quantification of the immobilized fraction in polymer ... of the immobilized fraction in polymer...

Quantification of the Immobilized Fraction

in Polymer Inorganic Nanocomposites

Dissertation

zur

Erlangung des akademischen Grades

doctor rerum naturalium (Dr. rer. nat.)

der Mathematisch-Naturwissenschaftlichen Fakultät

der Universität Rostock

vorgelegt von

Albert Sargsyan, geboren am 24. Juni 1980 in Jerewan,

Armenien

Rostock, 28 März 2007

Gutachter:

PD Dr. rer. nat. habil. Doris Pospiech, Leibniz-Institut für Polymerforschung Dresden Prof. Dr. Anahit Tonoyan, State Engineering University of Armenia Prof. Dr. Christoph Schick, Universität Rostock

Tag der Verteidigung:

11. Mai 2007

CONTENT

1. Introduction.................................................................................................. 5

2. Literature review .......................................................................................... 9

2.1. Polymer nanocomposites...................................................................... 9

2.1.1. Nanoparticles.................................................................................. 9

2.1.2. Preparation methods of polymer nanocomposites........................ 12

2.1.3. Morphology................................................................................... 17

2.1.4. Interfacial interactions................................................................... 19

2.1.5. Calorimetry ................................................................................... 23

2.2. Semicrystalline polymers .................................................................... 26

2.2.1. RAF in semicrystalline polymers................................................... 26

2.2.2. Vitrification of RAF........................................................................ 30

2.2.3. Devitrification of RAF.................................................................... 32

2.3. Heat capacity determination................................................................ 39

2.3.1. Linear scanning ............................................................................ 41

2.3.2. StepScan DSC ............................................................................. 44

3. Experimental.............................................................................................. 49

3.1. Materials ............................................................................................. 49

3.2. Preparation methods........................................................................... 53

3.2.1. Solution method............................................................................ 53

3.2.2. Shear mixing................................................................................. 54

3.2.3. Classical emulsion polymerization................................................ 54

3.2.4. Microemulsion polymerization ...................................................... 55

3.3. Characterization.................................................................................. 55

3.3.1. Gel permeation chromatography .................................................. 55

3.3.2. Electron Microscopy ..................................................................... 58

4

3.3.3. Small angle X-ray scattering..........................................................60

3.3.4. Thermogravimetry .........................................................................62

3.4. RAF determination ...............................................................................65

3.5. Annealing experiments.........................................................................66

4. Results........................................................................................................71

4.1. DSC measurements.............................................................................71

4.2. Specific heat capacity correction..........................................................75

4.3. RAF determination ...............................................................................84

4.4. Annealing experiments.........................................................................88

4.5. Devitrification of RAF at high temperature ...........................................94

4.5.1. StepScan DSC ..............................................................................94

4.5.2. High rate DSC ...............................................................................95

4.6. Plasticization experiments....................................................................95

5. Discussion ..................................................................................................99

6. Summary ..................................................................................................105

7. References ...............................................................................................107

Appendix........................................................................................................ A1

A1. Specific heat capacity data corrected .................................................. A1

A2. The calorimetric data from annealing experiments .............................. A3

A3. RAF layer thickness estimation ........................................................... A5

1. INTRODUCTION

Polymer nanocomposites have attracted a great deal of attention in

recent years due to their exceptional properties. Searching in SCOPUS™ [1]

for “polym* nanocompos* OR polymer inorganic hybrid” yields more than

8,000 hits from journals and more than 35,000 patents [2-17] and references

therein to name a few. Layered silicates [3], ceramic nanoparticles such as

silica and titania [18], carbon [19-21] and others are used as nanofillers.

Compared to conventional micro and macro composites the enormous surface

to volume ratio of the nanoparticles is the most important factor. The improved

properties of nanocomposites are related to the modification of the structure

and dynamics of the polymer at and near the particle surface. Because of the

large surface area this fraction of the polymer contributes significantly to the

properties of the whole nanocomposite, even at low filler content. In this

respect polymer nanocomposites are somehow similar to semicrystalline

polymers where the crystals can be considered as nanofillers too.

The glass transition, calorimetrically measured as well as the dynamic

glass transition studied by different probes (α-relaxation in amorphous

polymers), is often used to detect changes in molecular dynamics in polymers.

