composite materials based on polydimethylsiloxane and in situ generated silica by using the...

Composite Materials Based on Polydimethylsiloxaneand In Situ Generated Silica by Using theSol–Gel Technique

Mihaela Alexandru,1 Mariana Cristea,1 Maria Cazacu,1 Aurelia Ioanid,1 Bogdan C. Simionescu1,2

1‘‘Petru Poni’’ Institute of Macromolecular Chemistry, Aleea Gr. Ghica Voda 41A, 700487 Iasi, Romania

2Department of Natural and Synthetic Polymers, ‘‘Gh. Asachi’’ Technical University, 700050 Iasi, Romania

A polydimethylsiloxane-a,x-diol with molar mass Mn =43,000 has been synthesized by cationic polymerizationof octamethylcyclotetrasiloxane and reinforced withsilica. Two pathways were used for incorporation ofsilica in the polymeric matrix: ex situ by mechanicalblending of a pretreated fumed silica and in situ byadding tetraethyl-orthosilicate (TEOS) as silica precur-sor in the polymer matrix followed by their hydrolysisand condensation (sol–gel technique). The procedureoccurred in the absence of solvent. Composites withdifferent contents of silica were prepared and investi-gated by dynamic mechanical analysis (DMA), differen-tial scanning calorimetry (DSC), and scanning electronmicroscopy (SEM). The results were compared to thoseobtained on a model network based on the same poly-siloxane without silica. POLYM. COMPOS., 30:751–759,2009. ª 2008 Society of Plastics Engineers

INTRODUCTION

It is well known that, despite of a large range of inter-

esting and useful properties, polysiloxanes have two

major drawbacks, i.e., poor mechanical properties and

high costs. There are two main approaches to minimize

these disadvantages: (a) the chemical insertion of organic

sequences in the siloxane backbone, either in the main

chain or as pendant chains; (b) the incorporation of inor-

ganic powders, usually silica, as filler in the polysiloxane

matrix. Silica is a useful additive for the improvement of

the polymer capabilities (thermal stability, mechanical

strength, insulating properties, etc.). However, the effects

of the silica on the properties of the composites depend

on the preparation method. Typically, such fillers are

introduced in the polymeric matrix by mechanical blend-

ing followed by crosslinking. There are some inherent

problems concerning this type of ex situ blending process:

time- and energy-consuming, the limited possibilities to

control the morphology or the surface characteristics of

the filler, and the occurrence of curing by polymer-filler

interactions that leads to premature strengthening of the

polymer [1]. The surface of silica is generally highly re-

active and, in order to avoid the structuration of the com-

posite during processing or storage, the filler must be

hydrophobized by a treatment that usually is laborious.

For these reasons, in 1980s Mark et al. first applied a

method specific for inorganic chemistry, the sol–gel tech-

nique, to prepare polydimethylsiloxane (PDMS) networks

containing in situ precipitated silica [2, 3]. The major

advantage of the network filling by this in situ approach

is the avoidance of the problems associated with mechani-

cal blending [3]. In addition, in the case of in situ rein-

forcement some properties of the resulted materials are

improved, as compared to those obtained by classical

methods. Certain applications of the filled PDMS elasto-

mers, such as protective masks, contact lenses, medical

and industrial tubing, and light guides require transpar-

ency. Small domain sizes relative to the wavelength of

light are necessary in order to obtain a high transparency.

If the filler is a powder, then the primary particles them-

selves must be small and their aggregation should be

eliminated or minimized. The in situ generation of silica

by sol–gel procedure proved to be a good approach to

obtain unusually transparent PDMS nanocomposites [3].

In principle, the sol–gel reaction can be viewed as a

two-step network-forming polymerization process. In the

first step, metal alkoxides are hydrolyzed to generate

intermediary species of metal hydroxides. Then, the spe-

cies undergo a stepwise polycondensation reaction to

form a three-dimensional network [4–6]. However, it is

known that, once the hydrolysis begins, the condensation

Correspondence to: M. Cazacu; e-mail: [email protected]

Contract grant sponsor: Project CEEX-MATNANTECH; contract grant

number: 52/2006; Contract grant sponsor: Project RAINS/INCO-CT;

contract grant number: 2005-017142.

