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The effect of fillerfiller and fillerelastomerinteraction on rubber reinforcement

J. Frohlich*, W. Niedermeier, H.-D. Luginsland

Advanced Fillers and Pigments Division, Applied Technology Advanced Fillers, Degussa AG, Harry-Kloepfer-strasse 1, Koeln 50997, Germany

Abstract

The role of active fillers like carbon black and silica has been studied in the rubber matrix for a better understanding of the rubber

performance and the mechanism of reinforcement. In particular the influence of basic properties of carbon blacks, such as specific surfacearea, structure and surface activity on the Payne-effect, was investigated with the Rubber-Process-Analyzer (RPA) which allows a testing of

the strength of the filler network and the filler-polymer interaction in the green compound as well as in the vulcanizate in a wide range of

shear amplitudes. A comparison between carbon black and the silica-silane system leads to further scientific findings for the understanding of

the dynamic behavior of filled rubber compounds.

q 2004 Elsevier Ltd. All rights reserved.

Keywords: Carbon Black

1. Introduction

Despite the fact that carbon black is not only the most

widely used but also the oldest active filler in rubbercompounds and that a huge amount of research work[1] has

been carried out to characterize, describe and understand

this kind of filler, the reinforcement effect in filler-loaded

rubber is not satisfactorily explained.

Since carbon black is an essential ingredient of most

rubber formulations and has a considerable influence on the

performance of the final product, it is important to know how

rubber vulcanizates are affected by the nature of carbon

black. The aim ofthisreport is to givea deeper insight intothe

reinforcement mechanism of the carbon black filler system

by analyzing the strain dependence of the complex modulus

G

*

and the loss factor tand

. The Rubber-Process-Analyzer[2,3] (RPA, Alpha Technologies) allows a very reliable and

detailed investigation of the dynamic behavior on a much

easier and faster way than the common dynamic analysis of

vulcanizates. The main focus of this work lies on the

influence of the nature of carbon black (loading, surface area,

structure and surface activity) on the dynamic properties.

2. Background

Although the reinforcement of rubber by active fillers is a

well-recognized phenomenon the term reinforcement isnot well defined. Briefly it can be stated that reinforcement

means the pronounced increase in tensile strength, tear

resistance, abrasion resistance and modulus far beyond the

values expected on the basis of the Einstein-Guth and Gold

theory [4], taking into account the effects caused by

colloidal spherical particles (hydrodynamic effect) and

occlusion of rubber. The reinforcement of elastomers by

fillers has been studied in depth in numerous investigations

[5] and it is generally accepted that this phenomenon is

dependent, to a large extent, on polymer properties, filler

properties and processing.

Generally speaking, the primary filler factors influencing

elastomer reinforcement are:

The primary particle size or specific surface area, which,

together with loading, determines the effective contact

area between the filler and polymer matrix.

The structure or the degree of irregularity of the filler

unit, which plays an essential role in the restrictive

motion of elastomer chains under strain.

The surface activity, which is the predominant factor

with regard to fillerfiller and fillerpolymer interaction.

1359-835X/$ - see front matter q 2004 Elsevier Ltd. All rights reserved.

doi:10.1016/j.compositesa.2004.10.004

Composites: Part A 36 (2005) 449460

www.elsevier.com/locate/compositesa

* Corresponding author.

E-mail address: [email protected] (J. Frohlich).
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Most of the elastomeric components are deformed

dynamically and specified dynamic properties are required.Therefore the effect of strain amplitude on the dynamic

modulus was observed very intensively. The modulus of

filled rubbers decreases with increasing applied dynamic

strain up to intermediate amplitudes. A detailed study of the

low frequency dynamic properties of filled natural rubber

was carried out by Fletcher and Gent [6] and was later

extended by Payne [7,8]. In cyclic strain tests the shear

modulus can be simply expressed as a complex modulus G*:

G*ZG0CiG00 where G 0 is the in-phase modulus and G00 the

out-of-phase modulus. The phase angle d is given by:

tan dZG 00/G 0.

The addition of fillers to rubber compounds has a strongimpact on the static and dynamic behavior of rubber

samples. Fig. 1 shows the typical behavior of the complex

shear modulus of filled rubber samples versus dynamic

shear deformation. Similar to the model of Payne, we see

the strain-independent part of the modulus as a combination

of the polymer network, the contribution from the

hydrodynamic effect and the modulus resulting from the

in-rubber structure.