However, experimental results on polymer dynamics and the glass transition

in polymer nanocomposites are not conclusive concerning the mechanism and

the details of the modification near the particle surface. The glass transition

temperature of the nanocomposite was found to increase [22-28], to decrease

[27, 29-31], not to be influenced at all [22, 25, 27, 29, 32, 33] or the glass

transition disappeared totally [27, 34-36]. However, there are many

experimental results suggesting that the restriction of chain mobility caused by

the nanoparticles does not extend throughout the material but affects only the

chains within a few nanometers of the filler surface. The existence of such an

interfacial layer was shown for several filler polymer combinations by different

techniques [29, 32, 37-43]. In some cases the interfacial layer was identified

as totally immobilized [32, 35, 38] while in others a second glass transition [44,

45] was observed at higher temperature or at least a shoulder at the high

temperature flank of the relaxation peak [46]. The second peak observed in

6 Chapter 1

the mechanical tanδ curves by Eisenberg et al. [44, 45] was alternatively, as

an example to highlight the problem, interpreted as an indication for the

formation of a macroscopic gel in the studied nanocomposites [47] and not as

the glass transition of the interfacial layer as discussed in [44, 45]. Obviously a

peak in dynamic loss curves does not necessarily identify a glass transition.

Additional criteria must be fulfilled. A straight forward proof of a glass

transition is the observation of the typical step in heat capacity. This step like

change in heat capacity does not occur for local or normal mode relaxation

processes because of the missing contribution from entropy fluctuation [48-

50]. How important the length scale probed by the dynamic experiment for the

identification of a RAF is was demonstrated for semicrystalline poly(ethylene

terephthalate) (PET) [51, 52]. For the dynamic glass transition from dielectric,

dynamic mechanical and temperature modulated DSC an immobilized fraction

(rigid amorphous fraction (RAF)) was detected. In contrary, data from the

more local secondary ß-relaxation process were well described by a two

phase model not requiring the introduction of a RAF.

Even calorimetry, mainly Differential Scanning Calorimetry (DSC), is

routinely used to characterize nanocomposites often the glass transition

temperature was reported only. In a few other studies the shape of the glass

transition interval was investigated too [27, 32, 46, 53] or heat capacity was

measured quantitatively [42]. V.P. Privalko recognized very early the

importance of absolute heat capacity measurements for the thermodynamic

characterization of nanocomposites [25, 36]. Following these ideas heat

capacity measurements for poly(methyl methacrylate) (PMMA),

poly(butyl methacrylate) (PBMA) and polystyrene (PS) silicon oxide

nanocomposites of different morphology were performed. To identify an

immobilized interfacial fraction of the polymer we apply a formalism well

established for the determination of a rigid amorphous fraction (RAF) in

semicrystalline polymers as described by Wunderlich et al. [54, 55], was

applied.

For semicrystalline polymers there is an ongoing debate at what

temperature the immobilized fraction (RAF) devitrifies (relaxes), see e.g. [56-

58]. The question if the polymer crystals are melting first and simultaneously

Introduction 7

the RAF devitrifies or the RAF devitrifies first and later on the crystals melt can

not be answered easily on the example of semicrystalline polymers. This is

because the crystals, which are the reason for the immobilization of the

polymer, often disappear (melt) in the same temperature range as the RAF.

For polymer nanocomposites the situation is simpler. Silica nanoparticles do

not melt or undergo other phase transitions altering the polymer-nanoparticle

interaction in the temperature range where the polymer is thermally stable

(does not degrade). Therefore polymer silica nanocomposites are well suited

for a detailed study of the glass transition of an immobilized layer at the

interface between the polymer and the nanoparticle. Several authors claim to

observe such a second glass transition, see e.g. [44-46]. In all these cases the

second glass transition is detected as a separate peak or a shoulder of the α-

relaxation peak from dynamic measurements. But to the best of my

knowledge there is no evidence for a second glass transition in polymer

nanocomposites from calorimetric studies so far. It was therefore of interest to

obtain polymer nanocomposites with a significant amount of the immobilized

fraction and to measure heat capacity in order to detect a possible second

glass transition as an increase of heat capacity towards liquid heat capacity at

temperatures above the glass transition of the mobile polymer.

8 Chapter 1

2. LITERATURE REVIEW

2.1. Polymer nanocomposites

Filling polymers with inorganic particles is used to improve the stiffness

of the materials, to reinforce thermal and mechanical properties as well as the

chemical stability, to enhance the resistance to fire, decrease the gas

permeability etc. Due to the large surface area of the nanosized particles, its

dispersion in the polymers provides new properties or significantly improves

them in comparison to those of the pure polymer. The inorganic nanoparticles

uniform distribution in the polymer matrix generates a new class of materials

called polymer nanocomposites. The term hybrid composite or material is

commonly used as a synonym of organic inorganic nanocomposite. The

preparation of such materials dates back to 1990s when the first clay polymer

nanocomposite synthesis has been reported [59]. Kojima et al. found that

montmorillonite cation exchanged for 12-aminolauric acid was swollen by

epsilon-caprolactam to form a new intercalated compound. Caprolactam was

polymerized in the interlayer of montmorillonite, yielding a nylon 6-clay hybrid

(NCH). NCH is a nanocomposite of Nylon 6 and uniformly dispersed silicate

monolayers of montmorillonite. There are also many other nanoparticles used

to produce polymer nanocomposites depending on the properties which

should be improved [60-64]. Firstly the composites with layered silicates

(clays) are discussed, see [65] for a recent review.