DOI 10.1002/pc.20608

Published online in Wiley InterScience (www.interscience.wiley.com).

VVC 2008 Society of Plastics Engineers

POLYMER COMPOSITES—-2009

reactions occur simultaneously. The network structure

imposes restrictions to the molecular movements, thereby

making possible the preparation of a multicomponent sys-

tem without crystallization [4]. The use of sol–gel proc-

esses to generate SiO2 particles in PDMS polymers has

been studied extensively by Mark’s group [1, 2, 7–13].

This technique was applied to fill polysiloxanes in two

different ways:

Silica precipitated in situ after polymer curing [4]. This

consists in the swelling of the preformed PDMSs net-

works (prepared from PDMS crosslinked with tetraethyl-

orthosilicate (TEOS) in presence of stannous octoate as

catalyst) in a solution containing TEOS and 1 wt %

dibuthyltin diacetate (DBTDA) or dibuthyltin dilaurate

(DBTDL) [3];

Silica precipitated in situ during polymer curing [1]. The

process includes the mixing of the PDMS oligomer with

TEOS in a solvent mixture (e.g., tetrahydrofuran and 2-

propanol) in presence of added water as a hydrolysis

agent and hydrochloric acid as a catalyst [4]. In general,

low molecular polysiloxanes were used as polymeric ma-

trix, e.g., 1,700 and 550 [4], 4,200 [14], 8,000 [15],

1,700 and 18,000 [1], 21,100 and 22,600 [16].

In this study, we approached a modification of the sec-

ond protocol by extending the molar mass to a relatively

high value (43,000) and by working in the absence of sol-

vent. DBTDL was used as catalyst. No water was added as

hydrolysis agent, except for that absorbed by the sample

from the atmosphere. The polycondensation catalyst was

generated in situ from the tin salt (DBTDL) added to the

initial mixture [3]. The composites having different contents

of silica were prepared. The effects of the filler on the mac-

roscopic properties were investigated by dynamic mechani-

cal analysis (DMA), and the acquired data were correlated

with those obtained by differential scanning calorimetry

(DSC) measurements. The cross-section morphology of the

samples was analyzed by scanning electron microscopy

(SEM). The transparence of the samples was evaluated on

the basis of UV–vis transmittance spectra. The results were

compared to those obtained for a model network based on

the same polysiloxane but without silica.

EXPERIMENTAL

Materials

Octamethylcyclotetrasiloxane, [(CH3)2SiO]4, (D4), sup-

plied by Fluka AG, with the following characteristics: b.p.

¼ 1758C; n20D ¼ 1.396; d204 ¼ 0.955, purity [ 99% (GC),

was dried over Na wire and freshly distilled before use.

Purolite CT-175, a styrene-divinylbenzene ion

exchanger with ��SO3H groups (4.1 mequiv/g) was dehy-

drated by azeotrope distillation with toluene and vacuum-

ation at 1108C/10 mm Hg.

Fumed silica, Aerosil 380 (Degussa), 100% purity, spe-

cific surface 380 m2/g, particle diameter 0.003–0.015 lm,

was used after hydrophobization by treatment with D4 for

3 h at 1808C.Tetraethyl-orthosilicate (TEOS), purchased from Fluka

(purity[ 98%, b.p. ¼ 163–1678C, d204 ¼ 0.933) was used

as received.

Dibuthyltin dilaurate (DBTDL) was received from

Merck-Schuchardt, d204 ¼ 1.055 and was used as received.

Equipments

Gel permeation chromatographic analyses (GPC) were

carried out on a PL-EMD 950 Evaporative Mass Detector

instrument by using chloroform as eluant, after calibration

with standard polystyrene samples.

Fourier transform infrared (FTIR) spectra were

obtained on a Bruker Vertex 70 FTIR analyzer. Analyses

were performed on the films in reflectance mode (ATR),

in the 600–4,000 cm21 range, at room temperature with

2 cm21 resolution and accumulation of 32 scans.