(a) The polymer network contribution depends on the

crosslink density of the matrix and the nature of the

polymer.

(b) The hydrodynamic effectin this modelis nothingelse than the effect of strain amplification, resulting from

the fact that the filler is the rigid phase, which cannot be

deformed. As a consequence, the intrinsic strain of the

polymer matrix is higher than the external strain yielding

a strain-independent contribution to the modulus.

(c) The effect of the structure is attributed to the in-rubber

structure, which can be understood as a combination of

the structure of the filler in the in-rubber state (in-rubber

DBP) and the extent of fillerpolymer interaction. The

in-rubber structure is the measure for the occluded

rubber, which is shielded from deformation and there-

fore increases the effective filler content leading also to a

strain-independent contribution to the modulus. The

fillerpolymer interaction can be attributed to physical

(van der Waals) as well as to chemical linkages or a

mixture of both. In the case of the silicasilane system

this interaction is formed by chemical linkages.

(d) The stresssoftening at smallamplitudes is attributed to thebreakdown of the inter-aggregate association respectively

to the breakdown of the filler network. This stress

softening at small deformations, called Payne-effect

[9,10], plays an important role in the understanding of

reinforcement mechanism of filled rubber samples [11].

3. Carbon black characterization

3.1. Specific surface area

The primary particle size or specific surface area ischaracterized by different adsorption methods of specified

molecules. The most widely used adsorption methods are

based on [12]:

(a) The iodine number, which reflects a not true surface

area, because it is affected by porosity, surface

impurities and surface oxidation.

(b) The cetyltrimethyl ammonium bromide (CTAB)

method, which analyzes the so-called external surface

area which corresponds to the accessible surface area of

carbon black for an elastomer.

(c) The BET nitrogen adsorption surface area, which

provides the total surface area including porosity.(d) The Statistical Thickness Surface Area (STSA), which

is an alternative technique for the determination of the

external surface area.

All these measures surface areas are affected in different

ways by the nature of the carbon black surface and have

not to correlate exactly with the actual occupied surface by

an elastomer in a mixture. This should be kept in mind

when physical properties are deduced from the surface

area.

3.2. Structure

The primary particles formed during the initial carbon

black formation stage fuse together building up three-

dimensional branched clusters called aggregates. High-

structure blacks exhibit a high number of primary particles

per aggregate, which is called a strong aggregation, low-

structure blacks show only a weak aggregation. These

aggregates again may form loose agglomerates linked by

van der Waals interactions. The empty space (void volume)

between the aggregates and agglomerates, usually

expressed as the volume of dibutylphthalate (DBP)

absorbed by a given amount of carbon black, is described

Fig. 1. Idealized form of a typical elastic modulus curve.

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by the term structure of carbon black. It is assumed that the

DBP absorption reflects the total carbon black structure

which comprises both aggregates and agglomerates,

whereas the crushed DBP absorption test eliminates loose

agglomerates and easily destroyable aggregates. DBP and

CDBP absorption can be measured very precisely, but do

not give any information about the shape of the aggregates.Furthermore the DBP value does not reflect the actual

structure of the carbon black inside a compound since the

agglomerates change their shape during mixing.

It is also not surprising that the CDBP value does not

have to correlate automatically with the static dynamic

behavior because the agglomerates are separated into

aggregates by the shear forces during mixing and are not

firstly compressed and then dispersed as simulated by the

crushed DBP test. However the size as well as the shape of

the aggregates, which are not characterized by DBP or

CDBP absorption, are important primary carbon black

properties and should be analyzed in detail.

3.3. Surface activity

The third reinforcing parameter, surface activity, a poorly

defined term, but widely used in the filler field, can in a

chemical sense be related to different chemical groups on the

surface. Such are carboxyl, quinone, phenol and lactone

groups. In a physical sense, variations in surface energy

determine the capacity and energy of adsorption. The surface

chemistry of carbon blacks has a significant effect only on the

vulcanization behavior of filled compounds. No direct

correlation has been demonstrated between chemical groups

on the carbon black surface and rubber related properties[1315]. In the case of carbon black the fillerpolymer

interaction is mainly of physical nature (physisorption).

The surface activity can be influenced by heat treatment.