2.1.1. Nanoparticles

The commonly used clays for formation of nanocomposites consist of

nanoplates of silicates. The thickness of the layers is usually in a range of

several nanometers and the length can be up to 1 μm or even more.

Formation of polymer-clay nanocomposites depends on the type of dispersion

of the silicate layers within the polymer matrix. There are three particular

cases of clay distribution: agglomerated stacks of the layers within the

polymer matrix, intercalated and exfoliated structures. Intercalated

nanocomposites are formed when the polymer chains penetrate between clay

plates or are polymerized there. However the lamellar structures of the clay

10 Chapter 2

still remain unbroken. When the plates are completely separated and have

random orientation in the polymer matrix, exfoliated nanocomposites are

obtained. The exfoliated structure is of particular interest because it increases

the polymer-clay interactions. The specific surface area of exfoliated clays is

usually of about 700 m2/g compared to 2 m2/g for the not exfoliated structure

[66]. Fig. (2.1) illustrates schematically the situation for these cases of polymer

clay nanocomposites: agglomerated, intercalated, partly intercalated and

exfoliated, and fully exfoliated. The agglomerated system is just a stack of the

silicate layers without polymer in between the layers. Fig. (2.1b) corresponds

to the polymer intercalated into the interlayer space situation.

(a) (b)

(c) (d)

Figure 2.1. The schematic of (a) – agglomerated, (b) – intercalated, (c) –

partly intercalated and exfoliated and (d) – exfoliated polymer

clay nanocomposites. The heavy straight lines are silicate layers,

the random thin lines are the polymer chains [65].

The fully exfoliated clay nanocomposite is shown in Fig. (2.1d) when all

clay layers are deagglomerated and dispersed independently on each other in

Literature review 11

the polymer matrix. The situation when both cases are present is illustrated in

Fig. (2.1c) which is also called intercalated-flocculated.

Clay based nanocomposites are currently synthesized and studied

most frequently. But nanoparticles of different shapes than layered silicates

are used to produce polymer nanocomposites as well. Nanoparticles of

spherical shape have attracted a great attention nowadays due to

enhancement of polymer properties. How complex the situation is can be

explained on the example of barrier properties. For layered silicate

nanocomposites barrier properties are normally improved. But in special

cases the addition of 10–30 wt% of nanosized fumed spherical silica to a

number of high-permeability polymers increases small penetrants permeation

by up to an order of magnitude [65, 67-71]. Normally, the addition of low-

permeability fillers (such as silica) reduces penetrant diffusion simply by

volume fraction effects. It is believed that the anomalous behavior observed

for nanosized particles is associated with the greater specific interfacial area

for the same level of loading compared to conventional (i.e., micron-sized or

larger) filler particles.

Another example is the production of composite biomaterials such as

bioresorbable polymers filled with spherical calcium phosphate nanoparticles

[72-78]. Calcium phosphate nanospheres mixed with poly(d,l-lactide-

coglycolide) intensifies the activity of alkaline phosphatase, which is important

for the differentiation of osteoblasts that dictate the regeneration process

within the organism. The most used calcium phosphate in implant materials is

hydroxyapatite, Ca10(PO4)6(OH)2, since it is the most similar material to the

mineral component of bones. Here nanocomposites with biocompatible

polymers are of special interest. These nanocomposites exhibit good

properties of biomaterials, such as biocompatibility, bioactivity,

osteoconductivity, direct bonding to bone, etc. [79, 80].

Nanotubes of different elements, its oxides etc. are widely used to

enhance the properties of polymers too. A nanotube is a nanometer scale

wire-like structure that is most often composed of carbon. A single walled

carbon nanotube is a one-atom thick sheet of graphite rolled up into a

seamless cylinder with diameter of the order of a nanometer. This results in an

12 Chapter 2

essentially one-dimensional nanostructure where the length-to-diameter ratio

exceeds 10 000. Such cylindrical carbon molecules have novel properties that

make them potentially useful in a wide variety of applications in

nanotechnology, electronics, optics and other fields of materials science. They

exhibit extraordinary strength and unique electrical properties, and are

efficient conductors of heat. Inorganic nanotubes have also been synthesized.

