMat Inc

7,119,138 to InMat Inc

thereof. EP 1321489 A1 to Goodyear Tyre and Rubber Company.

US 2003/0144401 A1 to Goodyear Tyre and Rubber Company

containing same. WO 058874 A1 to Bridgestone Corporation.

6,232,389. to InMat and Michelin Recherche et technique

6,087,016 to InMat and Michelin Recherche et technique

Rubber Ind.

Engineering Company.

[138] Miyazaki T. (2006) Rubber composition for inner liner. EP 1726620 A1 to Sumitomo

[139] Elspass W.C., Peiffer G.D., Kresge E.N., Wright P.J., Wang H.C. (1998) Tactoidal

[140] Elspass C.W., Peiffer D.G. (2000) Nanocomposite materials formed from inorganic layered materials dispersed in a polymer matrix. US 6,034,164 to Exxon Research and

[142] Elspass C.W., Peiffer D.G., Kresge E.N., Hsieh D.T., Chludzinski J.J., Liang K.S., (1999) Nanocomposite materials. US 5,883,173 to Exxon Research and Engineering Company [143] Parker D.K., Larson B.K.F., Yang X. (2003) Preparation and use of a nanocomposite of elastomer and exfoliated clay platelets formed in situ within an elastomer host and articles of manufacture, including tires, having at least one component comprised

[144] Dias A.J., Tsou A. H., Chung D.Y.L. Weng W. (2005). Low permeability

[145] Kresge E.N., Lohse D.J. (1997) Tire inner-liners comprising a solid rubber and a complex of a reactive rubber and layered silicate clay. US 5,665,183 to Exxonmobil [146] Ajbani M., Geiser J.F., Parker D.K. (2003) Nanocomposite of elastomer and dispersion therein of intercalated clay prepared in an aqueous medium with funcional elastomer.

[147] Gong C., Dias A.J., Tsou A. H., Poole B.J., Karp K.R. (2004) Functionalized Elastomer

[150] Feeney C.A., Balzer, R.J. (1998) Barrier coating of an elastomer and a dispersed layered filler in a liquid carrier and coated compositions, particularly tires. WO 98/56598 to

[151] Feeney C. A., Goldberg,H. A., Farrell M., Karim D. P., Oree K. R. (2006) US Patent

[152] Feeney C. A., Goldberg H. A., Farrell M., Karim D. P., Oree K. R. (2006). US Patent

[153] Feeney C.A., Farrell M., Tannert K., Goldberg H.A., Lu M., Grah M.D. (2001) US Patent

[154] Feeney C.A., Farrell M., Tannert K., Goldberg H.A., Lu M., Grah M.D. (2000) US Patent

[148] Grah M.D. (2004) Tire with improved inner liner. US 2004/0194863 A1 to Michelin [149] Wang X., Fudemoto H., Hall J., Araki S., Hogan T., Foltz V., Sadhukhan P., Bohom, G.G.A. (2004) Method for clay exfoliation, compositions therefore, and modified rubber Pigments are a dynamically developinged area of chemical technology. This interesting scientific concept involves preparation of hybrid pigments by the formation of a stable dye sphere surrounding an inorganic substrate grain, such as silica, titanium dioxide, or aluminosilicates. The concept of combining properties of organic and inorganic pigments has been employed for a long time. At the end of the 19th century, this concept was applied to the production of a pigment lake where freshly precipitated aluminium hydroxide was employed as the support. These lakes were used for the production of graphic dyes due to their averaged properties. Silica, zinc oxide, carbon black, and titanium dioxide may be employed in rubber compounds as pigments and fillers. Fillers boast a wide range of industrial applications, such as in the pharmaceutical and textile industries, and are important as additives for paints and varnishes. Fillers are relatively inexpensive, solid substances that are added in fairly large volumes to polymers to adjust the volume, weight, cost, surface, colour, expansion coefficient, conductivity, permeability, and mechanical properties. They can be roughly divided into inactive or extender fillers and active or functional or reinforcing fillers. One of the most popular fillers is synthetic silica, which is used in the polymer industry as an active filler (Liu et al., 2008, Wang et al., 2009). The rubber industry utilises fine silicas, which have a large specific surface leading to favourable effects on the strength of the composites. One important processing challenge is that silica surfaces are hydrophilic in nature, which impedes their dispergation in polymer media. Nearly all of the silica forms tend to agglomerate, which is an undesirable behaviour from the point of view of utilising their functional properties in industrial applications. On-going research on this topic focuses on altering the characteristics of the filler surfaces through some type of modification (Ciesielczyk et al., 2007; Ciesielczyk et al., 2011; Klapiszewska et al., 2003; Kohno et al., 2011; Ladewig et al., 2012; Kishore et al., 2012; Natinee et al., 2011 ). For example, industry has developed hydrophobic silicas that easily dispergate in polymers. The

© 2012 Marzec and Zaborski, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

production of these silicas involves a silica silanisation reaction with alkoxysilane containing functional groups (Jesionowski et al., 2001, Krysztawkiewicz et al., 2000; Elham et al., 2012; Kister et al., 2012; Jianzhong et al., 2011; Wunpen et al., 2011; Jung-Woo et al., 2011). Over the past decade, numerous studies have employed a dye to modify the surface of silica to obtain hydrophobic silica surfaces and visually appealing end product colours (Ansarifar et al., 2004; Al Dwayya et al., 2012). The presence of amino groups in these supports enables the formation of ionic bonds with acid and direct dyes as well as the formation of covalent bonds with reactive dyes (Donia et al., 2009). Dyes, which were applied for this purpose, belonged to the textile class. Although they were combined with silica, they could not fulfill the requirements of modern pigments, such as light-resistance or resistance to solvents and to high temperature. High temperature during polymer processing can damage or destroy the dye causing changes in the shade or loss of colour. The present study was directed at obtaining composite pigments of indigothiazine and silicas with different surface areas. Indigothiazine consists of two thiazine residues conjugated to each other by a double bond. Because their cyclic carbonamide groups , thiazine can form hydrogen bond to silanol groups present on the silica surface. The thiazine pigments exhibited good heat stability, insolubility and exposure durability. An alternative method of producing coloured fillers is via a mechanical approach in which a coloured organic system is mechanically applied to the surface of a non-organic support. This method is a simple and cheap method of producing an organic-inorganic pigment encasing silica in the coloured chromophore with the additional benefit of eliminating the undesirable solvents from the process. The anti-ageing properties of indigothiazine were confirmed by oxidation-reduction potential measurements. The composites were produced by mechanically applying a coloured organic system (indigothiazine pigment) to a non-organic support (silica). The prepared pigment composites were characterised by the following properties: agglomerate size distribution, zeta potential, and specific surface area of the silica fillers before and after application of indigothiazine. The composite pigments were evaluated for their functional properties and subsequently employed as fillers in an acrylonitrile butadiene rubber (NBR). The mechanical properties as well as the spectrophotometric and DSC data of the vulcanisates produced with the composite pigment fillers were also studied before and after being aged under UV radiation.

### **2. Colourants in polymer composites**

#### **2.1. Fillers modification**

Compared to inorganic pigments, organic pigments exhibit greater vividness, higher colour intensity and higher staining potential. Their drawbacks include greater sensitivity to the action of chemical compounds, temperature and solvents (Binkowski et at., 2000). Several studies were conducted to obtain organic pigments that were permanently attached to a silica core to improve the resistance of the pigments to chemical and thermal treatments. Numerous studies have shown that the reaction mechanism that occurs on the surface of inorganic substances is extremely difficult to identify. For example, the reaction that occurs on the surface of silica with fluorescein results in the formation of covalent bonds between the reagents. The interactions may involve the formation of hydrogen bonds as well as van der Waals or electrostatic interactions (Jesionowski 2003; Wu et al., 1997). Silica is one of the most popular fillers employed as a pigment carrier. Silicas are extensively modifiable materials. Their modification results in products that have new functional groups on their surface and are capable of interacting with various organic compounds. For example, modification with aminosilane provided functional amino groups that can react with the carbonyl groups of aldehydes, ketones or esters. The aminosilane-modified silicas find multiple applications as polymer fillers in industry. Recently, they have been used with increasing frequency as coupling agents in pigment or organic systems. Until recently, studies of organic pigments involved modification of their basic properties, such as stability or intensity of the colours. However, in many cases, the pigments must satisfy additional requirements (e.g., well defined particle diameters). Pigments with silica cores can satisfy these requirements (Krysztafkiewicz et al., 2003). Bińkowski (Bińkowski et al., 2000) proposed a simple procedure for obtaining coloured silicas, which involves attaching the dye to a modified silica surface. The procedure is possible due to the availability of a wide range of commercial dyes. The resulting pigment exhibits a structure in which the silanol groups on the surface of the silica are chemically coupled to an organic dye (e.g., azo dye) through an aminosilane coupling agent. They represented the pigment structure using the following formula: S-(O-Si-R1-Z-D)n where R1 is the alkyl group, Z is the amine bridge fragment , D is the dye, n is a number not less than 1 and -Si-R1-Z- represents a group originating from the silane coupling agent. In this work, they employed precipitated silica, 3-aminopropylotriethoxysilane (U-13), 2-(aminoethyl)-3-aminopropylotriethoxysilane (U-15) and a dye obtained from Boruta-Kolor, C.I. Direct Red 81. The modifications resulted in uniform pigments. A subsequent report by the authors (Jesionowski et al., 2011) showed the adsorption of C.I. Mordant Red 3 on the surface of silica that was both unmodified and modified with *N*-2-(aminoethyl)-3-aminopropyltrimethoxysilane. They proposed two mechanisms for the dye addition to the silane modified silica surface including via a hydrogen bond between the hydroxyl group of C.I. Mordant Red 3 and the amine groups of the silane coupling agents and/or via electrostatic interaction between the dissociated anion of the dye and the cation appearing on the silica surface modified with silane. Sudam (Sudam et al., 2003) described the adsorption behaviour of selected styryl pyridinium dyes on silica gel. The adsorbents were styryl pyridinium dyes with a monochromatic group containing an alkyl chain at the pyridinium nitrogen and a bischromatic group containing methylene groups bridging at the pyridinium nitrogen atoms. They observed that the bischromophoric dyes anchor on the silica surface in a flat-on position in which the methylene units also come in contact with the silica surface. Andrzejewska (Andrzejewska et al., 2004) performed studies to obtain pigments by adsorption of organic dyes on the modified surface of a titanium dioxide pigment system. The titanium oxide surface was modified with silane coupling agents, such as 3-aminopropyltriethoxysilane (U-13), *N*-2- (aminoethyl)-3-aminopropyltriethoxysilane (U-15D), in various solvents (methanol, toluene, acetone and methanol-water mixture). C.I. Acid Orange 7 and C.I Reactive Blue 19 were employed to prepare the pigments. The organic dyes were deposited on the surface of the modified titanium white. The modification of the titanium white with both silanes (U-13 and U-15D) boosted the efficiency of adsorption of C.I. Acid Orange 7 dye on the surface. The improved adsorption efficiency of C.I Reactive Blue 19 on the titanium was obtained

122 Advanced Elastomers – Technology, Properties and Applications

radiation.

**2. Colourants in polymer composites** 

**2.1. Fillers modification** 

production of these silicas involves a silica silanisation reaction with alkoxysilane containing functional groups (Jesionowski et al., 2001, Krysztawkiewicz et al., 2000; Elham et al., 2012; Kister et al., 2012; Jianzhong et al., 2011; Wunpen et al., 2011; Jung-Woo et al., 2011). Over the past decade, numerous studies have employed a dye to modify the surface of silica to obtain hydrophobic silica surfaces and visually appealing end product colours (Ansarifar et al., 2004; Al Dwayya et al., 2012). The presence of amino groups in these supports enables the formation of ionic bonds with acid and direct dyes as well as the formation of covalent bonds with reactive dyes (Donia et al., 2009). Dyes, which were applied for this purpose, belonged to the textile class. Although they were combined with silica, they could not fulfill the requirements of modern pigments, such as light-resistance or resistance to solvents and to high temperature. High temperature during polymer processing can damage or destroy the dye causing changes in the shade or loss of colour. The present study was directed at obtaining composite pigments of indigothiazine and silicas with different surface areas. Indigothiazine consists of two thiazine residues conjugated to each other by a double bond. Because their cyclic carbonamide groups , thiazine can form hydrogen bond to silanol groups present on the silica surface. The thiazine pigments exhibited good heat stability, insolubility and exposure durability. An alternative method of producing coloured fillers is via a mechanical approach in which a coloured organic system is mechanically applied to the surface of a non-organic support. This method is a simple and cheap method of producing an organic-inorganic pigment encasing silica in the coloured chromophore with the additional benefit of eliminating the undesirable solvents from the process. The anti-ageing properties of indigothiazine were confirmed by oxidation-reduction potential measurements. The composites were produced by mechanically applying a coloured organic system (indigothiazine pigment) to a non-organic support (silica). The prepared pigment composites were characterised by the following properties: agglomerate size distribution, zeta potential, and specific surface area of the silica fillers before and after application of indigothiazine. The composite pigments were evaluated for their functional properties and subsequently employed as fillers in an acrylonitrile butadiene rubber (NBR). The mechanical properties as well as the spectrophotometric and DSC data of the vulcanisates produced with the composite pigment fillers were also studied before and after being aged under UV

Compared to inorganic pigments, organic pigments exhibit greater vividness, higher colour intensity and higher staining potential. Their drawbacks include greater sensitivity to the action of chemical compounds, temperature and solvents (Binkowski et at., 2000). Several studies were conducted to obtain organic pigments that were permanently attached to a silica core to improve the resistance of the pigments to chemical and thermal treatments. Numerous studies have shown that the reaction mechanism that occurs on the surface of inorganic substances is extremely difficult to identify. For example, the reaction that occurs on the surface of silica with fluorescein results in the formation of covalent bonds between

following modification with silane U-13. For the titanium modified with U-15D, the efficiency of C.I. Reactive Blue 19 dye adsorption decreased slightly. A subsequent report (Raha et al., 2012) proposed a new type of inorganic carrier, Na-montmorilllonite (MMT). If the carrier is a smectite clay or layered silicate, the most effective pathway found involved using the ion exchange process. For that reaction, the dye must exist in a cationic form, which is the case for basic or cationic dyes. Example for that intercalated products of Rhodamine dyes with hectorite, montmorilllonite and other smectite silicates. The synthesis of the water-soluble cationic dyes was performed using two azo dyes (i.e., Solvent Red 24 and Solvent Yellow 14) and one disperse dye (i.e., Red 60) intercalated with Na+-MMT to create nano-structured pigments. The dye-intercalated montmorillonite was successfully obtained. The modification of the dyes into their respective cationic species was confirmed by NMR study, and XRD indicated that these nanopigments had a nanostructured morphology. The authors also observed that the use of the nanopigments in PP reduced the migration of the dye, which results from the intercalation /adsorption of the dye within the clay-based pigments. The most common methods of modifications are presented in Table 1.



**Table 1.** Common methods for the preparation of pigments.

#### **2.2. Organic and inorganic pigments**

124 Advanced Elastomers – Technology, Properties and Applications

**Type of supporter** 

Sodiumaluminium silicate

following modification with silane U-13. For the titanium modified with U-15D, the efficiency of C.I. Reactive Blue 19 dye adsorption decreased slightly. A subsequent report (Raha et al., 2012) proposed a new type of inorganic carrier, Na-montmorilllonite (MMT). If the carrier is a smectite clay or layered silicate, the most effective pathway found involved using the ion exchange process. For that reaction, the dye must exist in a cationic form, which is the case for basic or cationic dyes. Example for that intercalated products of Rhodamine dyes with hectorite, montmorilllonite and other smectite silicates. The synthesis of the water-soluble cationic dyes was performed using two azo dyes (i.e., Solvent Red 24 and Solvent Yellow 14) and one disperse dye (i.e., Red 60) intercalated with Na+-MMT to create nano-structured pigments. The dye-intercalated montmorillonite was successfully obtained. The modification of the dyes into their respective cationic species was confirmed by NMR study, and XRD indicated that these nanopigments had a nanostructured morphology. The authors also observed that the use of the nanopigments in PP reduced the migration of the dye, which results from the intercalation /adsorption of the dye within the clay-based pigments. The most common methods of modifications are presented in Table 1.

**The type of coupling agent Type of dye Reference** 

Yellow 1, Acid Blue 25, Guinea Green B

bischromatic styryl piridinium dyes

C.I. Reactive Blue 19, C.I. Acid Red 18

C.I. Acid Red 18, C.I. Acid Violet 1

C.I. Reactive Blue 19, C.I. Acid Green 16, C.I. Acid red 18, C.I. Acid Violet 1, C.I. Direct Red 81

C.I. Acid Orange 7, C.I. Reactive Blue 19 Wu et al., 1997

Parida et al., 1997

C.I. Direct Red 81 Binkowski et al.,


C.I. Reactive Blue 19 Jesionowski et

C.I. Acid Red 18 Jesionowski et

2000

2001

al., 2002

al., 2003

Jesionowski 2003

Krysztafkiewicz et al., 2003

Jesionowski et al., 2004

Andrzejewska et

al., 2004

Silica **-** D&C Red 6, Acid

Silica - Monochromatic and

(aminoethyl)-3-aminopropyltriethoxysilane

(aminoethyl)-3-aminopropyltriethoxysilane,

Vinyltris(β-methoxyethoxy)silane; γ-methacryloxypropyltrimethoxysilane

aminopropyltrimethoxysilane

aminopropyltrimethoxysilane

aminopropyltrimethoxysilane

3-ureidopropyltriethoxysilane

aminopropyltrimethoxysilane

aminopropyltrimethoxysilane

Silica 3-aminopropyltriethoxysilane, *N*-2-

Silica 3-aminopropyltriethoxysilane, *N*-2-

Titanium dioxide 3-aminopropylotriethoxysilane, *N*-2- (aminoethyl)-3-

Silica *N*-2-(aminoethyl)-3-

Silica *N*-2-(aminoethyl)-3-

Silica *N*-2-(aminoethyl)-3-

Silica *N*-2-(aminoethyl)-3-

Colourants used in polymers are pigments or dyes (Drobny, 2007). Dyes are organic compounds that are soluble in the polymers forming a molecular solution. They produce bright, intense colours and are transparent and easy to disperse and process. In general, pigments are insoluble in polymers. They produce opacity or translucence in the final product. Pigments can be inorganic or organic compounds and are available in a variety of forms, such as dry powders, colour concentrates and liquids. Pigments and dyes produce colour in a resin due to the selective absorption of visible light with wavelengths ranging from ~380 (violet) to 760 nm (red). Because the dyes are in solution, the colour is produced only from light absorption, and the material is transparent. The colour shade is dependent on the particle size of the pigment. Ultramarine blue pigments are nonreflective due to their refractive index, which is similar to that of the polymer. Pigments must be adequately dispersed in the polymer for optimum scattering. Specks and uneven colouration can result from incomplete dispersion. If agglomerates are present, they may adversely affect the mechanical properties (i.e., tensile strength, impact strength, and flex fatigue) of the resulting product. The pigments should be compatible with the polymer because poor compatibility may result in part failure. High temperature during processing can damage or destroy the pigment resulting in shade variation or colour loss. The thermal sensitivity is related to both temperature and duration of exposure. Long cycles during injection moulding and rotational moulding can have a more adverse effect than high-speed extrusion. Some pigments can act as nucleating agents altering the mechanical properties and improving the clarity of the polymer. The most common inorganic pigments include oxides, sulphides, hydroxides, chromates, and other complexes based on metals, such as cadmium, zinc, titanium, lead, and molybdenum. In general, they are more thermally stable than organic pigments and are more opaque and resistant to migration, chemicals, and fading. They can cause wear on the processing equipment (e.g., extrusion machine screws and barrels). The use of heavy metal compounds (e.g., cadmium) has been restricted due to toxicity issues. The most widely used white pigment is titanium oxide (rutile),which is used either alone or in combination with other colourants to control the opacity and produce pastel shades. Other white pigments include zinc oxide, zinc sulphide, and lead carbonate (lead white).The most widely used black pigment is carbon black, which is essentially composed of pure carbon. When combined with white pigments, it produces various shades of grey depending on the particle size and tinting strength of the carbon black grade. Iron oxide, Fe3O4, is another black pigment that has a lower thermal stability and tinting strength. A variety of inorganic compounds are used as pigments to colour polymers. Drobny (Drobny, 2007) classified the most common types of pigments. The yellow pigments include chrome yellow, chrome-titanium yellow, iron oxides, and lead chromates. The orange pigments include molybdate orange and cadmium orange. The brown pigments include iron oxide or a combination of chrome/iron oxides. The red pigments include iron oxide and cadmium sulphide/selenide. The blue pigments include ultramarine (aluminosilicate with sodium ion and ionic sulfur groups) and mixed metal oxides that are primarily based on cobalt aluminate. The green pigments include chrome oxide and cobalt-based mixed oxides. In general, the organic pigments are brighter, stronger, and more transparent than inorganic pigments but are not as light resistant. They may be partially soluble in many polymers but exhibit a much greater tendency to migrate. The largest group of organic pigments are the azo pigments, which contain one or more azo chromophoric groups and form yellow, orange, and red pigments. Monoazo pigments, which have only one chromophore, exhibit low thermal and light stability and have a tendency to bleed. In addition, the monoazo pigments are not typically employed in plastics. Polyazo pigments, which have more than one chromophore, do not tend to bleed and have better thermal stability and excellent chemical stability. Nonazo pigments have a variety of structures including polycyclic and metal complexes. Phtalocyanine blues and greens, most of which are complexed with copper, are highly stable to light, heat, and chemicals and form highly transparent, intense colours with a high tinting strength. Other organic pigments include quinacridones (red, violet, orange), dioxazines (violet), isoindolines (yellow, orange, red), perylenes, flavanthrones, and anthraquinones. In addition, thiazine pigments are suitable for almost all organic pigment applications (Smith, 2002). In low concentrations, thiazine pigments provide high value-in-use due to their heat stability (up to 290 -300 0 C) in polyolefins enabling warpage-/deformation-free HDPE mouldings, high opacity coupled with very high colour saturation, ease of incorporation (dispersion, flow), chemical inertness, high insolubility in nearly all solvents, high performance in the intended applications (processability and durability), ability to form supramolecules with synergistic values, potential environmental compatibility, and cheap cost.

### **3. Experimental**

126 Advanced Elastomers – Technology, Properties and Applications

refractive index, which is similar to that of the polymer. Pigments must be adequately dispersed in the polymer for optimum scattering. Specks and uneven colouration can result from incomplete dispersion. If agglomerates are present, they may adversely affect the mechanical properties (i.e., tensile strength, impact strength, and flex fatigue) of the resulting product. The pigments should be compatible with the polymer because poor compatibility may result in part failure. High temperature during processing can damage or destroy the pigment resulting in shade variation or colour loss. The thermal sensitivity is related to both temperature and duration of exposure. Long cycles during injection moulding and rotational moulding can have a more adverse effect than high-speed extrusion. Some pigments can act as nucleating agents altering the mechanical properties and improving the clarity of the polymer. The most common inorganic pigments include oxides, sulphides, hydroxides, chromates, and other complexes based on metals, such as cadmium, zinc, titanium, lead, and molybdenum. In general, they are more thermally stable than organic pigments and are more opaque and resistant to migration, chemicals, and fading. They can cause wear on the processing equipment (e.g., extrusion machine screws and barrels). The use of heavy metal compounds (e.g., cadmium) has been restricted due to toxicity issues. The most widely used white pigment is titanium oxide (rutile),which is used either alone or in combination with other colourants to control the opacity and produce pastel shades. Other white pigments include zinc oxide, zinc sulphide, and lead carbonate (lead white).The most widely used black pigment is carbon black, which is essentially composed of pure carbon. When combined with white pigments, it produces various shades of grey depending on the particle size and tinting strength of the carbon black grade. Iron oxide, Fe3O4, is another black pigment that has a lower thermal stability and tinting strength. A variety of inorganic compounds are used as pigments to colour polymers. Drobny (Drobny, 2007) classified the most common types of pigments. The yellow pigments include chrome yellow, chrome-titanium yellow, iron oxides, and lead chromates. The orange pigments include molybdate orange and cadmium orange. The brown pigments include iron oxide or a combination of chrome/iron oxides. The red pigments include iron oxide and cadmium sulphide/selenide. The blue pigments include ultramarine (aluminosilicate with sodium ion and ionic sulfur groups) and mixed metal oxides that are primarily based on cobalt aluminate. The green pigments include chrome oxide and cobalt-based mixed oxides. In general, the organic pigments are brighter, stronger, and more transparent than inorganic pigments but are not as light resistant. They may be partially soluble in many polymers but exhibit a much greater tendency to migrate. The largest group of organic pigments are the azo pigments, which contain one or more azo chromophoric groups and form yellow, orange, and red pigments. Monoazo pigments, which have only one chromophore, exhibit low thermal and light stability and have a tendency to bleed. In addition, the monoazo pigments are not typically employed in plastics. Polyazo pigments, which have more than one chromophore, do not tend to bleed and have better thermal stability and excellent chemical stability. Nonazo pigments have a variety of structures including polycyclic and metal complexes. Phtalocyanine blues and greens, most of which are complexed with

#### **3.1. Characterization and preparation pigment and pigment composites**

The silicas used as the support medium of composite pigments in these studies are as follows: Aerosil 380 (Degussa S.A.) – a hydrophilic fumed silica with a specific surface area of 340 - 380 m2/g; Zeosil 175 (Rhodia) – a precipitated silica with a specific surface area of 175 m2/g; Silica Gel 60 (Merck) – a silica with a specific surface area of 470 – 540 m2/g. The indigothiazine pigment was used as the coloured chromophore. Below 180 ºC, the *cis* form of this pigment is yellow in colour and transforms into the highly durable *trans* form characterised by the red colour upon heating (Figure 1). The latter form is highly resistant to organic solvents, high temperatures, and light. The *cis* form of indigothiazine was obtained by according to the description in the patent Bansi Lal (Bansi Lal et al., 2002).

*cis*- indigothiazine *trans*- indigothiazine

**Figure 1.** Isomeric conversion of *cis*- to *trans*–indigothiazine – ([2,2']-bi(1,4-benzothiazynylidene)-3,3'- (4H, 4'H)-dione.

#### **3.2. Characterisation and preparation pigment and pigment composites**

The composite pigments were prepared by mechanically applying *cis*-indigothiazine on silica surfaces and subjecting them 180 ºC for 2 h to induce their transformation into the highly durable *trans* form. The mechanical application of the *cis* form on silica was accomplished by means of a high-speed blender (Warning Commercial) at ~14,000 rpm. The quantity of the pigment was measured relative to the weight of silica, and the modification required 15 min (Fig. 2).

**Figure 2.** Production process scheme of composite pigments (Copyright Lipińska at al.2011).

To assess the mechanism and kinetics of the electrochemical oxidation of the compounds under investigation, cyclic voltammetry (CV) and differential pulse (DPV) methods were employed on an Autolab analytical unit (EcoChemie, Holland). A three-electrode system was employed for the measurements. Platinum was used as the anode and the auxiliary electrode. The potential of the tested electrode was measured versus a ferricinium/ferrocene reference electrode (Fc+/Fc) where the standard potential is defined as zero and independent of the solvent used. Prior to the measurements, all of the solutions were deoxygenated with argon. During the measurements, an argon atmosphere was maintained over the solution. The effect of the scan rate on the electrooxidation of indigothiazine in an anhydrous medium was investigated. The differential scanning calorimeter (DSC) measurements of *cis–* and *trans–*indigothiazine and vulcanisates were performed on a DSC1 calorimeter (Mettler Toledo) with a heating rate of 10 0C/min. The morphology of the indigothiazine particles and the dispersion in the elastomer matrix containing either unmodified silica or the composite pigments was estimated using scanning electron microscopy with a LEO 1530 SEM microscope. The NBR vulcanisates were broken down in liquid nitrogen, and the fracture surfaces of the vulcanisate were examined. Prior to the measurements, the samples were coated with carbon. The prepared pigment composites were characterised to evaluate the agglomerate size distribution in the water medium using dynamic light scattering (Zetasizer Nano S90, Malvern) for a dispersion concentration of 0.2 g/l (the powders were pre-treated with ultrasound for 0.5 h). The specific surface area of the silica fillers before and after application of indigothiazine was determined. The measurements were performed using the Brunauer-Emmet-Teller (BET) nitrogen adsorption method (GEMINI 2360 V2.01). The zeta potential of the silicas and of pigment composite were evaluated (ZETASIZER 2000, the water dispersion with concentration 2 g/l)). The powders were pre-treated with ultrasound for 1 h.

#### **3.3. Preparation and vulcanization of rubber compounds**

The composite pigments were used as the rubber mixture fillers. The rubber mixture composition was made up of the following: NBR (Perbunan 28-45F containing 28 wt % acrylonitryle groups from Lanxess, 100 phr); sulphur crosslinking system (See Table 2) and optional: composite pigments or silica (30 phr).


aphr = parts per hundred of rubber by weight

**Table 2.** Rubber formulations.

128 Advanced Elastomers – Technology, Properties and Applications

required 15 min (Fig. 2).

highly durable *trans* form. The mechanical application of the *cis* form on silica was accomplished by means of a high-speed blender (Warning Commercial) at ~14,000 rpm. The quantity of the pigment was measured relative to the weight of silica, and the modification

**Figure 2.** Production process scheme of composite pigments (Copyright Lipińska at al.2011).

To assess the mechanism and kinetics of the electrochemical oxidation of the compounds under investigation, cyclic voltammetry (CV) and differential pulse (DPV) methods were employed on an Autolab analytical unit (EcoChemie, Holland). A three-electrode system was employed for the measurements. Platinum was used as the anode and the auxiliary The mixture was prepared in a laboratory rolling mill (Roll dimensions: D = 200 mm, L = 450), and the vulcanisation process was carried out in a hydraulic press (160 0C) in line with the curing time specifications (See Table 2). The vulcanisation kinetics was measured with evaluated using a Monsanto vulcameter with an oscillating rotor (ZACH Metalchem) in

compliance with the ISO-4317standard. The tests performed on the vulcanizates are summarised as follows: the mechanical properties were characterised using a Zwick 1435 strength testing machine and according to ISO-37. The crosslink density of vulcanizates was determined by equilibrium swelling in toluene, based on the Flory-Rehner (Flory et al., 1943) equation using the Huggins parameter of elastomer–solvent interaction µ = 0.381 + 0.671Vr [eq. (1)].

$$\upsilon\_{\varepsilon} = -\frac{\ln(1 - V\_r)V\_r + \mu V\_r^2}{\frac{1}{V\_0(V\_r^{\frac{\pi}{3}} - \frac{V\_r}{2})}} \tag{1}$$

where: �� - crosslink density, Vr - volume fraction of elastomer in swollen gel, Vo - molar volume of solvent [mol/cm3].

spectrophotometric data were taken using a KONICA–MINOLTA CM–3600 spectrophotometer under computer control with spectra analysis software (Colour Data Software CM – S 100 in Spectra Magic) according to the following standards: CIE No. 15; PN-EN ISO 105- J01; ASTM E1 164; and DIN 5033 Tell.7, resistance to ageing under UV radiation was tested by exposing the vulcanizates to a UV source (Atlas UV200) for a total of 5 days, equivalent to the ageing by approximately 3 months under normal conditions. The test timeline was as follows:day segment: UV power 0.7 W/m2; temperature 60 ºC; duration 8 h, night segment: UV power 0.0 W/m2; temperature 50 ºC; duration 4 h. The vulcanizates subjected to UV ageing were tested for tensile strength. Spectrophotometric measurements were performed before and after UV irradiation. For each vulcanizate being tested, the ageing coefficient, K, was computed using Equation 2:

$$\mathbf{K} = \frac{(\text{TS} \cdot \text{Eb})\_{\text{after aggregation}}}{(\text{TS} \cdot \text{Eb})\_{\text{before aginging}}} \tag{2}$$

where: K – ageing coefficients; TS – tensile strength; Eb – relative elongation.