It is well-known that heat treatment of carbon blacks at

temperatures ranging between 1500 and 2700 8C [15]

profoundly changes the carbon black influence on rubber

compound properties. The result is a low level of bound

rubber, a drastic reduction in high strain modulus, tensile

strength and abrasion resistance. The explanation for this

behavior is that during heat treatment the size of the nano-

crystallites, building up the primary particles of carbon

blacks, drastically increases and the arrangement of these

crystallites becomes more ordered. As a consequence the

microstructure is totally changed and the number of high-

energy sites, located at the edges of the crystallites,decreases drastically. This confirms the assumption that

the surface activity of carbon blacks, as far as the capacity of

the surface for adsorption is concerned, has its origin in the

disordered arrangement of the crystallographic structure.

Compared to the highly oriented pyrolytic graphite, the

graphitic structure of graphitized carbon blacks still remains

in a certain way imperfect [16].

4. Experimental

All investigations in the present study were performed

with the rubber process analyzer (RPA); the method and thespecific settings are described elsewhere [17].

All investigated rubber compounds are based on an

emulsion SBR 1500 formulation according to Table 1. The

mixing procedure is given in Table 2.

5. Results

5.1. Influence of carbon black loading

The influence of the carbon black loading on the dynamic

modulus of the green compound as well as of the

vulcanizate can be seen in Fig. 2. The black N 115 was

chosen as example, but in principle the behavior is similar

for all grades of carbon black. The unfilled rubber shows no

indication of non-linearity in the green compound as well as

in the vulcanizate. After adding the filler, the low strain

modulus G0 rises more that the high strain modulus GN,

resulting in a non-linear viscoelastic behavior, known as

Payne-effect G0KGN.

In both cases the overproportional increase of the Payne-

effect with filler loading can be seen. This increase is caused

by the formation of fillerfiller interactions, which are

formed with carbon black loading. This is due to the fact

Table 1

Basic formulation

Stage I

E-SBR 100

Carbon black 60

Other chemicals: ZnO 3; stearic acid 2; wax 1

Stage II

Batch stage I

Stage III

Batch stage II

CBS 1.5

Sulfur 1.5

Table 2

Basic mixing procedure

Stage I Intermix GK 1.5 N

010 Polymers

120 1/2 carbon black, ZnO, stearic acid

230 1/2 carbon black, 6PPD, wax

3 0 Sweep

34.50 Mix

4.50 Sweep

4.560 Mix and dump

Stage II

0.20 Batch stage I

250 Remill at 150 8C

Stage III

0.20 Batch stage II, accelerators, sulfur

2 0 Dump and homogenize on an open mill

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that the inter-aggregate distances become smaller with

rising filler content, and therefore the probability for the

formation of a filler network increases. The comparison of

the G*-curves of the vulcanizate and of the green compound

for fixed loading displays an increase of the Payne-effect

G0KGN during vulcanization. The crosslinkage of the

rubber matrix causes a parallel shift along the G*-axis, see

Fig. 1. As a consequence, the carbon black aggregates must

Fig. 2. Strain amplitude dependence of G* of N 115 at different loadings (a) green compound, (b) vulcanizate.

Fig. 3. Strain amplitude dependence of tan d of N 115 at different loadings (a) green compound, (b) vulcanizate.


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reagglomerate respectively flocculate during the vulcaniza-

tion process, which is possible because of the low viscosity

of the polymer matrix at high temperatures. The effect of

flocculation of fillers in green compounds annealing them atintermediate and high temperatures is of the same nature

and is well-known [18].

Fig. 3 gives the corresponding tan d-curves for the green

compounds and the vulcanizates. At small and intermediate

strains the hysteresis decreases with increasing filler content

in the case of the green compound. The unfilled sample

exhibits the highest tan d level, which is due to the viscous

nature of an unfilled raw polymer. With increasing carbon

black loading filler networking develops and stabilizes

the viscous polymer fluid yielding a diminished hysteresis

behavior at small and intermediate strains. At high strains,

when the filler network is broken down, the viscous polymer

is no longer stabilized and the tan d rises drastically.

The contribution of the breakdown of the fillerfiller

interactions to the hysteresis is not significant in the green

compound because of the already high level of tan d

resulting from the raw polymer.

The absolute values of tan d are considerably reduced in

the vulcanizates due to the crosslinkage of the polymer

matrix. The tan d-curve of the unfilled sample indicates that

there is essentially no hysteresis from the pure gum. Now it

can clearly be seen that the hysteresis increases with carbon

black loading. It can be stated that this hysteresis results

from the breakdown of the increased filler network whose

disruption during straining dissipates energy. So the concept

of filler networking yields a good interpretation of the

Payne-effect both for carbon clack filled green compounds

and vulcanizates, comprising the overproportional increaseof the Payne-effect with filler loading and the increase by

annealing.