UV–vis transmittance measurements on the films were

performed by using a SPECORD M42 spectrophotometer.

DMA were run on a Diamond Perkin Elmer apparatus,

in tension mode, at a frequency of 1 Hz. The temperature

scans were performed between 2150 and 3008C at 28C/min, in a nitrogen atmosphere. The films (10 3 10 3

0.8) mm were longitudinally deformed by small sinusoi-

dal stress and the resulting strain was measured. The force

amplitude used was well within the linear viscoelastic

range for the all investigated samples. The value of the

storage modulus (E0), the loss modulus (E00), and the ten-

sion loss tangent (tan d ¼ E00/E0) were obtained as a func-

tion of temperature. All samples were measured in the

same conditions. Moreover, the rate of cooling was main-

tained constant.

SEM was performed on a TESLA BS 301 SEM at

25 kV with a magnification of 300–15,000. The images

were recorded on freeze-fractured surfaces put on Al

support and coated by sputtering with Au thin films using

an EK 3135 EMITECH device.

DSC analysis was performed on Diamond Perkin

Elmer equipment between 2150 and þ1508C with a heat-

ing rate of 208C/min, in nitrogen. The glass transition

temperature was determined as the midpoint of the heat

capacity change in the heating scan.

Procedure

Synthesis of Polydimethylsiloxane-a,x-diol (HO-PDMS-

OH). In a thermostated reaction vessel equipped with

reflux condenser, thermometer, and mechanical stirrer,

100 ml D4, 2.5 g Purolite CT-175, and 0.5 ml water were

introduced. The temperature was increased to 708C and

the reaction mixture was maintained at this temperature

under stirring for 4 h. Then, the reaction was stopped by

removing the catalyst through filtration. The filtrate was

devolatilized in a rotavapor at 1508C/10 mm Hg to

remove the cyclic and linear low molecular weight com-

752 POLYMER COMPOSITES—-2009 DOI 10.1002/pc

pounds [17]. The molar mass of the remained polymer,

determined by GPC using toluene as eluant, was 43,000.

Polymer Filling with Silica

(a) Polymer filling by mechanical blending: HO-PDMS-

OH (100 g) was introduced in a blender with palettes.

A pre-established amount of treated silica was added

in small portions until the whole amount was incorpo-

rated. TEOS (0.005 ml) and DBTDL (0.0025 ml)

were then added and the blending was continued for

about 1/4 h. The resulted composite was poured into a

rectangular metallic frame and pressed between two

glass plates. This ensemble was maintained in this

manner for 7 days to allow the crosslinking. The

formed film was easily detached from the glass sur-

face.

(b) Polymer filling in situ by sol–gel procedure: 10 g

HO-PDMS-OH were introduced in a Teflon dish and

mixed with pre-established amounts of TEOS, accord-

ing to Table 1. After about 10 min of stirring,

0.17 mL of DBTDL was added and the stirring con-

tinued for another 10 min. The resulted mixture was

used to obtain thick films by pouring on a Teflon foil.

The films were maintained at room temperature for

48 h and another 24 h in vacuum at 508C. The

obtained colorless and transparent films (of about 0.8

mm thickness) were easily peeled off from the sub-

strate.

The films were then kept in the laboratory environment

about 2 months before investigations.

RESULTS AND DISCUSSION

The polydimethylsiloxane-a,x-diol with the numerical

molar mass of 43,000 was synthesized by cationic ring-

opening polymerization of D4 in the presence of Purolite

CT-175 as catalyst (Scheme 1). A certain amount of water

was added as chain transfer agent (or chain blocker) that

provides the terminal functional groups (��OH) and regu-

lates to some extent the molecular mass of the resulted

linear polymer [17].

This polymer served as matrix for both ex situ and

in situ prepared silica.

The used silica is a commercially available one having

particle diameter in the range 0.003–0.015 lm. To avoid

the strengthening of the compounds during storage, the

silica was hydrophobized by treatment with D4. Silica and

PDMS were mixed in a 0.55 molar ratio (PTMs, Table 1)

by mechanical blending.