Inorganic nanotube is a cylindrical molecule often composed of metal oxides,

and morphologically similar to the carbon nanotube. They are observed to be

contained naturally in some mineral deposits too [21, 81-83]. In recent years,

nanotubes have been synthesised of many inorganic materials, such as

vanadium oxide and manganese oxide, and are being used for such

applications as redox catalysts and cathode materials for batteries. Inorganic

nanotubes are heavier than carbon nanotubes and not as strong under tensile

stress, but they are particularly strong under compression, leading to potential

applications in impact resistant applications such as bullet proof vests. The

name “nanotube” is derived from their size, since the diameter of a nanotube

differs from its length by six orders of magnitude. There are two main types of

nanotubes: single-walled nanotubes (SWNTs) and multi-walled nanotubes

(MWNTs). The specific surface area of the nanotubes and other particles of

nanosize is in the range of several hundreds m2/g similar to that of clays.

Due to the large surface area the nanoparticles tend to form

agglomerates of much larger size which suppresses its ability to enhance the

properties of the polymers while producing the composites. Therefore one of

the most important prerequisites of fabricating the polymer nanocomposites is

the deagglomeration of the nanoparticles. And the deagglomeration becomes

more difficult with increasing specific surface area. This is the main

disadvantage of using nanotubes to obtain composite materials because of its

larger surface area in comparison to nanosized clays and spheres.

2.1.2. Preparation methods of polymer nanocomposites

Several possibilities of deagglomeration are known which differ from

one preparation method to another. The preparation methods could be divided

into three widely used groups: (i) dispersion of nanoparticles into the polymer


matrix, (ii) synthesis of the polymer in presence of nanofiller and (iii) the

synthesis of nanoparticles in presence of polymer.

To the first group belongs, for instance, the melt blending (or melt

compounding) of the polymer nanocomposites. It is processed typically in one

or two screw extruders at temperatures of the liquid state of polymer [23, 84-

87]. By control of mixing conditions the uniform distribution of the

nanoparticles in the polymer matrix can be achieved. The high shear forces

generated between extruder walls and screws make it sometimes possible to

obtain almost or full deagglomeration of the nanofiller added to the polymer.

This method is of technical importance because of the relatively simple

upgrade from laboratory to industrial scales. The disadvantage is that a

number of polymer chains may degrade due to the generated shear force or

high temperature. Control of molecular weight is therefore required. The

method is not technically available in our laboratory but the polymer

nanocomposites prepared by melt blending have been kindly provided by

colleagues at the Department of Polymer Structures, Leibniz Institute of

Polymer Research, Dresden.

Another way of producing polymer nanocomposites is the solution

preparation method (solution mixing) [88-91]. If the polymer can be solved and

the nanoparticles dispersed in a solvent, mixing brings reasonably good level

of dispersion of the nanoparticles in the polymer. In some cases even the

agglomeration tendency of the nanoparticles can be overcome. It is applicable

to polymers that can be dissolved or swelled by the solvent [92, 93]. The

deagglomeration of the nanoparticles may be reached by the mixing regimes

(intense stirring) as well as sonification of the solution.

Ultrasound is often used for the synthesis due to its influence on the

reaction [94-99]. The chemical effects of ultrasound derive primarily from

acoustic cavitation [100]. Bubble collapse in liquids results in an enormous

concentration of energy from the conversion of the kinetic energy of the liquid

motion into heating of the contents of the bubble. The high local temperatures

and pressures, combined with extraordinarily rapid cooling, provide a unique

means for dispersing nanoparticles or driving chemical reactions under

extreme conditions. A diverse set of applications of ultrasound to enhance

14 Chapter 2

chemical reactivity has been explored with important uses in synthetic

materials chemistry. For example, the sonochemical decomposition of volatile

organometallic precursors in low-volatility solvents produces nanostructured

materials in various forms with high catalytic activities. Nanostructured metals,

alloys, oxides, carbides and sulfides, nanometer colloids, and nanostructured

supported catalysts can all be prepared by this general route. But the main

advantage of this method in the context of the present work is the huge kinetic

energy which eventually is transferred to the nanoparticles resulting in

deagglomeration. Therefore the sonification is used to obtain polymer

nanocomposites during this work.