#### **4. Results**

The electrode reactions characterising the electrochemical oxidation of indigothiazine at the platinum electrode were studied by cyclic pulse voltammetry. The half-wave potential of the peak in the cyclic voltammogram is characteristic of each subsequent step in the investigated electrode reaction. Selected cyclic voltammograms recorded in a solution consisting of the indigothiazine and the supporting electrolyte are presented in Fig. 3. The cyclic voltammograms recorded for the indigothiazine solution exhibit two peaks, which involve at least three electrode steps of the pigment electrooxidation in the potential range prior to the potential at which electrolyte decomposition is initiated. The supporting electrolyte (in *N*-methyl-*2*-pyrrolidone) has no characteristic peaks except for charging the electrical double layer. The electrooxidation potentials were determined fallowing: the first peak - 0.71 V, second peak - 0.82 V, third peak - 1.01 V. In addition, the electrode reactions of the examined connection are becoming more and more irreparable. Polarisation towards the negative potentials of the reduction reaction was not observed (Fig. 4). The recorded voltammograms, under the linear diffusion of the first electrooxidation step, were employed to determine the half-wave potential (E1/2). Based on the results provided in Table 3, indigothiazine, E1/2 = 0.36 V, was easily oxidised. The energy of the highest filled molecular orbital (EHOMO) was estimated to determine the ease of electron backdonation (potential of the ionisation) and the enthalpy (ΔH). The EHOMO energy (-7.65 eV) and enthalpy (22.83) confirmed the good antioxidant properties and stability.


**Table 3.** Cyclic voltammetry and molecular orbital parameters for the studied compound

130 Advanced Elastomers – Technology, Properties and Applications

ageing coefficient, K, was computed using Equation 2:

0.671Vr [eq. (1)].

**4. Results** 

volume of solvent [mol/cm3].

compliance with the ISO-4317standard. The tests performed on the vulcanizates are summarised as follows: the mechanical properties were characterised using a Zwick 1435 strength testing machine and according to ISO-37. The crosslink density of vulcanizates was determined by equilibrium swelling in toluene, based on the Flory-Rehner (Flory et al., 1943) equation using the Huggins parameter of elastomer–solvent interaction µ = 0.381 +

�� = − �������)������

where: �� - crosslink density, Vr - volume fraction of elastomer in swollen gel, Vo - molar

spectrophotometric data were taken using a KONICA–MINOLTA CM–3600 spectrophotometer under computer control with spectra analysis software (Colour Data Software CM – S 100 in Spectra Magic) according to the following standards: CIE No. 15; PN-EN ISO 105- J01; ASTM E1 164; and DIN 5033 Tell.7, resistance to ageing under UV radiation was tested by exposing the vulcanizates to a UV source (Atlas UV200) for a total of 5 days, equivalent to the ageing by approximately 3 months under normal conditions. The test timeline was as follows:day segment: UV power 0.7 W/m2; temperature 60 ºC; duration 8 h, night segment: UV power 0.0 W/m2; temperature 50 ºC; duration 4 h. The vulcanizates subjected to UV ageing were tested for tensile strength. Spectrophotometric measurements were performed before and after UV irradiation. For each vulcanizate being tested, the

> K = ������)����� ������ ������)������ ������

The electrode reactions characterising the electrochemical oxidation of indigothiazine at the platinum electrode were studied by cyclic pulse voltammetry. The half-wave potential of the peak in the cyclic voltammogram is characteristic of each subsequent step in the investigated electrode reaction. Selected cyclic voltammograms recorded in a solution consisting of the indigothiazine and the supporting electrolyte are presented in Fig. 3. The cyclic voltammograms recorded for the indigothiazine solution exhibit two peaks, which involve at least three electrode steps of the pigment electrooxidation in the potential range prior to the potential at which electrolyte decomposition is initiated. The supporting electrolyte (in *N*-methyl-*2*-pyrrolidone) has no characteristic peaks except for charging the electrical double layer. The electrooxidation potentials were determined fallowing: the first peak - 0.71 V, second peak - 0.82 V, third peak - 1.01 V. In addition, the electrode reactions of the examined connection are becoming more and more irreparable. Polarisation towards the negative potentials of the reduction reaction was not observed (Fig. 4). The recorded voltammograms, under the linear diffusion of the first electrooxidation step, were employed to determine the half-wave potential (E1/2). Based on the

where: K – ageing coefficients; TS – tensile strength; Eb – relative elongation.

����� � � � �� � ) �

(1)

(2)

**Figure 3.** Voltammogram of indigothiazine oxidation on a Pt electrode; c = 1 mmol/dm3 in 0,1 mol/dm3 in *N*-methyl-*2*-pyrrolidone at various polarisation speeds.

**Figure 4.** Voltammogram of indigothiazine for negative potential on the Pt electrode; c = 1 mmol/dm3 in 0,1 mol/dm3 in *N*-methyl-*2*-pyrrolidone.

The thermal strength of *cis –* and *trans–* indigothiazine (Fig.6) was measured applying the DSC method (DSC 1 Mettler Toledo). The resulting diagrams show the melting bands of the pigment's crystalline phase at around 440° C (Fig. 5). The exothermic peak on the *cis–* indigothiazine curve is in the regime where the *cis*- to *trans-* conversion took place most intensively.

**Figure 5.** DSC curves of indigothiazine pigment – *cis* and *trans* conformation.

**Figure 6.** SEM image of trans-indigothiazine pigment .

The fumed silica, Aerosil 380, used in the tests is characterised by the high numerical percentage (99.67 %) of aggregates in the range of 116-330 nm in size; however, these aggregates account for only 28.83 % of the filler's volume in the water medium. The outstanding fraction (71.17 %) consists of agglomerates whose sizes are in the range of 1056- 3000 nm. The application of indigothiazine pigment on the surface of Aerosil 380 silica reduced the average particle size and fragmented agglomerates, effectively producing the composite pigments with sizes ranging from 82-131 nm (100 %). For the precipitated silica, Zeosil 175, the application of pigment helped produce the composite pigments with 99.92 % of the aggregates in the size range from 116-262 nm and the rest the size range of 2379-3000 nm. Although the modification of Zeosil 175 reduced the average particle size and increased the particles uniformity, the particles still formed agglomerates in the order of a few µm in size. For Silica Gel 60, applying indigothiazine only slightly increased the particles tendency to agglomerate; that is, the aggregates size increased from 104-147 nm (unmodified) to 116- 208 nm (modified). The results are summarized in Table 4.

132 Advanced Elastomers – Technology, Properties and Applications

**Figure 5.** DSC curves of indigothiazine pigment – *cis* and *trans* conformation.

**Figure 6.** SEM image of trans-indigothiazine pigment .

intensively.

The thermal strength of *cis –* and *trans–* indigothiazine (Fig.6) was measured applying the DSC method (DSC 1 Mettler Toledo). The resulting diagrams show the melting bands of the pigment's crystalline phase at around 440° C (Fig. 5). The exothermic peak on the *cis–* indigothiazine curve is in the regime where the *cis*- to *trans-* conversion took place most


**Table 4.** Aggregates size and specific surface area of the silicas and composite pigments.

The specific surface area was measured using the BET nitrogen adsorption method before and after applying the indigothiazine pigment. Each composite pigment was characterised by reduced specific surface area as compared with the unmodified silicas. The decrease in the specific surface proceeded to a larger extent in the case of Aerosil 380 and Silica Gel 60 with a large specific surface.

Figures 7-9 shows the change in zeta potential as a function of pH for both unmodified silicas and composite pigments. The largest change in the surface charge was caused by the application of indigothiazine to Zeosil 175 silica. The difference in zeta potential between the composite and unmodified silica Zeosil 175 was approximately 30 mV (pH 4), indicating that the modification reduced the acidic characteristics of the Zeosil 175 surface, which is a favourable result for its use as a polymer composite filler. For the A 380/ indigothiazine pigment, the zeta potential was slightly higher than for the unmodified Aerosil 380. For Silica Gel 60 and the Silica Gel 60/ indigothiazine composite, the two zeta potential curves overlapped one another, indicating that the application of the pigment induced practically no change to the surface charge in this filler.

**Figure 7.** Zeta potential diagram for the composite pigments in which the supports was Aerosil 380.

**Figure 8.** Zeta potential diagram for the composite pigments in which the support was Zeosil 175.

134 Advanced Elastomers – Technology, Properties and Applications



Zeosil 175

Zeosil 175/indigothiazine

Aerosil 380/indigothiazine





0

10

**Figure 7.** Zeta potential diagram for the composite pigments in which the supports was Aerosil 380.

0 2 4 6 810 12

0 2 4 6 810 12

pH

pH

**Figure 8.** Zeta potential diagram for the composite pigments in which the support was Zeosil 175.

**Figure 9.** Zeta potential diagram for the composite pigments in which the support was Silica Gel 60.

The effects of silicas and the composite pigments on the properties of rubber mixtures were evaluated with respect to the curing kinetics and tensile strength of the vulcanizates (Table 5). Rheometric tests showed that adding composite pigments to the rubber mixture did not significantly affect the vulcanisation time, t09, torque increase, Δ*M*, and the initial viscosity values (*M*min) in comparison to the mixtures containing unmodified silicas. The only exception was the Zeosil 175/ indigothiazine composite, which showed a lower torque increase than the original Zeosil 175. Tg values from DSC curves for vulcanizates filled modified and unmodified silicas have similar glass transitios temperature in range -28 0C – 30 0C (Table 5) The vulcanisation process produced red-coloured vulcanizates, which were then subjected to tensile strength tests. The results showed that the application of the composite pigments had a similar effect on the mechanical properties of the vulcanizates as did the application of the unmodified fillers, as indicated by the tensile strength TS (Table 5) and the Scanning Electron Microscope SEM observations. The SEM images revealed agglomerates in the vulcanizates filled with unmodified silicas (Figs. 11, 13, 15) and composite pigments (Figs. 12, 14, 16). Even though the modification of the test fillers reduced the average size of agglomerates in the water medium, it did not significantly affect the dispersion level of silicas in the elastomer. The vulcanizates containing modified silicas had also similar values of curing density, module values at 100% stress and elongations at break EB in comparison to composites filled unmodified silicas.

**Figure 10.** Ageing coefficient value (K) of vulcanizates filled with composte pigments and unmodified silicas.

Next, the functional properties of the vulcanizates after undergoing the 120 h of UV ageing were re-examined. After ageing, all samples exhibited lower TS and EB values and a higher curing density; however, the vulcanizates containing either the A 380/ indigothiazine, Z 175/ indigothiazine and S60/ indigothiazine composites exhibited less UV degradation. Based on the K values (Fig.10), it is possible to conclude that the pigment protects against ageing most effectively when applied to the A380 and Z175 silica supports. The closer the K value is to unity, the less ageing is considered to have taken place. Colorimetric tests of the aged vulcanizates were also performed using the non-aged samples as benchmarks (Table 6). The colorimetric measurements were taken in the CIE-Lab space that expressed the colours in the Cartesian L-a-b coordinate system of the brightness L, the red-greenness a, and the yellow-blueness b. The L value ranges from 0 (black) to 100 (white), and all shades of grey are expressible by the values in between. The colorimetric data reveal that the UV radiation not only degrades the functional properties but also changes the colour of the vulcanizates (Table 6). The parameter dE\*ab depicts the degree of change in colour as a result of ageing. All the vulcanizates containing composite pigments exhibited a lesser degree of change in colour than those containing unmodified silica. The vulcanizates with either the A 380/ indigothiazine composite or the Silica Gel 60/ indigothiazine composite exhibited three times less the degree of change in colour than those with ordinary silicas.


silicas.

**Figure 10.** Ageing coefficient value (K) of vulcanizates filled with composte pigments and unmodified

NBR/A380

azine NBR/Z175

azine NBR/S60

**K**

NBR/A380/indigothi

NBR/Z175/indigothi

0 0.2 0.4 0.6 0.8 1

Next, the functional properties of the vulcanizates after undergoing the 120 h of UV ageing were re-examined. After ageing, all samples exhibited lower TS and EB values and a higher curing density; however, the vulcanizates containing either the A 380/ indigothiazine, Z 175/ indigothiazine and S60/ indigothiazine composites exhibited less UV degradation. Based on the K values (Fig.10), it is possible to conclude that the pigment protects against ageing most effectively when applied to the A380 and Z175 silica supports. The closer the K value is to unity, the less ageing is considered to have taken place. Colorimetric tests of the aged vulcanizates were also performed using the non-aged samples as benchmarks (Table 6). The colorimetric measurements were taken in the CIE-Lab space that expressed the colours in the Cartesian L-a-b coordinate system of the brightness L, the red-greenness a, and the yellow-blueness b. The L value ranges from 0 (black) to 100 (white), and all shades of grey are expressible by the values in between. The colorimetric data reveal that the UV radiation not only degrades the functional properties but also changes the colour of the vulcanizates (Table 6). The parameter dE\*ab depicts the degree of change in colour as a result of ageing. All the vulcanizates containing composite pigments exhibited a lesser degree of change in colour than those containing unmodified silica. The vulcanizates with either the A 380/ indigothiazine composite or the Silica Gel 60/ indigothiazine composite exhibited three times less the degree of change in colour than those with ordinary silicas.

*M*min. – minimum torque [dNm], Δ*M* – torque increase [dNm], *t09* – vulcanisation time [min]**,** *SE100* – stress at 100 % elongation [MPa], *TS* – tensile strength [MPa], *EB* – relative elongation at break [%], *ν<sup>T</sup>* – curing density of vulcanizates calculated using the Flory Rehner equation [mol/cm3]

**Table 5.** Effects of composite pigments and unmodified silicas on selected properties of NBR.


L\* – brightness parameter, a\* – redness – greenness reading, b\* – yellowness – blueness reading, dl\*, da\*, db\* – the difference in colour between trial and standard readings; dE\*ab – vector of overall errors in colour and colour darkness;

**Table 6.** Colorimetric data for the vulcanizates containing unmodified silicas or composite pigment.

**Figure 11.** SEM image of NBR vulcanizate filled with Aerosil 380.

**Figure 12.** SEM image of NBR vulcanizate filled with A 380/indigothiazine.

**Figure 13.** SEM image of NBR vulcanizate filled with Zeosil 175.

**Figure 14.** SEM image of NBR vulcanizate filled with Z 175/indigothiazine.

**Figure 11.** SEM image of NBR vulcanizate filled with Aerosil 380.

**Figure 12.** SEM image of NBR vulcanizate filled with A 380/indigothiazine.

**Figure 13.** SEM image of NBR vulcanizate filled with Zeosil 175.

**Figure 15.** SEM image of NBR vulcanizate filled with Silica Gel 60.

**Figure 16.** SEM image of NBR vulcanizate filled with S60/indigothiazine.

### **5. Conclusions**

Indigothiazine exhibited antioxidant properties, which allows them to be successfully employed as anti-ageing substances. Red-colour composite pigments were produced by modifying silica Aerosil 380, Zeosil 175, and Gel 60. The modification method employed in the present work eliminated harmful solvents employed in the conventional process. The newly developed method is much simpler, cheaper, and easier to use on an industrial scale compared to the conventional process, which requires a solvent. The silica modification process using the blender reduced the specific surface area of the tested fillers. For the composites with Aerosil 380, Zeosil 175 and Silica Gel 60, this process reduced the agglomerate sizes. For the samples with Aerosil 380 and Zeosil 175, the zeta potential also increased, which indicated that their surface became less polarised and their acceptor characteristic was partially lost. The addition of the composite pigments to the polymer mixtures did not significantly change their curing kinetics or processing properties. The vulcanisation of NBR in the presence of the modified silicas resulted in vulcanisates with a vivid red colour were produced. Vulcanisates filled with the composite pigments and those filled with unmodified silicas exhibited similar mechanical strengths. The strength tests after UV ageing of the vulcanisates showed that the A 380/ indigothiazine composite and the Zeosil 175/ indigothiazine composite possess antiageing properties and have a favourable impact on the plastic's mechanical properties and colour stability. The present work led to the successful production of vulcanisates with enhanced ageing resistance.

### **Author details**

Anna Marzec and Marian Zaborski *Institute of Polymer and Dye Technology, Technical University of Lodz, Poland* 

#### **6. References**


Bansi Lal, K. & Bruno, P. (2002). Production of thiazine-indigo pigments. 6339084, United States.

140 Advanced Elastomers – Technology, Properties and Applications

Indigothiazine exhibited antioxidant properties, which allows them to be successfully employed as anti-ageing substances. Red-colour composite pigments were produced by modifying silica Aerosil 380, Zeosil 175, and Gel 60. The modification method employed in the present work eliminated harmful solvents employed in the conventional process. The newly developed method is much simpler, cheaper, and easier to use on an industrial scale compared to the conventional process, which requires a solvent. The silica modification process using the blender reduced the specific surface area of the tested fillers. For the composites with Aerosil 380, Zeosil 175 and Silica Gel 60, this process reduced the agglomerate sizes. For the samples with Aerosil 380 and Zeosil 175, the zeta potential also increased, which indicated that their surface became less polarised and their acceptor characteristic was partially lost. The addition of the composite pigments to the polymer mixtures did not significantly change their curing kinetics or processing properties. The vulcanisation of NBR in the presence of the modified silicas resulted in vulcanisates with a vivid red colour were produced. Vulcanisates filled with the composite pigments and those filled with unmodified silicas exhibited similar mechanical strengths. The strength tests after UV ageing of the vulcanisates showed that the A 380/ indigothiazine composite and the Zeosil 175/ indigothiazine composite possess antiageing properties and have a favourable impact on the plastic's mechanical properties and colour stability. The present work led to the successful production of vulcanisates with

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Natinee, L., Dolmalik, J. & Manus, S. (2011). Hybridized reinforcement of natural rubber with silica-modified short cellulose fibers and silica. *J. Appl. Polym. Sci*., 120, 3242-

containing mesoporous silicate. *Rev. Mineral. Geochem*., 141, 77–80.

modified with 3-aminopropyltriethoxysilane. *Colloids Surf.,* 173, 73–84.

	- Jung-Woo, P., Jun, P. Y. & Chul-Ho, J. (2011). Post-grafting of silica surfaces with prefunctionalized organosilanes: new synthetic equivalents of conventional trialkoxysilanes. *Chem Commun*., 47, 4860-4871.

## **Smart Elastomers**

144 Advanced Elastomers – Technology, Properties and Applications

trialkoxysilanes. *Chem Commun*., 47, 4860-4871.

Jung-Woo, P., Jun, P. Y. & Chul-Ho, J. (2011). Post-grafting of silica surfaces with prefunctionalized organosilanes: new synthetic equivalents of conventional

**Chapter 6** 

## **Microstructure and Properties of Magnetorheological Elastomers**

Anna Boczkowska and Stefan Awietjan

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/50430

### **1. Introduction**

Magnetorheological elastomers (MREs) belong to the new group of the functional materials called "smart". Although smart materials are known since long time, their intensive development started in the end of the XXth century. The term *smart materials, intelligent materials* or less frequently used *adaptive materials* or *multifunctional materials,* was introduced in the eighties of the twentieth century, when some materials, which were included in the group were already known. Till today there is no accepted universal definition of smart material, it is also not included in the encyclopedia devoted to these materials, published in 2002 [1, 2].

The term smart material generally refers to material which changes its properties under the influence of various external stimuli [3]. However, it seems that this approach is too simplistic. Much more precise description of smart material is given by Takagi in his work [4]. According to him, smart material is capable of reacting to external stimuli by changing its material properties for the desired way and effectively responds to these stimuli. Such material should therefore be some kind of sensor, processor, and actuator. These attributes should be looped, and the effect of changing the properties of the material should be done in real time.

Among smart materials, magnetorheological materials (MR) are an important group. They are a class of materials with rheological properties rapidly varied by the application of a magnetic field (Fig. 1). The change in their properties is in the proportion to the magnitude of the magnetic field applied and is immediately reversible. The Bingham plastic model is often used to model the behaviour of MR materials [1].

Research focussing on magnetorheological materials was initiated by Jacob Rabinow at the National Bureau of Standards, now the National Institute for Science and Technology, in the

© 2012 Boczkowska and Awietjan, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

United States in the 1940s and early 1950s. His early works on magnetorheological fluids led to a host of devices and products based on dry magnetic powders, for example the magnetic, powder brake. In addition to MR fluids, the MR materials also contain magnetic field responsive gels [5], foams, powders, and elastomers. The comparison of the most typical magnetorheological materials is shown in Table 1.

**Figure 1.** Magnetorheological material A- before, B – and after the application of an external magnetic field.


**Table 1.** Composition and properties of typical magnetorheological materials [6].

The interest in magnetorheological elastomers (MREs) has recently increased because of their prospects for application in various smart systems. They are still much less known than magnetorheological fluids. The MREs are smart materials, analogues of the magnetorheological fluids (MRFs), in which the fluid component is replaced by a crosslinked material like rubber or silicone. They consist of micron sized magnetically permeable particles in non-magnetic matrix. In a similar way, as in the case of MRFs, the particles tend to align themselves in the direction of the magnetic field, but after curing of the matrix such microstructure is fixed. Magnetorheological materials change their rheological properties under the influence of an external magnetic field. Their rheological properties can be changed continuously, rapidly and reversibly by changing an external magnetic field. This behaviour is described by the magnetorheological effect.

In this chapter the overview on magnetorheological elastomers is given, as well as basic knowledge about rheology and magneto-rhelogy. Also the effect of the amount, size and orientation of the particles on the microstructure and properties of MREs is discussed, as well as the influence of magnetic field on the compressive characteristics and rheological properties of developed MREs.

## **2. Magnetorheological elastomers**

The MREs consist of magnetically permeable particles (such as iron or other ferromagnetic particles) added to a viscoelastic polymeric material prior to crosslinking. The MREs contain ferromagnetic particles having sizes from few to few hundreds of μm. Pure iron has the highest saturation magnetization of known elements and it has also high permeability and low remanent magnetization, providing high, short-term inter-particle attraction. It is known from available literature that the amount and the shape of the particles, as well as the type of the polymer matrix influence the MREs properties.

MREs can be classified according to several parameters like: particles type, matrix, structure and distribution of particles [7].

MREs classification:

148 Advanced Elastomers – Technology, Properties and Applications

field.

Viscosity without external magnetic field [mPa∙s]

Change of properties in an external magnetic field

typical magnetorheological materials is shown in Table 1.

United States in the 1940s and early 1950s. His early works on magnetorheological fluids led to a host of devices and products based on dry magnetic powders, for example the magnetic, powder brake. In addition to MR fluids, the MR materials also contain magnetic field responsive gels [5], foams, powders, and elastomers. The comparison of the most

**Figure 1.** Magnetorheological material A- before, B – and after the application of an external magnetic

100 - 1000 2 - 200 none

The interest in magnetorheological elastomers (MREs) has recently increased because of their prospects for application in various smart systems. They are still much less known than magnetorheological fluids. The MREs are smart materials, analogues of the magnetorheological fluids (MRFs), in which the fluid component is replaced by a crosslinked material like rubber or silicone. They consist of micron sized magnetically permeable particles in non-magnetic matrix. In a similar way, as in the case of MRFs, the particles tend to align themselves in the direction of the magnetic field, but after curing of the matrix such microstructure is fixed. Magnetorheological materials change their rheological properties under the influence of an external magnetic field. Their rheological properties can be changed continuously, rapidly and reversibly by changing an external magnetic field. This

**Ferrofluid Magnetorheological** 

surfactants none

Relative viscosity

Δη/η ~1

**elastomer MRE** 

Storage modulus ΔG' ~20 kPa

**Magnetorheological** 

Particles type iron magnetite iron

Yield stress τy up to

**Table 1.** Composition and properties of typical magnetorheological materials [6].

Particles size 0,1 – 10 μm 2 –10 nm 10 – 50 μm Carrier oils oils, water elastomers Volume fraction 0,1 – 0,5 0,02 – 0,2 0,1 – 0,5

**fluid MRF** 

Additives surfactants, thicsotropic

agents

~100 kPa

behaviour is described by the magnetorheological effect.

	- a. Soft magnetic particles
	- b. Hard magnetic particles
	- c. Magnetostrictive particles
	- d. Magnetic shape-memory particles
	- a. Solid matrix
	- b. Porous matrix
	- a. Isolating matrix
	- b. Conductive matrix
	- a. Isotropic
	- b. Anisotropic

Usually magnetic field is applied to the polymer composites during crosslinking of the matrix. Such treatment locks the columnar particle chain structures during the final cure giving special anisotropic properties. The formation of columnar particle structures within the elastomers corresponds to a low dipolar energy state. Shearing of the cured composite under applied magnetic field requires additional energy because of the particles displacement from this low energy state. The amount of work required and the field dependent shear modulus increase steadily with the applied magnetic field [8, 9, 10].

Changes of the properties in MREs under the influence of the magnetic field depend strongly on the microstructure formed during the curing of elastomer matrix, also as a result

of magnetic field. Interactions between particles in a magnetic field bring them closer, resulting in increased stiffness of the material. This changes the stress-strain curves [11]. The magnetic field causes a shift towards higher stresses. The same mechanism is responsible for the effect of "magnetostriction", shortening of length of the cylindrical sample in the presence of an external magnetic field [12], as schematically shown in Figure 2A, or its extension, the MRE when the particles are distributed isotropically [13]. Under the influence of the field the particles move in the direction of the field, as shown in Figure 2B. This effect is possible when the adhesion between the particles and the matrix is large enough so that the movement of particles leads to deformation of the elastomeric matrix, which is has such low stiffness that such deformations can occur. A similar phenomenon occurs in the MREs, in which the particles are spherical and have an elongated shape [14]. Then, under the influence of the magnetic field they become dipoles, rotating in the direction of the magnetic field vector, which also leads to deformation of the elastomer and thus increase in the length of the sample (Fig. 2C). In all cases, these changes are rapid and fully reversible.

**Figure 2.** Magnetostriction effect in MREs under external magnetic field.

Under the influence of magnetic field rheological properties of MREs change [15], mainly shear modulus, but also the modulus of elasticity determined in a compression test [11]. These properties are strongly dependent on a magnetic field strength. Shearing of MREs in the presence of a magnetic field displaces the particles from the position of minimum energy state, which requires additional work increasing monotonically with increasing magnetic field strength. Therefore, shear modulus depends on the field and is a characteristic feature of the MREs [10, 16]. Changes of the module in the magnetic field also depend on the content of ferromagnetic particles in the elastomer matrix.

The vast majority of MREs described in the literature contain soft magnetic particles, mainly iron, cobalt and their oxides [17,18,19,20]. There is a group of magnetorheological elastomers with hard magnetic fillers like BaFe12O19 or SrFe12O19 [21]. These materials, like permanent magnets remain magnetized after turning off the external magnetic field [22,23]. For the fabrication of MREs also magnetostrictive particles are used, usually Terfenol D - a material with giant magnetostriction [24,25] and the particles with magnetic shape memory (MSM), such as Ni-Mn-Ga [26,27,28].

150 Advanced Elastomers – Technology, Properties and Applications

of magnetic field. Interactions between particles in a magnetic field bring them closer, resulting in increased stiffness of the material. This changes the stress-strain curves [11]. The magnetic field causes a shift towards higher stresses. The same mechanism is responsible for the effect of "magnetostriction", shortening of length of the cylindrical sample in the presence of an external magnetic field [12], as schematically shown in Figure 2A, or its extension, the MRE when the particles are distributed isotropically [13]. Under the influence of the field the particles move in the direction of the field, as shown in Figure 2B. This effect is possible when the adhesion between the particles and the matrix is large enough so that the movement of particles leads to deformation of the elastomeric matrix, which is has such low stiffness that such deformations can occur. A similar phenomenon occurs in the MREs, in which the particles are spherical and have an elongated shape [14]. Then, under the influence of the magnetic field they become dipoles, rotating in the direction of the magnetic field vector, which also leads to deformation of the elastomer and thus increase in the length

of the sample (Fig. 2C). In all cases, these changes are rapid and fully reversible.

**Figure 2.** Magnetostriction effect in MREs under external magnetic field.

content of ferromagnetic particles in the elastomer matrix.

Under the influence of magnetic field rheological properties of MREs change [15], mainly shear modulus, but also the modulus of elasticity determined in a compression test [11]. These properties are strongly dependent on a magnetic field strength. Shearing of MREs in the presence of a magnetic field displaces the particles from the position of minimum energy state, which requires additional work increasing monotonically with increasing magnetic field strength. Therefore, shear modulus depends on the field and is a characteristic feature of the MREs [10, 16]. Changes of the module in the magnetic field also depend on the The matrix of magnetorheological elastomers can be solid or porous. MREs with porous matrix, also called magnetorheological foams, have foamed matrix in order to increase the abilities to change their properties. Majority of magnetorheological elastomers have a matrix with electrically insulating properties. In some cases magnetorheological elastomers, which were doped with electrically conductive particles such as graphite or silver and their percolation threshold has been reached, become electric conductors [29]. Conductive polymers such as polypyrole, polyacetylene, polyaniline can also be used for matrix [30,31]. However, using a conductive matrix has no significant effect on changes in the rheological or mechanical properties in a magnetic field, and raises costs.

The spatial distribution of the particles is determined during curing process. In the presence of a magnetic field during curing, the obtained elastomer has anisotropic, oriented structure, consisting of chains of particles [10,6,7,8,32,33,34,35,36] as schematically shown in Figure 3B. Magnetorheological elastomer cured without the magnetic field allows to receive isotropic structure, which is shown on a schematic diagram in Figure 3A [11,20,21,37,38,39,40,41,42].

**Figure 3.** Schematic microstructure of MRE: A - isotropic, B - anisotropic spatial distribution of particles.

The distribution of particles in MREs microstructure is influenced by many factors, such as the magnetic interaction forces between particles, orientation and magnetic field strength, sample size, the volume fraction of particles, temperature. Understanding the influence of all these factors on the MREs microstructure is important, but extremely difficult for both experimental and computational methods.

## **3. The rheological properties and magnetorheological effect of magnetorheological elastomers**

Many materials, especially polymers and their composites are characterized by viscoelastic properties [43]. This means that they combine the features of elastic solids and viscous liquids, as schematically shown in Figure 4. Their behavior is between the ideal solid described by Hooke's law, in which the stress is always directly proportional to the strain and is independent of strain rate, and a viscous liquid, in which according to Newton's law, stress is always directly proportional to the strain rate and does not depend on the strain. Viscoelastic materials under rapid deformation behave more like elastic body, and under very slow - as viscous liquid. Rheology describes the flow and deformation of solids and liquids under the influence of an external forces.

**Figure 4.** Viscoelastic properties of materials.

For a perfectly elastic solid when the force is applied, strain occurs immediately, and it is linearly proportional to the applied force. The ratio of stress to strain is a measure of elasticity of the material. After unloading deformed body immediately returns to its initial state. The applied force can cause the shear stress (τ) and shear modulus (G). G modulus determines the resistance of the solid to deformation and is expressed by the ratio of shear stress (τ) to the shear strain (γ):

$$G = \frac{\pi}{\mathcal{Y}}.\tag{1}$$

For an elastic solid, both stress and strain are independent on time. For the viscoelastic solid, rheological parameters are dependent on the time and described by Kelvin-Voight model for linear viscoelasticity:

$$
\tau = \mathbf{G} \cdot \boldsymbol{\gamma} + \eta \cdot \frac{d\boldsymbol{\gamma}}{dt} \,\tag{2}
$$

where: η – dynamic viscosity, t – time.

Deformed material can undergo relaxation when the applied force is maintained, which results in a decrease in the stress in time, until its complete disappeared, as schematically shown in Figure 5. When the force is removed, the disappearance of deformation is delayed. This delay is given by the relaxation time λ:

Microstructure and Properties of Magnetorheological Elastomers 153

$$
\lambda = \frac{\eta}{G} \tag{3}
$$

In the area of linear viscoelasticity delay times during creep and recovery are the same. Usually solids are more complicated and to describe the viscoelastic behavior during the creep and recovery it is necessary to use the whole spectrum of relaxation times. In most testing methods of viscoelastic materials, instead of constant stress, dynamic strain measurements in the form of an oscillating sinusoidal function of time is used (Fig. 6):

**Figure 5.** Stress relaxation (τ) at time (t).

152 Advanced Elastomers – Technology, Properties and Applications

liquids under the influence of an external forces.

**Figure 4.** Viscoelastic properties of materials.

stress (τ) to the shear strain (γ):

for linear viscoelasticity:

η – dynamic viscosity,

This delay is given by the relaxation time λ:

where:

t – time.

**magnetorheological elastomers** 

**3. The rheological properties and magnetorheological effect of** 

Many materials, especially polymers and their composites are characterized by viscoelastic properties [43]. This means that they combine the features of elastic solids and viscous liquids, as schematically shown in Figure 4. Their behavior is between the ideal solid described by Hooke's law, in which the stress is always directly proportional to the strain and is independent of strain rate, and a viscous liquid, in which according to Newton's law, stress is always directly proportional to the strain rate and does not depend on the strain. Viscoelastic materials under rapid deformation behave more like elastic body, and under very slow - as viscous liquid. Rheology describes the flow and deformation of solids and

For a perfectly elastic solid when the force is applied, strain occurs immediately, and it is linearly proportional to the applied force. The ratio of stress to strain is a measure of elasticity of the material. After unloading deformed body immediately returns to its initial state. The applied force can cause the shear stress (τ) and shear modulus (G). G modulus determines the resistance of the solid to deformation and is expressed by the ratio of shear

> *G* .