The sol–gel technique was applied to generate the

in situ silica during silicone curing. One of the most pop-

ular precursors for inorganic polymerization [18] by the

sol–gel method, TEOS, was used. TEOS plays both silica

generator and polysiloxane crosslinker roles. It is known

that Si(OR)4 and its homologues are rapidly hydrolyzed

by water in the presence of acid or base as catalyst. The

hydrated tetrahedral silanol undergoes a polycondensation,

resulting in SiO2 networks [19]. The reaction mechanism

of the sol–gel process that involves hydrolysis and con-

densation of tetraalkoxysilane is variable and depends on

factors such as catalyst type, water content, and solvent

[20]. In the sol–gel process, it was proved that the acidity

of the environment and the water content of the system

display a critical effect on the structure of the final prod-

ucts. This happens because each of them influences the

hydrolysis rate and the chemical equilibrium of the sys-

tem [4]. Mainly, the type of a catalyst determines the

course of the polycondensation reaction and the final

structure of the silica in the organic matrix [20]. In the

base catalyzed sol–gel process, the initial hydrolysis is

slow, but each subsequent hydrolysis and condensation

occur quickly resulting in a mixture of highly branched

clusters and unreacted monomer. According to the litera-

ture data [19, 21], in the acid catalysis, hydrolysis and the

first condensation are rapid, but each subsequent conden-

sation takes place gradually slower. In the case of the

acid catalysis, the gel time is longer and the resulted

structures are more ramified. However, some additional

factors become important due to the addition of the poly-

TABLE 1. The prepared siloxane-silica composites.

Sample Preparation pathway

Feed ratio

Masic ratio TEOS:

HO-PDMS-OH

Molar ratio SiO2:a

HO-PDMS-OH

Corresponding SiO2,a

wt % total mixture

PTMo Pure crosslinked PDMS – – –

PTMs Crosslinked PDMS ex situ filled with silica – 0.55 31.0

PT1 Crosslinked PDMS in situ filled with silica 1:1 0.36 22.5




a SiO2 corresponding to the initial added TEOS.

SCHEME 1.

DOI 10.1002/pc POLYMER COMPOSITES—-2009 753

meric component: the amount of the added polymer, the

polymer molar mass, the difference between the solubility

parameters of the polymer and of the glassy component, as

well as the used solvent(s) and the reaction temperature.

All these variables may affect the miscibility of the system

during the reaction and, consequently, the structure and the

properties of the final products [4]. The tetraalkoxysilanes,

Si(OR)4, are soluble in common organic solvents and

therefore react efficiently with different organic com-

pounds. Silanol-terminated PDMS is often chosen as poly-

meric matrix due to the similarity of its backbone structure

(Si��O��Si) with the sol–gel glass matrix of TEOS [4]. In

this work, we used a polydimethylsiloxane-a,x-diol that,

unlike other literature reports, has relatively high molar

mass. The hydrolysis of TEOS relied on atmosphere hu-

midity. DBTDL was used as a catalyst for condensation.

Most probable, a structure close to the one presented

in Scheme 2 is formed.

The postprocessing of the samples, consisting in keep-

ing at room temperature for 48 h and for another 24 h in

vacuum at 508C, followed by their maintaining in the

laboratory environment about 2 months before investiga-

tions, seems to permit the condensation of the OH groups

to a high degree. FTIR spectra for some representative

samples are presented in Fig. 1, where the specific bands:

1258, 1268 (Si-CH3 sym.), 1009, 1007 (Si-O-Si asym.),

788, 768, 786 cm21 (CH3 rocking asym. Si-C) can be seen.