To the second group of preparation methods belong chemical

processes, in which polymerization is performed directly in the presence of the

inorganic particles. Examples of emulsion [35, 99, 101], miniemulsion [98,

102, 103], microemulsion [104, 105], suspension or dispersion [24, 106, 107]

polymerization, as well as differently performed free radical polymerization

[30, 108, 109] and ionic polymerization [110] etc. can be found in the literature

but emulsion polymerization is by far the technique most frequently used.

Heterogeneous polymerization, especially emulsion polymerization,

provides an effective way of synthesizing nanocomposites with various

architectures and forms [111]. Seeded (in the presence of nanoparticles)

emulsion polymerization technology is commonly used in the production of

nanocomposite emulsions. The seeded emulsion polymerization occurs

beforehand in the presence of water, emulsifier (surfactant), water soluble

initiator and a small amount of monomer, in which the emulsion has a large

number of particles with very small size. Then the polymerization reaction

continues in emulsion with the presence of the seeds (nanoparticles). This

method can control the reaction rate, particle size and morphology effectively

[112]. A number of papers deal with the encapsulation of sol–gel type metal

oxide particles (SiO2, TiO2) and other inorganic pigments [113] to give

organic–inorganic hybrid dispersions, where the polymer shell is built in situ by

means of conventional emulsion [35, 114], miniemulsion [115, 116] and

related dispersed-phase (Huang and Brittain 2001; Hwu, Ko et al. 2004)

polymerization processes. When the hydrophobic coat layer is simply


adsorbed on the hydrophilic inorganic particle surface the poor chemical

interaction between the three phases (inorganic particle–hydrophobic

surfactant–organic polymer) can result in dewetting of the cover-forming

organic polymer. While the encapsulant in these organic–inorganic hybrids is

an organic polymer, the core inorganic particle usually presents a more

hydrophilic surface. Therefore, either the adsorption of the polymerization

initiator onto the particle surface through electrostatic interaction [117] or

special polymerization techniques are generally required when polymerization

onto an unmodified (i.e. not by surface-graft reaction hydrophobically

modified) inorganic particle is carried out by conventional emulsion process to

decrease the tendency to agglomerate.

Another way to keep the emulsion stable without using hydrophobic

modification or emulsifier is non-surfactant emulsion polymerization under

sonification. The idea is that the emulsion is formed by applying the

sonification and kept stable until the polymerization finished. Sonification of

the reaction media gives also the other advantage: inorganic nanoparticles are

kept deagglomerated up to the encapsulation by the polymer.

Frontal polymerization in the presence of nanofillers also belongs to the

second group of the preparation methods [118-120]. Frontal polymerization is

a process in which the polymerization propagates through the reaction vessel.

This approach allows producing the polymer nanocomposites by the

deagglomeration of the nanoparticles by an emulsifier in the reaction media

and then followed by radical polymerization in frontal regime. Frontal

polymerization is carried out usually in tubular reactors. Thermal frontal

polymerization begins when a heat source contacts a side of the tube and the

heat released by the exothermal polymerization initiates continuously next

portions of tube-like reaction media. This method allows the synthesis of

polymer nanocomposite materials which may have varying properties on the

product length scale of product.

The synthesis of nanoparticles in presence of monomers, oligomers or

polymers can be performed by different techniques. For instance, inorganic

(CdS, Ag)-polyacrylamide (PAM) nanocomposites can be prepared

successfully using a convenient ultraviolet irradiation technique; the initiation

16 Chapter 2

of reaction and polymerization are carried out by means of ultraviolet

irradiation [121]. It was found that the inorganic nanoparticles could be well

homogeneously dispersed in the polymer matrix because polymerization of

organic monomer and formation of inorganic nanoparticles were

simultaneous. It is very interesting that the presence of inorganic ions may be

favorable for the polymerization of the organic monomer. At the same time the

organic polymer matrices can efficiently prevent the produced inorganic

nanoparticles from agglomeration. The polymer nanocomposites can be

prepared also by synthesis of the nanoparticles in the polymer matrix [122-

124]. For instance, well-dispersed titanium dioxide (TiO2) nanoparticles were

synthesized utilizing a block copolymer as a template [125]. The nanoparticles

were confined within microphase separated domains of sulfonated styrene-b-

(ethylene-ran-butylene)-b-styrene (S-SEBS) block copolymers. Another

possibility is described in [126]. The formation of nanosized lanthanum

hydroxide particles in aqueous medium was carried out in the presence of

double-hydrophilic block copolymers. These copolymers contain a polyacrylic

acid block as an ionizable block, and a polyacrylamide (PAM) or a

polyhydroxyethylacrylate (PHEA) block as a neutral block. The nanoparticles

were synthesized by a two-step procedure. Firstly, the complexation of

lanthanum ions in water by the polyacrylate blocks induced the formation of

star-shaped micelles stabilized by the PAM or PHEA blocks. Secondly, the

inorganic polycondensation of lanthanum ions led to the formation of organic-

inorganic nanohybrids. Also here the organic polymer matrices efficiently

prevent the formed inorganic nanoparticles from agglomeration.