For an elastic solid, both stress and strain are independent on time. For the viscoelastic solid, rheological parameters are dependent on the time and described by Kelvin-Voight model

, *<sup>d</sup> <sup>G</sup> dt* 

Deformed material can undergo relaxation when the applied force is maintained, which results in a decrease in the stress in time, until its complete disappeared, as schematically shown in Figure 5. When the force is removed, the disappearance of deformation is delayed.

 

(1)

(2)

$$
\tau = \tau\_0 \cdot \sin(o \cdot t) \tag{4}
$$

where: τ0 - applied stress, ω - angular velocity [1/s or rad/s],

$$
\alpha = 2\pi \cdot f,\text{ f } - \text{ frequency} \left[ \text{Hz} \right]. \tag{5}
$$

**Figure 6.** Dynamic measurements: a - deformation of the angle φ, b - sinusoidal strain or stress.

The course of deformation and the induced stress in the perfectly elastic solid is in-phase, as shown schematically in Figure 7. However, in the viscoelastic material phase shift angle δ occurs in the range 0 - 90°, between the stress and strain that induces it (Fig. 8). It is the result of delayed reaction of the material to an applied strain. The response to an applied strain consistent with the phase is called the elastic (material entirely elastic - behaves according to Hooke's law), shifted in phase by 90° - viscous (material entirely viscous behaves like a Newtonian fluid), and between 0 and 90° - viscoelastic (viscoelastic material).

**Figure 7.** Dynamic measurement: strain and stress in an elastic solid.

**Figure 8.** Dynamic measurement: strain and stress in the viscoelastic body, δ - phase shift angle.

To combine the elastic and viscous properties the complex shear modulus G\* has been introduced. It is the total resistance of the solid to the applied strain and is defined as:

$$G^\* = \frac{\tau\_0}{\mathcal{V}\_0}.\tag{6}$$

For viscoelastic materials, both complex modulus and phase angle is frequency dependent. For easy distinction between elastic and viscous behavior of the material complex numbers were introduced, so that the complex shear modulus can be described as follows:

$$\text{G}^\* = \text{G}^\* + i\text{G}^\*,\tag{7}$$

where:

154 Advanced Elastomers – Technology, Properties and Applications

**Figure 7.** Dynamic measurement: strain and stress in an elastic solid.

**Figure 8.** Dynamic measurement: strain and stress in the viscoelastic body, δ - phase shift angle.

introduced. It is the total resistance of the solid to the applied strain and is defined as:

To combine the elastic and viscous properties the complex shear modulus G\* has been

\* 0 0 *G* . 

(6)

material).

The course of deformation and the induced stress in the perfectly elastic solid is in-phase, as shown schematically in Figure 7. However, in the viscoelastic material phase shift angle δ occurs in the range 0 - 90°, between the stress and strain that induces it (Fig. 8). It is the result of delayed reaction of the material to an applied strain. The response to an applied strain consistent with the phase is called the elastic (material entirely elastic - behaves according to Hooke's law), shifted in phase by 90° - viscous (material entirely viscous behaves like a Newtonian fluid), and between 0 and 90° - viscoelastic (viscoelastic

G'- storage modulus, or elastic (the real part of complex modulus G\*), G"- loss modulus, or viscous (imaginary part of complex modulus G\*).

The term storage modulus indicates that the strain energy is temporarily stored during deformation, and can be later recovered, and the loss modulus term means that the energy is irretrievably lost and converted to shear heat. Modules G' and G" are described by the following relations:

$$\boldsymbol{G}^{\cdot} = \boldsymbol{G}^{\ast} \cdot \cos \delta = \frac{\tau\_0}{\varkappa\_0} \cdot \cos \delta \,\tag{8}$$

$$\boldsymbol{G}^\* = \boldsymbol{G}^\* \cdot \sin \delta = \frac{\tau\_0}{\gamma\_0} \cdot \sin \delta. \tag{9}$$

Storage modulus G' and loss modulus G" are parameters describing the rheological properties of viscoelastic polymeric materials. Storage modulus G' defined by the ratio of elastic stress to strain, describes the amount of stored (saved) energy during shear, and refers to the elastic properties of the material. Loss modulus G" expressed by the ratio of viscoelastic stress to strain, shows how much energy was dissipated in a cycle of deformation in the form of heat and defines the viscous properties of the material [44].

If the complex modulus value is equal to the storage modulus G\* = G', the loss modulus G'' is equal to zero and the material is completely elastic - the whole cycle of deformation energy is stored and then released, and the deformation is completely reversible. When the complex modulus is equal to the loss modulus G\* = G'', storage modulus G' is equal to zero and the material is completely viscous. Modules G' and G" are dependent on the frequency of oscillation and the temperature. The higher the frequency of oscillation, the higher the value of the modules G' and G". A rise in temperature causes a decrease in storage modulus G' and increase of loss modulus G" [44].

The ratio of loss modulus G" to the storage modulus G' is called the loss angle (damping) and is expressed by the tangent of the phase angle δ [45].

$$\tan(\delta) = \frac{\stackrel{\circ}{G}^{\cdot}}{\stackrel{\circ}{G}^{\cdot}}\tag{10}$$

The loss angle tan(δ) is the ratio between the energy lost and stored during the deformation and determines the viscoelastic material's ability to dissipate (damp) energy [45]. In the

case of completely viscous material value of tan (δ) is greater than 1, and the loss modulus values are higher than the storage modulus G" > G'. When the value of tan(δ) is less than 1 it means supremacy of the elastic properties over the viscous, which means higher values of storage modulus in relation to the loss modulus G' > G". The lower the value of the tan(δ), the more elastic the material. A zero value of tan (δ) indicates that the material is perfectly elastic [46].

To describe the damping properties of the material parameters such as the loss angle tan (δ) and the complex shear modulus G\* are used [47,48]. Damping of vibration or the damping of cyclic deformation is related to a non-reversible absorption of the deformation energy and its conversion into heat that is dissipated by the material [49].

Magnetorheological elastomers are viscoelastic materials, which because of their potential applications are subjected to dynamic loads. Characteristic MREs work area is in the range of small deformations below the yield stress. Under the influence of the magnetic field MREs change their stiffness over a wide range, allowing customization for effective vibration damping. MREs properties are studied and described by the equations typical for viscoelastic materials, with the parameter describing the influence of the external magnetic field on the rheological properties of MREs.

Change in the properties of MREs under the magnetic field is called the magnetorheological (MR) effect and is the most important and most extensively studied property of magnetorheological composites.

The term "magnetorheological" is used only in relation to a group of magnetorheological materials. It describes the reversible properties changes under the magnetic field in comparison to the material properties without the presence of the field. Changes in the rheological properties are usually described by the shear modules: storage G' and loss G" [10,50,51,52,53,54,55]. The magnetorheological effect is also associated with increased stiffness (modulus of elasticity, measured in compression test for example) and the change in damping of magnetorheological materials [11].

Changes in the properties of MREs under the magnetic field, described as the MR effect are related to the magnetic particles tendency to change their position under the influence of an applied magnetic field. The magnetic field induces dipole moments in the ferromagnetic particles, which tend to obtain the positions of minimum energy state. Particles movement introduce deformations in the elastomer matrix, resulting in the increase of shear modulus and stiffness of the MREs. Secondly, the interactions between particles in a magnetic field causes their attraction to each other, which also introduces a deformation in the elastomeric matrix and, consequently, increases the stiffness of the material and shear modules [56]. We can conclude that due to efforts of magnetic particles to change the position in a magnetic field, the elastic matrix is deformed, which leads to an increase in the shear modules (G\*) [20].

The creation of oriented chains of particles in the elastomer leads to a state of minimum magnetic energy [10]. Such MREs under shear deformation in the presence of an external magnetic field losses the low-energy state and requires additional work associated with the movement of entire chains of particles that oppose the deformation and tend to return to equilibrium state. Deformation of the matrix leads to an increase in the value of shear modules G' and G'' and increase in the stiffness of the elastomer. MREs properties strongly depend on the magnetic field strength. Studies of the MREs magnetorheological effect are usually carried out at relatively small frequencies (from 1 to 20 Hz) [8,10,57,58,59,60].

156 Advanced Elastomers – Technology, Properties and Applications

its conversion into heat that is dissipated by the material [49].

field on the rheological properties of MREs.

in damping of magnetorheological materials [11].

an increase in the shear modules (G\*) [20].

magnetorheological composites.

perfectly elastic [46].

case of completely viscous material value of tan (δ) is greater than 1, and the loss modulus values are higher than the storage modulus G" > G'. When the value of tan(δ) is less than 1 it means supremacy of the elastic properties over the viscous, which means higher values of storage modulus in relation to the loss modulus G' > G". The lower the value of the tan(δ), the more elastic the material. A zero value of tan (δ) indicates that the material is

To describe the damping properties of the material parameters such as the loss angle tan (δ) and the complex shear modulus G\* are used [47,48]. Damping of vibration or the damping of cyclic deformation is related to a non-reversible absorption of the deformation energy and

Magnetorheological elastomers are viscoelastic materials, which because of their potential applications are subjected to dynamic loads. Characteristic MREs work area is in the range of small deformations below the yield stress. Under the influence of the magnetic field MREs change their stiffness over a wide range, allowing customization for effective vibration damping. MREs properties are studied and described by the equations typical for viscoelastic materials, with the parameter describing the influence of the external magnetic

Change in the properties of MREs under the magnetic field is called the magnetorheological (MR) effect and is the most important and most extensively studied property of

The term "magnetorheological" is used only in relation to a group of magnetorheological materials. It describes the reversible properties changes under the magnetic field in comparison to the material properties without the presence of the field. Changes in the rheological properties are usually described by the shear modules: storage G' and loss G" [10,50,51,52,53,54,55]. The magnetorheological effect is also associated with increased stiffness (modulus of elasticity, measured in compression test for example) and the change

Changes in the properties of MREs under the magnetic field, described as the MR effect are related to the magnetic particles tendency to change their position under the influence of an applied magnetic field. The magnetic field induces dipole moments in the ferromagnetic particles, which tend to obtain the positions of minimum energy state. Particles movement introduce deformations in the elastomer matrix, resulting in the increase of shear modulus and stiffness of the MREs. Secondly, the interactions between particles in a magnetic field causes their attraction to each other, which also introduces a deformation in the elastomeric matrix and, consequently, increases the stiffness of the material and shear modules [56]. We can conclude that due to efforts of magnetic particles to change the position in a magnetic field, the elastic matrix is deformed, which leads to

The creation of oriented chains of particles in the elastomer leads to a state of minimum magnetic energy [10]. Such MREs under shear deformation in the presence of an external Changes of the properties of magnetorheological composites in the magnetic field are usually described by absolute and relative magnetorheological effect [61].

The absolute magnetorheological (MR) effect is the difference between the maximum value of shear modulus Gmax achieved in a magnetic field, and the value obtained without a magnetic field G0 (called zero field modulus). The absolute magnetorheological effect is described by the equation:

$$
\Delta \mathbf{G} = \mathbf{G}\_{\text{max}} - \mathbf{G}\_0 \left[ \mathbf{M} \mathbf{P} \mathbf{a} \right]. \tag{11}
$$

The relative magnetorheological (MR) effect is the ratio of the absolute magnetorheological effect to the zero field modulus G0, expressed in percents:

$$
\Delta G\_r = \frac{\Delta G}{G\_0} \cdot 100\%. \tag{12}
$$

As it is shown in the literature the absolute and relative MR effects depend on the content of magnetic particles, the frequency of oscillations and the magnetic field strength [20,32,40,42,57,61].

According to reports relative magnetorheological effect can be of 30-40% on average [62], and even 60% for MRE containing 30% vol. iron particles [63]. With the increase in the frequency in the range from 100 to 1400 Hz, shear modulus increases up to four times for MRE containing 27% vol. iron particles. An increase in shear modulus (about 2 MPa) under the influence of the magnetic field of 0.56 T, as well as significant MR effect has been obtained [8,58,60,64].

In most studies, the magnetic field is applied parallel to the chains with iron particles, in order to increase the interaction between particles inside chains [8,57,62].

The MR effect is also influenced by the amplitude of the applied strain, and this is because forces of the magnetic interaction strongly depend on the distance between the dipoles (the iron particles) [20,58]. Magnetorheological elastomers are characterized by a greater strain at which the MR effect is significant than the corresponding MR fluids [10,20,57].

In [65], an increase in storage modulus of 100% in a magnetic field of 670 mT for the MRE based on silicone rubber matrix containing 20 vol.% of iron particles has been reported. The authors of [66] recorded the relative MR effect reaching 878% under 1 T magnetic field (800 kA/m), the MRE was based on silicone rubber (10 wt.%), plasticized with silicone oil (10 wt.%), containing 80 wt.% of particles and cured in 1.5 T magnetic field (1200 kA/m).

Beneficial effects of plasticizers, causing a reduction of material matrix stiffness, on the magnetorheological effect is shown in other papers [20].

Most of the results presented in the literature describes the MREs with isotropic distribution of particles. In [20,61] it was found that larger relative MR effect can be achieved in materials with softer matrix. Therefore, plasticizers are often added in order to increase the relative MR effect. The increase in the relative MR effect is obtained in this case, but the zero field modulus (without magnetic field) is lower.

Also reports can be found about the influence of the anisotropic microstructure on the properties of magnetorheological elastomers. The static and dynamic tensile testing of MREs with particles oriented in the form of chains have shown that they exhibit much higher Young's modulus in comparison to MREs with the isotropically distributed particles [67], which is explained by the presence of column-like structure formed by the particles. Stiffness of the chains depends on the distance between the particles within them, and increases with the applied magnetic field due to the dipole formation and dipole-dipole interactions.

There is also a significant effect of particle size on the MR effect. In order to achieve a significant MR effect, particles should contain at least several magnetic domains [10]. The shape of particles can also have influence on the MR effect, it is better if they have an asymmetric shape with the main axis of anisotropy [63]. Few works have reported positive effect of surface modification of particles on the growth of magnetorheological effect [68], although others show clearly that due to the better adhesion particle-matrix boundary region occurs and the ability of particles to move is reduced, which leads to a reduction of MR effect [69,70].

### **4. Experimental results**

### **4.1. Materials characterisation**

In our study, the selection of composite matrix was carried out to produce MREs with mechanical properties, which can be varied widely under the influence of the magnetic field.

Magnetorheological elastomers were manufactured using:


PU substrates were mixed in the weight ratio, respectively 30:70:23. Mixing and curing process were conducted at room temperature.

EPU with molar ratio of MDI:(OAE+DCDA) equal to 2.5 was synthesized by one-shot method. The curing process was carried out at temperature of 150oC, what makes technology aspects of MRE manufacturing more complicated. The existence in every short hard segment of strong polar urea group and strong polar nitrilimide side-group increases the urea-urethane thermal and mechanical properties, as well as stiffness and hardness in comparison to soft polyether polyurethane (see Table 2).


**Table 2.** Selected physical and mechanical properties of elastomers used for MREs fabrication.

As shown in Table 2 polyurethane PU 70/30 is characterized by lower density and mechanical properties than EPU 2.5. Low hardness and stiffness of the polyurethane matrix can lead to the higher relative property changes of the MRE under an external magnetic field. On the other hand, the EPU 2.5 is distinguished by lower viscosity of reactive mixture of substrates and good mechanical properties. Low viscosity during the processing of the MRE makes the arrangement of the particles into aligned chains very easy.

During the fabrication of MRE, two polyols were mixed first and then carbonyl iron was added. This mixture was subsequently put under vacuum to remove trapped air bubbles. After de-gassing the isocyanate compound was added and the reactive mixture was poured into moulds. The mixing of substrates and curing processes were carried out at room temperature, without or with application of magnetic field.

Three iron powders, which differed in size and shape of the particles, were used as magnetoactive component of MREs:


Images of the particles and particle size distribution histograms are presented in Fig. 9, particles characteristics are shown in Table 3.

All the particles used for the study are characterized by their high saturation magnetization, which has a positive effect on the properties of MREs by increasing the strength of interaction between particles in a magnetic field. Moreover, they are commercially available and relatively inexpensive. A variety of MRE samples were produced:

a. pure elastomers (without particles),

158 Advanced Elastomers – Technology, Properties and Applications

modulus (without magnetic field) is lower.

interactions.

MR effect [69,70].

**4. Experimental results** 

Chemical Company,

process were conducted at room temperature.

**4.1. Materials characterisation** 

Magnetorheological elastomers were manufactured using:

magnetorheological effect is shown in other papers [20].

Beneficial effects of plasticizers, causing a reduction of material matrix stiffness, on the

Most of the results presented in the literature describes the MREs with isotropic distribution of particles. In [20,61] it was found that larger relative MR effect can be achieved in materials with softer matrix. Therefore, plasticizers are often added in order to increase the relative MR effect. The increase in the relative MR effect is obtained in this case, but the zero field

Also reports can be found about the influence of the anisotropic microstructure on the properties of magnetorheological elastomers. The static and dynamic tensile testing of MREs with particles oriented in the form of chains have shown that they exhibit much higher Young's modulus in comparison to MREs with the isotropically distributed particles [67], which is explained by the presence of column-like structure formed by the particles. Stiffness of the chains depends on the distance between the particles within them, and increases with the applied magnetic field due to the dipole formation and dipole-dipole

There is also a significant effect of particle size on the MR effect. In order to achieve a significant MR effect, particles should contain at least several magnetic domains [10]. The shape of particles can also have influence on the MR effect, it is better if they have an asymmetric shape with the main axis of anisotropy [63]. Few works have reported positive effect of surface modification of particles on the growth of magnetorheological effect [68], although others show clearly that due to the better adhesion particle-matrix boundary region occurs and the ability of particles to move is reduced, which leads to a reduction of

In our study, the selection of composite matrix was carried out to produce MREs with mechanical properties, which can be varied widely under the influence of the magnetic field.



PU substrates were mixed in the weight ratio, respectively 30:70:23. Mixing and curing

about 2000 g/mol and dicyandiamide (DCDA) as a chain extender.

b. elastomers with randomly dispersed iron particles,

	- c. elastomers with aligned iron particles.

**Figure 9.** SEM images and particle size distribution histograms A - carbonyl iron Fluka, B - carbonyl iron HQ, C - Iron PYRON.


**Table 3.** Characteristics of ferromagnetic particles used for MREs fabrication.

The carbonyl iron particle volume fractions were 1.5, 11.5, 18, 25 and 33.0 vol. %. The samples were subjected to a magnetic field during curing, to produce aligned chains in the elastomer matrices. Two magnetic field strengths were used: 80 and 240 kA/m. Samples with different orientation of the particle chains to the long sample axis (30, 45, 60 and 90 degree) were produced.

### **4.2. Microstructure-properties relationship in MRE**

160 Advanced Elastomers – Technology, Properties and Applications

A

B

C

**Figure 9.** SEM images and particle size distribution histograms A - carbonyl iron Fluka,

B - carbonyl iron HQ, C - Iron PYRON.

c. elastomers with aligned iron particles.

The spatial distribution of particles in the magnetorheological elastomer matrix depends on the curing conditions. It is known [6,7,8,10,32,33,34,35,36] that curing in the presence of a magnetic field, allows to obtain magneto-anisotropic structure, consisting of chains of particles, as shown in the example of MRE with PU matrix and 11.5 vol.% carbonyl iron particles (Fig. 10A and C). The cross-section parallel to the direction of the magnetic field shows the arranged paths (chains) of Fe particles, clearly indicating the orientation of the particles in a magnetic field (MF). The curing of magnetorheological elastomer without application of magnetic field, leads to the isotropic structure of Fe particles uniformly distributed throughout the volume of the matrix (Fig. 10D).

As it is seen in the cross-section of the composite, perpendicular to the direction of the magnetic field (Fig. 10B), the distribution of chains is relatively uniform, moreover, most of them consist of more than one row of particles.

The viscosity of the reactive mixture is a technological parameter, which turned out to have an significant influence on the ability of particles to create a structure oriented along the magnetic field direction. The viscosity is not constant over time, but increases as the reaction follows, to the principle of doubling the length of the molecules in each act of addition, often starting during mixing of liquid substrates. The lifetime of the mixture is the period in which it is liquid, until the gel formation. The movement of particles in a reactive mixture of substrates is possible until it is liquid.

**Figure 10.** Microstructure of MRE obtained from PU matrix containing 11.5 vol. % Fe particles: A anisotropic, cross-section parallel to the direction of the MF, SEM image, B - anisotropic cross section perpendicular to the direction of the MF, SEM image, C - anisotropic, cross-section parallel to the field, the image in polarized light, D - isotropic, SEM image. Arrows show the direction of the magnetic field during the curing.

Observations of the brittle fractures of produced MREs, as shown in Figure 11, show that too high viscosity of the reactive mixture inhibits the movement of particles and the applied magnetic field does not allow to obtain the oriented structure of particles chains. This is the case for the silicone rubber matrix (Fig. 11C), which has a viscosity before curing of 30 000 mPa∙s and is nearly four times greater than the viscosity of the PU 70/30 mixture and up to 20 times than EPU. Even a relatively short time of gelling (about 10 min for the PU 70/30) is not an obstacle for obtaining oriented structure of particle chains. It is sufficient for particles to shift the position to obtain the lowest Zeeman energy and arrange along the lines of magnetic field.

Another important parameter that affects the microstructure and consequently the properties of MREs is the magnetic field strength during curing. At higher magnetic field strengths the formed chains are thicker, consisting of more than 1-2 rows of particles.

during the curing.

magnetic field.

**Figure 10.** Microstructure of MRE obtained from PU matrix containing 11.5 vol. % Fe particles: A anisotropic, cross-section parallel to the direction of the MF, SEM image, B - anisotropic cross section perpendicular to the direction of the MF, SEM image, C - anisotropic, cross-section parallel to the field, the image in polarized light, D - isotropic, SEM image. Arrows show the direction of the magnetic field

Observations of the brittle fractures of produced MREs, as shown in Figure 11, show that too high viscosity of the reactive mixture inhibits the movement of particles and the applied magnetic field does not allow to obtain the oriented structure of particles chains. This is the case for the silicone rubber matrix (Fig. 11C), which has a viscosity before curing of 30 000 mPa∙s and is nearly four times greater than the viscosity of the PU 70/30 mixture and up to 20 times than EPU. Even a relatively short time of gelling (about 10 min for the PU 70/30) is not an obstacle for obtaining oriented structure of particle chains. It is sufficient for particles to shift the position to obtain the lowest Zeeman energy and arrange along the lines of

Another important parameter that affects the microstructure and consequently the properties of MREs is the magnetic field strength during curing. At higher magnetic field

strengths the formed chains are thicker, consisting of more than 1-2 rows of particles.

**Figure 11.** Microstructure of MREs with 11.5 vol.% of Fe particles, cured in the 80 kA/m, matrix: A - EPU 2.5, B - PU 70/30, C - silicone rubber. Arrows show the direction of the magnetic field during curing.

**Figure 12.** Model of MRE microstructure dependently on the magnetic field strength during the matrix curing: A - 80 kA/m, B - 240 kA/m.

The structural anisotropy was confirmed by the Vibration Sample Magnetometer (VSM) studies. Tests were carried out parallel (II) and perpendicular (L) to the direction of particles

chains, corresponding to the magnetic field direction during curing. The example of the hysteresis loops obtained for the MREs is shown in Fig. 13A. From the hysteresis loops the anisotropy coefficient (Ab) was calculated at the selected value of magnetic field strength of 160 kA/m. It was expressed by the ratio of magnetization measured at 160 kA/m, respectively parallel and perpendicular to the particles alignment direction. The Ab values are shown in Fig. 13B.

**Figure 13.** (A) Magnetization curves as a function of a magnetic field in the direction parallel (II) and perpendicular (L) to the chains of particles for the MRE containing 11.5 vol. % Fe particles, cured in the 240 kA/m magnetic field. (B) A change in the magnetic anisotropy Ab(160) for MRE with different particle amount and magnetic field during curing (80 and 240 kA/m).

It was found that there is a maximum value of the anisotropy coefficient corresponding to 11,5 vol.% of carbonyl-iron. The anisotropy coefficient decreases with further increase in the volume fraction of particles. Moreover, for the sample with the highest Ab a significant effect of the aligning magnetic field is visible; higher magnetic field leads to a better alignment. Structural and magnetic anisotropy has not been found in the specimens having 33 vol. % of carbonyl-iron, therefore, the anisotropy coefficient is equal to 1 for these specimens. The result can be attributed to the formation of the network of particles, in contrast to the chains of particles that characterize anisotropic composites. This was also observed on SEM images.

The anisotropy of MRE microstructure, obtained due to the magnetic field during curing has a significant influence on the changes of their mechanical properties in the magnetic field. Also the field strength applied during manufacture influences the mechanical properties of the MREs. This was confirmed by the results of mechanical and rheological tests presented in this chapter.

#### *4.2.1. Compression tests*

Compression tests were performed using a specially designed magnetic coil device for the MTS tensile tester. The study was conducted on samples with different particle volume content (1.5, 11.5 and 33%), produced in a magnetic field of 80 and 240 kA/m. Samples with dimensions of φ = 20 mm and h = 25 mm, with particle chains parallel to the axis of the sample, were compressed uniaxially without and with a magnetic field strength of 240 kA/m. The obtained stress-strain curves are shown in Figure 14.

164 Advanced Elastomers – Technology, Properties and Applications

amount and magnetic field during curing (80 and 240 kA/m).

are shown in Fig. 13B.

in this chapter.

*4.2.1. Compression tests* 

chains, corresponding to the magnetic field direction during curing. The example of the hysteresis loops obtained for the MREs is shown in Fig. 13A. From the hysteresis loops the anisotropy coefficient (Ab) was calculated at the selected value of magnetic field strength of 160 kA/m. It was expressed by the ratio of magnetization measured at 160 kA/m, respectively parallel and perpendicular to the particles alignment direction. The Ab values

**Figure 13.** (A) Magnetization curves as a function of a magnetic field in the direction parallel (II) and perpendicular (L) to the chains of particles for the MRE containing 11.5 vol. % Fe particles, cured in the 240 kA/m magnetic field. (B) A change in the magnetic anisotropy Ab(160) for MRE with different particle

It was found that there is a maximum value of the anisotropy coefficient corresponding to 11,5 vol.% of carbonyl-iron. The anisotropy coefficient decreases with further increase in the volume fraction of particles. Moreover, for the sample with the highest Ab a significant effect of the aligning magnetic field is visible; higher magnetic field leads to a better alignment. Structural and magnetic anisotropy has not been found in the specimens having 33 vol. % of carbonyl-iron, therefore, the anisotropy coefficient is equal to 1 for these specimens. The result can be attributed to the formation of the network of particles, in contrast to the chains of particles that characterize anisotropic composites. This was also observed on SEM images. The anisotropy of MRE microstructure, obtained due to the magnetic field during curing has a significant influence on the changes of their mechanical properties in the magnetic field. Also the field strength applied during manufacture influences the mechanical properties of the MREs. This was confirmed by the results of mechanical and rheological tests presented

Compression tests were performed using a specially designed magnetic coil device for the MTS tensile tester. The study was conducted on samples with different particle volume

**Figure 14.** Stress-strain curves during the compression of MREs cured in a field intensity of 80 and 240 kA/m, for various Fe particle volume fractions: A - 1.5%, B - 11.5%, C - 33%.

The compressive curves recorded without magnetic field are marked with a dashed line, while with MF using the solid line. It is clear that in each case MREs in a magnetic field are characterized by higher stiffness and compressive strength than without magnetic field. MREs with a higher content of particles are characterized by higher stiffness and

compressive strength both without and with the magnetic field. Moreover, in a magnetic field, obtained stress-strain curves are at a higher level for MREs with higher structural and magnetic anisotropy, those cured in a magnetic field of greater intensity, even when no field curve was at a similar level, as for the MRE containing 1.5% vol. Fe particles (Fig. 14A).

Changes of the properties of magnetorheological elastomers in the magnetic field are described by absolute and relative MR effect, which in this case is expressed by the relative and absolute change in compressive modulus. Based on the compressive curves, the absolute and relative change of the compressive modules in the field of 240 kA/m, at the deformation of 0.1 mm/mm has been calculated and presented in Table 4.


**Table 4.** Comparison of absolute (ΔEc (240)) and relative (ઢ۳܋ሺሻ ሺሻ ) changes in the compressive modules in the magnetic field of 240 kA/m, under the deformation of 0.1 mm/mm.

The absolute MR effect is greater with the higher content of iron particles and depends on the applied magnetic field strength during curing. The higher the magnetic field during curing, the larger the absolute MR effect. The relative MR effect is not directly proportional to the particle content and reaches the highest value for contents of 11.5 vol.% Fe. Relative MR effect for the largest particle content, as well as magnetic anisotropy coefficient is practically the same regardless of the magnetic field during manufacture. Moreover, it is worth noting that the relative MR effect reaches its maximum for the MRE with the highest values of the coefficient of magnetic anisotropy.

### *4.2.2. Rheological properties*

The knowledge of dynamic properties is needed especially for materials that are used in equipment subjected to vibration such as in aerospace and automotive industries. To describe the efficiency of damping, parameters such as the loss angle tan (δ) and shear modulus G\*, consisting of the storage modulus G' and loss modulus G" are used, determined by rheological tests [47,48].

(Fig. 14A).

Table 4.

Fe particles fraction [vol. %]

1,5

11,5

33

compressive strength both without and with the magnetic field. Moreover, in a magnetic field, obtained stress-strain curves are at a higher level for MREs with higher structural and magnetic anisotropy, those cured in a magnetic field of greater intensity, even when no field curve was at a similar level, as for the MRE containing 1.5% vol. Fe particles

Changes of the properties of magnetorheological elastomers in the magnetic field are described by absolute and relative MR effect, which in this case is expressed by the relative and absolute change in compressive modulus. Based on the compressive curves, the absolute and relative change of the compressive modules in the field of 240 kA/m, at the deformation of 0.1 mm/mm has been calculated and presented in

**ΔEc(240)**

ઢ۳܋ሺሻ ሺሻ

) changes in the compressive modules

Ψ כ

**[%]** 

**[MPa]** 

80 0,03 25 240 0,05 51

80 0,17 86 240 0,23 100

80 0,72 56 240 0,75 55

ሺሻ

The absolute MR effect is greater with the higher content of iron particles and depends on the applied magnetic field strength during curing. The higher the magnetic field during curing, the larger the absolute MR effect. The relative MR effect is not directly proportional to the particle content and reaches the highest value for contents of 11.5 vol.% Fe. Relative MR effect for the largest particle content, as well as magnetic anisotropy coefficient is practically the same regardless of the magnetic field during manufacture. Moreover, it is worth noting that the relative MR effect reaches its maximum for the MRE with the highest

The knowledge of dynamic properties is needed especially for materials that are used in equipment subjected to vibration such as in aerospace and automotive industries. To describe the efficiency of damping, parameters such as the loss angle tan (δ) and shear modulus G\*, consisting of the storage modulus G' and loss modulus G" are used,

**Magnetic field during curing [kA/m]** 

**Table 4.** Comparison of absolute (ΔEc (240)) and relative (ઢ۳܋ሺሻ

values of the coefficient of magnetic anisotropy.

*4.2.2. Rheological properties* 

determined by rheological tests [47,48].

in the magnetic field of 240 kA/m, under the deformation of 0.1 mm/mm.

**Figure 15.** The change of rheological properties in the magnetic field as a function of frequency at a constant angular deformation of 0.1% for the MRE with PU matrix containing 11.5 vol.% Fe particles, cured in the 80 and 240 kA/m: A - storage modulus G', B - the loss modulus G", C - damping factor tan(δ).

MREs cured at different magnetic field strengths, tested without magnetic field have a similar storage (elasticity) modulus, loss (viscosity) modulus and the loss angle, which increases with increasing oscillation frequency of deformation. This pattern is characteristic for viscoelastic materials for which the higher the frequency of oscillation, the higher the value of the modules G' and G" [44]. The sharp increase in tan(δ) as a function of oscillation frequency shows an increase in the share of the viscous properties of MREs.

The increase in the magnetic field causes an increase of both modules G' and G", much higher values were obtained in 160 kA/m for the MRE with a higher magnetic and structural anisotropy (cured in the 240 kA/m). For larger magnetic field the difference is not so evident. In a magnetic field damping is characterized by smaller changes in the function of frequency, but its value clearly depends on the magnetic field strength (tan (δ) increases with increasing field) and to a lesser extent, but also depends on the microstructure formed during curing. MREs cured in the field of higher intensity are characterized by higher values of tan(δ). With increasing magnetic field, the share of the viscous properties is greater, and increases with the structural and magnetic anisotropy. The material absorbs more energy which is dissipated as heat.