It should be noted that the existence of filler networking

is not necessarily corresponding to the existence of a filler

network which is percolated through the whole specimen.

Fillerfiller contacts in local subnetworks are also denoted

as filler networking and yield the Payne-effect.

5.2. Influence of carbon black surface area

Fig. 4 shows the RPA-curves of three carbon black

grades with varied CTAB surface area, but similar structure(Table 3).

First of all it can be recognized that the GN

level at high

strain is comparable for all grades. On the other hand the

initial modulus G0 differs drastically from the variation

Fig. 4. Strain amplitude dependence of G* (a) and tan d (b) of three grades with different surface area but similar structure.

Table 3

Analytical characteristics of the blacks N 115, N 220 and N 339

CTAB (m2/g) DBP (ml/100 g) CDBP (ml/100 g)

N 115 128 113 97

N 220 111 113 101

N 339 92 121 99


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in surface area, the higher the surface area the higher

the Payne-effect which corresponds to the extent of filler

networking. The same tendency is displayed by the graph

tan d versus strain. The influence of the CTAB surface area

on the formation of a filler network can be explained as

follows: the primary particle size diameter decreases with

increasing surface area. At fixed levels of structure and

carbon black loading, both the aggregate size and the inter-

aggregate distance decrease with reduced primary particle

diameter respectively increasing surface area [19]. The

smaller the inter-aggregate distance the higher the prob-ability for the formation of a filler network. Consequently,

the extent of filler networking is more pronounced, when the

surface area rises.

5.3. Influence of carbon black structure

In order to investigate the effect of the carbon black

structure on the complex modulus two ASTM grades N 326

and N 330, which show strong deviations in DBP level, but

almost identical surface areas (Table 4), were selected.

In Fig. 5a, a comparison of the modulus curves of the

blacks N 326 and N 330 demonstrates higher modulus

values for the initial as well as for the final strain

deformation for the black with the higher structure.

Strictly speaking the higher structure of the black N 300imparts a parallel shift of the curve of the black N 326 to

higher G* values. Two consequences can be stated: first an

increase in structure leads to an enhancement in high strain

modulus; second, this increase is independent of the

deformation and is also valid for small strain amplitudes.

The parallel shift indicates that the Payne-effect G0KGN is

not significantly affected by structure. The expected

increase of the Payne-effect by a higher structure might be

compensated to some extent by a better microdispersion

behavior. The macrodispersion for all samples is between 8

and 9 in the well-known Phillips dispersion rating. The

influence of structure alone on the modulus curve can beseen by changes of the high strain modulus, because it is not

affected by surface area (see Fig. 4). The influence of the

surface area on the high strain modulus can be analyzed in

more detail by adding the blacks N 347 and N 375 for

comparison to Fig. 5a (Fig. 5b). Despite the fact that the

surface area of N 375 is higher than for N 347 ( Table 4) the

G* level for the final modulus is higher for N 347. The only

explanation for this behavior is the higher structure of N

347. Studying the order of the blacks regarding the final

Table 4

Analytical characteristics of blacks of the N 300 series


N 326 81 71 68

N 330 82 102 88

N 347 86 124 97

N 375 95 113 96

Fig. 5. (a) Influence of the structure on the dynamic modulus. (b) Influence of four different blacks of the N 300 series on the dynamic modulus.


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modulus it can be recognized that the order is in line with

the DBP numberthe higher the structure the higher the

final modulus. As seen before, above 57% strain the carbon

black network is destroyed to a certain extent, may be not in

individual aggregates but in the form of unconnected

subnetworks of carbon black, hindering the polymer chain

movement upon straining.

5.4. Influence of carbon black surface activity

One of the objectives of this investigation was to

determine the relative contribution of carbon black

surface activity to the complex modulus curve. In order

to judge the effect of surface activity two carbon blacks

of different structure, but having comparable particle sizeand surface area, were selected for surface modification

by high-temperature treatment (2700 8C) in a nitrogen

atmosphere. Heat treatment is known to change surface

properties without altering the particle size and structure

to any important degree [20]. Thus, with the different

structure levels chosen and the varying surface activity

produced by heat treatment, the effects of these variables

could be studied separately. Despite the fact of similar

surface area and structure levels of graphitized and non-

graphitized grades, the stressstrain behavior is distinctly

different. At low strain amplitudes a stronger filler

networking can be observed in the case of the graphitized

blacks, whereas at high strain the complex modulus is

significantly lowered compared to the corresponding

ASTM grade (Fig. 6).