Consequently the polymer nanocomposites can be obtained by various

preparation methods. The properties of the composites depend on different

factors, such as nanoparticle size, the polymer type and most important on the

interaction between the polymer and the nanoparticles. One of the important

factors influencing on the composite properties is the degree of agglomeration

of nanoparticles in the polymer matrix. Therefore characterization of the

morphology of the nanocomposites obtained is an important task. Morphology

of the nanocomposites could be investigated by X-ray scattering or electron

microscopy and many other techniques.


2.1.3. Morphology

Generally, the structure of polymer clay nanocomposites has typically

been studied using wide angle X-ray diffraction (WAXD) analysis and

transmission electron microscopic (TEM) observations. Due to its easiness

regarding sample preparation and availability WAXD is one of the commonly

used methods to probe nanocomposite structure [34, 127]. By monitoring the

position, shape, and intensity of the basal reflections from the distributed

silicate layers, the nanocomposite structure (intercalated or exfoliated) may be

identified. For example, in an exfoliated nanocomposite, the extensive layer

separation associated with the delamination of the original silicate layers in the

polymer matrix results eventually in the disappearance of any coherent X-ray

diffraction from the distributed silicate layers. On the other hand, for

intercalated nanocomposites, the finite layer expansion associated with the

polymer intercalation results in the appearance of new basal reflections

corresponding to the larger gallery height. Although WAXD offers a convenient

method to determine the interlayer spacing of the silicate layers in the original

layered silicates and in the intercalated nanocomposites (within 1–4 nm), little

can be said about the spatial distribution of the silicate layers or any structural

non-homogeneities in the nanocomposites.

Additionally, some layered silicates initially do not exhibit well-defined

basal reflections. Thus, peak broadening and intensity decrease are very

difficult to study systematically. Therefore, conclusions concerning the

mechanism of nanocomposites formation and their structure based solely on

WAXD patterns are only tentative. On the other hand, TEM allows a

qualitative understanding of the morphology, spatial distribution of the various

phases, and views of the defect structure through direct visualization.

Moreover TEM is used to characterize not only nanocomposites with layered

silicates but with the nanoparticles of any shape. However, special care must

be exercised to guarantee a representative cross-section of the sample and to

avoid artefacts. The WAXD patterns and corresponding TEM images of three

different types of nanocomposites are presented in Fig. (2.2). Both TEM and

WAXD are essential tools [128] for evaluating nanocomposite structure.

However, TEM is time-intensive, and only gives qualitative information on the

18 Chapter 2

sample as a whole, while low-angle peaks in WAXD allow quantification of

changes in layer spacing.

Figure 2.2. (left) WAXD patterns and (right) TEM images of three different

types of nanocomposites [65]

Typically, when layer spacing exceed 6–7 nm in intercalated

nanocomposites or when the layers become relatively disordered in exfoliated

nanocomposites, associated WAXD features weaken to the point of not being

useful. However, recent simultaneous small angle X-ray scattering (SAXS)

and WAXD studies yielded quantitative characterization of nanostructure and

crystallite structure in Nylon 6 based nanocomposites [129, 130].

As an example of electron microscopic characterization of

polymer/nanospheres composites TEM images of PMMA filled with silica are

represented in Fig. (2.3) [87]. Fig. (2.3a) shows the PMMA nanocomposites


containing organically modified (PMMA grafted onto silica surface) while

Fig. (2.3b) non-modified silica nanoparticles dispersed in a PMMA matrix.

Figure 2.3. TEM images of PMMA/silica-gPMMA nanocomposites; (a) –

deagglomerated and uniformly dispersed, (b) – agglomerated

nanoparticles in polymer matrix (d(silica)=15-20 nm) [87]

It is seen that by grafting the PMMA onto silica surface one gets a

totally deagglomerated system while the unmodified nanoparticles appeared

to agglomerate into bigger particles.

The mentioned above reveals that there are several possibilities to

characterize the polymer nanocomposites morphology. Another important

parameter for polymer nanocomposites is the interfacial interaction between

the inorganic filler and the polymer. Different techniques are available for an

investigation of the properties of the polymer near the interface.