C - damping factor tan(δ).

which is dissipated as heat.

**Figure 15.** The change of rheological properties in the magnetic field as a function of frequency at a constant angular deformation of 0.1% for the MRE with PU matrix containing 11.5 vol.% Fe particles, cured in the 80 and 240 kA/m: A - storage modulus G', B - the loss modulus G",

frequency shows an increase in the share of the viscous properties of MREs.

MREs cured at different magnetic field strengths, tested without magnetic field have a similar storage (elasticity) modulus, loss (viscosity) modulus and the loss angle, which increases with increasing oscillation frequency of deformation. This pattern is characteristic for viscoelastic materials for which the higher the frequency of oscillation, the higher the value of the modules G' and G" [44]. The sharp increase in tan(δ) as a function of oscillation

C

The increase in the magnetic field causes an increase of both modules G' and G", much higher values were obtained in 160 kA/m for the MRE with a higher magnetic and structural anisotropy (cured in the 240 kA/m). For larger magnetic field the difference is not so evident. In a magnetic field damping is characterized by smaller changes in the function of frequency, but its value clearly depends on the magnetic field strength (tan (δ) increases with increasing field) and to a lesser extent, but also depends on the microstructure formed during curing. MREs cured in the field of higher intensity are characterized by higher values of tan(δ). With increasing magnetic field, the share of the viscous properties is greater, and increases with the structural and magnetic anisotropy. The material absorbs more energy

**Figure 16.** Changes in modules G' and G'' and damping factor tan(δ) as a function of oscillation frequency (at constant strain of 0.1%) for the MREs with PU matrix and different content of Fe particles without and in the field 160 kA/m. Particle chains parallel to the MF in rheometer.

Selected plots of rheological properties as a function of frequency of oscillation, described by the storage modulus G', loss modulus G'' and loss angle tan(δ) without field and in a magnetic field of 160 kA/m, for MREs with a polyurethane matrix and different content of carbonyl iron particles, are shown in Fig 16. The samples were cured in the magnetic field of 240 kA/m, and the particle chains were oriented parallel to the direction of the MF applied in rheometer.

It can be concluded that the increase in the particle content causes a slight increase in both modules in the initial state (without magnetic field) and they are dependent on the frequency of oscillation. Modules (storage and loss) in the initial state increase with increasing frequency of oscillation. Loss angle also depends on the content of particles and is greater with the higher particle content. The higher the value of tan (δ), the greater the ability to absorb energy. However, note that the values obtained for the studied MREs are well below 1, which means a significant advantage of the elastic properties over the viscous.

In the magnetic field storage and loss modulus increase, and this increase depends on the content of the particles. The more particles the greater the value of the modules, and both increase with increasing frequency. For high frequencies, a decrease in loss modulus is observed, inversely proportional to the particle content.

Increase of loss angle as a function of frequency in the magnetic field is much smaller, except for the MREs with very low amount of particles 1.5 vol.%, for which the tan(δ) in the MF is the same as without magnetic field. The values of tan(δ) clearly increase for the low oscillation frequency, and rise with the rising particle content. The higher the value of tan(δ), the greater the share of energy dissipated, and the material exhibits better damping properties.

The influence of the size of particles on the microstructure and rheological properties of MREs has also been investigated. The particle size does not affect the possibility of obtaining anisotropic structure consisting of chains of particles. In case of irregularly shaped particles (PYRON iron) during curing in magnetic field, they turn so their long axis is parallel to the direction of the magnetic field, which is consistent with the direction of easy magnetization and the tendency to minimize the energy associated with the shape anisotropy.

Rheological parameters were measured for a selected frequency of 10 Hz, the parameters G', G", and their absolute (ΔG', ΔG") and relative (ΔG'/G'0, ΔG"/G"0) changes describing absolute and relative MR effect are shown in Figure 18 as a function of magnetic field strength. Compared MREs have parallel arrangement of particles chains to the magnetic field in rheometer, the same content of particles but different sizes.

Storage modulus, loss modulus, and their absolute and relative changes increase with increasing magnetic field to the value of 320 kA/m, above the increase is weaker, which is associated with reaching a saturation magnetization of particles. In case of loss modulus G" and the absolute and relative changes, significantly higher values as a function of field were obtained for the largest particles, with an average size of ~ 74 microns and irregular shapes. Storage modulus G' values were clearly lower for MRE with particles of small diameter (~ 1.4 μm), while for other types of particles are similar. Noticeable difference in favor of larger particles occurs in the field of 320 kA/m, which comes from the fact that particle size affects the strength of magnetic interactions between them. The bigger they are the stronger the interactions, the particles are attracted to each other in a chain with higher force, which ultimately leads to a greater MR effect. The maximum relative MR effect for the MREs with 11.5 vol.% of particles is 760% of the relative change in storage modulus, and 770% for the loss modulus. For this field strength these values are much higher than the values found in the literature.

170 Advanced Elastomers – Technology, Properties and Applications

observed, inversely proportional to the particle content.

in rheometer.

over the viscous.

properties.

Selected plots of rheological properties as a function of frequency of oscillation, described by the storage modulus G', loss modulus G'' and loss angle tan(δ) without field and in a magnetic field of 160 kA/m, for MREs with a polyurethane matrix and different content of carbonyl iron particles, are shown in Fig 16. The samples were cured in the magnetic field of 240 kA/m, and the particle chains were oriented parallel to the direction of the MF applied

It can be concluded that the increase in the particle content causes a slight increase in both modules in the initial state (without magnetic field) and they are dependent on the frequency of oscillation. Modules (storage and loss) in the initial state increase with increasing frequency of oscillation. Loss angle also depends on the content of particles and is greater with the higher particle content. The higher the value of tan (δ), the greater the ability to absorb energy. However, note that the values obtained for the studied MREs are well below 1, which means a significant advantage of the elastic properties

In the magnetic field storage and loss modulus increase, and this increase depends on the content of the particles. The more particles the greater the value of the modules, and both increase with increasing frequency. For high frequencies, a decrease in loss modulus is

Increase of loss angle as a function of frequency in the magnetic field is much smaller, except for the MREs with very low amount of particles 1.5 vol.%, for which the tan(δ) in the MF is the same as without magnetic field. The values of tan(δ) clearly increase for the low oscillation frequency, and rise with the rising particle content. The higher the value of tan(δ), the greater the share of energy dissipated, and the material exhibits better damping

The influence of the size of particles on the microstructure and rheological properties of MREs has also been investigated. The particle size does not affect the possibility of obtaining anisotropic structure consisting of chains of particles. In case of irregularly shaped particles (PYRON iron) during curing in magnetic field, they turn so their long axis is parallel to the direction of the magnetic field, which is consistent with the direction of easy magnetization

Rheological parameters were measured for a selected frequency of 10 Hz, the parameters G', G", and their absolute (ΔG', ΔG") and relative (ΔG'/G'0, ΔG"/G"0) changes describing absolute and relative MR effect are shown in Figure 18 as a function of magnetic field strength. Compared MREs have parallel arrangement of particles chains to the magnetic

Storage modulus, loss modulus, and their absolute and relative changes increase with increasing magnetic field to the value of 320 kA/m, above the increase is weaker, which is associated with reaching a saturation magnetization of particles. In case of loss modulus G" and the absolute and relative changes, significantly higher values as a function of field were obtained for the largest particles, with an average size of ~ 74 microns and irregular shapes.

and the tendency to minimize the energy associated with the shape anisotropy.

field in rheometer, the same content of particles but different sizes.

**Figure 17.** Microstructure ofs MRE with 11.5 vol.% of particles of different sizes: A - Fe ~ 9 μm, B - Fe HQ ~ 1.4 μm, C - Fe PYRON ~ 74 μm. Arrows shows the direction of the magnetic field during curing.

The alignment of particle chains within the matrix may have a significant effect on the performance of the MREs. In order to investigate this effect, samples with different alignment of particles were produced and tested. A schematic representation of the samples, which contained 11.5 vol.% of Fe particles (Fluka, 9 μm), is shown in Fig. 19.

172 Advanced Elastomers – Technology, Properties and Applications

**Figure 18.** Changes of the storage modulus G' (A), loss modulus G" (B) and the absolute (C, D) and relative (E, F) magnetorheological effects as a function of magnetic field for the MRE with polyurethane matrix and content of 11.5 vol.% of particles. Sizes of particles: Fe ~ 9 μm, Fe HQ ~ 1.4 μm, Fe PYRON ~ 74 μm. Magnetic field in rheometer parallel to the chains of particles.

**Figure 18.** Changes of the storage modulus G' (A), loss modulus G" (B) and the absolute (C, D) and relative (E, F) magnetorheological effects as a function of magnetic field for the MRE with polyurethane matrix and content of 11.5 vol.% of particles. Sizes of particles: Fe ~ 9 μm, Fe HQ ~ 1.4 μm, Fe PYRON ~

74 μm. Magnetic field in rheometer parallel to the chains of particles.

**Figure 19.** Schematic representation of the particle distribution with respect to the magnetic field direction.

**Figure 20.** The effects of different orientation of carbonyl iron particles in elastomer matrix, changes of relative MR effect as a function of the magnetic field.

It can be concluded that the values of storage and loss modules, and their absolute and relative changes strongly depend on the orientation of the particles chains. The largest absolute and relative MR effects were obtained for the arrangement of particles at an angle of 60° to the direction of the magnetic field. Under the influence of a magnetic field the direction of easy magnetization comes along the chains of ferromagnetic particles. To minimize the Zeeman energy the chains tend to set to the direction of the magnetic field, which involves the displacement of particles. Because they are embedded in the elastomer matrix and restrained by interactions at the interface, their displacements will introduce additional shear stress into the matrix, causing increased stress concentration in the matrix between the particles [71]. This does not take place when the particles chains are perpendicular to the direction of the field, and the resulting MR effect is much higher than for the isotropic structure. The lowest MR effect as a function of the field, both absolute and relative has been obtained in isotropic MREs.

The properties of the elastomeric matrix have also a great influence on MREs rheological properties. Rheological properties measured for the MREs fabricated from elastomers with different stiffness and hardness are shown in Fig. 21. Elastomer matrix with higher stiffness and hardness (EPU 2.5) leads to the higher storage and loss modulus of the MREs measured without magnetic field, while under the magnetic field the values are significantly lower in comparison to MREs obtained from soft elastomer (PU 70/30). As a result the absolute and relative MR effect is much higher in MREs obtained from the soft elastomer.

**Figure 21.** Absolute and relative MR effect of the MREs based on PU 70/30 and EPU 2.5 elastomers vs. the magnetic field strength (H).

MREs with the stiffer matrix (EPU 2.5) exhibit significant lower MR effect because the stiffer matrix makes impossible the particles to displace when they are subjected to the magnetic field [72].

Developed magnetorheological composites belong to a group of smart materials, and under the influence of the magnetic field changes of the properties should be fully reversible. It is a complete reversibility of change in modules G' and G", as shown in Fig. 22 for the MRE containing 11.5 vol.% Fe particles oriented at an angle of 45 degrees in the increasing magnetic field, which was alternating turned on and off. Each time after switching off the magnetic field the modules immediately returned to initial state.

**Figure 22.** The changes of rheological properties of MRE in the magnetic field alternating switched on and off.

### **5. Conclusions**

174 Advanced Elastomers – Technology, Properties and Applications

relative has been obtained in isotropic MREs.

the magnetic field strength (H).

field [72].

It can be concluded that the values of storage and loss modules, and their absolute and relative changes strongly depend on the orientation of the particles chains. The largest absolute and relative MR effects were obtained for the arrangement of particles at an angle of 60° to the direction of the magnetic field. Under the influence of a magnetic field the direction of easy magnetization comes along the chains of ferromagnetic particles. To minimize the Zeeman energy the chains tend to set to the direction of the magnetic field, which involves the displacement of particles. Because they are embedded in the elastomer matrix and restrained by interactions at the interface, their displacements will introduce additional shear stress into the matrix, causing increased stress concentration in the matrix between the particles [71]. This does not take place when the particles chains are perpendicular to the direction of the field, and the resulting MR effect is much higher than for the isotropic structure. The lowest MR effect as a function of the field, both absolute and

The properties of the elastomeric matrix have also a great influence on MREs rheological properties. Rheological properties measured for the MREs fabricated from elastomers with different stiffness and hardness are shown in Fig. 21. Elastomer matrix with higher stiffness and hardness (EPU 2.5) leads to the higher storage and loss modulus of the MREs measured without magnetic field, while under the magnetic field the values are significantly lower in comparison to MREs obtained from soft elastomer (PU 70/30). As a result the absolute and

**Figure 21.** Absolute and relative MR effect of the MREs based on PU 70/30 and EPU 2.5 elastomers vs.

MREs with the stiffer matrix (EPU 2.5) exhibit significant lower MR effect because the stiffer matrix makes impossible the particles to displace when they are subjected to the magnetic

Developed magnetorheological composites belong to a group of smart materials, and under the influence of the magnetic field changes of the properties should be fully reversible. It is a complete reversibility of change in modules G' and G", as shown in Fig. 22 for the MRE

relative MR effect is much higher in MREs obtained from the soft elastomer.

Studies on fabrication of MREs were carried out using different elastomers as a matrix, ferromagnetic particles with various shape and size. Samples with isotropic and anisotropic particles arrangement, were examined. Particles were oriented into chains under external magnetic field. Special attention was put on fabrication of samples with different orientation of chains to the magnetic field direction.

It was found that the microstructure of the MREs depends on the amount of ferrous particles and manufacturing conditions. The orientation of the iron particles into aligned chains is possible for lower volume content of the ferromagnetic fillers. High carbonyl-iron volume content in the matrix leads to the formation of more complex three-dimensional lattices. Also the magnetic measurements confirmed the existence of the microstructure anisotropy for the lower volume content of the iron particles. The structural and magnetic anisotropy has not been found in the MREs with 33 vol. % of particles. To evaluate the effect of the external magnetic field on the magnetorheological properties compressive strength, storage and loss modulus, as well as loss factor were measured. Both, the content of particles and their arrangement have significant effect on the properties of magnetorheological

urethane elastomers. Compression test results showed that under external magnetic field samples are distinguished by a higher compressive strength.

Rheological properties of magnetorheological urethane elastomers also depend on the content of particles and their arrangement. Application of an external magnetic field leads to a significant increase in elastic modulus. Absolute and relative changes of storage modulus, calculated from obtained curves, show that the microstructure of the samples has a significant effect on their magnetorheological effect. Magnetorheological effect, expressed by relative change of storage modulus under magnetic field, is the highest for the sample with the highest magnetic and structural anisotropy.

The aligned particle network structure has a significant influence on the elastic properties of the composite material. Inside the chains the effective filler content is higher than the average filler content. By optimizing the content of particles and their alignment, either the stiffness or the damping properties of MREs can be increased by applying a magnetic field.

As a result of the studies it was also found, that the MREs having anisotropic microstructure exhibit, for the same particles content, much higher magnetorheological effect in comparison to the isotropic ones. The non-linear change of the rheological properties versus particles fraction was found. It is due to structural and magnetic anisotropy of the MREs, which has the greatest influence on the changes of the properties under the magnetic field, i.e. magnetorheological effect. Moreover, it was found that the magnetorheological effect can be controlled by the particles alignment to the magnetic field lines. It means that it is possible to obtain high magnetorheological effect not by increasing of the particles volume fraction, but by the formation of appropriate microstructure. It can be achieved for lower particles volume fraction, what advantageously decreases weight of devices based on the MREs. As a result of this work the new MREs, with application capabilities, characterized by extremely high magnetorheological effect (900% in MF of 480 kA/m), were elaborated.

### **Author details**

Anna Boczkowska and Stefan Awietjan *Warsaw University of Technology, Faculty of Materials Science and Engineering, Warsaw, Poland* 

### **6. References**


[6] Ginder J.M., Rheology controlled by magnetic fields, in: Encyclopedia of Applied Physics, vol. 16, VCH Publisher Inc., New York (1996)

176 Advanced Elastomers – Technology, Properties and Applications

the highest magnetic and structural anisotropy.

**Author details** 

**6. References** 

Anna Boczkowska and Stefan Awietjan

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samples are distinguished by a higher compressive strength.

urethane elastomers. Compression test results showed that under external magnetic field

Rheological properties of magnetorheological urethane elastomers also depend on the content of particles and their arrangement. Application of an external magnetic field leads to a significant increase in elastic modulus. Absolute and relative changes of storage modulus, calculated from obtained curves, show that the microstructure of the samples has a significant effect on their magnetorheological effect. Magnetorheological effect, expressed by relative change of storage modulus under magnetic field, is the highest for the sample with

The aligned particle network structure has a significant influence on the elastic properties of the composite material. Inside the chains the effective filler content is higher than the average filler content. By optimizing the content of particles and their alignment, either the stiffness or the damping properties of MREs can be increased by applying a magnetic field. As a result of the studies it was also found, that the MREs having anisotropic microstructure exhibit, for the same particles content, much higher magnetorheological effect in comparison to the isotropic ones. The non-linear change of the rheological properties versus particles fraction was found. It is due to structural and magnetic anisotropy of the MREs, which has the greatest influence on the changes of the properties under the magnetic field, i.e. magnetorheological effect. Moreover, it was found that the magnetorheological effect can be controlled by the particles alignment to the magnetic field lines. It means that it is possible to obtain high magnetorheological effect not by increasing of the particles volume fraction, but by the formation of appropriate microstructure. It can be achieved for lower particles volume fraction, what advantageously decreases weight of devices based on the MREs. As a result of this work the new MREs, with application capabilities, characterized by extremely

high magnetorheological effect (900% in MF of 480 kA/m), were elaborated.

*Warsaw University of Technology, Faculty of Materials Science and Engineering, Warsaw, Poland* 

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magnetorheological elastomers, Smart. Mater. Struct. 16 (2007) 2645-2650



## **Thermo-Shrinkable Elastomers**

Magdalena Maciejewska and Alicja Krzywania-Kaliszewska

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/48209

### **1. Introduction**

180 Advanced Elastomers – Technology, Properties and Applications

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Modern Phys. B, Vol. 16, No. 17 i 18, (2002) 2447-2453

The ability of polymeric materials to sense and respond to external stimuli has great scientific and technological significance. It enables these materials to change their properties, such as shape, colour and electrical conductivity, as a result of pH, temperature, chemicals, light, or stimulation by an electric or magnetic field. Materials that respond dynamically to external stimuli are called intelligent or smart materials, and it is important that their response should be repeatable and controllable (Landlein, 2010). One of the most important classes of smart materials is shape memory polymers (SMPs), which can change their shape in a predetermined way upon the application of an external stimulus. The shape memory effect in polymers depends primarily on the existence of separated phases that are related to the coiled structure, crosslinks (covalent bonds), hydrogen or ionic bonding or physical intermolecular interactions of the polymer (Hu, 2007). Covalent crosslinks are formed during suitable crosslinking of the polymer, whereas physical crosslinks are obtained when the polymer morphology consists of segregated domains, such as crystalline and amorphous phases or hard and soft segments (Landlein, 2010) (e.g., linear block copolymers). In multiphase polymers, the hard segments act as the frozen phase, which is usually semicrystalline or physically crosslinked and provides stiffness and reinforcement to the material (Hu, 2007), while the soft segments are responsible for the thermo-elastic behaviour of polymers and act as the reversible phase. In this case, the shape memory effect is produced by the reversible phase transformation of the soft segments (Hu, 2007). Thermo-shrinkable polymers, which change their shape as the temperature changes, are a unique class of SMP materials with interesting properties and many potential applications. The shape transition temperature (Ttrans) for this type of SMP can be the melting point (Tm) or glass transition temperature (Tg) of the soft phase (Ratna & Karger-Kocsis, 2008). Melting points are preferred because the transition is sharper than the glass transition; therefore, the temperature of the shape recovery can be better determined. Heating the SMP above the Tm or Tg of the hard segment enables its processing. This original (permanent) shape can be

© 2012 Maciejewska and Krzywania-Kaliszewska, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

memorised by cooling the material below the Tm or Tg of the hard phase. Cooling the SMP below the Tm or Tg of the soft segment while the shape is deformed allows a temporary shape to be fixed. The permanent shape of the SMP is recovered by heating it above the Tm or Tg of the soft phase (Hu, 2007). Another way to fix the temporary shape of a SMP is to deform it at a temperature lower than the Tm or Tg of the soft phase, which causes stress and strain absorption by the soft segments. When the material is heated above the Tm or Tg of the soft segments, the stresses and strains are relieved, causing the material to return to its original shape (Leng & Du, 2010). The thermally induced shape memory effect in polymers is schematically presented in Figure 1. SMPs that use Tm as the shape transition temperature are represented by polymers such as polyurethanes, block copolymers of polyethyleneterephthalate and polyethyleneoxide, and copolymers consisting of polystyrene and poly(1,4-butadiene) (Wang et al., 1998). Tg is the shape transition temperature for thermoset SMPs; e.g., styrene-based resins (Ivens et al., 2011). SMPs can also be single phase materials with a certain number of crosslinks between their polymer chains. In this case, the crosslinks are the net points that enable fixing and storage of the permanent shape of the polymer, whereas the free polymer chains between the crosslinks act as switching segments that possess increased mobility above Tg (Ratna & Karger-Kocsis, 2008). Stretching the polymer chain segments in a certain direction reduces their entropy. At the same time, the net points deform elastically. The overall result of this process is an increase in the SMP enthalpy. The loss of polymer chain segment mobility in the switching segments stabilises the temporary shape of the SMP when the material is allowed to cool in the deformed state (Ivens et al., 2011).

**Figure 1.** Thermally induced shape memory effect in polymers

Recently, research activities related to the development of elastomeric composites that exhibit a thermally induced shape memory effect have intensified. Thermoplastic elastomers (TPEs) are an interesting class of materials. TPEs behave like cured elastomers at room temperature, and they can be processed as plastics at higher temperatures because of their special multiphase morphology. TPEs consist of plastic phases embedded in a continuous elastomer phase, and they form physical crosslinks. Among the TPEs, segmented polyurethane elastomers and ionomers or blends of elastomers with thermoplastic polymers have many potential applications. In shape memory polyurethanes, the hard segment phases with the highest thermal transition temperatures (Tperm) acts as the physical crosslink points and control the permanent shape of the polymer. When the polymer is heated above Tperm, the physical crosslinks between the hard segments are destroyed. The molecular chains melt, and the polymer can be processed to a permanent shape like a thermoplastic material. When the shape memory polyurethane is cooled below Tperm and above Tg or Tm of its soft segments, the polymer becomes relatively soft, but it cannot flow because of the physical crosslinks. Consequently, it can be easily deformed to a temporary shape by stretching or compression. Reheating the polymer above Tg or Tm of its soft segments but below Tperm induces shape recovery (Hu, 2007). In the case of polyolefin/elastomer blends, the crosslinked elastomeric phase causes an enhancement of the blend shrinkability upon heating. Crosslinked points in the elastomer network are believed to serve as memory points, increasing the heat shrinkability (Patra & Das, 1997). Weiss et al. designed a new type of SMP based on blends of an elastomeric ionomer and low molar mass fatty acids or their salts (Weiss et al., 2008). Nanophase separation of the ionomer was used to develop the permanent network, and the fatty acids or salts were used to produce a secondary network. The role of the ionomer was to provide a strong intermolecular bond between crystals of the fatty acids or salts and the polymer, acting as a physical crosslink below the fatty acids or salt Tm and allowing reshaping of the material above Tm. The polymer films were heated to 100°C and stretched to 47% strain, then cooled to room temperature to fix an elongated temporary shape. When the samples were reheated to 100°C, the film recovered to the permanent shape with a length recovery of approximately 92%. A two-way temperatureinduced shape memory effect was observed for polymer laminates. SMP-laminated composites were prepared with SMP polyurethane films and elastic polymers films. Their two-way shape memory behaviour was produced by bending upon heating and reverse bending upon cooling. The shape memory mechanism was ascribed to the release of the elastic strain of the SMP layer upon heating and the recovery of the elastic strain induced by the bending force of the substrate layer upon cooling (Chen et al., 2008). Mishra et al. have studied the heat shrinkability behaviour of grafted low-density polyethylene/polyurethane elastomers. They have suggested that the interchain crosslinking between the grafted polyethylene and the elastomer improves their shrinkability (Mishra et al., 2004).

182 Advanced Elastomers – Technology, Properties and Applications

deformed state (Ivens et al., 2011).

permanent shape

deformation

T>Ttrans

**Figure 1.** Thermally induced shape memory effect in polymers

shape programming

Recently, research activities related to the development of elastomeric composites that exhibit a thermally induced shape memory effect have intensified. Thermoplastic elastomers (TPEs) are an interesting class of materials. TPEs behave like cured elastomers at room temperature, and they can be processed as plastics at higher temperatures because of their special multiphase morphology. TPEs consist of plastic phases embedded in a continuous elastomer phase, and they form physical crosslinks. Among the TPEs, segmented polyurethane elastomers and ionomers or blends of elastomers with thermoplastic polymers

shape fixation

T<Ttrans

shape recovery

shape

temporary shape permanent

T>Ttrans

memorised by cooling the material below the Tm or Tg of the hard phase. Cooling the SMP below the Tm or Tg of the soft segment while the shape is deformed allows a temporary shape to be fixed. The permanent shape of the SMP is recovered by heating it above the Tm or Tg of the soft phase (Hu, 2007). Another way to fix the temporary shape of a SMP is to deform it at a temperature lower than the Tm or Tg of the soft phase, which causes stress and strain absorption by the soft segments. When the material is heated above the Tm or Tg of the soft segments, the stresses and strains are relieved, causing the material to return to its original shape (Leng & Du, 2010). The thermally induced shape memory effect in polymers is schematically presented in Figure 1. SMPs that use Tm as the shape transition temperature are represented by polymers such as polyurethanes, block copolymers of polyethyleneterephthalate and polyethyleneoxide, and copolymers consisting of polystyrene and poly(1,4-butadiene) (Wang et al., 1998). Tg is the shape transition temperature for thermoset SMPs; e.g., styrene-based resins (Ivens et al., 2011). SMPs can also be single phase materials with a certain number of crosslinks between their polymer chains. In this case, the crosslinks are the net points that enable fixing and storage of the permanent shape of the polymer, whereas the free polymer chains between the crosslinks act as switching segments that possess increased mobility above Tg (Ratna & Karger-Kocsis, 2008). Stretching the polymer chain segments in a certain direction reduces their entropy. At the same time, the net points deform elastically. The overall result of this process is an increase in the SMP enthalpy. The loss of polymer chain segment mobility in the switching segments stabilises the temporary shape of the SMP when the material is allowed to cool in the

> Zhang et al. reported a novel type of shape memory polymer blend that consisted of two immiscible components, an elastomer and a switch polymer. The elastomer could be a rubber or thermoplastic elastomer, and the switch polymer could be an amorphous or crystalline polymer. Styrene–butadiene–styrene tri-block copolymer (SBS) was chosen as the elastomer, and poly(ε-caprolactone) (PCL) was used as the switch polymer (Zhang et al., 2009). The SBS/PCL blends demonstrated good shape recovery performance, with a shape recovery ratio of approximately 100%. The shape memory effect was also observed in elastomer networks containing reversibly associating side-groups. The supramolecular shape-memory elastomer consisted of a lightly crosslinked polymer network that was covalently bonded to reversibly associating side-groups. These elastomers exhibited shapememory effects arising from reversible hydrogen bond association, and the shape memory

recovery rate was strongly dependent on the temperature. Hydrogen-bonding interactions could stabilise mechanically strained states in these elastomers, and the thermo-mechanical cycling produced a strain fixity of approximately 90% and a strain recovery of approximately 100% (Li et al., 2007).

Thermal shrinkability was also reported for elastomer blends containing carboxylated groups, which could form ionic labile bonds during crosslinking, and low-density polyethylene (Mishra et al., 2000) or poly(ethylene-vinyl-acetate) (Raychowdhury et al., 2000).

In our previous studies, we proved that the thermally induced shape memory effect was also exhibited by a carboxylated nitrile elastomer cured with zinc oxide because it contained labile ionic crosslinks, which were able to rearrange upon external deformation (Przybyszewska & Zaborski, 2009).

In this work, the thermo-shrinkable properties of carboxylated acrylonitrile – butadiene elastomer (XNBR) and hydrogenated acrylonitrile – butadiene elastomer (HNBR) containing ionic crosslinks were studied. The XNBR was vulcanised with nanosized calcium, magnesium oxide or zinc oxide to ensure the formation of ionic crosslinks during vulcanisation. In the case of HNBR, nanosized calcium and magnesium oxides were coated with unsaturated carboxylic acids (itaconic, 2,4-pentadienoic, oleic, linoleic and linolenic acids) and applied as coagents in the peroxide vulcanisation of the elastomer. The application of these coagents led to formation of ionic crosslinks in the elastomer network.

Heat-shrinkable polymers are widely used in packaging and in the cable industry; therefore, the shrinkability of XNBR and HNBR vulcanisates is technologically important.

### **2. Experimental section**

#### **2.1. Materials**

Carboxylated nitrile elastomer XNBR (Krynac X7.50) containing 27 wt % acrylonitrile and 6.7 wt % carboxylic groups was obtained from Bayer C.O. The Mooney viscosity was (ML1+4 (100oC):47). Nanosized calcium oxide CaO (Aldrich), magnesium oxide MgO (Nanostructured & Amorphous Materials Inc., Houston, USA) and zinc oxide ZnO (Nanostructured & Amorphous Materials Inc., Houston, USA) were used as crosslinking agents. Hydrogenated acrylonitrile-butadiene elastomer HNBR (Therban 3407) containing 34 wt % acrylonitrile and 0.9 wt % of residual double bonds after hydrogenation was obtained from Bayer C.O. The Mooney viscosity was (ML1+4 (100oC):70). It was vulcanised with dicumyl peroxide DCP (Aldrich). Nanosized calcium oxide (Aldrich), magnesium oxide (Nanostructured & Amorphous Materials Inc., Houston, USA), itaconic acid IA (Fluka), 2-4-pentadienoic acid (Aldrich), oleic acid, linoleic acid and linolenic acid (Aldrich) were applied as coagents.

#### **2.2. Preparation of coagents**

The nanosized metal oxides were mixed with a solution of modifying agent (unsaturated carboxylic acid) in acetone for 30 minutes during ultrasonic treatment (BANDELIN DT 255) at a frequency of 35 kHz. The mixture was left for 24 hours. Then, the solvent (acetone) was evaporated using a vacuum evaporator at 50oC. The coagents obtained were dried in a vacuum drier at 70oC for 96 hours.

#### **2.3. Preparation and characterisation of rubber compounds**

184 Advanced Elastomers – Technology, Properties and Applications

approximately 100% (Li et al., 2007).

(Przybyszewska & Zaborski, 2009).

**2. Experimental section** 

were applied as coagents.

**2.2. Preparation of coagents** 

**2.1. Materials** 

recovery rate was strongly dependent on the temperature. Hydrogen-bonding interactions could stabilise mechanically strained states in these elastomers, and the thermo-mechanical cycling produced a strain fixity of approximately 90% and a strain recovery of

Thermal shrinkability was also reported for elastomer blends containing carboxylated groups, which could form ionic labile bonds during crosslinking, and low-density polyethylene

In our previous studies, we proved that the thermally induced shape memory effect was also exhibited by a carboxylated nitrile elastomer cured with zinc oxide because it contained labile ionic crosslinks, which were able to rearrange upon external deformation

In this work, the thermo-shrinkable properties of carboxylated acrylonitrile – butadiene elastomer (XNBR) and hydrogenated acrylonitrile – butadiene elastomer (HNBR) containing ionic crosslinks were studied. The XNBR was vulcanised with nanosized calcium, magnesium oxide or zinc oxide to ensure the formation of ionic crosslinks during vulcanisation. In the case of HNBR, nanosized calcium and magnesium oxides were coated with unsaturated carboxylic acids (itaconic, 2,4-pentadienoic, oleic, linoleic and linolenic acids) and applied as coagents in the peroxide vulcanisation of the elastomer. The application of these coagents led to formation of ionic crosslinks in the elastomer network. Heat-shrinkable polymers are widely used in packaging and in the cable industry; therefore,

Carboxylated nitrile elastomer XNBR (Krynac X7.50) containing 27 wt % acrylonitrile and 6.7 wt % carboxylic groups was obtained from Bayer C.O. The Mooney viscosity was (ML1+4 (100oC):47). Nanosized calcium oxide CaO (Aldrich), magnesium oxide MgO (Nanostructured & Amorphous Materials Inc., Houston, USA) and zinc oxide ZnO (Nanostructured & Amorphous Materials Inc., Houston, USA) were used as crosslinking agents. Hydrogenated acrylonitrile-butadiene elastomer HNBR (Therban 3407) containing 34 wt % acrylonitrile and 0.9 wt % of residual double bonds after hydrogenation was obtained from Bayer C.O. The Mooney viscosity was (ML1+4 (100oC):70). It was vulcanised with dicumyl peroxide DCP (Aldrich). Nanosized calcium oxide (Aldrich), magnesium oxide (Nanostructured & Amorphous Materials Inc., Houston, USA), itaconic acid IA (Fluka), 2-4-pentadienoic acid (Aldrich), oleic acid, linoleic acid and linolenic acid (Aldrich)

The nanosized metal oxides were mixed with a solution of modifying agent (unsaturated carboxylic acid) in acetone for 30 minutes during ultrasonic treatment (BANDELIN DT 255)

(Mishra et al., 2000) or poly(ethylene-vinyl-acetate) (Raychowdhury et al., 2000).

the shrinkability of XNBR and HNBR vulcanisates is technologically important.