The deviations are so marked that any fluctuation in

DBP or CTAB number cannot cause these differences.

As a consequence, these changes in complex modulus

behavior can only be attributed to the lowering in surface

activity for the graphitized carbon blacks. It can be

argued that low surface activity results in a decreased

interaction with the rubber phase, providing a reduction

of the high strain modulus and therefore a reduction of

the strain-independent part of G*. On the other hand,

fillerfiller and fillerpolymer interaction are competitive

in a certain way. So it can be understood that for

graphitized blacks, which exhibit a lower fillerpolymer

interaction, as seen before, filler networking should be

more pronounced. This is confirmed by the higher Payne-

effect for the compounds containing the graphitized

grades. Up to now only a deactivation of the surface

activity by heat treatment could be achieved and

investigated. The novel nanostructure blacks, produced

by a physical modification of the carbon black pro-

duction process, are characterized by a high surface

roughness and a large number of high-energy patches or

sites equivalent to a high surface activity. The origin of

the roughness found at the carbon black surface is

explained by a reduced lateral extension of the nano-

crystallites linked with a high degree of geometrical

disarrangement. These crystallites build up the primary

particles [2124]. The active sites of carbon black

particles are attributed to the edges of the graphitic

layers respectively crystallites [15]. Both facts, the

reduced size of the crystallites, leading to a larger

number of crystallites to form the same particle diameter,

as well as the high degree of disarrangement, yield a

large number of edges respectively high-energy sites.

This conclusion is valid in both directions. It is well-known that under heat treatment a growth of the

crystallographic dimension (LA and LC) of the carbon

black takes place at constant primary particle sizes. As

consequence in particular the high-energy sites are

reduced. These sites play a very important role in

elastomer reinforcement [25].

The surface activity of the nanostructure blacks therefore

should be higher compared to conventional ASTM grades.

In order to investigate only variations in surface activity the

ASTM black N 234, graphitized and non-graphitized, and

the nanostructure black EB 171, with almost the same

surface area respectively primary particle diameter and

structure (Table 5), were chosen.

In Fig. 7 the modulus curves of these three blacks show

the expected trends regarding the differences in surface

activity. The high surface activity of the nano-structure

black results in high interaction with the rubber phase,

preventing carbon black network formation and reducing

the low strain modulus.

The differences in high strain modulus are not so

considerable for EB 171 compared to N 234. Nevertheless

the trend of high strain modulus enhancement with

increasing surface activity of vulcanizates containing an

nanostructure black can clearly be seen.

Fig. 6. Complex modulus versus strain for graphitized (N 326 g, N 347 g)

and untreated grades.

Table 5

Analytical data of nanostructure (EB 171) and corresponding ASTM black


N 234 117 125 102

EB 171 118 125 103


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As expected the tan d curve versus strain gives for all

deformation distinctly lower values for the nanostructure

black regarding the lower filler network, as a consequence

of the strong mechanical/physicochemical interaction withthe polymer resulting from the high surface activity.

5.5. Payne-effect of silica and silicasilane system

It is well-known that silica as a filler itself shows a very

strong filler networking due to poor compatibility to

hydrocarbon rubber, its polar character and the ability to

form hydrogen bonds [26]. In order to study this strong

network two silicas with different surface areas (Table 6)

were investigated in a solution SBR/BR formulation

(Table 7) and compared to the carbon black N 234.

The following explanations based on the so far obtained

findings for carbon black can be also applied to silica. As

can be seen in Fig. 8, the fillerfiller interaction of Ultrasil

7000 GR with the high surface area of 160 m2/g (CTAB) is

significantly higher than the one of Ultrasil VN 2 GR

(Table 7). Based on the G*-curve the following can be

observed: The silica network starts to break down at higherstrain amplitudes due to a stronger filler network formed by

hydrogen bonds compared to carbon black.