The quantitative UV–vis transmittance spectra for rela-

tively thick films based on pure crosslinked HO-PDMS-

OH (PTMo), crosslinked HO-PDMS-OH reinforced with

31% added silica (PTMs), and 52.8% in situ prepared

silica (PT3) are comparatively illustrated in Fig. 2. As

can be seen, the sample reinforced by in situ technique

has practically the same very good transmittance (close to

90%) as the pure crosslinked HO-PDMS-OH (PTMo), but

differs from the sample reinforced with ex situ prepared

silica that has lower transparence (about 70%). This can

be due to the better compatibilizing between silica and

PDMS in the first case, where the silica is found as a net-

work interpenetrated with PDMS one. In the ex situ rein-

forced sample, the silica is as particulate only.

Two types of restrictions may be imposed when PDMS

is incorporated into the network: (1) the one caused by

coupling its ends through chemical bonding (it may con-

nect to either a TEOS species or to another oligomer

through the silanol functionality); (2) the one caused by

the local dense structure of the three-dimensional network

developed through the self-condensation of the hydrolyzed

TEOS. As a result, the thermal energy needed to mobilize

a constrained polymeric chain would be higher than that

for an unconstrained one. This would be expected to

result in an increase of the glass transition temperature of

SCHEME 2.

FIG. 1. Illustrative FTIR-ATR spectra of the samples: (a) PTMo; (b)

PTMs; (c) PT3.

FIG. 2. Comparative UV–vis transmittance spectra of the samples: (a)

PTMo; (b) PTMs; (c) PT3.

FIG. 3. DSC scans (second heating) of the reinforcing samples: (a)

PTMo; (b) PTMs; (c) PT4; (d) PT1; (e) PT2; (f) PT3.


the polymeric chain. Because of polydimethylsiloxane-

a,x-diol self-condensation, two phases can be developed

within the silica network: dimethylsiloxane rich phases

and dimethylsiloxane poor phases. The phase-separated

PDMS should display a glass transition temperature (Tg)near that of the pure oligomer (about 21238C). If the

PDMS is better incorporated into the silica network in the

form of relatively short chains, the chain extending

through reaction with other oligomers can not occur. The

chain motion is restricted by the presence of the network

and, as a result, Tg would be shifted to a higher tempera-

ture [4]. In our case, the Tg values for the silica filled

PDMS samples evaluated from the DSC curves (Fig. 3)

are very close to that of pure PDMS. This can be

explained by the presence of pure PDMS domains in the

network. SEM was employed to study the fractured surfa-

ces (Fig. 4). Cryo-fractured specimens have been used for

SEM study.

In the case of PTMo sample, the general view of the

fractured surface revealed a homogeneous structure char-

acteristic to the crosslinked polymers and resins. Small

globular polymeric domains formed as a result of con-

straints imposed by crosslinks are visible on the detailed

image (Fig. 4a). As the TEOS amount increases the tex-

ture changes. Thus, in PT4 sample, the texture becomes

cylindrical with very rarely spherical silica domains (Fig.

4b). In PT2 the very dense silica domains having a high

dimensional distribution are developed (Fig. 4c). A strati-

fication of the silica domains near the surface is observed

in the fracture. The image is completely different in the

FIG. 4. Scanning electron micrographs of the broken surfaces deposed on Al supports and coated with Au: (a) PTMo (32,500); (b) PT4 (32,500);

(c) PT2 (32,500); (d) PTMs (33,100).


case of the sample consisting in PDMS reinforced with

silica prepared ex situ, PTMs (Fig. 4d). The fracture sur-

face presents a globular texture with the globule size

lower than in the case PTMo but with interstices between

them. It is presumed that powerful hydrophobic interac-

tions occur between pretreated silica and PDMS, the for-

mer being embedded in the polymeric matrix (bound rub-

ber). On this background, silica with particle dimension

of few tens nanometers or lower, are dispersed. This is

probably the silica remained hydrophilic, that appears as a

separated phase. By comparing samples with the same

silica content but different in the preparing procedure

(PTMs and PT2), the morphologies are also completely

different.

If one analyses at a glance the viscoelastic behavior of

the samples (Figs. 5–7), some similarities are noticeable:

a glassy region (T \ 21258C) with a storage modulus

higher than 109Pa (with the exception of the sample

PTMs, whose storage modulus goes under 109Pa), a one

or two-step descent of the storage modulus till 2508C,

followed by an abrupt descent until a plateau character-

ized by a specific modulus value for each sample.