Rubber compounds with the formulations given in Table 1 were prepared using a laboratory two-roll mill. The samples were cured at 160oC until they developed a 90% increase in torque, as measured by an oscillating disc rheometer.

The crosslink densities (νT) of the vulcanisates were determined by their equilibrium swelling in toluene, based on the Flory-Rehner equation (Flory & Rehner, 1943). The Huggins parameter of the XNBR-solvent interaction (χ) was calculated from the equation χ = 0.487 + 0.228Vr (Equation 1) (Przybyszewska & Zaborski, 2008), where Vr is the volume fraction of elastomer in the swollen gel, and χ = 0.501 + 0.273Vr for HNBR-solvent interaction (Equation 2) (Przybyszewska & Zaborski, 2009). To determine the content of ionic crosslinks in the elastomer network, samples were swollen in toluene in a dessicator with saturated ammonia vapour (25% aqueous solution). The ionic crosslink content (Δν) was calculated from Equation 3, where νA is the crosslink density determined for samples treated with ammonia vapour.

$$
\Delta \nu = \frac{\nu\_T - \nu\_A}{\nu\_T} \bullet 100\% \tag{3}
$$

The tensile properties of the vulcanisates were determined according to ISO-37 with a ZWICK 1435 universal machine.


**Table 1.** Composition of the XNBR and HNBR-based rubber compounds [phr]

### **2.4. Dynamic-mechanical analysis**

Dynamic - mechanical measurements were carried out in the tension mode using a DMA/SDTA861e analyser (Mettler Toledo). Measurements of the dynamic moduli were performed over the temperature range (-60 - 120oC) for XNBR and (-80 - 100oC) for HNBR with a heating rate of 2oC/min, a frequency of 1 Hz and a strain amplitude of 4 µm. The temperature of the elastomer glass transition was determined from the maximum of tanδ = f(T), where tanδ is the loss factor and T is the measurement temperature.

#### **2.5. Shrinkability measurements**

To measure the shrinkability of the XNBR and HNBR vulcanisates, the samples were stretched above their Tg at a temperature of 100oC until they reached an elongation of 200%

and left in the stretched form for 48 h. They were then stabilised in the stretched form for 4 h at (-7°C). Finally, the stretched samples were allowed to shrink above the ionic transition temperature at 70°C for 48 h. The length of the samples at each state of study was measured using the digital callipers (Preisser) with the measurement error 1 mm. The lengthwise shrinkage was calculated according to Equation 4, in which Sh is the percentage of shrinkability, Lstr is the length of the sample after stretching, and Lshr is the length of the shrunk sample. The maximum shrinkage, Shmax, was calculated from the length of the sample before stretching according to Equation 5, in which L0 is the original length of the sample before stretching. The physical properties of the vulcanisates were studied before and after the thermal treatment.

$$\mathbf{S}\_h\text{\%} = \frac{L\_{str} - L\_{str}}{L\_{str}} \mathbf{\bullet}100\tag{4}$$

$$\mathbf{S}\_{h\,\mathrm{max}}\,\%=\frac{L\_{str}-L\_{0}}{L\_{str}}\bullet 100\,\tag{5}$$

The continuous increase in temperature causes the recovery of the stretched sample deformation, which reflects the memory effect of the vulcanisate. The percentage recovery R is the ratio of the lengthwise shrinkage to the maximum shrinkage (Equation 6) (Khonakdar et al., 2007).

$$R\% = \frac{\mathcal{S}\_h}{\mathcal{S}\_{h\,\mathrm{max}}} \bullet 100\tag{6}$$

#### **2.6. Scanning Electron Microscopy (SEM)**

The morphology of the metal oxide particles and their dispersion in the elastomer matrix were estimated using scanning electron microscopy with a LEO 1530 SEM. The vulcanisates were broken down in liquid nitrogen, and the surfaces of the vulcanisate fractures were examined. Prior to the measurements, the samples were coated with carbon.

#### **3. Results and discussion**

#### **3.1. Thermo-shrinkable XNBR vulcanisates**

#### *3.1.1. Crosslink density and ionic crosslink content of XNBR vulcanisates*

The carboxylated nitrile elastomer XNBR reacts with metal oxides to form carboxylic salts, which behave as ionic crosslinks. These salts are able to associate, forming multiplets and clusters. This association is caused by the electrostatic interactions between multiplets, and it is impaired by the retractive elastic forces of the backbone chains. The restricted elastomer chain mobility in the proximity of the ionic clusters produces a hard phase surrounded with the soft elastomer matrix (Mishra et al., 2000). The biphasic structure of the XNBR crosslinked with metal oxide and the presence of labile ionic crosslinks provide possible routes to obtain thermo-shrinkable composites.

Because the ionic crosslinks play a crucial role in the return of the sample to the original shape and serve as memory points, the crosslink density and ionic crosslinks content in the elastomer network were determined in the first stage of the study. These results are presented in Figs. 2 and 3.

The nanosized calcium, magnesium and zinc oxides exhibited high crosslinking activity in XNBR. Their application led to the formation of ionic crosslinks in the elastomer network. The content of ionic crosslinks was in the range of 55% to 76%. Nanosized zinc oxide, for which the highest vulcanisate crosslink density and ionic crosslink content was observed (approximately 76%), appeared to be the most active, whereas the lowest activity was exhibited by CaO (with an ionic crosslink content of approximately 55%). Therefore, it could be supposed that the highest shape recovery would be obtained for vulcanisates crosslinked with nanosized ZnO.

**Figure 2.** Crosslink density of XNBR vulcanisates

186 Advanced Elastomers – Technology, Properties and Applications

**2.6. Scanning Electron Microscopy (SEM)** 

**3.1. Thermo-shrinkable XNBR vulcanisates** 

**3. Results and discussion** 

the thermal treatment.

et al., 2007).

and left in the stretched form for 48 h. They were then stabilised in the stretched form for 4 h at (-7°C). Finally, the stretched samples were allowed to shrink above the ionic transition temperature at 70°C for 48 h. The length of the samples at each state of study was measured using the digital callipers (Preisser) with the measurement error 1 mm. The lengthwise shrinkage was calculated according to Equation 4, in which Sh is the percentage of shrinkability, Lstr is the length of the sample after stretching, and Lshr is the length of the shrunk sample. The maximum shrinkage, Shmax, was calculated from the length of the sample before stretching according to Equation 5, in which L0 is the original length of the sample before stretching. The physical properties of the vulcanisates were studied before and after

% 100 *str shr*

*L L*

The continuous increase in temperature causes the recovery of the stretched sample deformation, which reflects the memory effect of the vulcanisate. The percentage recovery R is the ratio of the lengthwise shrinkage to the maximum shrinkage (Equation 6) (Khonakdar

> max % 100 *<sup>h</sup> h*

The morphology of the metal oxide particles and their dispersion in the elastomer matrix were estimated using scanning electron microscopy with a LEO 1530 SEM. The vulcanisates were broken down in liquid nitrogen, and the surfaces of the vulcanisate fractures were

The carboxylated nitrile elastomer XNBR reacts with metal oxides to form carboxylic salts, which behave as ionic crosslinks. These salts are able to associate, forming multiplets and clusters. This association is caused by the electrostatic interactions between multiplets, and it is impaired by the retractive elastic forces of the backbone chains. The restricted elastomer chain mobility in the proximity of the ionic clusters produces a hard phase surrounded with the soft elastomer matrix (Mishra et al., 2000). The biphasic structure of the XNBR

*<sup>R</sup> S S*

examined. Prior to the measurements, the samples were coated with carbon.

*3.1.1. Crosslink density and ionic crosslink content of XNBR vulcanisates* 

0 max% <sup>100</sup> *str <sup>h</sup> str*

*<sup>L</sup> <sup>S</sup>* (4)

*<sup>L</sup> <sup>S</sup>* (5)

(6)

*L L*

*<sup>h</sup> str*

The stretched samples were allowed to shrink above the ionic transition temperature at 70°C to initiate their return to the original shape. The influence of this thermal treatment on the ionic crosslink content and crosslink density of the vulcanisates was studied. From the data presented in Fig. 3, it follows that heating the samples to 70°C caused decomposition of the ionic crosslinks. The number of ionic crosslinks in the elastomer network was reduced by 21-34% compared to the vulcanisates before thermal treatment. The greatest number of ionic crosslinks was decomposed in the case of vulcanisates containing nanosized ZnO.

**Figure 3.** Ionic crosslink content in XNBR vulcanisates

#### *3.1.2. Thermo-shrinkability of XNBR vulcanisates*

Having determined the number of ionic crosslinks in the elastomer network and confirmed that these crosslinks could decompose during thermal treatment, we then examined the thermo-shrinkability of the vulcanisates. In Fig. 4, the percentage of thermo-shrinkability and the percentage shape recovery of the vulcanisates are presented.

The vulcanisates of XNBR crosslinked with nanosized calcium, magnesium and zinc oxides exhibited heat shrinkability. The greatest shrinkage upon heating (50%) was achieved for vulcanisates containing ZnO nanoparticles. The lower shrinkability of the vulcanisates with calcium and magnesium oxides (42% and 41%, respectively) was produced by the lower crosslink density and ionic crosslink content in their elastomer networks. Because the crosslinked points in the elastomer network serve as shape memory sites, a greater crosslink density improves the shrinkability of the XNBR. The vulcanisates demonstrated good shape recovery performance. The percentage shape recovery was in the range of 90-100% (Fig. 4). The highest percentage recovery was observed for vulcanisates containing ZnO, which had the greatest ionic crosslink content. The stretched XNBR samples shrunk upon heating above the temperature of the ionic transition because of the occurrence of ionic clusters in the elastomer network, which could rearrange or decompose. The results described in the previous section confirm that the decomposition of the ionic crosslinks is one of the reasons for the heat shrinkability of the XNBR vulcanisates containing nanosized CaO, MgO and ZnO.

**Figure 4.** Percentage thermo-shrinkability and shape recovery of XNBR vulcanisates

#### *3.1.3. Dynamic mechanical properties of XNBR vulcanisates*

188 Advanced Elastomers – Technology, Properties and Applications

**Figure 3.** Ionic crosslink content in XNBR vulcanisates

*3.1.2. Thermo-shrinkability of XNBR vulcanisates* 

ZnO.

and the percentage shape recovery of the vulcanisates are presented.

Having determined the number of ionic crosslinks in the elastomer network and confirmed that these crosslinks could decompose during thermal treatment, we then examined the thermo-shrinkability of the vulcanisates. In Fig. 4, the percentage of thermo-shrinkability

The vulcanisates of XNBR crosslinked with nanosized calcium, magnesium and zinc oxides exhibited heat shrinkability. The greatest shrinkage upon heating (50%) was achieved for vulcanisates containing ZnO nanoparticles. The lower shrinkability of the vulcanisates with calcium and magnesium oxides (42% and 41%, respectively) was produced by the lower crosslink density and ionic crosslink content in their elastomer networks. Because the crosslinked points in the elastomer network serve as shape memory sites, a greater crosslink density improves the shrinkability of the XNBR. The vulcanisates demonstrated good shape recovery performance. The percentage shape recovery was in the range of 90-100% (Fig. 4). The highest percentage recovery was observed for vulcanisates containing ZnO, which had the greatest ionic crosslink content. The stretched XNBR samples shrunk upon heating above the temperature of the ionic transition because of the occurrence of ionic clusters in the elastomer network, which could rearrange or decompose. The results described in the previous section confirm that the decomposition of the ionic crosslinks is one of the reasons for the heat shrinkability of the XNBR vulcanisates containing nanosized CaO, MgO and Dynamic - mechanical analysis was performed to confirm the biphasic structure of XNBR crosslinked with nanosized metal oxides, as well as the existence of ionic clusters in the elastomer network. The values of the glass transition temperature (Tg) are given in Table 2. The loss factor, tanδ, is presented in Fig. 5 as a function of temperature for the vulcanisates before thermal treatment.

**Figure 5.** Tan δ versus temperature for XNBR vulcanisates before thermal treatment

The results of the DMA analysis confirm the biphasic structure of the XNBR crosslinked with nanosized CaO, MgO and ZnO. The existence of two phase transitions was observed. The first transition is the glass transition of the XNBR at low temperatures, the maximum of which represents Tg. The determined glass transition temperatures for the vulcanisates with MgO and ZnO were approximately (-11.1oC) and (-10.5oC), respectively, whereas the Tg value was (-13.0oC) for the vulcanisate with CaO nanoparticles. The CaO vulcanisate most likely exhibits the lowest Tg because it contains the lowest crosslink density of the vulcanisates. The second peak, which is fuzzy and has a low-intensity, was observed in the temperature range of (50-100oC). This peak corresponds to the ionic transition at high temperature that is caused by the occurrence of a hard phase arising from ionic associations (ionic clusters or aggregates). These transitions were observed for all XNBR vulcanisates. Therefore, the existence of a biphasic structure in XNBR crosslinked with metal oxides was confirmed.


**Table 2.** Glass transition temperature of XNBR vulcanisates

**Figure 6.** Tan δ versus temperature for XNBR vulcanisates after thermal shrinking

DMA measurements were also performed for the vulcanisates after the thermal shrinking process. These results are presented in Fig. 6. Heating the samples to the ionic transition temperature decreased the glass transition temperatures of the vulcanisates. This reduction was most likely caused by the decomposition of the ionic crosslinks, leading to a reduction of the crosslink density of the vulcanisates. The highest decrease in the Tg value was observed for the vulcanisate containing nanosized ZnO, for which the greatest amount of ionic crosslinks decomposed during the thermal treatment. The second fuzzy peak corresponding to the ionic transition disappeared, which suggests that the return of the vulcanisates to their original shape was caused by the decomposition of ionic crosslinks, along with changes in or the disappearance of the biphasic structure of the crosslinked elastomer.

#### *3.1.4. Mechanical properties of XNBR vulcanisates*

190 Advanced Elastomers – Technology, Properties and Applications

**Table 2.** Glass transition temperature of XNBR vulcanisates

confirmed.

The results of the DMA analysis confirm the biphasic structure of the XNBR crosslinked with nanosized CaO, MgO and ZnO. The existence of two phase transitions was observed. The first transition is the glass transition of the XNBR at low temperatures, the maximum of which represents Tg. The determined glass transition temperatures for the vulcanisates with MgO and ZnO were approximately (-11.1oC) and (-10.5oC), respectively, whereas the Tg value was (-13.0oC) for the vulcanisate with CaO nanoparticles. The CaO vulcanisate most likely exhibits the lowest Tg because it contains the lowest crosslink density of the vulcanisates. The second peak, which is fuzzy and has a low-intensity, was observed in the temperature range of (50-100oC). This peak corresponds to the ionic transition at high temperature that is caused by the occurrence of a hard phase arising from ionic associations (ionic clusters or aggregates). These transitions were observed for all XNBR vulcanisates. Therefore, the existence of a biphasic structure in XNBR crosslinked with metal oxides was

Vulcanisate Tg before thermal treatment, oC Tg after thermal shrinking, oC

CaO -13.0 -15.0 MgO -11.1 -13.5 ZnO -10.5 -16.0

**Figure 6.** Tan δ versus temperature for XNBR vulcanisates after thermal shrinking

DMA measurements were also performed for the vulcanisates after the thermal shrinking process. These results are presented in Fig. 6. Heating the samples to the ionic transition temperature decreased the glass transition temperatures of the vulcanisates. This reduction was most likely caused by the decomposition of the ionic crosslinks, leading to a reduction of the crosslink density of the vulcanisates. The highest decrease in the Tg value was The mechanical properties of the vulcanisates (especially their tensile strength) are technologically important. Therefore, the effect of the thermal shrinking process on the tensile strength and elongation at break of the vulcanisates was studied. These results are presented in Table 3.


**Table 3.** Tensile strength and elongation at break of XNBR vulcanisates before thermal treatment (TS0, EB0) and after thermal shrinking (TSshr, EBshr)

Regarding the properties of the vulcanisates before the thermal treatment, the greatest tensile strength was exhibited by the vulcanisate with nanosized MgO, whereas the lowest was produced by CaO. One possible reason for the different activities of the metal oxides is their tendency to agglomerate in the elastomers (see Fig. 7).

The CaO nanoparticles were poorly dispersed in the elastomer matrix (Fig. 7a). They created microsized agglomerates with complex structures, which displayed poor adhesion to the elastomer. These agglomerates acted as centres for stress concentration in the vulcanisates during the deformation and initiate breakage of the sample under external stress. As a result, the tensile strength of the vulcanisates decreased. The ZnO nanoparticles also created microsized agglomerates that were smaller than the CaO particles and surrounded by an elastomer film (Fig. 7c). The wetting of the ZnO agglomerates with the elastomer probably produced the better mechanical properties of the ZnO vulcanisates, despite the heterogeneous dispersion of the nanoparticles. The MgO nanoparticles revealed the weakest ability to agglomerate in the XNBR, creating clusters of approximately 3 µm in size that were tightly bound to the elastomer matrix (Fig. 7b). The highest tensile strength was observed for the vulcanisate with MgO. The elongation at break was the lowest for the vulcanisate containing ZnO nanoparticles, and this result was correlated with the crosslink density of the examined vulcanisates.

The stretching of the samples at high temperature, their stabilisation in the stretched state and the shrinkage above the ionic transition temperature reduced the tensile strength and elongation at break of the vulcanisates by decomposing the ionic crosslinks and reducing the crosslink density of the vulcanisates. However, these mechanical properties remained

c) XNBR/ZnO

**Figure 7.** SEM images of XNBR vulcanisates

satisfactory, especially in the case of the vulcanisates containing MgO and ZnO (TSshr, EBshr in Table 3).

### **3.2. Thermo-shrinkable HNBR vulcanisates**

### *3.2.1. Crosslink density and ionic crosslink content for HNBR vulcanisates*

Studies performed on the XNBR elastomer crosslinked with nanosized metal oxides confirmed the existence of a biphasic structure and the presence of ionic crosslinks in the elastomer network, which are able to rearrange or decompose upon heating to produce thermo-shrinkable vulcanisates. Therefore, to obtain thermo-shrinkable vulcanisates from the HNBR elastomer, multifunctional crosslinking coagents based on nanosized calcium and magnesium oxides were used in combination with unsaturated carboxylic acids (UCAs). UCAs containing easily abstractable hydrogen atoms and readily accessible double bonds were grafted onto the powder surface during the modification process. Because of their multifunctionality, this type of coagent based on zinc oxide was proven to be able to react directly with the elastomer and form effective covalent bonds. Moreover, they can form ionic crosslinks that increase the vulcanisate crosslink density and improve tensile strength (Przybyszewska & Zaborski, 2009). Because the presence of ionic crosslinks is crucial for the thermo-shrinkability of vulcanisates, the effect of these coagents on the crosslink density of the vulcanisates and their ionic crosslink content was studied. These results are given in Figs. 8 and 9.

192 Advanced Elastomers – Technology, Properties and Applications

**Figure 7.** SEM images of XNBR vulcanisates

**3.2. Thermo-shrinkable HNBR vulcanisates** 

in Table 3).

c) XNBR/ZnO

satisfactory, especially in the case of the vulcanisates containing MgO and ZnO (TSshr, EBshr

Studies performed on the XNBR elastomer crosslinked with nanosized metal oxides confirmed the existence of a biphasic structure and the presence of ionic crosslinks in the elastomer network, which are able to rearrange or decompose upon heating to produce thermo-shrinkable vulcanisates. Therefore, to obtain thermo-shrinkable vulcanisates from the HNBR elastomer, multifunctional crosslinking coagents based on nanosized calcium and magnesium oxides were used in combination with unsaturated carboxylic acids (UCAs). UCAs containing easily abstractable hydrogen atoms and readily accessible double bonds were grafted onto the powder surface during the modification process. Because of their multifunctionality, this type of coagent based on zinc oxide was proven to be able to react

*3.2.1. Crosslink density and ionic crosslink content for HNBR vulcanisates* 

a) XNBR/CaO b) XNBR/MgO

The application of coagents based on nanosized CaO and MgO grafted with UCA caused the formation of ionic crosslinks in the elastomer network. The greatest content of ionic crosslinks was obtained for vulcanisates containing nanosized CaO in combination with oleic and linoleic acid, as well as for MgO with itaconic and linoleic acid. The results of these studies confirmed that these coagents contributed to the increased vulcanisation efficiency. The ionic crosslink content should affect the thermo-shrinkability of the vulcanisates considerably. The crosslinked points in the elastomer network are believed to serve as memory points, enhancing the heat shrinkability.

To confirm that the decomposition and/or rearrangement of the ionic crosslink aggregates is one of the causes of heat shrinkability in the examined HNBR vulcanisates, the crosslink density and ionic crosslink content were determined for the vulcanisates after thermal shrinking and compared with the values of the vulcanisates before thermal treatment.

The data presented in Figs. 8 and 9 suggest that the ionic crosslinks that decomposed during the shrinking of vulcanisates at 70oC allowed the samples to return to their original shape. As a result of the ionic crosslink decomposition, the crosslink density of the vulcanisates decreased. The most considerable reduction of the ionic crosslink number after shrinkage above the ionic transition temperature was observed for the vulcanisates containing CaO-OA, CaO-LA, MgO-IA and MgO-OA as coagents, which were characterised by the highest ionic crosslink content before the heat treatment. These vulcanisates may be expected to show the greatest thermal shrinkage and shape recovery.

**Figure 8.** Crosslink density of HNBR vulcanisates

**Figure 9.** Ionic crosslink content of HNBR vulcanisates

#### *3.2.2. Thermo-shrinkability of HNBR vulcanisates*

As mentioned before, the shrinkage of polymer is caused by an internal rearrangement of the structural elements within the stretched sample (Mishra et al., 2000). In contrast to the covalent crosslinks formed during conventional vulcanisation with peroxides, ionic crosslinks are multifunctional and labile. Ionic crosslinks group together, forming clusters that are immersed in the elastomer matrix. Moreover, ionic clusters can rearrange in the elastomer matrix upon external deformation or temperature change.

**Figure 10.** Percentage thermo-shrinkability of HNBR vulcanisates

The crosslinked samples were heated and stretched at 100oC (a temperature above the glass transition temperature). Cooling these samples in this stretched state stabilised the temporary shape. Finally, the stretched samples were allowed to shrink above the ionic transition temperature at 70°C. The heat shrinkability values and percentage recovery of the HNBR samples containing 7 phr of the coagents are presented in Figs. 10 and 11.

194 Advanced Elastomers – Technology, Properties and Applications

**Figure 9.** Ionic crosslink content of HNBR vulcanisates

*3.2.2. Thermo-shrinkability of HNBR vulcanisates* 

elastomer matrix upon external deformation or temperature change.

**Figure 10.** Percentage thermo-shrinkability of HNBR vulcanisates

As mentioned before, the shrinkage of polymer is caused by an internal rearrangement of the structural elements within the stretched sample (Mishra et al., 2000). In contrast to the covalent crosslinks formed during conventional vulcanisation with peroxides, ionic crosslinks are multifunctional and labile. Ionic crosslinks group together, forming clusters that are immersed in the elastomer matrix. Moreover, ionic clusters can rearrange in the The vulcanisates of the HNBR elastomers crosslinked in the presence of coagents based on nanosized CaO and MgO in combination with unsaturated carboxylic acids exhibited heat shrinkability. The greatest shrinkage upon heating (approximately 35%) was achieved for vulcanisates containing the CaO-OA, CaO-LA and MgO-IA coagents (Fig. 10). The lower shrinkability of the vulcanisates with other coagents is a result of the lower crosslink density and ionic crosslink number in their elastomer networks (see Figs. 8 and 9). Because the crosslinked points in the elastomer network serve as shape memory sites, a higher crosslink density improves the shrinkability of the vulcanisate.

The HNBR-based vulcanisates demonstrated good shape recovery performance, with a shape recovery percentage in the range of 91-100% (Fig. 11). The greatest percentage recovery was obtained for the vulcanisates with the highest ionic crosslink content. Moreover, the most considerable reduction of the ionic crosslink number during the thermal shrinking process was observed for these vulcanisates. This result confirms the assumption that the decomposition of ionic crosslinks is the most important reason for the shrinkability and shape recovery of HNBR vulcanisates. It can be concluded that the ionic crosslinks formed in the elastomer by the coagents serve as memory points in the elastomer network.

**Figure 11.** Percentage recovery of HNBR vulcanisates

#### *3.2.3. Dynamic mechanical properties of HNBR vulcanisates*

Dynamic - mechanical analysis was performed to confirm the existence of ionic clusters in the elastomer network. The loss factor (tanδ) of the HNBR vulcanisates with coagents is presented as a function of temperature in Figs. 12 and 13 as an example. The values of the glass transition temperature (Tg) are given in Table 4.

**Figure 12.** Tan δ versus temperature for HNBR vulcanisates before thermal treatment

**Figure 13.** Tan δ versus temperature for HNBR vulcanisates after thermal shrinking


**Table 4.** Glass transition temperature of HNBR vulcanisates

glass transition temperature (Tg) are given in Table 4.

*3.2.3. Dynamic mechanical properties of HNBR vulcanisates* 

Dynamic - mechanical analysis was performed to confirm the existence of ionic clusters in the elastomer network. The loss factor (tanδ) of the HNBR vulcanisates with coagents is presented as a function of temperature in Figs. 12 and 13 as an example. The values of the

**Figure 12.** Tan δ versus temperature for HNBR vulcanisates before thermal treatment

**Figure 13.** Tan δ versus temperature for HNBR vulcanisates after thermal shrinking

The DMA analysis revealed the biphasic structure of the HNBR crosslinked in the presence of coagents based on nanosized CaO and MgO grafted with UCA. Two phase transitions were observed. The glass transition of the HNBR, which occurs at Tg, was observed in the range from (-21.2oC) to (-19.1oC) for vulcanisates with CaO and from (-20,6oC) to (-17.8oC) for vulcanisates containing MgO. The determined Tg values were correlated with the crosslink density of the vulcanisates. A fuzzy low-intensity peak was observed in the temperature range of (40-100oC) (corresponding to the ionic transition) because of the occurrence of a hard phase arising from the existence of the ionic aggregates.

Heating to the ionic transition temperature decreased the glass transition temperature of the vulcanisates by decomposing ionic crosslinks, reducing the crosslink density of the vulcanisates. The greatest decrease of the Tg value occurred in the case of the vulcanisates for which the greatest number of ionic crosslinks decomposed during heating (CaO-OA, CaO-LA, MgO-IA, MgO-OA). The fuzzy peak corresponding to the ionic transition disappeared; consequently, it can be concluded that the shape recovery of the HNBR vulcanisates was caused by the decomposition of the ionic crosslinks and changes in the biphasic structure of the crosslinked elastomer, similarly to XNBR.

### *3.2.4. Mechanical properties of HNBR vulcanisates*

Having established the ability of the HNBR vulcanisates to shrink upon exposure to the ionic degradation temperature and return to their original shape, we then examined their mechanical properties.

It is known that the formation of coagent bridges, which are labile ionic crosslinks inside the elastomer network formed during vulcanisation, improved the tensile properties of the vulcanisates (Przybyszewska & Zaborski 2008, 2009). However, it is reasonable to investigate the effect of the heat treatment of the vulcanisates on their mechanical properties.

As could be supposed, the thermal treatment of the vulcanisates during their shrinkage deteriorated their tensile strength as a result of the decomposition of their ionic crosslinks. Their tensile strength was reduced by 5% to 30% in comparison with that of the vulcanisates before thermal treatment (Fig. 14). The vulcanisates subjected to shrinking at elevated temperature also exhibited a lower elongation at break (Fig. 15).

**Figure 14.** Tensile strength of HNBR vulcanisates

**Figure 15.** Elongation at break of HNBR vulcanisates

#### **4. Conclusions**

The thermo-shrinkable properties of carboxylated (XNBR) and hydrogenated (HNBR) acrylonitrile butadiene elastomer were studied. XNBR was crosslinked with nanosized calcium, magnesium and zinc oxide to ensure the formation of ionic crosslinks in the elastomer matrix, which can serve as memory points and enhance its heat shrinkability. Similarly, HNBR was crosslinked in the presence of coagents based on nanosized calcium and magnesium oxides grafted with unsaturated carboxylic acids to achieve ionic crosslinks. The examined samples were allowed to shrink at a temperature above the ionic transition temperature. The XNBR vulcanisates revealed thermo-shrinkability in the range of 41% for MgO to 50% for ZnO, whereas the percentage of shape recovery was in the range of 90- 100%. Good shape recovery performance was also observed for the HNBR vulcanisates. The greatest shrinkage upon heating (approximately 35%) was achieved for the vulcanisates containing CaO-OA, CaO-LA and MgO-IA coagents. The percentage of shape recovery of these vulcanisates was in the range of 91-100%.

The thermo-shrinkability value and shape recovery ratio were strongly correlated with the ionic crosslink content of the elastomer network, and the number of ionic crosslinks was reduced by heating the samples above their ionic transition temperature. Greater ionic crosslink contents and their more significant decomposition corresponded to vulcanisates with greater thermo-shrinkability. DMA measurements confirmed the presence of ionic crosslinks and the existence of a biphasic structure in both the XNBR and HNBR elastomers. A fuzzy low-intensity peak was observed in the (tan δ) curve, in the temperature range of (50-100oC), which corresponds to the ionic transition that occurs at high temperatures as a result of the occurrence of a hard phase arising from ionic aggregates. This peak disappeared when the samples were heated above their ionic transition temperature. Therefore, it could be concluded that the thermo-shrinkability of the XNBR and HNBR vulcanisates and their shape recovery was a result of the decomposition of ionic crosslinks and changes in or the disappearance of the biphasic structure of the crosslinked elastomer.

### **Author details**

198 Advanced Elastomers – Technology, Properties and Applications

**Figure 14.** Tensile strength of HNBR vulcanisates

**Figure 15.** Elongation at break of HNBR vulcanisates

**4. Conclusions** 

temperature also exhibited a lower elongation at break (Fig. 15).

Their tensile strength was reduced by 5% to 30% in comparison with that of the vulcanisates before thermal treatment (Fig. 14). The vulcanisates subjected to shrinking at elevated

The thermo-shrinkable properties of carboxylated (XNBR) and hydrogenated (HNBR) acrylonitrile butadiene elastomer were studied. XNBR was crosslinked with nanosized calcium, magnesium and zinc oxide to ensure the formation of ionic crosslinks in the elastomer matrix, which can serve as memory points and enhance its heat shrinkability. Similarly, HNBR was crosslinked in the presence of coagents based on nanosized calcium Magdalena Maciejewska and Alicja Krzywania-Kaliszewska *Technical University of Lodz; Institute of Polymer and Dye Technology, Poland* 

### **Acknowledgement**

The authors wish to acknowledge the Polish Ministry of Science and Higher Education as well as the National Centre of Research and Development for supporting this study.

### **5. References**


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### **Chapter 8**

## **Elastomer Application in Microsystem and Microfluidics**

Shuang (Jake) Yang and Kunqiang Jiang

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/48121

### **1. Introduction**

Elastomer has been widely used in academic and industry since it was invented in the nineteen century 1. By definition, elastomer is a polymer which is viscoelastic and able to regain its shape after deformation. Compared to other materials, elastomer has notably low Young's modulus and high yield strain. In general, amorphous elastomer is made of carbon, hydrogen, oxygen and/or silicon whose glass transition temperature is well below the application temperature, e.g., room temperature. As a result, it has excellent properties in elasticity 2, transparency 3, permeability 4, and insulation 5.