2.1.4. Interfacial interactions

The large specific surface area of the nanoparticles is important for

polymer properties in composites due to the interfacial interaction between

nanofiller and polymer matrix. Therefore it is of interest to investigate the

nature of such interactions. Several possibilities for the investigation of the

interface are known such as solid state nuclear-magnetic resonance (NMR)

[131-134], dynamic mechanical analysis (DMA) [37, 44, 133, 135-140],

nanoindentation [37, 141-143], infrared spectroscopy (IR) [35, 144], positron

annihilation spectroscopy [70, 145, 146], calorimetry (mainly differential

scanning calorimetry, DSC) [85, 87, 147, 148] and others.

20 Chapter 2

To investigate the interaction between nanofiller and polymer NMR

analysis is often used [131-133, 149-151]. In hybrid nylon 6/silica composites

nuclear relaxation measurements using low field NMR were applied [131]. The

NMR results showed that with up to 20% of silica there is compatibility due to

a weak intermolecular interaction. This may be concluded from the values of

spin-lattice relaxation time, which are in between the values of each initial

composite’s component. The data from NMR measurements also show that in

the Nylon 6/Si composites there is some interaction, as only two values of this

parameter were found for each Nylon 6/Si composition, and they were

different from pure nylon 6 and silica. The interphase behaviour of the cured

poly(amic acid) with polyhedral oligomeric silsesquioxane (POSS epoxide),

namely the octa(ethylcyclohexylepoxidedimethylsiloxy) silsesquioxane was

investigated by solid state NMR [133]. The results show that properties of the

interphase are varied systemically by adjusting the nanotether structure of the

epoxide molecules. Solid-state NMR data can be used also as a powerful tool

to characterize not only surfactant loading of the clays in polymer/layered

silicate composites but to obtain further insight into temperature-dependent

surfactant dynamics and structure of the surfactant layer too [132]. The NMR

measurements aimed at surfactant headgroups and tail ends supply

complementary information on the structure of that layer. It was found that two

microphases with different mobility and probably with different strength of the

attachment of surfactant headgroups to the silicate surface coexist over a

broad range of surfactant loadings and temperatures. By NMR analysis it was

possible to show that an excess of surfactant with respect to the cation

exchange capacity of the silicate causes plasticization of the surfactant layer

in pure organoclays and diminishes the tendency for intercalation.

Consequently NMR appears to be a very consistent tool to investigate the

interface interaction between polymer and nanofiller in composite materials.

But there is also a disadvantage of NMR use because this method

investigates mainly a behavior on a very local length scale and not necessarily

on the length scale representative for the glass transition.

Dynamic mechanical analysis is generally used to detect the property

enhancement of the polymer nanocomposites [139, 140] but also information


about the influence of nanofillers on glass transition can be obtained.

Eisenberg et al. have investigated the interface for the different

polymer/inorganic filler nanocomposites using DMA [45]. The results showed

that interactions of polymer chains with silica nanoparticles restrict the mobility

of the chains at the interphase. Two peaks in the tanδ curves, one

corresponding to the common glass transition and a second peak at higher

temperature (Fig. 2.4a) were observed [45]. Fagiadakis et al. [46] found a

shoulder at the high temperature flank of the relaxation peak. Both findings

were interpreted as the glass transition of the immobilized polymer close to

the interphase.

(a) (b)

Figure 2.4. tan δ versus T - Tg curves for (a) - poly(styrene-co-4.5 mol %

sodium methacrylate) ionomer and the following polymers filled

with 10 wt% of 7 nm silica particles: PS, PMMA,

poly(4-vinylpyridine) and poly(vinylacetate) (the curves of the

filled polymers have been shifted for clarity by 0.2, 0.9, 1.4, and

1.9, respectively) [45] and (b) – styrene-4-vinylpyridine (S4VP)

copolymers of different vinylpyridine (VP) contents, but

containing 20 wt% silica [44].

The second peak observed in the mechanical tanδ curves by Eisenberg

et al. [44, 45] was alternatively, as an example to highlight the problem,

22 Chapter 2

interpreted as an indication for the formation of a macroscopic gel in the

studied nanocomposites [47] and not as the glass transition of the interfacial

layer as in [44, 45]. Obviously a peak in a dynamic loss curve does not

necessarily identify the occurrence of a glass transition, additional criteria

must be fulfilled. The first peak in Fig. (2.4) for all systems is assigned as a

conventional glass transition because it is present also in pure polymers. It is

important to mention that the second peak in the tan δ curve for PS/10 wt%

silica nanocomposites in Fig. (2.4a) does not appear for the similar system but

filled with 20 wt% filler (the curve which corresponds to 0 wt% VP in

Fig. (2.4b)). Anyway in Fig. (2.4b) the second peak in the dynamic loss curve

is present only for the S4VP copolymers. This means that this peak appears

most likely due to the interaction between nanoparticles and VP, but not PS.