Nevertheless, a detailed examination reveals peculiar

features for each system. The glassy storage modulus of

the PTMo sample (4.7 3 109 Pa) decreases in one step

till 1.3 3 109 Pa during the glass transition process in the

interval range between 2125 and 21058C. This fall rep-

resents less than one order of magnitude and usually sug-

gests the presence of some kind of constraints in the poly-

mer system, i.e., physical crosslinkings or crystalline

domains. Recalling that PDMS chains have unusually low

intermolecular forces, the physical crosslinks are

excluded. DMA offers few clues to the presence of poten-

tial crystalline domains. This is why additional DSC

measurements were carried out in order to ascertain

whether any crystallization processes take place (Fig. 3).

DSC experiment evidences for PDMS the step of heat

capacity (DCp), associated with the glass transition, an

exothermic peak and an endothermic peak attributed to a

cold crystallization phenomenon and to the melting of the

crystalline phase, respectively. The temperatures of these

processes are 2124.41, 2102.84, and 246.528C. The fact

that the melting peak is much more prominent than the

crystallization one (DHm/DHc [ 10) evidences that an im-

portant amount of crystalline phase is developed during

cooling scan. This is an outcome of the great mobility of

PDMS chains that make possible a three-dimensional

chain arrangement required for crystalline phase develop-

ment. Therefore, DSC experiment supports the assumption

that the small drop of E0 in the glass transition range is

justified by the presence of crystalline domains in PDMS.

These domains formed during cooling scan behave simi-

larly to crosslinkings narrowing the E0 fall. The glass

transition appears as a peak on E00 and tan d plots (Figs. 6

and 7, respectively), the former comes out at lower tem-

perature (21218C) than the tan d peak (21198C). For

clarity, each curve of Figure 7 was shifted in relation to

each other. Some significant thermal characteristics

resulted for all samples from DSC and DMA (Table 2).

The E0 variation with temperature levels off till

FIG. 5. Storage modulus vs. temperature plots for: (a) PTMo; (b)

PTMs; (c) PT4; (d) PT1; (e) PT2; (f) PT3.

FIG. 6. Loss modulus vs. temperature plots for: (a) PTMo; (b) PTMs;

(c) PT4; (d) PT1; (e) PT2; (f) PT3.

FIG. 7. tan d vs. temperature plots for: (a) PTMo; (b) PTMs; (c) PT4;

(d) PT1; (e) PT2; (f) PT3.


� 2508C, point where an abrupt descent on E0 (more

than three orders of magnitude) over of rather narrow

range of temperature (158) marks the melting of the crys-

talline phase. The sign on tan d plot for this melting is a

single peak at � 2408C. Beyond 2508C, PDMS pre-

serves a good thermal stability until over 1508C (the limit

of the experimental temperature), in spite of the quite low

value of E0 (5.3 3105 Pa), characteristic for an elastomer.

The DSC thermogram of the PT4 sample (Fig. 3, curve

c) detects no cold crystallization. However, the high and

sharp melting peak, with the biggest melting enthalpy,

denotes that this sample contains the largest proportion of

crystalline domains. Hence, one can presume that all the

crystalline phase was formed during the cooling scan.

This indicates that the macromolecular chains are unex-

pectedly more inclined than in the case of PTMo sample

to acquire the three-dimensional order required for crys-

tallization. Structurally, the PDMS macromolecular chains

might be visualized as a macromolecular coil that con-

fines the silanol groups in the inner part. Once in the sys-

tem, the natural tendency of TEOS molecules is to

migrate inside the coil, toward hydrophilic silanol groups.

As a result, the macromolecular coil expands and

becomes perceptibly more flexible. Moreover, the self-

condensation of HO-PDMS-OH cannot be excluded, but

the consequence is the disentanglement of PDMS chains

accompanied by flexibility growth, too. Accordingly, the

smaller decrease of E0 in the glassy region (2.2 3 109 Pa)

as compared to PTMo is not any more unusual since

the storage modulus is considered a gauge of sample

stiffness.