One of the most commonly used elastomers is silicone elastomer such as poly(dimethylsiloxane) (PDMS), which has been widely used in electronics and microfluidics due to its elasticity, optical transparency, UV transmission, permeability, biocompatibility, and availability. PDMS generally consists of repeat unit dimethylsiloxane. The repeat methyl is modifiable using functional groups such as urea or bis-urea for development of high mechanical strength of elastomer. The modified PDMS presents both a high mechanical strength and elasticity at room temperature due to functionality of the cross-linking domains and intrinsic reversibility of the supramolecular chains. PDMS is optically transparent and is usually inert, non-toxic and non-flammable 6. PDMS is viscoelastic and is normally measured using dynamic mechanical analysis (DMA), in which a specialized instrument is used to determine the material's flow characteristics over a wide range of temperature, flow rates, and deformations. In general, the shear modulus of PDMS is in the range of 100 kPa to 3 MPa, varied with the preparation conditions.

Characterization of elastomer by dynamic-mechanical testing includes temperature sweep, frequency sweep, strain sweep, damping and friction properties. The dynamic mechanical measurement is particularly useful to characterize conductive particles in elastomer for

© 2012 Yang and Jiang, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

understanding the particles effect on overall elastomer mechanical properties. It is very common to measure damping factor (tan) change with temperature for understanding the elastomer degree of crosslinking 7. The dynamic mechanical testing provides information on both uncrosslinked and crosslinked phases of the elastomer, as well as viscoelasticity. To this end, nanoindentation plays a significant role on viscoelastic characterization of elastomers for their creep and relaxation properties 8, 9.

Elastomer is an excellent matrix which can mix with other materials such as conductive particles and thermoplastics to enhance mechanical properties. For example, silver particles are added into silicone elastomer matrix to formation of conductive elastomer for highdensity electrical contacts in microelectronics. Polypropylene and ethylene-propylene copolymer elastomer significantly improves the dynamic mechanical, tensile and impact properties and even changes the modes of failure 10. The size, shape and spatial packing of elastomer inclusions are the most significant factors in controlling the mechanical behavior of blends. Since PDMS can be cured with different ratio of dimethylsiloxane and curing reagent, it is very useful to fabricate microfluidic devices for a variety of applications, including cell culture 11, ELISA assay 12, capillary electrophoresis 13, DNA sequencing 14.

In this chapter, we briefly discussed elastomer elastic and viscoelastic properties, as well as permeability. Elastomer is widely used in microelectronics and microfluidics due to all or part of these characteristics. Then we discussed several applications in microelectronics and microfluidics using elastomer as substrate. Finally, we introduce potential development of elastomer.

### **2. Elastomer property and its applications**

Elasticity is the ability of a material to return to its original shape and size after being stretched, compressed, twisted or bent 15. Elastic deformation (change of shape or size) lasts only as long as a deforming force is applied to the object, and disappears once the applied force is removed. The elasticity of elastomer makes it extremely useful in microelectronics and microfluidics. For example, the Intel CPU uses land grid array sockets (LGA), in which the socket has over thousands of elastomer contacts that consist of silver particles in elastomer matrix 16, 17. When socket, the CPU and printed circuit board (PCB) are assembled by compressing force, the silver-embedded elastomer is compressed to increase contact area between socket and the CPU, and between socket and the PCB. The excellent elastic deformation of elastomer can keep the electrically tight contact for minimum contact resistance. Elasticity of PDMS minimizes its stress relaxation and creep effect, which is severe in other polymer with low elasticity. The excellent elasticity of PDMS allows to release from microfluidic masters without damaging PDMS device 18, 19 and significant deflection to occur when a pressure difference develops 20. Therefore, it is important to characterize PDMS elasticity and understand how to improve such property for various applications.

The most common elasticity experiment for many decades involved mechanically blending a reinforcing filler such as carbon black into natural rubber, crosslinking the elastomer with sulfur, and characterizing it network structure by stress-strain measurements in simple extension, or by extents of equilibrium swelling 21. However, in most instances, such as thinfilm elastomer, land grid array elastomer contact, and PDMS microfluidics, these methods may not be applicable due to effects of supporting materials. To characterize elastomer in those applications, nanoindentation has been an emerging technology for elastomer characterization and is an indispensable tool in many instances where other techniques are inapplicable. Nanoindentation is such tool for characterization of polymer on its viscoelasticity. Indentation technique was first introduced as advent of indentation fracture mechanics 22 in 1975 and ever since it has been developed for characterization of elasticplastic materials using the Oliver-Pharr's method 23. No valid method was reported for polymer nanoindentation analysis based on stress-strain experiments until the viscoelastic model was proposed in 2004 and other emerging studies that followed 8, 24, 25. Because of viscoelasticity, elastic-plastic model cannot be directly used to derive modulus and hardness. Also these measurements are dependent on the experimental parameters, such as loading rate, maximum load, holding duration, unloading rate, and minimum unloading force 26. To minimize the effects by those experimental parameters during indentation, the experimental conditions were proposed in such a way that fast loading was applied for minimizing viscoelastic deformation, long dwell time for full viscoelastic deformation, and quick unloading rate. The fast loading and unloading showed that only elastic-plastic deformation was observed, while viscoelastic deformation was observed in the dwell duration. Also, it was observed that the viscoelastic deformation can be measured by keeping a minimum force on the indenter so that the instrument can record the deformation progress with time. The modulus and viscosity coefficient of elastomer can be derived using flat-ended punch-tip 8, the elastic-viscoelastic-viscous deformation is described by following equation,

204 Advanced Elastomers – Technology, Properties and Applications

elastomers for their creep and relaxation properties 8, 9.

**2. Elastomer property and its applications** 

such property for various applications.

elastomer.

understanding the particles effect on overall elastomer mechanical properties. It is very common to measure damping factor (tan) change with temperature for understanding the elastomer degree of crosslinking 7. The dynamic mechanical testing provides information on both uncrosslinked and crosslinked phases of the elastomer, as well as viscoelasticity. To this end, nanoindentation plays a significant role on viscoelastic characterization of

Elastomer is an excellent matrix which can mix with other materials such as conductive particles and thermoplastics to enhance mechanical properties. For example, silver particles are added into silicone elastomer matrix to formation of conductive elastomer for highdensity electrical contacts in microelectronics. Polypropylene and ethylene-propylene copolymer elastomer significantly improves the dynamic mechanical, tensile and impact properties and even changes the modes of failure 10. The size, shape and spatial packing of elastomer inclusions are the most significant factors in controlling the mechanical behavior of blends. Since PDMS can be cured with different ratio of dimethylsiloxane and curing reagent, it is very useful to fabricate microfluidic devices for a variety of applications, including cell culture 11, ELISA assay 12, capillary electrophoresis 13, DNA sequencing 14.

In this chapter, we briefly discussed elastomer elastic and viscoelastic properties, as well as permeability. Elastomer is widely used in microelectronics and microfluidics due to all or part of these characteristics. Then we discussed several applications in microelectronics and microfluidics using elastomer as substrate. Finally, we introduce potential development of

Elasticity is the ability of a material to return to its original shape and size after being stretched, compressed, twisted or bent 15. Elastic deformation (change of shape or size) lasts only as long as a deforming force is applied to the object, and disappears once the applied force is removed. The elasticity of elastomer makes it extremely useful in microelectronics and microfluidics. For example, the Intel CPU uses land grid array sockets (LGA), in which the socket has over thousands of elastomer contacts that consist of silver particles in elastomer matrix 16, 17. When socket, the CPU and printed circuit board (PCB) are assembled by compressing force, the silver-embedded elastomer is compressed to increase contact area between socket and the CPU, and between socket and the PCB. The excellent elastic deformation of elastomer can keep the electrically tight contact for minimum contact resistance. Elasticity of PDMS minimizes its stress relaxation and creep effect, which is severe in other polymer with low elasticity. The excellent elasticity of PDMS allows to release from microfluidic masters without damaging PDMS device 18, 19 and significant deflection to occur when a pressure difference develops 20. Therefore, it is important to characterize PDMS elasticity and understand how to improve

The most common elasticity experiment for many decades involved mechanically blending a reinforcing filler such as carbon black into natural rubber, crosslinking the elastomer with sulfur, and characterizing it network structure by stress-strain measurements in simple

$$h = \frac{P\_0 h\_{ln}}{E\_0 A\_0} + \sum\_{1}^{n} \frac{P\_0 h\_{ln}}{E\_l A\_0} (1 - e^{-E\_l t}/\eta\_l) + \frac{h\_{ln}}{\eta\_o} t^{\frac{1}{2}}$$

where *P* represents load, *h* is deformation and subscript *e* is elastic, coefficient of viscosity, *E* modulus, *t* time. Elastic modulus is thus derived by *P0hin/A0he*. This model provides a means for characterization of elastomer, such as investigation of effect of crosslinking, work of adhesion and fluid environment on elastic modulus 24. Nanoindentation on PDMS indicated that it is capable of measuring several order of magnitudes of elastic modulus of PDMS, with expected increase of the elastic modulus at higher crosslinked PDMS (2.80 MPa at 10:1 ratio, and 0.88 MPa at 30:1 ratio; monomer:crosslinker, wt%) 24, 27.

Elastomer is permeable to solvents, gases, moisture or even some small molecules 4, 28-30. Permeability has been employed to fabricate non-contact pump for the manipulation of aqueous solutions within PDMS microfluidic devices 20 or supply oxygen for large-scale culture of hepatocytes 31. On the other hand, PDMS permeability limits its application to microfluidic technology, in which coating such as glass-like layer using sol-gel chemistry is developed on PDMS to minimize permeability 32. The permeability of a PDMS film or membrane to a penetrant (i.e., gas or solvent) is defined as 28, 33,

$$P\_A = \frac{N\_A l}{p\_{2A} - p\_{1A}}$$

Where *PA* is the permeability coefficient, *NA* is the steady-state penetrant flux through the film, *p1A* and *p2A* are the downstream and upstream partial pressures of component *A*, and l is the membrane or film thickness. In general, PDMS permeability depends strongly on the feeding gas/solvent composition and temperature: an increase in vapor concentration or liquid pressure leads to increased chain mobility at constant feeding pressure and temperature, resulting in high diffusivity and permeability. The increase of permeability with pressure is attributed to the swelling of PDMS.

Due to the permeability nature of PDMS, many researchers synthesized co-polymer which contained PDMS as a permeable medium for gas/solvent. For example, the 'smart' tricomponent amphiphilic membranes consisting of poly(ethylene glycol) (PEG), PDMS and polypentamethylcyclopentasiloxane (PD5) domains were synthesized for biological applications, including immunoisolation of cells 29, 30. The high oxygen permeability of PDMS makes it feasible for the application in microdevices in cell culture 34. Compared to large-scale system / platform, microfluidics provides suitable environments for sample interaction, such as cell culture, with the significant increase of surface-to-volume ratio while the fluidic behavior is similar to the environments in vivo 29. Integration of PDMS into microfluidic platform can implement studies on cell-to-cell interactions and understanding of cell behaviors in vitro, emulation of situations observed in vivo. Several studies have demonstrated that microfluidic devices consisted of two PDMS layers were capable for cell culture 35. They further fabricated a complex microfluidics which consisted of stacking ten PDMS layers, in which four cell culture chambers and one-oxygen chamber inserts between.

### **3. Elastomer as electrical contact matrix in microelectronics**

Insulation elastomer in microelectronics plays a significant role to prevent a short-circuit 36. Electrical insulating polymers have been introduced nearly one hundred years ago and proved to be excellent long-term physical and electrical properties, including weatherability, moisture-sealing, corrosion resistant and durability. The polymer insulating materials are used in a wide range of applications such as surge arresters, insulators, insulation enhancement, and bushings 5. The initial polymers used in insulating materials were mainly on radiation-crosslinked and semicrystalline polyolefin copolymers. With development of manufacturing and materials technologies, elastomer has been increasingly designed for insulating applications 5, 37.

The elastomer formulations have been optimized so that an exceptional electrical and weathering performance is achieved, which are equal or above that of polyolefin copolymer materials. The silicone elastomer has many advantages such as hydrophobicity, hydrophobic recovery, weathering resistance, processability, and elastomeric mechanical properties. These advantages make silicone elastomer suitable for outdoor insulating applications. For example, PDMS has been used for high-voltage (HV) outdoor insulation in the real world as a replacement over conventional porcelain and glass insulation. As one silicone elastomer, PDMS as HV insulating materials is light weight, has superior vandal resistance, and contains better contamination performance. However, many research works have been conducted on molecular modification in order to improve PDMS properties, including degradation under exposure to discharge and arcing, mechanical strength etc. It has been demonstrated that hydrophobicity of PDMS can be rendered by exposure to discharges, which can be recovered when exposure ceases. Their studies showed that the lowly crosslinked PDMS can readily reorient between hydrophilic and hydrophobic states, therefore improving its weathering capabilities 38.

206 Advanced Elastomers – Technology, Properties and Applications

with pressure is attributed to the swelling of PDMS.

insulating applications 5, 37.

Where *PA* is the permeability coefficient, *NA* is the steady-state penetrant flux through the film, *p1A* and *p2A* are the downstream and upstream partial pressures of component *A*, and l is the membrane or film thickness. In general, PDMS permeability depends strongly on the feeding gas/solvent composition and temperature: an increase in vapor concentration or liquid pressure leads to increased chain mobility at constant feeding pressure and temperature, resulting in high diffusivity and permeability. The increase of permeability

Due to the permeability nature of PDMS, many researchers synthesized co-polymer which contained PDMS as a permeable medium for gas/solvent. For example, the 'smart' tricomponent amphiphilic membranes consisting of poly(ethylene glycol) (PEG), PDMS and polypentamethylcyclopentasiloxane (PD5) domains were synthesized for biological applications, including immunoisolation of cells 29, 30. The high oxygen permeability of PDMS makes it feasible for the application in microdevices in cell culture 34. Compared to large-scale system / platform, microfluidics provides suitable environments for sample interaction, such as cell culture, with the significant increase of surface-to-volume ratio while the fluidic behavior is similar to the environments in vivo 29. Integration of PDMS into microfluidic platform can implement studies on cell-to-cell interactions and understanding of cell behaviors in vitro, emulation of situations observed in vivo. Several studies have demonstrated that microfluidic devices consisted of two PDMS layers were capable for cell culture 35. They further fabricated a complex microfluidics which consisted of stacking ten PDMS layers, in which four cell culture chambers and one-oxygen chamber inserts between.

**3. Elastomer as electrical contact matrix in microelectronics** 

Insulation elastomer in microelectronics plays a significant role to prevent a short-circuit 36. Electrical insulating polymers have been introduced nearly one hundred years ago and proved to be excellent long-term physical and electrical properties, including weatherability, moisture-sealing, corrosion resistant and durability. The polymer insulating materials are used in a wide range of applications such as surge arresters, insulators, insulation enhancement, and bushings 5. The initial polymers used in insulating materials were mainly on radiation-crosslinked and semicrystalline polyolefin copolymers. With development of manufacturing and materials technologies, elastomer has been increasingly designed for

The elastomer formulations have been optimized so that an exceptional electrical and weathering performance is achieved, which are equal or above that of polyolefin copolymer materials. The silicone elastomer has many advantages such as hydrophobicity, hydrophobic recovery, weathering resistance, processability, and elastomeric mechanical properties. These advantages make silicone elastomer suitable for outdoor insulating applications. For example, PDMS has been used for high-voltage (HV) outdoor insulation in the real world as a replacement over conventional porcelain and glass insulation. As one silicone elastomer, PDMS as HV insulating materials is light weight, has superior vandal resistance, and contains better contamination performance. However, many research works Recently, PDMS has been micropatterned using a photoresist lift-off technique for selective electrical insulation in microelectrode array applications 39. The micropatterning technique is able to manufacture PDMS patterns with feature as small as 15 µm and various thicknesses down to 6 µm. With micropatterned PDMS insulation layer, the electrical resistance between adjacent electrodes is within the specification. Additionally, the micropatterned PDMS insulation can be applied to biosensing microdevices such as in an extensive neuronal network. Other researchers used PDMS to develop a replaceable insulator for a single-use planar microelectrode array (MEA) in the study of electrogenic tissues 40. They demonstrated applications using microstencils for rejuvenation of an old MEA and the fabrication of a single-use MEA.

Conductive materials are used to produce shielding gaskets for military, aerospace, electronics, and communications. The conductive elastomers are designed to balance requirements for electrical conductivity, thermal management, and cost performance 41. Conductive elastomer materials are ideal for customer applications requiring both excellent electromagnetic interference shielding and environmental sealing across a wide range of temperature. The conductive elastomers could be blended polyacetylene with different types of thermoplastic elastomers: styrene-butadiene-styrene, styrene-isoprene-styrene, and styrene-ethylenebutylene-styrene tri-block copolymers 42. The polyacetylene is a highly conjugated polymer which contains high concentration of unsaturated sites. Those sites are easily attacked by ozone / UV light and reacted with elastomer, formation of conductive elastomer composites. Or insulation elastomers are coated with compliant electrode material on both sides of the elastomer film 43. The dielectric elastomers were actuated by means of electrostatic forces applied via compliant electrodes. The performance of conductive elastomers is dependent on many factors, e.g., elastomer crosslinking method and type of monomer. In an electrically conductive heterogeneous binary polymer blends, consisting of ethylene-propylene-diene-monomer (EPDM) and polyaniline (PAni), the performance was significantly affected by the crosslinking method 44. In the blend, PAni underwent doping upon exposure to a protonic acid and became electrically conductive. Several methods have been studied for monomer crosslinking, such as the use of a phenolic resin for unsaturated rubbers and high service temperature products, inter-chain C-C bonds formation after the reticulation reaction, electron beam induced crosslinking for 3-D network. Results indicated that electron beam irradiation crosslinking method was not affected by the presence of the acid necessary for doping the conductive elastomers.

Insulation elastomer has been used as a matrix to hold conductive particles in land grid array socket since it significantly improves contact resistance and increases interconnect density. The need for high density electronics has led to the development of semiconductor array packages, such as ball grid array (BGA) and land grid array (LGA) packages. However, soldering a package to a printed circuit board (PCB) is challenging for input/output (I/O) counts greater than a thousand, and solder joint defects can induce reliability problems in the application. These problems can be solved by using an LGA socket, which eliminates the soldering process, provides a separable interconnect between the component and the PCB, and also enables easy product rework, repair, and upgrade. The conductive particles in elastomer matrix has been introduced into electronics in 1980's 16, 45. One such conductive elastomer composite is called Z-direction anisotropically conductive materials which are made of the magnetic alignment of conductive particles in elastomers 46. The composites contain many vertically aligned but laterally isolated chains of ferromagnetic metal spheres, and the ends of which protrude from the surfaces for better electrical contact. These materials exhibit Z-direction only electrical conduction, in combination with the compliant nature of the materials, can be exploited advantageously for a variety of electronic applications, including fine-pitch, area-array, circuit interconnects, circuit-testing, heat sink interfacing, and sensor devices 47, 48. The schematic diagram of this design structure is shown in Figure 1. This design has been used in a tactile shear sensor for applications such as robotic skins or grippers, touch sensitive actuators, and finger operated controls 49.

**Figure 1.** Schematic diagram of a Z-direction only electrical conduction, in which the component is electrical connected to the printed circuit board through vertically aligned and laterally isolated conductive materials. The contact is made by external contact force. Reprint and permission from 16

An LGA socket assembly consists of component, socket, and board as shown in Figure 2. The contact force is applied through a fixture 16, 45. The socket design can vary in terms of target component, contact, and housing. One class of LGA sockets incorporates conductive particles in elastomer (polymer) as electrical interconnects for LGA. This class of socket has comparatively low cost, ease of assembly, and short signal path; the elastomer has a sealing effect. However, elastomer-based interconnection is susceptible to failure due to creep and stress relaxation, whenever a force or deformation is provided. A decrease in the contact force due to stress relaxation may lead to the degradation of contact resistance. Increase of elasticity can effectively mitigate creep and stress relaxation, in that crosslinked elastomer has instantaneous recovery after removal of stress.

208 Advanced Elastomers – Technology, Properties and Applications

controls 49.

array packages, such as ball grid array (BGA) and land grid array (LGA) packages. However, soldering a package to a printed circuit board (PCB) is challenging for input/output (I/O) counts greater than a thousand, and solder joint defects can induce reliability problems in the application. These problems can be solved by using an LGA socket, which eliminates the soldering process, provides a separable interconnect between the component and the PCB, and also enables easy product rework, repair, and upgrade. The conductive particles in elastomer matrix has been introduced into electronics in 1980's 16, 45. One such conductive elastomer composite is called Z-direction anisotropically conductive materials which are made of the magnetic alignment of conductive particles in elastomers 46. The composites contain many vertically aligned but laterally isolated chains of ferromagnetic metal spheres, and the ends of which protrude from the surfaces for better electrical contact. These materials exhibit Z-direction only electrical conduction, in combination with the compliant nature of the materials, can be exploited advantageously for a variety of electronic applications, including fine-pitch, area-array, circuit interconnects, circuit-testing, heat sink interfacing, and sensor devices 47, 48. The schematic diagram of this design structure is shown in Figure 1. This design has been used in a tactile shear sensor for applications such as robotic skins or grippers, touch sensitive actuators, and finger operated

**Figure 1.** Schematic diagram of a Z-direction only electrical conduction, in which the component is electrical connected to the printed circuit board through vertically aligned and laterally isolated conductive materials. The contact is made by external contact force. Reprint and permission from 16

An LGA socket assembly consists of component, socket, and board as shown in Figure 2. The contact force is applied through a fixture 16, 45. The socket design can vary in terms of target component, contact, and housing. One class of LGA sockets incorporates conductive particles in elastomer (polymer) as electrical interconnects for LGA. This class of socket has comparatively low cost, ease of assembly, and short signal path; the elastomer has a sealing effect. However, elastomer-based interconnection is susceptible to failure due to creep and stress relaxation, whenever a force or deformation is provided. A decrease in the contact force due to stress relaxation may lead to the degradation of contact resistance. Increase of

**Figure 2.** Schematic diagram of elastomer land grid socket assembly. The interconnection part, elastomer socket, can be replaced with different type of sockets, including metalized particle interconnect (MPI), PariPoser film, and conductive rubber sockets (CRS). Reprint and permission from 16

Several mechanisms can lead to the elastomer failure, such as electrochemical migration 50-52, stress relaxation and creep 45, 53 and corrosion 16. This section will extensively emphasize these mechanisms and use respective models to predict the lifecycle of the components. The metal electrochemical migration was first identified in the short-circuit failure of Sun Microsystems Sun Fire 6800. Investigation on the failed parts did not observe any dendrite to bridge adjacent components and/or contacts. In general, dendrites were considered to short the electrical circuit. This finding indicated that even without dendrite formation a short-circuit can happen. The design of the experiment, which mimicked the device environment with accelerated conditions, was conducted to identify the failure mechanism and electrochemical migration processes based on the failure products 50-52. The new failure mechanism in the metal-in-elastomer socket was proposed on the facts that surface insulation resistance has catastrophically decreased and unknown materials were formed on

surface after a certain period of time incubation in the temperature-humidity-bias conditions. The metal-in-elastomer socket consists of silver particles in elastomer matrix. Previous studies showed that silver can migrate in the humidity-temperature-bias conditions 54. However, those studies added the external conductive ions, e.g., contaminants or silver ions, into the system to initialize the migration process 55. In recent studies, several substantial results exploited different electrochemical migration, including silver electrochemical migration (ECM) processes, migration products, and failure criterion on surface insulation resistance (105 ohm).

The silver ECM could occur when bias was applied to silver-in-elastomer socket in the humidity-temperature. Initially, the water layer absorbed on the surface is electrolyzed to hydrogen and hydroxide ions, which further migrate to cathode and anode respectively. Silver oxides to silver ions by accepting electrons through electrolyzed water solution under bias, and silver ions are accumulated in absorbed water layer. The silver ions are also migrated to cathode through water layer. Overall, the surface insulation resistance (SIR) decreases when ions are generated in the absorbed water layer, or formation of conductive electrolyte. The short-circuit failure occurs when silver ion concentration increases to the specific value when the surface insulation resistance decreases to 105 ohm.

Both film and dendrites were observed on the film after permanent failure when the surface insulation resistance of the silver-in-elastomer was less than 1000 ohm. Studies on this regard showed that the film growth was developed in deionized water, while the dendrite growth was developed in high conductivity water. It was found that the film grew on the surface without distinctive direction while dendrites grew from cathode to anode. The analysis of film and dendrites from electron scanning spectrometry and X-ray photoelectron microscopy showed that the film consists of silver oxide and dendrites are silver only. The process of the silver electrochemical migration is progressed as follow: silver oxidation at anode to formation of silver ion; silver ion migration from anode to cathode through absorbed water layer (6-7 water monolayers); partial silver ion reduction to silver oxide with hydroxide from water electrolysis (film); other silver ion reduction to formation of silver at the cathode (dendrites) 50-52.

New failure criterion on surface insulation resistance was proposed. In the IPC standard (IPC-TR-476A), the failure was defined as the surface insulation resistance below 106 ohm. New experimental findings showed that the component didn't fail when resistance below the resistance, no electrochemical migration related products detected either 50-52. In addition, the surface insulation resistance immediately decreased to the order of 106 ohm when the metal-in-elastomer was conditioned in the highly accelerated stress testing (130C/85% RH/bias). Until SIR was below 105 ohm, the silver dendrites were observed. The resistance ultimately was within hundreds of ohms when the dendrites bridged the adjacent contacts.

The new ECM process was proposed as: presence of moisture, anodic metal dissolution or ion generation and ion migration to cathode (electrochemical reaction), ion accumulation, and metal dendritic growth. The ion accumulation was found to be the rate controlling step in which the surface insulation resistance degraded significantly when the ion concentration accumulated to a specific value. This added step explained the failure occurrence without dendritic formation/bridging and perfected the process. Time-to-failure (TTF) model considered these factors to estimate the failure time, including water adsorption, ion generation, ion migration, and ion accumulation. The dendritic growth was excluded from the model based on the experimental findings. Furthermore, incorporation of dendritic growth model with TTF model can be used to estimate time-to-permanent component failure, since the dendritic bridge creates a permanent conductive path between conductors.

210 Advanced Elastomers – Technology, Properties and Applications

surface insulation resistance (105 ohm).

the cathode (dendrites) 50-52.

contacts.

surface after a certain period of time incubation in the temperature-humidity-bias conditions. The metal-in-elastomer socket consists of silver particles in elastomer matrix. Previous studies showed that silver can migrate in the humidity-temperature-bias conditions 54. However, those studies added the external conductive ions, e.g., contaminants or silver ions, into the system to initialize the migration process 55. In recent studies, several substantial results exploited different electrochemical migration, including silver electrochemical migration (ECM) processes, migration products, and failure criterion on

The silver ECM could occur when bias was applied to silver-in-elastomer socket in the humidity-temperature. Initially, the water layer absorbed on the surface is electrolyzed to hydrogen and hydroxide ions, which further migrate to cathode and anode respectively. Silver oxides to silver ions by accepting electrons through electrolyzed water solution under bias, and silver ions are accumulated in absorbed water layer. The silver ions are also migrated to cathode through water layer. Overall, the surface insulation resistance (SIR) decreases when ions are generated in the absorbed water layer, or formation of conductive electrolyte. The short-circuit failure occurs when silver ion concentration increases to the

Both film and dendrites were observed on the film after permanent failure when the surface insulation resistance of the silver-in-elastomer was less than 1000 ohm. Studies on this regard showed that the film growth was developed in deionized water, while the dendrite growth was developed in high conductivity water. It was found that the film grew on the surface without distinctive direction while dendrites grew from cathode to anode. The analysis of film and dendrites from electron scanning spectrometry and X-ray photoelectron microscopy showed that the film consists of silver oxide and dendrites are silver only. The process of the silver electrochemical migration is progressed as follow: silver oxidation at anode to formation of silver ion; silver ion migration from anode to cathode through absorbed water layer (6-7 water monolayers); partial silver ion reduction to silver oxide with hydroxide from water electrolysis (film); other silver ion reduction to formation of silver at

New failure criterion on surface insulation resistance was proposed. In the IPC standard (IPC-TR-476A), the failure was defined as the surface insulation resistance below 106 ohm. New experimental findings showed that the component didn't fail when resistance below the resistance, no electrochemical migration related products detected either 50-52. In addition, the surface insulation resistance immediately decreased to the order of 106 ohm when the metal-in-elastomer was conditioned in the highly accelerated stress testing (130C/85% RH/bias). Until SIR was below 105 ohm, the silver dendrites were observed. The resistance ultimately was within hundreds of ohms when the dendrites bridged the adjacent

The new ECM process was proposed as: presence of moisture, anodic metal dissolution or ion generation and ion migration to cathode (electrochemical reaction), ion accumulation, and metal dendritic growth. The ion accumulation was found to be the rate controlling step

specific value when the surface insulation resistance decreases to 105 ohm.

The ion-accumulation based model assumed that the silver discharged at the anode was directly proportional to the amount of electrical charge that passed through the electrode, or Faraday's Law, the total mass of metal dissolution equaled the total mass of metal ions discharged into the electrolyte. Unlike the bulk resistance of conductive particles in elastomer 56, by use of Ohm's law and Brunauer-Emmett-Teller (BET) model 57, the time-tofailure model was given by,

$$TTF = nF \times \frac{m\_0}{M} \times \beta \times \frac{1}{V} \times \frac{(1 - RH)[1 + (c - 1)RH]}{cRH} \times e^{\left(\frac{E\_\alpha}{RT}\right)t}$$

Where *<sup>s</sup>* is the surface insulation conductivity, *R* is the resistance, *m0* is the critical mass loss required to dissolve to reach the critical ion concentration when resistance decreases to value of failure criterion, *V* is potential, *RH* is relative humidity, *n* is chemical valence, *M* is the molecular weight, *F* and *R* are constant, , *c* and *E* are coefficient relevant to relative humidity, temperature and materials properties. Using this model, the time-to-failure is estimated for elastomer sockets. For silver in elastomer socket, the *TTF* can be simplified as,

$$TTF = 1.11 \times 10^{-5} \times \frac{1}{V} \times \frac{(1 - RH)}{RH} \times e^{\left(\frac{7846}{T}\right)t}$$

Based on this model, the silver-in-elastomer socket fails in 150 hours after incubation in 90C/90%/20 V conditions; while it can resist up to 27000 h in service at conditions of 25C/50%/110 V.

Another failure mechanism is called creep and stress relaxation due to viscoelastic elastomer as matrix in socket. As shown in Figure 2, the mechanical support fixed elastomer deformation as constant. As a result, the stress applied to the socket is expected to decrease gradually due to the stress relaxation of elastomer. To understand failure mechanism and its effect on contact resistance, the dynamic mechanical analyzer (DMA) was used to characterize elastomer stress relaxation. The chamber of DMA can control temperature by air gas and liquid nitrogen. Creep is performed by constant stress and stress relaxation by constant strain. The creep process is often categorized into three stages: primary creep (a stage of decreasing creep rate), secondary creep (a stage of constant creep rate), and tertiary creep (a stage of increasing creep rate). When the load is initially applied, there is an instantaneous elongation, a primary stage of a transient nature during which slip and work hardening occur in the most favorably oriented grains. Then there is a secondary stage of steady-state creep during which the deformation continues at an approximately constant rate. The third stage (tertiary stage) takes place when the stress is high enough that the creep

$$R\_c = \frac{\rho}{2} \sqrt{\frac{\pi H}{F\_0[1 - B \ln\left(\frac{t}{C} + 1\right)]}} + \frac{\sigma\_f H}{F\_0[1 - B \ln\left(\frac{t}{C} + 1\right)]}$$

generates micropatterns by contact printing and microstructures by embossing and replica molding, has been used to manufacture blazed grating optics, stamps for chemical patterning, and microfluidics devices 19, 62, 63. Different from photolithography technique 64, the soft lithography is not subject to the limitations set by optical diffraction and optical transparency, with feature size down to 10-100 nm. Soft lithography techniques is termed as 'multilayer soft lithography' and developed by Quake et al. 59 in 2000. Briefly, multilayer structures were formed by bonding PDMS layers from separated cast in a micromachining mold. Since each layer has excess of one of the components, the reactive molecules remained at the interface form permanent bonding after further curing. They used such technique for development of microfluidic large-scale integration chips that contain plumbing networks with thousands of micromechanical PDMS valves and hundreds of individually addressable chambers. Their chips can be used to construct the microfluidic analog of a comparator array and a microfluidic memory storage device similar to the behavior of random-access memory.

212 Advanced Elastomers – Technology, Properties and Applications

secondary stage was also a function of temperature.