This breakdown takes place over a much wider range of

amplitudes. The tan d versus strain curves reflect the same

tendency. In the case of the carbon black the maximum

energy dissipation, due to the breakdown of the filler

Table 7

Solution SBR/BR formulation

Stage I

S-SBR Styrene 25% 96Vinyl 50%

Arom.oil 37.5 pbw

BR cis-1,4O96% 30

Filler 80

Silane Only for silica 6.4

Other chemicals: ZnO 3; stearic acid 2; arom. oil 10; 6PPD 1.5; wax 1

Stage II

Batch stage I

Stage III

Batch stage II

DPG 2

CBS 1.5

Sulfur 1.5

Fig. 7. Complex modulus (a) and tan d (b) versus strain as a function of surface activity.

Table 6

Analytical data of the silicas Ultrasil VN 2 GR, Ultrasil 7000 GR and the

carbon black N 234

CTAB (m2/g) DBP (ml/100 g)

Ultrasil VN 2 GR 125 z185

Ultrasil 7000 GR 160 z235

N 234 125 119


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Fig. 8. Comparison of Ultrasil 7000 GR, Ultrasil VN 2 GR and N 234 regarding complex modulus and tan d.

Fig. 9. Shear modulus G* and tan d of vulcanizates in S-SBR/Br formulation of Ultrasil 7000 GR, Ultrasil 7000 GRCPTES and Uitrasil 7000 GRCSi 69.


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network, is reached at 57% strain, whereas even at 42%

strain the maximum is not reached in the case of the silicas.

The breakdown of the inter-aggregate bonding of the

silica over a broad range of strain amplitudes indicates that

the silicasilica interaction differs strongly from the van der

Waals bonding of carbon black aggregates [27].

The hydrophobation of the silica surface by a mono-

functional alkylsilane (PTES) makes the silica more

compatible with the rubber matrix reducing the filler

networking respectively agglomeration. This can be seen

on the drastically reduced low strain modulus (Fig. 9). Byadding a bifunctional silane TESPT (Si 69) instead of PTES

not only an additional hydrophobation but also a chemical

bonding of filler-to-rubber takes place. This is indicated by

the once again lowered G* level at low strain and a higher

G* level at high strain for Si 69 instead of PTES. The higher

G* value at high strain can clearly be attributed to the

bounded rubber (in-rubber structure) given by chemical

linkages between the silica surface and the rubber matrix via

Si 69 and in the silica structure [28]. Thus it can be stated

that an enhanced high strain modulus indicates a better

fillerpolymer interaction, this is also valid for carbon black.

The effect of the lower filler networking can also be seen in

the lowered tan d curve of Fig. 9.

6. Discussion

6.1. In-rubber structure

As mentioned above in-rubber structure was interpreted

as a combination of the effects caused by the real carbon

black structure in the in-rubber state and the fillerpolymer

interaction. Medalia has proposed [29] that in a carbon

blackrubber system the rubber which fills the void space

within each aggregate is occluded and shielded fromdeformation and thus acts as part of the filler rather than

as part of the deformable matrix. Due to this phenomenon,

the effective volume of the filler is increased considerably

and should of course enhance the high strain modulus

significantly. As can be seen for the graphitized black the

modulus level of the vulcanizate at high strain is lowered

compared to the non-graphitized one (Fig. 6). Due to

the fact that the DBP number of the graphitized and

non-graphitized blacks is on the same level, it can be stated

that the structure in the in-rubber state is nearly ineffective,

if the surface activity of the filler is low. On the other hand,

without enough structure in the in-rubber state, the surface

activity is also relatively ineffective (see Fig. 5, comparisonof N 326 and N 330). Hence, it can be concluded, in-rubber

structure is a measure for the occluded rubberrubber,

which is unable to participate fully in the macrodeforma-

tion. The nature of surface activity in this context can be

attributed to different featureschemical-type (silica

silane), physical-type (carbon black) bondings or both.

Fig. 11. Dynamic modulus G* and tan d vs. temperature in E-SBR (0.1% dynamic strain, 10 Hz, Gabo Explexor).

Fig. 10. Proposed model for the fillerfiller contacts, specific surface area

high (1) and low (2), equal loading.


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As a consequence, the term in-rubber structure explains

immediately the term of reinforcement, defined as modulus

300% divided by modulus 100% in a stressstrain

measurement. Reinforcement is a measure for the strength

of a filler to keep its in-rubber structure respectively to

restrain the occluded rubber at increasing deformation.