Keeping in mind that the water for the TEOS hydroly-

sis reaction comes only from the environmental moisture

and due to the small amount of initial TEOS, it is very

unlikely that an important amount of TEOS ethoxy groups

is rapidly converted to hydroxyl groups in order to per-

form the co-condensation reaction. However, it is still

possible that even the limited cocondensation reaction

develops few joints between TEOS and PDMS chains

generating a very loose network. The tan d peak that is

the mark of the glass transition of PDMS shifts discerni-

bly to 21158C (Fig. 7). Noteworthy, the right-side peak

of the a relaxation in tan d plot broadens substantial, indi-

cating a dispersion in the distribution of the relaxation

times associated with the PDMS chains. This shoulder

could be associated with the PDMS segments neighbored

upon network joints that are constrained PDMS segments.

This position of the PDMS characterized by lower mobil-

ity is known as bound rubber [22–25] Because of this

phenomenon the crystallinity degree, reflected in melting

endotherm intensity, decreases until the disappearance

with silica content increasing (Fig. 3, curves c–f).

Through a scrupulous examination of storage modulus

curve one discerns that after the break around 21058C,the second E0 declining becomes less steep because the

content of mobile chain segments diminishes. This argu-

ment reinforces the preceding DSC rationale. The DSC

thermogram of the PT1 sample indicates likewise that all

the crystalline domains are formed during cooling scan

and the peak corresponding to the cold crystallization is

absent (Fig. 3, curve d). The DSC melting peak is less

important beside the ones of PTMo and PT4 sample and

this is a sign for lower quantity of crystalline domains in

the material. This means that more TEOS included in the

network increases the stiffness—DMA confirms this fact

(E0glassy ¼ 4.1 3 109 Pa)—at the expense of the macromo-

lecular segments proportion prone to crystallization. The

steepness of the storage modulus increases in the second

descent step of the glass transition due to the reduction of

crystalline domains. Rising up TEOS amount, more joints

are formed between HO-PDMS-OH and TEOS and this

augment the quantity of bound rubber. On that account

the tan d plot registers the displacement of the shoulder

associated to the glass transition of constrained PDMS to

higher temperature (Fig. 7). The behavior changes in the

case of PT2 sample, but all the transformations reflect the

trend induced by the extra-TEOS added to the system.

According to the DHm value (Table 2) obtained in the

DSC experiment, it is evident that a less crystalline phase

is present in the system as compared to the previous sam-

ple (PT1). In addition, the ratio DHm/DHc � 2 reflects

that the crystalline domains are formed, in equal parts,

during cooling and heating steps. During the sol–gel pro-

cess, more joint points are formed in the network and

thus the probability of the three-dimensional ordering

decreases. The tan d shoulder attributed to the glass tran-

sition of the constrained PDMS moves to higher tempera-

tures and becomes part of the right-side melting peak

(Fig. 7, curve e). A special situation is encountered for

the PT3 sample. The crystallization process did not occur

during cooling scan as can be inferred from the absence

of the melting peak on the DSC thermogram (Fig. 3,

curve f). The crosslinking rate is so high that the bound

rubber increment renders into the immobility of the net-

work that inhibits any crystallization. As a consequence,

the polymer is frozen in the disordered glassy state.

Unfortunately, due to the brittleness of the sample, the

experiment was not successful in all the range of the neg-

TABLE 2. The main parameters of DSC and DMA curves.