�� <sup>=</sup> � 2 �

Where the *Rc* is contact resistance,

socket for scheduled socket replacement.

rate accelerates until fracture occurs 58. The results for the creep of the silicone elastomer are shown in Figure 3 (a). Both primary and secondary creeps are observed. The results show that the creep deformation increases with temperature. The constant creep rate during the

Stress relaxation refers to the time-dependent contact force decrease under a constant strain. Figure 3 (b) plotted behavior of stress relaxation of elastomer under various temperatures. The significance of stress relaxation on elastomer socket is that contact resistance increases at a decreased contact force. It is thus important to derive the relationship time-dependent contact resistance due to stress relaxation of elastomer. Combination of stress relaxation and

� + ���

is initial contact force, *t* is time, and *B*, *C* are temperature dependent parameters, given by,

The *Rc – t* can be used for estimation of contact resistance degradation of metal-in-elastomer

**Figure 3.** Elastomer creep (a) and stress relaxation (b) behavior at various temperatures. Dots:

Elastomer is also one of popular materials for microfluidics application as substrate, microvalves, pumps 59-61. Elastomer can be easily patterned by curing on a micromachining mold as an alternative microfabrication technique. The soft lithography technique, which

experimental data; solid line: fitted curve. Reprint and permission from 45

**4. Elastomer as substrate in microfluidics** 

+

��� ���� � ��� ��

� + ���

� (for silicon elastomer)

is electrical resistivity (ohm.m), *H* is contact hardness, *F0*

contact resistance, the time-dependent contact resistance model is derived 45,

�� ���� � ��� ��

� = ��2�� � � �����, and � = ���2� � � ����������

PDMS microfluidics have great potential to be biomedical applications which only require small amounts of sample, routine operation by untrained personnel and low cost 65. One of such examples was developed on a simple, inexpensive PDMS microfluidic diagnostic device, which performs sandwich immunoassays for medicine and biological studies 12, 63. Screws are used in the device as virtual valves to on and off fluids. The low-cost on soft lithography microfabrication and raw materials makes this device great potential for portable healthcare delivery and monitoring. PDMS microfluidics can be used for reagent mixing in microscale channel 66. The challenge for flow mixing in a small dimension is mainly due to the difficult liquid turbulence inside chip. By use of PDMS as microfluidics substrate, a topological structure can be fabricated to exploit the laminarity of the flow to repeatedly fold the flow and double the lateral concentration gradient in a compact chip. Two different solutions are contained in a T-junction, and fluid flow is repeatedly split, rotated, and merged 67. The chip design was obtained by a rapid-prototyping master mold, and results showed that effective mixing can be achieved on short length scales.

PDMS microfluidics has applied to a variety of applications such as cell sorting, DNA sizing and sorting. DNA sequence is successfully performed in PDMS devices including sample preparation and electrophoresis analysis 14. PDMS has substantial advantages over glass or silicon dioxide substrates, since these latter materials were adhesive for the cells of interest. In addition, Chou et al. 68 developed a PDMS device for the sizing and sorting of restriction fragments of DNA based on single DNA molecule detection. The device sizes DNA on the basis of fluorescence intensity in which the DNA was labeled with fluorescence dye.

PDMS device can be easily fabricated by non-covalent bonding or plasma treated permanent bonding after replica by molding. The PDMS microfluidics for capillary electrophoresis was developed by Effenhauser et al. 69 in 1997. They used replica molding to obtain PDMS device and bonded reversibly the device with a flat piece of PDMS. The performance of PDMS CE is not ideal on unmodified surface, especially for proteins/peptides analysis, due to hydrophobicity of the unmodified PDMS surface. Later, Whitesides et al. developed PDMS microfluidics device by a rapid prototyping and irreversible bonding through plasma

oxidation19. Different PDMS devices were developed to demonstrate separation performance of fluorescently labeled amino acids and proteins using capillary zone electrophoresis (CZE) and gel electrophoresis (GE).

One-dimensional sodium dodecyl sulfate (SDS) capillary gel electrophoresis (CGE) can be also performed in PDMS microfluidics due to its versatile surface chemistry. PDMS surface is treated by UV-Ozone so that it could immobilize other hydrophilic molecules to prevent electroosmotic flow and protein/peptide non-specific binding. For complex cell lysate analysis, isoelectric focusing (IEF) can be interfaced with SDS-CGE as the first dimension separation. What makes PDMS special for IEF-SDS/CGE is because of conventional procedures for fabricating 3-D microfluidic channels 70, and the facility with which PDMSbased devices can be readily assembled and disassembled. For example, after IEF separation, the PDMS was disassembled and connected with a 3-D channel for SDS/CGE. Recent development on 2-D IEF-SDS/CGE showed that other thermoplastic would be better substrate of choice for integrated 2-D complex cell lysate analysis 71, 72.

The design of elastomer microfluidics can be made based on these studies 71-73. The electrokinetic injection of defined sample plugs within single dimension microfluidic systems has been extensively studied, with the most widely used injector configurations including the cross, double-T, and triple-T topologies 74, 75, and extended configurations employing continuous sample injection leveraging flow switching techniques. Floating injection and pinched-injection methods 76, 77 have been employed in simple cross-injectors for the controlled definition of small sample plugs, while double- and triple-T designs are generally employed for the introduction of larger sample volumes. Sample leakage 78 was observed in a 2-D device during sample injection which was the central issue for separation performance as shown in Figure 4. After the desired sample plug has been transferred to the separation channel, additional samples enter the injection region due to diffusion and fringing of the electrical field during the electrokinetic transfer process. The excess sample tailing degrades separation performance. To solve this issue, backbiasing is a commonly used to eliminate sample leakage. Practically, bias voltages can be applied to sample inlet and waste reservoirs, electrokinetically pulling back excess sample from the injection zone. The effective of this approach depended on the accurate selection of backbiasing bias. This requires characterization of device geometry, medium in device, and buffer resistivity. A few groups 49, 72, 79 have proposed 2-D chip designs consisting of an array of second dimension separation channels aligned with an identical number of injection channels on the opposite side of the first dimension microchannel, a more efficient design for spatially multiplexed 2-D separation chips is shown in Figure 4 (a). The design consists of a single first dimension separation channel intersected by multiple second dimension separation channels on one side, and sample injection channels on the opposite side, with the injection channels staggered with respect to the separation channels to ensure complete and simultaneous sampling of species separated in the first dimension channel. In this staggered design, the parallel second dimension microchannels may be regarded as an array of double-T injectors operating in parallel. However, unlike the single-channel case, the double-T injectors are not electrically isolated in the multidimensional system.

electrophoresis (CZE) and gel electrophoresis (GE).

substrate of choice for integrated 2-D complex cell lysate analysis 71, 72.

oxidation19. Different PDMS devices were developed to demonstrate separation performance of fluorescently labeled amino acids and proteins using capillary zone

One-dimensional sodium dodecyl sulfate (SDS) capillary gel electrophoresis (CGE) can be also performed in PDMS microfluidics due to its versatile surface chemistry. PDMS surface is treated by UV-Ozone so that it could immobilize other hydrophilic molecules to prevent electroosmotic flow and protein/peptide non-specific binding. For complex cell lysate analysis, isoelectric focusing (IEF) can be interfaced with SDS-CGE as the first dimension separation. What makes PDMS special for IEF-SDS/CGE is because of conventional procedures for fabricating 3-D microfluidic channels 70, and the facility with which PDMSbased devices can be readily assembled and disassembled. For example, after IEF separation, the PDMS was disassembled and connected with a 3-D channel for SDS/CGE. Recent development on 2-D IEF-SDS/CGE showed that other thermoplastic would be better

The design of elastomer microfluidics can be made based on these studies 71-73. The electrokinetic injection of defined sample plugs within single dimension microfluidic systems has been extensively studied, with the most widely used injector configurations including the cross, double-T, and triple-T topologies 74, 75, and extended configurations employing continuous sample injection leveraging flow switching techniques. Floating injection and pinched-injection methods 76, 77 have been employed in simple cross-injectors for the controlled definition of small sample plugs, while double- and triple-T designs are generally employed for the introduction of larger sample volumes. Sample leakage 78 was observed in a 2-D device during sample injection which was the central issue for separation performance as shown in Figure 4. After the desired sample plug has been transferred to the separation channel, additional samples enter the injection region due to diffusion and fringing of the electrical field during the electrokinetic transfer process. The excess sample tailing degrades separation performance. To solve this issue, backbiasing is a commonly used to eliminate sample leakage. Practically, bias voltages can be applied to sample inlet and waste reservoirs, electrokinetically pulling back excess sample from the injection zone. The effective of this approach depended on the accurate selection of backbiasing bias. This requires characterization of device geometry, medium in device, and buffer resistivity. A few groups 49, 72, 79 have proposed 2-D chip designs consisting of an array of second dimension separation channels aligned with an identical number of injection channels on the opposite side of the first dimension microchannel, a more efficient design for spatially multiplexed 2-D separation chips is shown in Figure 4 (a). The design consists of a single first dimension separation channel intersected by multiple second dimension separation channels on one side, and sample injection channels on the opposite side, with the injection channels staggered with respect to the separation channels to ensure complete and simultaneous sampling of species separated in the first dimension channel. In this staggered design, the parallel second dimension microchannels may be regarded as an array of double-T injectors operating in parallel. However, unlike the single-channel case, the double-T injectors are not electrically isolated in the multidimensional system.

**Figure 4.** Schematic diagram of two-dimensional capillary electrophoresis microfluidics design with (a) simplified device of a spatially-multiplexed separation platform with five second dimension microchannels, and (b) image during the sample transfer process using a fabricated chip without backbias channels. Reprint and permission from 71

This interconnected design can result in significant variations in performance between the different injectors, as depicted in Figure 4 (b). In this image, the sample within the first dimension channel is being electrokinetically transferred into the second dimension channel array by applying a uniform bias voltage in the injection channel reservoirs while grounding the second dimension reservoirs. Three main features are evident in this injection process. The substantial tailing of sample occurs at the head of each second dimension channel, resulting in sample dispersion during transfer; in all but the center second dimension channel, the injection is highly asymmetric, reflecting a non-uniform electric field distribution within the first dimension channel; sample from the outermost regions of the first dimension channel

continually leaks into the second dimension array due to a combination of diffusion and electric field fringing, leading to additional tailing which can continue long after sample from the center region of the first dimension channel has been fully transferred separation devices .

The solution for unsymmetrical injection and tailing was proposed by adding backbias channel s parallel to the separation channels as shown in Figure 4 (b) 70. Analytical calculation indicated that two-fold of resistance in the backbias channel generated uniform current and potential distribution in the 2-D channel network, not only in 1-D channel, but also injection and separation channels. Additionally, the angled 1-D channel can level the electric field distribution in the local region of T intersection. As a result, the angled 1-D channel substantially reduced sample tailing with negligible effect on the first dimension separation. These design perspectives have been used for complex biological sample analysis.

Use of PDMS for liquid chromatography has been limited largely due to its low bonding strength 80 and high gas/liquid permeability 81. It was reported to fabricate column in a device using PDMS 82-84. Since PDMS is inexpensive and thus it can be fabricated as a disposable column, it can be used as a stationary phase in liquid chromatography columns. The siloxane monomer was added into a molding to formation of desired features. The featured PDMS was reversibly sealed with cover PDMS by placing the molded chip in contact with cover PDMS, or irreversibly sealed by plasma oxidation. The PDMS chip was treated with silanes to minimize sample non-specific absorption. In general, PDMS is not the good materials of choice for high performance liquid chromatography. In such instance, other thermoplastics such as polyimide (PI) 85, fuse silica capillary 75, cyclic olefin copolymer (CoC) 86, are recommended.

Droplets formed by mixing one fluid in another or emulsions are very useful in a wide range of applications including personal care products, foods, microbiology, cell biology, and drug delivery vehicles 87, 88. It has been recognized that a successful droplet device for applications is able to control over size and distribution in microscale and nanoscale emulsions 88. Droplet formation can be obtained by either turbulence to break apart an immiscible mixture or continuous flow. PDMS devices for micro-droplet formation are fabricated by soft lithography in a molding. The native PDMS surface is usually hydrophobic and excellent surface for water-in-oil emulsions. The hydrophobic PDMS surface ensures that the oil droplets encapsulation of water phase inside without being contacted the PDMS channel walls 89, 90; while hydrophilic PDMS surface can be used for oilin-water droplets and biological sample analysis 91, 92. Several surface modifications have been reported by the sequential layer-by-layer deposition of polyelectrolytes yielding hydrophilic microchannels. Others used oxygen-plasma and oxygen-C2F6 to modify PDMS surface in order to enhance hydrophilicity and reduce channel electroosmotic flow (EOF). Alternatively, use of chemical modification, e.g., 2-hydroxyethyl methacrylate (HEMA), can permanently enhance PDMS hydrophilicity 91. In layer-by-layer (LbL) surface modification approach, segments of sodium chloride, poly(allylamine hydrochloride), and poly(sodium 4-styrenesulfonate) solutions are separated by air plugs and loaded into a piece of tubing and sequentially flushed through the channel at a constant flow rate. This LbL modification provides long storage time up to 5 months without noticeable behavior change of droplet formation. The LbL modified PDMS device can be used for water-oil-water double emulsion system for applications such as pharmaceutical compound delivery.

An interesting PDMS device for static microdroplet arrays is fabricated to droplet trapping, incubation, and release of enzymatic and cell-based assays 93. They simply used a singlelayered PDMS microfluidic structure. The aqueous droplets are trapped en masse and optically monitored for extended periods of time. The array droplet approach is used to characterize droplet shrinkage, aggregation of encapsulated biological cells and enzymatic reactions. On the other hand, to optimize the droplet formation, a design which use feedback control loop is developed 88. They successfully demonstrated that the closed loop control system with negative feedback can produce the required droplet size with minimal noise and no systematic steady sate error. This close-loop design is utilized for a number of applications, including micro-chemistry and emulsion production.

## **5. Outlook of elastomer application**

216 Advanced Elastomers – Technology, Properties and Applications

continually leaks into the second dimension array due to a combination of diffusion and electric field fringing, leading to additional tailing which can continue long after sample from the center region of the first dimension channel has been fully transferred separation devices .

The solution for unsymmetrical injection and tailing was proposed by adding backbias channel s parallel to the separation channels as shown in Figure 4 (b) 70. Analytical calculation indicated that two-fold of resistance in the backbias channel generated uniform current and potential distribution in the 2-D channel network, not only in 1-D channel, but also injection and separation channels. Additionally, the angled 1-D channel can level the electric field distribution in the local region of T intersection. As a result, the angled 1-D channel substantially reduced sample tailing with negligible effect on the first dimension separation.

Use of PDMS for liquid chromatography has been limited largely due to its low bonding strength 80 and high gas/liquid permeability 81. It was reported to fabricate column in a device using PDMS 82-84. Since PDMS is inexpensive and thus it can be fabricated as a disposable column, it can be used as a stationary phase in liquid chromatography columns. The siloxane monomer was added into a molding to formation of desired features. The featured PDMS was reversibly sealed with cover PDMS by placing the molded chip in contact with cover PDMS, or irreversibly sealed by plasma oxidation. The PDMS chip was treated with silanes to minimize sample non-specific absorption. In general, PDMS is not the good materials of choice for high performance liquid chromatography. In such instance, other thermoplastics such as polyimide

Droplets formed by mixing one fluid in another or emulsions are very useful in a wide range of applications including personal care products, foods, microbiology, cell biology, and drug delivery vehicles 87, 88. It has been recognized that a successful droplet device for applications is able to control over size and distribution in microscale and nanoscale emulsions 88. Droplet formation can be obtained by either turbulence to break apart an immiscible mixture or continuous flow. PDMS devices for micro-droplet formation are fabricated by soft lithography in a molding. The native PDMS surface is usually hydrophobic and excellent surface for water-in-oil emulsions. The hydrophobic PDMS surface ensures that the oil droplets encapsulation of water phase inside without being contacted the PDMS channel walls 89, 90; while hydrophilic PDMS surface can be used for oilin-water droplets and biological sample analysis 91, 92. Several surface modifications have been reported by the sequential layer-by-layer deposition of polyelectrolytes yielding hydrophilic microchannels. Others used oxygen-plasma and oxygen-C2F6 to modify PDMS surface in order to enhance hydrophilicity and reduce channel electroosmotic flow (EOF). Alternatively, use of chemical modification, e.g., 2-hydroxyethyl methacrylate (HEMA), can permanently enhance PDMS hydrophilicity 91. In layer-by-layer (LbL) surface modification approach, segments of sodium chloride, poly(allylamine hydrochloride), and poly(sodium 4-styrenesulfonate) solutions are separated by air plugs and loaded into a piece of tubing and sequentially flushed through the channel at a constant flow rate. This LbL modification provides long storage time up to 5 months without noticeable behavior change of droplet formation. The LbL modified PDMS device can be used for water-oil-water double emulsion

These design perspectives have been used for complex biological sample analysis.

(PI) 85, fuse silica capillary 75, cyclic olefin copolymer (CoC) 86, are recommended.

system for applications such as pharmaceutical compound delivery.

The application of elastomer is being continuously exploited. The future of elastomer is closely related to the specific requirement on applications. One of the hot topics is development of nano- and microfluidics for assisted reproductive technologies 94. For example, the studies have demonstrated that thickness of PDMS layers at 100 µm in chips produced significantly fewer embryos due to drastically shift in osmolality of embryo-culture media. It could be the solution by development of a hybrid membrane consisting of PDMS-parylene-PDMS for maintaining a stable cell growth environment. Therefore, investigation of new type of PDMS materials or hybrid PDMS can provide means for such applications.

Due to the low bonding strength, PDMS has seldom used for high performance liquid chromatography application (HPLC). The HPLC can be directly interfaced with detection system such as electrospray ionization or mass spectrometry. As a result, it becomes research of interest to synthesize PDMS which has strong bonding strength (via covalent bond) and low permeability, while PDMS optical transparency and elasticity are largely remained. In summary, PDMS is becoming more and more popular as the materials of choice for biological analysis due to its good physical, mechanical, and chemical properties, as well as its excellent biocompatibility.

A thermoplastic rubber material has recently been developed that can completely mend itself when the fracture interfaces are rejoined and left to heal for a moderate time 95. This "smart" rubber is easy to synthesize and displays excellent mechanical properties. This material could be used for microelectronics, microfluidics and biology system.

### **Author details**

Shuang (Jake) Yang\* *Department of Pathology, Johns Hopkins University, Baltimore, USA* 

Shuang (Jake) Yang and Kunqiang Jiang *Department of Mechanical Engineering, University of Maryland, USA Department of Chemistry and Biochemistry, University of Maryland, USA* 

<sup>\*</sup> Corresponding Author

#### **6. References**


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[57] Brunauer, S., Emmett, P.H. & Teller, E. Adsorption of gases in multimolecular layers. *J* 

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## **Elastomeric Electronics: A Microfluidic Approach**

Shi Cheng

222 Advanced Elastomers – Technology, Properties and Applications

[81] Kuncová-Kallio, J. & Kallio, P.J. 2486-2489 (IEEE, 2006).

molded in poly (dimethyl siloxane). 22, 3736-3743 (2001).

capillary electrophoresis devices. 21, 107-115 (2000).

plasma and chemical treatments. 83, 1277-1279 (2006).

in-oil-in-water double emulsions. 10, 1814-1819 (2010).

assisted reproduction. 57, 125-135 (2002).

1200, 55-61 (2008).

(2010).

(2002).

in microchannels. 82, 364 (2003).

[79] Li, Y., Buch, J.S., Rosenberger, F., DeVoe, D.L. & Lee, C.S. Integration of isoelectric focusing with parallel sodium dodecyl sulfate gel electrophoresis for multidimensional

[80] Gent, A. & Tobias, R. Effect of interfacial bonding on the strength of adhesion of elastomers. III. Interlinking by molecular entanglements. 22, 1483-1490 (1984).

[82] Weng, X., Chon, C.H., Jiang, H. & Li, D. Rapid detection of formaldehyde concentration in food on a polydimethylsiloxane (PDMS) microfluidic chip. 114, 1079-1082 (2009). [83] Slentz, B.E., Penner, N.A., Lugowska, E. & Regnier, F. Nanoliter capillary electrochromatography columns based on collocated monolithic support structures

[84] Ocvirk, G. et al. Electrokinetic control of fluid flow in native poly (dimethylsiloxane)

[85] Levkin, P.A. et al. Monolithic porous polymer stationary phases in polyimide chips for the fast high-performance liquid chromatography separation of proteins and peptides.

[86] Mair, D.A., Geiger, E., Pisano, A.P., Fréchet, J.M.J. & Svec, F. Injection molded

[87] Anna, S.L., Bontoux, N. & Stone, H.A. Formation of dispersions using "flow focusing"

[88] Miller, E., Rotea, M. & Rothstein, J.P. Microfluidic device incorporating closed loop feedback control for uniform and tunable production of micro-droplets. 10, 1293-1301

[89] Link, D.R. et al. Electric control of droplets in microfluidic devices. 45, 2556-2560 (2006). [90] Fujii, T. PDMS-based microfluidic devices for biomedical applications. 61, 907-914

[91] Bodas, D. & Khan-Malek, C. Formation of more stable hydrophilic surfaces of PDMS by

[92] Bauer, W.A.C., Fischlechner, M., Abell, C. & Huck, W.T.S. Hydrophilic PDMS microchannels for high-throughput formation of oil-in-water microdroplets and water-

[93] Huebner, A. et al. Static microdroplet arrays: a microfluidic device for droplet trapping, incubation and release for enzymatic and cell-based assays. 9, 692-698 (2009). [94] Beebe, D., Wheeler, M., Zeringue, H., Walters, E. & Raty, S. Microfluidic technology for

[95] Wietor, J.L. & Sijbesma, R.P. A Self‐Healing Elastomer. 47, 8161-8163 (2008).

microfluidic chips featuring integrated interconnects. 6, 1346-1354 (2006).

protein separations in a plastic microfludic network. 76, 742-748 (2004).

Additional information is available at the end of the chapter

http://dx.doi.org/10.5772/50208

### **1. Introduction**

Thousands of millions of electronic units have been produced since the first invention of electronics. These devices and systems, in which stiff, solid metals, insulators, and semiconductors have dominated for ages, are exclusively made in the rigid format, and usually retain a static shape once fabricated. Things have however changed recently. Researchers are currently attempting to shape a soft and rubbery future for electronics, where elastic materials like elastomers would be widely used to replace conventional rigid materials. The reason is rather simple. We as human beings are soft, just like most other living creatures in nature, but existing electronics are not. Shouldn't we create some kind of new electronics that resemble ourselves?

Extensive research activities on developing elastomeric electronics that may withstand severe bending, twisting, and straining, with maintained electronic functions, have been recently carried out. The mechanical characteristics of gold thin films on polydimethylsiloxane (PDMS) surfaces: wrinkles and cracks, have been investigated by Whitesides and Hutchinson for the first time in 1998 [1]. Subsequently, wrinkled electrodes on pre-stretched elastic substrates were employed to implement electroactive polymer actuators [2], and studies on wrinkled gold electrodes on PDMS substrates were conducted in the Wagner's and Suo's groups at Princeton [3],[4]. Micromechanics and manufacturing processes based on relaxed and pre-stressed PDMS slabs were developed in-depth. The most significant contributions were made by the Rogers's group at University of Illinois, Urbana-Champaign, where stretchable and foldable silicon integrated circuits (ICs) on "wavy" silicon ribbons (from Silicon On Insulator (SOI)) in PDMS were proposed for fully integrated stretchable electronics [5],[6]. Dozens of appealing devices based on stretchable silicon ICs, including electronic eye sensor [7], smart gloves/skins [8], implanted medical devices [9], and wearable ergonomic biomedical sensors [10], have been demonstrated. An alternative approach utilizing anisotropic etching of bulk wafers was also introduced [11],[12]. Later, thin meandered stretchable interconnects encased in silicone rubber substrates, operating at

© 2012 Cheng, licensee InTech. This is an open access chapter distributed under the terms of the Creative Commons Attribution License (http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited. © 2012 The Author(s). Licensee InTech. This chapter is distributed under the terms of the Creative Commons Attribution License http://creativecommons.org/licenses/by/3.0), which permits unrestricted use, distribution, and reproduction in any medium, provided the original work is properly cited.

different frequencies, were presented [13]-[15]. New electrical nanocomposite materials containing sliver nano-particles or carbon nanotubes/graphene that can flex and strain showed promising results as well [16],[17].

Microfluidics based stretchable electronics were initiated by Whitesides, who first studied low temperature metling solder filled microstructured elastomeric channels [18],[19]. Recently, this concept was extended to implement elastic direct current (DC) circuits, by incorporating eutectic gallium and indium alloy (EGaIn) into microfluidic channels in thin sheets of elastomers [20]-[22]. However, all these previously mentioned studies deal with either low-frequency ICs or relatively simple interconnects. Stretchable electronics at radio frequencies remained an unexploited field until early 2009, at which time Cheng demonstrated the first stretchable fluidic antennas, for enabling wireless communication and remote sensing [23],[24]. Similar work was reported afterwards using EGaIn alloy as conductors for the antennas instead of Galinstan [25]. Significantly enhanced elasticity of the resulting antennas could be achieved by introducing a different type of siloxane [26], and mechanically reconfigurable antennas could also be realized [27]. In 2010, Cheng proposed a hybrid integration strategy for the first demonstration of active microfluidic stretchable RF electronics (μFSRFEs), an integrated RF radiation sensor, *cf*. Fig. 1.a. [28]. Cheng further developed the μFSRFEs to the multi-layer configurations, and reported an integrated stretchable large-area wireless strain sensor, as seen in Fig. 1.b. [29].

**Figure 1.** Foldable, flexible, stretchable elastomeric electronic devices: a). Microfluidic stretchable radiation sensor. b). Elastomeric, reversibly stretchable, large-area wireless strain sensor.

The following sections in this chapter address the recent progresses in the emerging field of microfluidics based elastomeric electronics. Fabrication processes, integration techniques, application examples, as well as future perspectives are presented and discussed.

## **2. Fabrication of single-layer microfluidics based passive elastomeric electronic devices**

In brief, the manufacturing process steps can be summarized as follows: master construction, molding, liquid-alloy injection or filling, and encapsulation, *cf*. Fig. 2.

showed promising results as well [16],[17].

different frequencies, were presented [13]-[15]. New electrical nanocomposite materials containing sliver nano-particles or carbon nanotubes/graphene that can flex and strain

Microfluidics based stretchable electronics were initiated by Whitesides, who first studied low temperature metling solder filled microstructured elastomeric channels [18],[19]. Recently, this concept was extended to implement elastic direct current (DC) circuits, by incorporating eutectic gallium and indium alloy (EGaIn) into microfluidic channels in thin sheets of elastomers [20]-[22]. However, all these previously mentioned studies deal with either low-frequency ICs or relatively simple interconnects. Stretchable electronics at radio frequencies remained an unexploited field until early 2009, at which time Cheng demonstrated the first stretchable fluidic antennas, for enabling wireless communication and remote sensing [23],[24]. Similar work was reported afterwards using EGaIn alloy as conductors for the antennas instead of Galinstan [25]. Significantly enhanced elasticity of the resulting antennas could be achieved by introducing a different type of siloxane [26], and mechanically reconfigurable antennas could also be realized [27]. In 2010, Cheng proposed a hybrid integration strategy for the first demonstration of active microfluidic stretchable RF electronics (μFSRFEs), an integrated RF radiation sensor, *cf*. Fig. 1.a. [28]. Cheng further developed the μFSRFEs to the multi-layer configurations, and reported an integrated

stretchable large-area wireless strain sensor, as seen in Fig. 1.b. [29].

**Figure 1.** Foldable, flexible, stretchable elastomeric electronic devices: a). Microfluidic stretchable

The following sections in this chapter address the recent progresses in the emerging field of microfluidics based elastomeric electronics. Fabrication processes, integration techniques,

In brief, the manufacturing process steps can be summarized as follows: master

radiation sensor. b). Elastomeric, reversibly stretchable, large-area wireless strain sensor.

application examples, as well as future perspectives are presented and discussed.

construction, molding, liquid-alloy injection or filling, and encapsulation, *cf*. Fig. 2.

**electronic devices** 

**2. Fabrication of single-layer microfluidics based passive elastomeric** 

**Figure 2.** Manufacturing process steps of single-layer passive microfluidics based elastomeric electronics.

First of all, the design patterns are transferred to a 100 μm thick SU-8 100 (MicroChem, Newton, MA) layer on a silicon wafer with a standard soft lithography, and then developed and thermally stabilized at 150°C for 30 min to enhance the adhesion between SU-8 layer and silicon substrate. Subsequently, the PDMS prepolymer and cross linker (Elastosil RT601A and B, Wacker Chemie, Munich, Germany) are thoroughly mixed at a ratio of 9:1 (wt:wt) and poured onto the SU-8 master. After degassing, the PDMS mixture is cured at 70°C in an oven for 30 min. Then the cured thin PDMS replica is peeled off and a couple of holes are punched out for liquid alloy injection or filling. Meanwhile a thin blank PDMS lid is prepared using blank silicon wafer. Later, the PDMS replica and blank lid are bonded, using corona discharging (ETP, Chicago, IL, USA) activation. Through the punched holes in the PDMS, the liquid metal alloy (Galinstan, 68.5 % Ga, 21.5 % In, 10 % Sn, σ=3.46106 S/m) is manually injected into the channels. This alloy remains in liquid state from -19°C to 1300°C. The ventilation outlets are encapsulated with uncured PDMS mixture as mentioned above.

### **3. Liquid metal, elastic, microfluidic unbalanced loop antenna**

The fabrication process described in the previous section has been successfully employed to implement a single-layer, liquid metal, stretchable, unbalanced loop antenna, which consists of a radiating element and a semi-circular ground plane, as depicted in Fig. 3 [23]. Several cylindrical reservoirs are aligned along the upper tube of the antenna prototype to ensure

good electrical connections while being flexed. The presence of these reservoirs introduces slight resonance frequency decrease of the antenna. The lower semi-circular area acts as a ground plane for the antenna. A number of posts are aligned to space the top and bottom PDMS membranes in the ground area, which has negligible influence on the electrical characteristics of the antenna.

**Figure 3.** Geometrical schematic of stretchable unbalanced loop antenna: *R*=18.1 mm, *R*sub=47.5 mm, *R*reservior=1 mm, *W*S=400 μm, *L*grid=1.5 mm, *W*grid=1.5 mm, *D*grid=1.9 mm, *h*=1.0 mm, *h*metal=100 μm, and *θ*=30°.

Mechanical properties of the antenna prototype were characterized experimentally. Fig. 4 shows that the resulting antenna is reversibly foldable, flexible, and stretchable. Severe bending and twisting in the tests did not cause any mechanical damages. In addition, the antenna can withstand extreme levels of strains of up to 40% along multiple axis. As a result of the uneven thickness of the PDMS substrate and the heterogeneous pattern of the antenna, slight mechanical inhomogeneity was observed while stressing along different orientations, *cf*. Figs. 4. b and c. In principle, elasticity of up to 100% should be in reach. Nevertheless, the openings in the membranes for electrical tests limit the mechanical deformability, reliability, and robustness, as the antenna can be easily torn from these openings.

Electrical properties of the antenna in its original state were obtained from numerical simulations. Experiments on port impedance and radiation characteristics of the antenna in both relaxed and flexed states were performed. Fig. 5 presents the simulated and measured reflection coefficients (*S*11). The non-strained antenna features good impedance matching at 2.4 GHz, with an input impedance of approximately 75+15j Ω, Fig. 5.a. Applying strains on the antenna introduces an increase of the length of the upper radiating loop so that leads to a decrease on its resonance frequency, Fig. 5.b. In the case of stressing along *y*-axis, the antenna input resistance decreases due to the stronger coupling between the upper antenna arm and the lower ground plane, whereas this resistances goes up while stretching along *x*-axis.

characteristics of the antenna.

good electrical connections while being flexed. The presence of these reservoirs introduces slight resonance frequency decrease of the antenna. The lower semi-circular area acts as a ground plane for the antenna. A number of posts are aligned to space the top and bottom PDMS membranes in the ground area, which has negligible influence on the electrical

**Figure 3.** Geometrical schematic of stretchable unbalanced loop antenna: *R*=18.1 mm, *R*sub=47.5 mm, *R*reservior=1 mm, *W*S=400 μm, *L*grid=1.5 mm, *W*grid=1.5 mm, *D*grid=1.9 mm, *h*=1.0 mm, *h*metal=100 μm, and *θ*=30°.

and robustness, as the antenna can be easily torn from these openings.