6.2. Fillerfiller/fillerpolymer interaction

Silica alone is building up a strong filler network and

shows a low interaction with the polymer. But if a

bifunctional silane is used chemical linkages between the

silica surface and the polymer are formed. Through the

hydrophobation the fillerfiller interaction is lowered.

The bondings are of chemical nature, and therefore polymer

segments cannot undergo surface displacements by mol-

ecular slippages. In the case of carbon black linkages

between the surface and polymer chains are formed only by

physical adsorption. In the past as well as in the presentmolecular slippages along the carbon black surface as well

as adsorption and desorption mechanisms are still under

discussion.

Up to now, it appears almost impossible to clarify exactly

the molecular mechanism of the fillerfiller contacts. All

experimental results known, those reported in this paper and

results from investigations of the dynamic modulus as a

function of temperature might be explained both by direct

filler to filler contacts and fillerfiller contacts which are

separated by a very thin polymer layer with a thickness in

the nanometer range. The existence of such an elastomeric

layer between the particles could at least be confirmed bydielectrical investigations done by Goritz and Lanzl [30].

Therefore the authors prefer the idea of a thin polymer

layer between the filler particles where the mobility of the

molecular polymer chains is restricted to some extent due to

fillerpolymer interaction, depending on the type of the

polymer and the filler to filler distance which is mainly

controlled by the specific surface area and the filler loading

(Fig. 10). In this case, the amount of overlapping areas

containing polymer with lower mobility controls the

hardness of the contacts and as a consequence the dynamic

modulus and the hysteresis during breaking off the contacts

under deformation.

From a steady decrease of the chain mobility near the

boundary of the filler surface one might expect a glass

transition, which has an onset at elevated temperatures. This

should especially be valid if the glass transition is measured

under dynamic conditions at very low strains where the filler

networking is fully active.

Fig. 11 shows the dynamic modulus G* and tan d in

dependence of the temperature and compares an unfilled E-

SBR compound and a compound filled with 50 phr N 234

(0.1% dynamic strain, 10 Hz, Gabo Eplexor). As expected,

an earlier onset of the glass transition of the carbon black

filled sample can be observed in G*, whereas the position of

the tan d maximum is not affected. But is has to be

mentioned that the measured effect is rather small.

7. Conclusion

The present study was undertaken to give a more detailedanalysis of the influence of carbon black morphological

parameters (surface area, structure and surface activity) on

viscoelastic properties of filled rubber compounds. Using

the Rubber-Process-Analyzer the phenomenon, reinforce-

reinforcement of active fillers can be studied with high

accuracy and repeatability. Two essential parameters of

filled rubber are investigated with this apparatus. For the

following conclusions two prerequisites have to be fulfilled:

the crosslink density and the filler loading have to be

constant.

7.1. High strain modulus (O30%)

The complex modulus at high strain indicates the real in-

rubber structure of fillers.

The high strain modulus of carbon black filled

compounds is enhanced only if both, the filler structure

in the in-rubber state and the surface activity of the filler

are high, which means that the in-rubber structure is high

as well. Consequently, the in-rubber structure is com-

posed of two parameters, the actual filler structure in

the in-rubber state and the surface activity of the filler.

The strong enhancement of the high strain modulus by

incorporating carbon black into rubber must be attributed

to the in-rubber structure, which is a direct measure forthe occluded rubberrubber, which is shielded from the

deformation of the matrix leading to an increase in

effective filler volume. Surface activity in this context

can be attributed to chemical-type, physical-type bond-

ings or both.

7.2. Low strain modulus (/5%)

The complex modulus at low strain, exactly G0KGN,

indicates the fillerfiller network.

The higher the filler loading the stronger the filler

networking is marked. Increasing surface area respectively

decreasing primary particle size provide also a higher filler

networking. Both findings can be explained by the fact that

the inter-aggregate distance becomes smaller with increas-

ing surface area as well as loading and thus the probability

of forming a network rises. Filler networking respectively

the Payne-effect G0KGN is decisively responsible for the

hysteresis level at amplitudes up to 10% strain. Therefore

filler networking plays a major role in the dynamic

properties of a tire, namely the rolling resistance, as well

as in the electrical resistance. On the other hand, a higher

surface activity leads to a decrease in filler networking due

to a better fillerpolymer interaction. It can be stated, if


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the affinity of the filler surface to the polymer is increased

respectively decreased, a reduced respectively enhanced filler

networking is the consequence. Therefore fillerfiller and

fillerpolymer interaction are two competitive processes.

References

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