Sample

Tg(DSC)a (8C)

Tg(DMA)b (8C)

Tcc

(8C)Tm

d

(8C)DHm/

DHce

PTMo 2124.4 2118.8 2102.84 246.52 10.6

PTMs 2119.6 2111 – 247.72 –

PT4 2119.24 2115 – 246.6 –

PT1 2119.88 2114 299.91 248.8 230

PT2 2124.59 2113 296.59 250.6 1.89

PT3 2128.10 2123 – – –

a Tg, glass transition temperature evaluated from DSC curves.b Tg, glass transition temperature evaluated from DMA curves (read

as tan d peak).c Tc, crystallization temperature evaluated from DSC curves.d Tm, melting temperature evaluated from DSC curves.e DHm/DHc, ratio between the melting and crystallization enthalpies.


ative temperatures. Around 2758C the sample starts to

slipper between the clamp chucks, although repeated

experiments were performed. Even so, we chose to repre-

sent the data for two temperature ranges: 21458C and

2758C; 258C and 2508C. In tan d plot (Fig. 7, curve f)

still comes out the peak at 21238C. The appearance of

this peak characteristic for pure PDMS is consistent with

the existence of quite long PDMS segments uninvolved in

the matrix. Maybe, in spite of the high molecular weight,

this PDMS is embedded in the dense network and no

crystallization happens. The melting point comes out at

the same temperature range for all the samples (2408C),because the nature of crystalline domains is alike.

Important to mention is that the storage modulus val-

ues reached after melting increase with the temperature

(Fig. 5) as well as with the initial quantity of TEOS, that

is with network density. During the sol–gel process,

incomplete condensation of TEOS can not be excluded.

As the temperature is raised to high positive values the

sol–gel process may perfect the condensation, the catalyst

being present in the system. The increase in storage mod-

ulus is consistent with a completion of reaction. The

impact of raising the TEOS quantity in the system on the

storage modulus values is well reflected by the Fig. 8.

A distinctive approach should be considered when

examining the HO-PDMS-OH/SiO2 mixture. Any types of

reactions are excluded because SiO2 has no reactive

groups for HO-PDMS-OH and the initial treatment elimi-

nates the possibility of H-bonding with the PDMS end-

chain hydroxyl groups. Solely, the presence of the melting

peak on DSC thermogram signifies that all the crystalline

domains are formed during cooling scan (Fig. 3, curve b).

Comparing DHm values (Table 2), it is evident that the

fraction of crystalline domains is commensurate with PT1

sample, but lower than PTMo and PT4 sample. Addition-

ally, SiO2 loosens the PDMS macromolecular coil and

generates the smallest E0 value (7.3 3 108Pa). Several

peaks and shoulders are noticed on DMA thermogram till

the melting. There is an environment wherein SiO2 rein-

forces the PDMS matrix therefore inducing the largest

increasing of Tg. Miscellaneous morphologies prompted

by the SiO2/PDMS proximity indicate that SiO2 influences

in different ways the PDMS segments. However, a precise

correlation of viscoelastic behaviour with specific mor-

phologies is at least risky. Notable are the strength and

the thermal stability of the network throughout the posi-

tive temperature interval, suggesting that SiO2 might act

as a reinforcing agent.

CONCLUSIONS

A relative high molar mass crosslinked PDMS rein-

forced with either ex situ or in situ silica has been pre-

pared and processed by casting as films of about 0.8 mm

thickness. Transparent composite samples were obtained

with an advanced condensation degree of the OH groups.

The thermomechanical tests performed by DMA empha-

sized differences in macroscopic properties of the

obtained materials depending on composition. The me-

chanical properties of the in situ reinforced materials

proved to be higher than those reinforced by ex situmethod. The DMA results correlate well with the DSC

ones.

Based on the DMA data, it can be appreciated that

storage modulus increases by increasing silica content.

This increasing is more significantly in the case of the

samples reinforced with silica prepared in situ as com-

pared to those containing silica prepared ex situ.Because of using of a PDMS with relatively high mo-

lecular weight, the Tg values evaluated from the DSC

curves are very close to that of pure crosslinked PDMS.

This can be explained by the formation of the dimethyl-

siloxane domains that behave as such.

The cross-section morphology evaluated by SEM on the

sample fractured in liquid nitrogen revealed a typical granular

or interpenetrated networks aspect, depending on the compo-

sition. SEM images support the DSC and DMA results.

FIG. 8. Storage modulus vs. feed molar ratio TEOS/PDMS plot: (a) at 258C; (b) 21408C.


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