Mechanical properties of the antenna prototype were characterized experimentally. Fig. 4 shows that the resulting antenna is reversibly foldable, flexible, and stretchable. Severe bending and twisting in the tests did not cause any mechanical damages. In addition, the antenna can withstand extreme levels of strains of up to 40% along multiple axis. As a result of the uneven thickness of the PDMS substrate and the heterogeneous pattern of the antenna, slight mechanical inhomogeneity was observed while stressing along different orientations, *cf*. Figs. 4. b and c. In principle, elasticity of up to 100% should be in reach. Nevertheless, the openings in the membranes for electrical tests limit the mechanical deformability, reliability,

Electrical properties of the antenna in its original state were obtained from numerical simulations. Experiments on port impedance and radiation characteristics of the antenna in both relaxed and flexed states were performed. Fig. 5 presents the simulated and measured reflection coefficients (*S*11). The non-strained antenna features good impedance matching at 2.4 GHz, with an input impedance of approximately 75+15j Ω, Fig. 5.a. Applying strains on the antenna introduces an increase of the length of the upper radiating loop so that leads to a decrease on its resonance frequency, Fig. 5.b. In the case of stressing along *y*-axis, the antenna input resistance decreases due to the stronger coupling between the upper antenna arm and the lower ground plane, whereas this resistances goes up while stretching along *x*-axis.

**Figure 4.** Photographs of the elastomeric fluidic unbalanced loop antenna: a). relaxed state, b). with 40% *x*-axis elongation, c). with 40% *y*-axis elongation, d). with biaxial elongation, e). bent state, and f). twisted state.

**Figure 5.** Simulated and measured port impedance of the elastic unbalanced loop antenna operating between 2.0 and 2.8 GHz: a). simulated (solid) and measured (dashed) *S*11 of the relaxed antenna; b). measured *S*11 of the flexed antenna with 20% *x*-axis (solid), 40% *x*-axis (dashed) and 20% *y*-axis (dotted), 40% *y*-axis (dashdot) elongation. The *S*11 at 2.4 GHz is marked by a cross sign on each curve.

Fig. 6 shows simulated and measured radiation patterns of the antenna prototype at 2.4 GHz.

As expected, the relaxed antenna resembles conventional unbalanced loops that exhibit broad beam coverage, especially in the *yz*-plane where nearly perfect omnidirectionality is seen. The maximum antenna gain is around 2.7 dBi, and the measured cross-polarization (*Gφ* in the *xz*-plane and *Gθ* in the *yz*-plane) is approximately 15 dB lower than the corresponding co-polarization. Variations of the radiation patterns of the strained antenna can be found, *cf*. Fig. 6.b. The omnidirectionality of the *yz*-plane radiation patterns degrades, particularly while the antenna is stressed along *y*-axis. The increasing cable influence caused by the *y*-axis strain explains the ripples within the angle of 45o-135o on the measured *xz*plane radiation pattern. Also, the level of the measured cross-polarization of the stretched antenna slightly increases compared to that of the relaxed one.

**Figure 6.** Simulated and measured radiation patterns of the a). relaxed and b). flexed antenna at 2.4 GHz. The corresponding coordinate system is depicted in Fig. 3. The antenna gain along *φ* and *θ* is defined as *Gφ* and *Gθ*, respectively.

High conductivity of Galinstan and large cross dimensions of the microfluidic channels attribute to low conductive losses and high radiation efficiency of the antenna. The measured data indicate a resonance frequency decrease of up to 18% while stretching, but the radiation efficiency at 2.4 GHz remains greater than 80%. In the experiments, no significant radiation efficiency decrease can be seen. This fact implies that galvanic connections in the microstructured elastomeric channels are not interrupted by straining. Good electrical continuity of the liquid alloy owes to its excellent wettability on PDMS substrate surfaces. Though the presented unbalanced loop antenna achieves good radiation properties, its resonance frequency detuning introduced by stretching leads to relatively poor total efficiency at 2.4 GHz. Other antenna concepts with more robust port impedance and radiation characteristics in response to deformation would be feasible alternative solutions.

### **4. Bendable, stretchable, fluidic UWB antenna**

The concept of microfluidics based elastomeric electronics is then extended to realize a planar inverted cone antenna (PICA) for the ultrawideband (UWB) frequency range of 3.1- 10.6 GHz [24]. The reason for choosing the PICA is that its impedance matching and radiation characteristics are expected to be insensitive to its mechanical deformation. Also, its uniplanar configuration makes it a suitable antenna type for a design that should be bendable and stretchable. Fig. 7 shows the design schematically, with a leaf-shaped radiation and a large ground plane.

Resembling to the previously demonstrated elastic unbalanced loop, the implemented PICA prototype features excellent reversible deformability in the mechanical tests, as shown in Fig. 8. Extreme levels of straining of up to 40% along either *x*- or *y*-axis do not cause any mechanical failures, *cf*. Figs. 8.b and d. Moreover, the resulting prototype also withstands severe mechanical folding and twisting. After removal of applied mechanical forces, the antenna returns to its original state.

228 Advanced Elastomers – Technology, Properties and Applications

a). b).

defined as *Gφ* and *Gθ*, respectively.

radiation and a large ground plane.

**Figure 6.** Simulated and measured radiation patterns of the a). relaxed and b). flexed antenna at 2.4 GHz. The corresponding coordinate system is depicted in Fig. 3. The antenna gain along *φ* and *θ* is

in response to deformation would be feasible alternative solutions.

**4. Bendable, stretchable, fluidic UWB antenna** 

High conductivity of Galinstan and large cross dimensions of the microfluidic channels attribute to low conductive losses and high radiation efficiency of the antenna. The measured data indicate a resonance frequency decrease of up to 18% while stretching, but the radiation efficiency at 2.4 GHz remains greater than 80%. In the experiments, no significant radiation efficiency decrease can be seen. This fact implies that galvanic connections in the microstructured elastomeric channels are not interrupted by straining. Good electrical continuity of the liquid alloy owes to its excellent wettability on PDMS substrate surfaces. Though the presented unbalanced loop antenna achieves good radiation properties, its resonance frequency detuning introduced by stretching leads to relatively poor total efficiency at 2.4 GHz. Other antenna concepts with more robust port impedance and radiation characteristics

The concept of microfluidics based elastomeric electronics is then extended to realize a planar inverted cone antenna (PICA) for the ultrawideband (UWB) frequency range of 3.1- 10.6 GHz [24]. The reason for choosing the PICA is that its impedance matching and radiation characteristics are expected to be insensitive to its mechanical deformation. Also, its uniplanar configuration makes it a suitable antenna type for a design that should be bendable and stretchable. Fig. 7 shows the design schematically, with a leaf-shaped

Resembling to the previously demonstrated elastic unbalanced loop, the implemented PICA prototype features excellent reversible deformability in the mechanical tests, as shown in Fig. 8. Extreme levels of straining of up to 40% along either *x*- or *y*-axis do not cause any mechanical failures, *cf*. Figs. 8.b and d. Moreover, the resulting prototype also withstands

**Figure 7.** Geometry of the 2-D stretchable PICA. Dimensions are: *R*=10 mm, *Rsub*=47.5 mm, *Lg*=25 mm, *Wg*=40 mm, *W1*=1.25 mm, *W2*=1.75 mm, *G*=300 μm, *h*=1 mm, and *hmetal*=100 μm.

**Figure 8.** Photographs of the bendable, flexible, stretchable PICA: a). and c). relaxed state, b). strained antenna with 40% *x*-axis elongation, d). strained antenna with 40% *y*-axis elongation, e). folded antenna, and f). twisted antenna. The corresponding coordinate system is presented in Fig. 7.

Figs. 9 and 10 present simulated and measured reflection coefficients of the elastic, microfluidic PICA. The relaxed antenna shows good impedance match (*S*11 < -10 dB), within 3-11 GHz, both in numerical simulations and experiments.

**Figure 9.** Simulated and measured *S*11 of the non-stretched antenna.

Applying strains along the *x*-axis results in the increased height of the antenna radiator so that its first resonance frequency decreases, *cf*. Fig. 10. Port impedance of the antenna is somewhat sensitive to its geometry, and consequently the antenna exhibits slightly varying impedance matching while stressed. However, good impedance match maintains at the entire UWB frequency band even if the antenna is strained to 40%.

**Figure 10.** Measured reflection coefficients of the flexed antenna.

Measured radiation patterns at 2.5 GHz of the antenna in its relaxed and flexed states are displayed in Figs. 11 and 12. Similar to conventional fat monopole antennas, the nonstrained antenna features broad coverage, especially in the *yz*-plane. The maximum antenna gain at 2.5 GHz was measured to be 2.2 dBi. The cross-polarization discrimination is very good. Numerical simulations are in line with the corresponding experimental data. Stretching the antenna along either *x*- or *y*-axis up to 40% introduces slight variations in the measured radiation patterns at 2.5 GHz, but without any significant gain reduction.

Similar simulations and experiments are also performed at 5 GHz, where ripples and slight asymmetry occur in the radiation patterns caused by disturbance from feed cable. The presence of higher order modes at 5 GHz together with the increasing cable influence degrades the cross-polarization discrimination. Compared to the experimental data at 2.5 GHz, larger variations on the measured radiation patterns at 5 GHz are observed while the antenna is in its strained states, particularly in the *yz*-plane.

**Figure 9.** Simulated and measured *S*11 of the non-stretched antenna.

entire UWB frequency band even if the antenna is strained to 40%.

**Figure 10.** Measured reflection coefficients of the flexed antenna.

antenna is in its strained states, particularly in the *yz*-plane.

Applying strains along the *x*-axis results in the increased height of the antenna radiator so that its first resonance frequency decreases, *cf*. Fig. 10. Port impedance of the antenna is somewhat sensitive to its geometry, and consequently the antenna exhibits slightly varying impedance matching while stressed. However, good impedance match maintains at the

Measured radiation patterns at 2.5 GHz of the antenna in its relaxed and flexed states are displayed in Figs. 11 and 12. Similar to conventional fat monopole antennas, the nonstrained antenna features broad coverage, especially in the *yz*-plane. The maximum antenna gain at 2.5 GHz was measured to be 2.2 dBi. The cross-polarization discrimination is very good. Numerical simulations are in line with the corresponding experimental data. Stretching the antenna along either *x*- or *y*-axis up to 40% introduces slight variations in the

Similar simulations and experiments are also performed at 5 GHz, where ripples and slight asymmetry occur in the radiation patterns caused by disturbance from feed cable. The presence of higher order modes at 5 GHz together with the increasing cable influence degrades the cross-polarization discrimination. Compared to the experimental data at 2.5 GHz, larger variations on the measured radiation patterns at 5 GHz are observed while the

measured radiation patterns at 2.5 GHz, but without any significant gain reduction.

**Figure 11.** Measured a). *xz*- and b). *yz*-plane (according to the coordinate system in Fig. 7) radiation patterns at 2.5 GHz of the non-stretched and stretched antenna with 40% *x*-axis elongation.

**Figure 12.** Measured a). *xz*- and b). *yz*-plane (according to the coordinate system in Fig. 7) radiation patterns at 2.5 GHz of the non-stretched and stretched antenna with 40% *y*-axis elongation.

The measured radiation efficiency within the frequency range of 3-10 GHz is shown in Fig. 13. Although the radiation efficiency at the lower end of the frequency range decreases when the antenna is stretched, it is still greater than 70%. This, together with the experimental data on port impedance, indicates that the PICA prototype achieves good total antenna efficiency regardless of stretching.

Much work however remains. First of all, comprehensive environmental reliability and durability tests, e.g. vibration, temperature cycling, and aging, are needed, since it is indeed a new way of fabricating antennas, with new materials. Electrical characteristics of the antenna at extreme temperature condition, e.g. below the melting point of the liquid alloy, should be evaluated when a special measurement setup is established. Studies on antenna radio interfaces are also necessary. Development of fully integrated stretchable wireless electronic systems consisting of flexible thin embedded active chips, stretchable interconnects, and highly efficient stretchable antennas are the end objective, with a need for much research efforts.

**Figure 13.** Experimental results on the radiation efficiency of the relaxed and flexed antennas.

### **5. Single-layer stretchable elastomeric integrated active RF electronics**

The concept of microfluidics based passive stretchable elastomeric electronics has been further developed to the integrated device level utilizing localized stiff cells (LSCs), as illustrated in Fig. 14 [28]. A microfluidic, elastic, large-area antenna is realized in the same manner as the previous stretchable antennas, by incorporating liquid alloy into microstructured elastomeric channels. Established IC chips associated with passive components are assembled on small flexible laminates. Subsequently, a few tin-plated contact pins resembling cantilevers are soldered to the flexible circuits, and a semi-spherical solder ball is then mounted on the bottom surface of each contact pin on the other end to improve galvanic connection to the liquid fluid, *cf*. Fig. 14.e. Whereafter, the flexible circuits are embedded into the antenna substrate, with each contact pin immersing in the liquid alloy. In the end, uncured PDMS mixture droplets are deposited on top of the flexible circuits to locally enhance the stiffness of the elastomeric substrate and encapsulate the flexible circuits as well as the flex-to-stretch interfaces. The LSCs with enhanced stiffness than surrounding areas ensures that nearly zero stress and displacement between the rigid and the elastic structures would arise inside the LSCs when applying strains to the heterogeneously integrated device. The overall elasticity of the hybrid device is degraded compared to the standalone stretchable antennas due to the presence of the LSCs. But the reliability of the integrated device is greatly improved. Also, good wetting of the liquid metal alloy on tin-plated pins and solder balls ensures reliable electrical connections between contact pins and the antenna.

The proposed hybrid integration strategy was employed to implement a 900 MHz microfluidic stretchable RF radiation sensor, as seen in Fig. 15. The resulting sensor prototype is capable of performing real-time monitoring on the human exposure level to electromagnetic fields (EMF). Once the exposure level to EMFs exceeds the threshold power, the integrated light emission diode (LED) will be switched on. This sensor is of importance to human health as more and more EMFs are generated by today's telecom networks and mobile terminals, and might cause health issues. The integrated radiation sensor contains three sub-modules fully encased in a large-area elastomeric substrate, including a stretchable unbalanced loop antenna for capturing RF radiation from ambient environments, an RF power detection unit for converting received RF energy to the corresponding DC voltages, and a LED indicator for visualizing.

between contact pins and the antenna.

and a LED indicator for visualizing.

**Figure 13.** Experimental results on the radiation efficiency of the relaxed and flexed antennas.

**5. Single-layer stretchable elastomeric integrated active RF electronics** 

The concept of microfluidics based passive stretchable elastomeric electronics has been further developed to the integrated device level utilizing localized stiff cells (LSCs), as illustrated in Fig. 14 [28]. A microfluidic, elastic, large-area antenna is realized in the same manner as the previous stretchable antennas, by incorporating liquid alloy into microstructured elastomeric channels. Established IC chips associated with passive components are assembled on small flexible laminates. Subsequently, a few tin-plated contact pins resembling cantilevers are soldered to the flexible circuits, and a semi-spherical solder ball is then mounted on the bottom surface of each contact pin on the other end to improve galvanic connection to the liquid fluid, *cf*. Fig. 14.e. Whereafter, the flexible circuits are embedded into the antenna substrate, with each contact pin immersing in the liquid alloy. In the end, uncured PDMS mixture droplets are deposited on top of the flexible circuits to locally enhance the stiffness of the elastomeric substrate and encapsulate the flexible circuits as well as the flex-to-stretch interfaces. The LSCs with enhanced stiffness than surrounding areas ensures that nearly zero stress and displacement between the rigid and the elastic structures would arise inside the LSCs when applying strains to the heterogeneously integrated device. The overall elasticity of the hybrid device is degraded compared to the standalone stretchable antennas due to the presence of the LSCs. But the reliability of the integrated device is greatly improved. Also, good wetting of the liquid metal alloy on tin-plated pins and solder balls ensures reliable electrical connections

The proposed hybrid integration strategy was employed to implement a 900 MHz microfluidic stretchable RF radiation sensor, as seen in Fig. 15. The resulting sensor prototype is capable of performing real-time monitoring on the human exposure level to electromagnetic fields (EMF). Once the exposure level to EMFs exceeds the threshold power, the integrated light emission diode (LED) will be switched on. This sensor is of importance to human health as more and more EMFs are generated by today's telecom networks and mobile terminals, and might cause health issues. The integrated radiation sensor contains three sub-modules fully encased in a large-area elastomeric substrate, including a stretchable unbalanced loop antenna for capturing RF radiation from ambient environments, an RF power detection unit for converting received RF energy to the corresponding DC voltages,

**Figure 14.** Schematic of the hybrid integration of the single-layer stretchable elastomeric RF electronics.

**Figure 15.** Photograph of the stretchable elastomeric RF radiation sensor prototype.

Prior to the hybrid integration, experimental evaluations on electrical characteristics of the standalone antenna and the RF power detection sub-module assembled on a flex foil are conducted. The implemented stretchable fluidic antenna exhibits similar mechanical and electrical performance as the previously presented unbalanced loop, but operates at a lower frequency of approximately 900 MHz. The power conversion sub-module in the integrated RF radiation sensor involves an off-the-shelf power detector IC chipset (Linear Technology, LT 5534), two decoupling and one coupling capacitors, an inductor for impedance matching at the RF input, and a green LED indicator, all assembled on a small flex foil with a size of 10 mm × 18 mm, as seen in Fig. 15. When the input RF energy at the detector exceeds the threshold power of the integrated device, the LED indicator is switched on, and vice versa. The entire integrated power detection unit can be powered by four serially connected AA rechargeable batteries with a DC supply voltage of 5.23 V. The power conversion behavior of the integrated device is first characterized using a signal generator and a digital

multimeter. Fig. 16 presents the measured DC voltages at the output of the power detection unit in response to varying input RF power at 900 MHz. A dynamic range of -60 dBm to 0 dBm can be achieved, and the threshold power for turning on/off the LED indicator is found to be slightly higher than -30 dBm.

**Figure 16.** Measured output DC voltages of the RF power detection sub-module and the LED on/off states versus varying input RF power.

The integrated elastic radiation sensor device is tested in a demonstration setup in Fig. 17. The RF radiation source consists of an RF signal generator and a horn antenna placed 5 m away from the sensor device in the line-of-sight.

**Figure 17.** Schematic illustration of the RF radiation sensing demonstration setup.

Experiments presented in Fig. 18 verify that the resulting microfluidics based elastomeric integrated device maintains its radiation sensing capabilities even if it is strained along multiple axis. Extreme levels of twisting do not cause any failures in its operation either.

to be slightly higher than -30 dBm.

states versus varying input RF power.

either.

away from the sensor device in the line-of-sight.

multimeter. Fig. 16 presents the measured DC voltages at the output of the power detection unit in response to varying input RF power at 900 MHz. A dynamic range of -60 dBm to 0 dBm can be achieved, and the threshold power for turning on/off the LED indicator is found

**Figure 16.** Measured output DC voltages of the RF power detection sub-module and the LED on/off

**Figure 17.** Schematic illustration of the RF radiation sensing demonstration setup.

The integrated elastic radiation sensor device is tested in a demonstration setup in Fig. 17. The RF radiation source consists of an RF signal generator and a horn antenna placed 5 m

Experiments presented in Fig. 18 verify that the resulting microfluidics based elastomeric integrated device maintains its radiation sensing capabilities even if it is strained along multiple axis. Extreme levels of twisting do not cause any failures in its operation

**Figure 18.** Photographs of the elastomeric RF radiation sensor device operating in ordinary office environment: a). relaxed state, b). with 15% strain applied along *y*-axis, e). with manually applied strain along multiple directions, and f). with severe twisting. The sensor was directly illuminated by a radiation source 5 m away. The coordinate system is shown in Fig. 15.

### **6. Multi-layer stretchable elastomeric integrated active RF electronics**

Very recently, the emerging field of microfluidics based elastomeric electronics has been further advanced to multi-layer μFSRFEs, with the demonstration of a microfluidic, reversibly stretchable, large-area wireless strain sensor [29]. The manufacturing process for multi-layer elastomeric passive components resembles the single-layer fabrication, but with a few minor modifications and process steps added, *cf*. Fig. 19.

In summary, the upper and lower microfluidic channels are respectively constructed in the top and the bottom PDMS sheets, using standard soft lithography techniques. Then a few inlets and outlets are punched. In addition, the blank middle PDMS slab is prepared. The microstructured top elastomer layer is bonded to the blank silicone rubber slab using plasma boning, and addition inlets are punched on the bonded PDMS sheet. The bottom elastomer substrate is bonded to the previously bonded PDMS layers afterwards, using corona discharging activation. Prior to filling the upper microfluidic channels with galinstan fluid metal, the inlet 3 is sealed with a piece of Scotch® tape. Later, the ventilation outlets in the top silicone rubber sheet are encapsulated using PDMS prepolymer. Whereafter, both in the inlets 1 and 3 are taped, and galinstan alloy is injected into the lower microfluidic channels from the bottom side. All remaining ventilation outlets and the inlet 2 are then encapsulated, and the inlets 1 and 3 are reserved for connecting active circuitry in a heterogeneously integrated device.

**Figure 19.** Fabrication process of a multi-layer microfluidics based stretchable elastomeric passive electronic device.

The subsequent fabrication process steps are active circuit assembly and hybrid device integration, which resemble the manufacturing and the integrated processes presented in the previous section. Schematic illustrations describing the entire assembly and integration procedure are displayed in Fig. 20, and not discussed in detail in this section.

**Figure 20.** Schematic drawings of the integration procedure for a multi-layer stretchable, microfluidic integrated electronic device.

The first resulting integrated device based on multi-layer μFSRFEs is a self-contained wireless strain sensor, fully encased in a thin large-area silicone elastomer, as shown in Fig. 21. A stretchable liquid metal microstrip patch antenna comprising an upper rectangular meshed patch and a lower ground plane constructed in the same manner, takes up the major area of the hybrid device. It is worthy mentioning that this antenna not only serves as a radiator like any other conventional antennas, but also acts as a sensing device owing to its varying electrical characteristics in response to mechanical deformation. The entire integrated sensor device is slightly larger than its meshed fluid metal ground with a size of 100 mm × 80 mm, and approximately four times as big as the antenna patch. Though the liquid metal patch is the actual strain sensing element, the large ground plane can also serve an effective sensing area, and extend the strain sensing functionality to almost the entire device. Apart from the self-contained wireless strain sensor, a custom-designed personal computer (PC)-assisted radio receiver for remotely collecting, processing, and storing the measured data wirelessly transmitted from the integrated sensor device is also implemented. It removes the need for costly RF measurement facilities, and significantly reduces the cost for building up a complete system.

236 Advanced Elastomers – Technology, Properties and Applications

electronic device.

integrated electronic device.

**Figure 19.** Fabrication process of a multi-layer microfluidics based stretchable elastomeric passive

procedure are displayed in Fig. 20, and not discussed in detail in this section.

The subsequent fabrication process steps are active circuit assembly and hybrid device integration, which resemble the manufacturing and the integrated processes presented in the previous section. Schematic illustrations describing the entire assembly and integration

**Figure 20.** Schematic drawings of the integration procedure for a multi-layer stretchable, microfluidic

**Figure 21.** Demonstrated microfluidic reversibly stretchable wireless strain sensor: a). schematic illustration, and b). optical photograph of the resulting device.

Port impedance and radiation characteristics of the standalone, mechanically reconfigurable, elastomeric patch antenna in its relaxed and flexed states are measured. Excellent port impedance matching around 1.46 GHz along with very good correlation between the simulated and measured reflection coefficients of the non-stretched antenna are seen in Fig. 22.a. Applying increasing strain of up to 15% along its *x*-axis introduces persistent downshift of its resonant frequency of the patch antenna. The lowest resonant frequency of approximately 1.33 GHz is achieved at the maximum *x*-axis elongation of 15%.The measured resonant frequency of the relaxed antenna is slightly less than the minimum operational frequency of the integrated transmitter circuit. This negative offset is crucial to strain sensing, since persistently rising antenna mismatch losses at the operational frequencies in response to increasing stressing along the *x*-axis of the hybrid device.

Furthermore, this frequency offset should be as little as possible to avoid too high mismatch losses so that reasonably long remote sensing ranges as well as sufficient sensing sensitivity can be attained. Placing the antenna original resonant frequency above the highest operational frequency of the transmitter is not an option, as inconsistent variations of mismatch losses would occur if the antenna based strain sensor is stressed from its relaxed state to a high tensile strain.

Measured radiation patterns, including mismatch losses, at 1.46 GHz, of the relaxed and flexed antennas are shown in Figs. 22.b and c. The realized peak gain is measured along the +*z*-axis (according to the coordinate system depicted in Fig. 21.a.) of the mechanically reconfigurable antenna, and greatly degraded from 2.0 dBi to -10.7 dBi while applying an increasing strain from 0% to 15% along its *x*-axis. The meshed ground plane directs the radiation forward, with a front-to-back ratio as high as 10.0 dB, regardless of stretching. Moreover, the cross-polarization discrimination is very good in both relaxed and strained cases.

**Figure 22.** a). Simulated and measured reflection coefficients of the elastomeric fluid metal patch antenna in its relaxed state, and with different strains along the *x*-axis (according to the coordinate system in Fig. 21). b). *xz*- and c). *yz*-plane radiation patterns (including mismatch losses) at the original resonant frequency of the elastic antenna in its relaxed state, and with 15% elongation along its *x*-axis. The antenna gain along *φ* and *θ* orientations is defined as *Gφ* and *Gθ*, respectively. The antenna co- and cross-polarization along +*z*-axis are in parallel to the *x*- and *y*-axis.

Mechanical properties of the elastomeric patch antenna are evaluated after characterizations on its electrical performance. The implemented antenna prototype is more than twice as thick as the previously demonstrated stretchable RF electronic devices, and thus features degraded mechanical deformability. Yet moderate twisting or folding do not cause any mechanical damages to the fluidic patch antenna during experiments.

238 Advanced Elastomers – Technology, Properties and Applications

state to a high tensile strain.

cases.

Furthermore, this frequency offset should be as little as possible to avoid too high mismatch losses so that reasonably long remote sensing ranges as well as sufficient sensing sensitivity can be attained. Placing the antenna original resonant frequency above the highest operational frequency of the transmitter is not an option, as inconsistent variations of mismatch losses would occur if the antenna based strain sensor is stressed from its relaxed

Measured radiation patterns, including mismatch losses, at 1.46 GHz, of the relaxed and flexed antennas are shown in Figs. 22.b and c. The realized peak gain is measured along the +*z*-axis (according to the coordinate system depicted in Fig. 21.a.) of the mechanically reconfigurable antenna, and greatly degraded from 2.0 dBi to -10.7 dBi while applying an increasing strain from 0% to 15% along its *x*-axis. The meshed ground plane directs the radiation forward, with a front-to-back ratio as high as 10.0 dB, regardless of stretching. Moreover, the cross-polarization discrimination is very good in both relaxed and strained

**Figure 22.** a). Simulated and measured reflection coefficients of the elastomeric fluid metal patch antenna in its relaxed state, and with different strains along the *x*-axis (according to the coordinate system in Fig. 21). b). *xz*- and c). *yz*-plane radiation patterns (including mismatch losses) at the original resonant frequency of the elastic antenna in its relaxed state, and with 15% elongation along its *x*-axis. The antenna gain along *φ* and *θ* orientations is defined as *Gφ* and *Gθ*, respectively. The antenna co- and

Mechanical properties of the elastomeric patch antenna are evaluated after characterizations on its electrical performance. The implemented antenna prototype is more than twice as

cross-polarization along +*z*-axis are in parallel to the *x*- and *y*-axis.

The measured total efficiency of the standalone stretchable patch antenna including mismatch losses is 36.9% in its relaxed state at 1.46 GHz. Straining the patch to 15% along its *x*-axis leads to considerable total efficiency drop of 33.8%. In the intermediate states with the strains between 5% and 10%, the total efficiency of 11.5% and 5.8% is achieved according to the experiments.

The system demonstration setup is first calibrated by characterizing the output DC voltages of the RF power detector with different static stresses applied to the integrated wireless strain sensor, as presented in Fig. 23. A fairly linear decline of the measured DC voltages versus increasing mechanical elongation along the *x*-axis of the self-contained sensor device is seen, and can be explained by decreased resonant frequencies and increased mismatch losses of the transmitting, elastomeric antenna as a result of incremental stretching. As copolarization components dominate in both microfluidic stretchable patch antenna in the hybrid sensor and the receiving horn in the custom-designed RF receiver, the measured output voltages with respect to the cross-polarization are considerably lower than that of the co-polarization, and also exhibit smaller variations versus different strains. The steepest voltage decline in response to increasing stretch is found in the case of the co-polarization measured in an anechoic chamber. When it comes to ordinary office environment, the presence of reflections and scatterings limits the range of voltage variations.

**Figure 23.** Measured output DC voltages at the receiver versus varying mechanical stress along the *x*axis (according to the coordinate system in Fig. 21) applied to the integrated strain sensor. Its co- and cross-polarizations along the +*z*-axis were in parallel to the *x*- and *y*-axis, and the receiving standard horn was horizontally polarized.

Fig. 24 shows the demonstration of remote sensing of repeated body motion in a corridor of ordinary office environment, using the resulting wireless strain sensor prototype. Periodically repeated dynamic strains of up to 15% along its *x*-axis, with a period of 10s and a duty cycle of 50%, are manually applied to the two shorter edges of the integrated, elastic, strain sensor. The output DC voltages in the PC-assisted RF receiver varying with mechanical strains of the integrated, body-worn, sensor device, are continuously monitored and recorded, *cf*. Fig. 24.c. The measured data of six cycles, well correlating with the applied varying tensile strains, are presented in the subplot in Fig. 24.c, in which six fairly uniform,

quasi-rectangular waves with varying amplitude between 1.28V and 1.55V are recorded in an overall period of 60s. This experimental data also verifies that the implemented wireless strain sensor can rapidly return to its relaxed state without any hysteresis, once removing the applied stress. This great feature reflects reversible deformability as well as high degree of elasticity of the multi-layer μFSRFEs based sensor device.

**Figure 24.** Integrated elastomeric strain sensor a). in its original state and b). with 15% vertical strain. c). system demonstration (periodical manual straining). Real-time recorded data at the PC-assisted RF receiver 5m away is presented in the subplot.

### **7. Conclusion**

Recent advances in the emerging field of elastomeric electronics that are able to be conformed into complex curvilinear shapes, or be compressed, twisted, and stressed to extreme degrees, have been briefly reviewed. Various techniques and strategies for realizing bendable, flexible, stretchable elastomeric electronic devices and system have been discussed. As the main focus, elastomeric electronics based on microfluidic approaches have been addressed in detail. Fabrication processes, hybrid integration techniques, as well as appealing application examples involving single- and multi-layer μFSRFEs have also been presented.

Microfluidics based elastic electronics together with other members in the family of elastomeric electronics are introducing a revolution to the world of electronics, and shaping the future for electronics so as to change our tomorrow's daily life and contribute to our networked society. It is anticipated that more than 50 billion devices will be wirelessly connected by 2020, which would involve units as intelligent as smartphones/tablets, and as soft as elastomeric electronics.

### **Acknowledgement**

240 Advanced Elastomers – Technology, Properties and Applications

receiver 5m away is presented in the subplot.

soft as elastomeric electronics.

**7. Conclusion** 

of elasticity of the multi-layer μFSRFEs based sensor device.

quasi-rectangular waves with varying amplitude between 1.28V and 1.55V are recorded in an overall period of 60s. This experimental data also verifies that the implemented wireless strain sensor can rapidly return to its relaxed state without any hysteresis, once removing the applied stress. This great feature reflects reversible deformability as well as high degree

**Figure 24.** Integrated elastomeric strain sensor a). in its original state and b). with 15% vertical strain. c). system demonstration (periodical manual straining). Real-time recorded data at the PC-assisted RF

Recent advances in the emerging field of elastomeric electronics that are able to be conformed into complex curvilinear shapes, or be compressed, twisted, and stressed to extreme degrees, have been briefly reviewed. Various techniques and strategies for realizing bendable, flexible, stretchable elastomeric electronic devices and system have been discussed. As the main focus, elastomeric electronics based on microfluidic approaches have been addressed in detail. Fabrication processes, hybrid integration techniques, as well as appealing application

Microfluidics based elastic electronics together with other members in the family of elastomeric electronics are introducing a revolution to the world of electronics, and shaping the future for electronics so as to change our tomorrow's daily life and contribute to our networked society. It is anticipated that more than 50 billion devices will be wirelessly connected by 2020, which would involve units as intelligent as smartphones/tablets, and as

examples involving single- and multi-layer μFSRFEs have also been presented.

The author greatly appreciates the Editor Board in InTech for providing the free article processing support.

### **Author details**

Shi Cheng *Ericsson AB, Stockholm, Sweden* 

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**Chapter 10** 