However the second peak in Fig. (2.4a) cannot be explained by this

assumption because the polymer used for the PS/SiO2 nanocomposites

preparation was pure PS which seems to exhibit no interaction with silica

nanoparticles. Anyway, the situation here could be clarified having the results

of the measurement at the exactly same conditions for silica used. The silica

nanoparticles might have an organic cover on its surface which may exhibit

the peak in the tanδ curve. Consequently it is difficult to draw final conclusions

about the mobility of the interfacial polymer layer for the pure polystyrene

silica nanocomposites by DMA. This problem will be discussed in Chapter (4).

The study of organic phase mobility in organic-inorganic coatings by

DMA is reported in [139]. The analysis of the hybrids coated on a PET film

(coating thickness 10 μm and 40 μm) shows an additional up-shift of glass

transition temperature, more markedly in the case of the thinner hybrid

coating. This result is attributed to molecular interactions at the

substrate-coating interface that locally hinder molecular mobility. The

consequent increase of glass transition temperature is more evident when the

coating layer is thin.

The chain mobility in the polymer-clay nanocomposites is greatly

reduced as studied by dynamic mechanical analysis (DMA) and dielectric

analysis (DEA) in [140]. The modulus of the composite increases significantly.

The modulus enhancement strongly relates to the volume of the added clay as


well as the volume of the constrained polymer. This modulus enhancement

follows a power law with the content of the clay in the composite. This

study [140] also indicates that the structure of clay nanocomposites with

strong interfacial interactions is analogous to that of semicrystalline polymers.

In the case of polymer-clay nanocomposites, the intercalated clay phase

serves as an unmeltable crystalline phase which results in improvement in

mechanical and thermal properties. The same consideration can be applied

for the other polymer/inorganic nanofiller systems.

The nanoindentation measurements could be revealed as an

appropriate technique to characterize hybrid organic–inorganic thin films [37,

142] which obviously exhibits very similar mechanical and thermal properties

to those of interfacial polymer layer in composite materials. The indentation

study shows that the extent of the hybrid interface could be adjusted by the

use of preformed silica nanoparticles. It was also shown that the mechanical

response was governed by the size of the hybrid interface since the

mechanical properties of materials based on sol–gel silica are more elevated

than those obtained from materials formed from silica nanoparticles which

exhibit a more defined interface. Therefore to allow the precise quantification

of the nanofiller surface area in polymer nanocomposites spherical silica

nanoparticles has been used for the present work.

2.1.5. Calorimetry

Privalko et al. have consistently investigated the interfacial organic

layer mobility for different hybrid materials using calorimetric methods [22, 25,

152-155]. The heat capacity of polyurethane filled with finely dispersed Aerosil

(specific surface area is 175 m2/g) was studied as a function of temperature in

[156]. The addition of filler was found to decrease the crystallinity of the filled

polyurethane and significantly reduce enthalpy of the polymer. This is

explained by the appearance of macromolecules with reduced mobility in the

amorphous zones at the polymer-particle interface. The calorimetric study of

oligo-ethylene glycol adipate (OEGA) filled with Aerosil and colloidal graphite

(calculated specific surface area is 0.67 m2/g) showed an interesting result

(Fig. (2.5)) [36]. In Fig. (2.5) the specific heat capacity as a function of

temperature for quenched samples of differently filled OEGA is presented.

24 Chapter 2

Figure 2.5. The heat capacity of quenched samples of filled OEGA. Filler

contents (wt %): 1 – 0, 2 (filled triangle) – 1, 3 – 50 graphite, 4 –

10 Aerosil [36].

It is seen that at 50 wt% graphite loading the glass transition of the

OEGA disappears. The same can be observed for Aerosil filled sample but

already at much lower filler content. The reason for that could be the

difference in specific surface area of the inorganic fillers. Consequently,

Fig. (2.5) [36] clearly demonstrates that the interfacial organic layer shows

much lower mobility in comparison to that of the pure (unfilled) substance,

which at relatively high filler concentrations can result in disappearance of the

step in heat capacity at glass transition.

Later authors confirmed [25] that the properties of filled polymer

systems are determined by the amount of the interfacial layer. The heat

capacity data from the calorimetric measurements indicated that an increase

in Aerosil content results in a more or less reduction of calorimetric relaxation

strength at the glass transition temperature. It has to be mentioned that in [25]

only the polymer was varied in polymer/filler hybrids to get comparable data

for each system. A calorimetric study of PMMA and PS filled with powdered

glass [22] confirmed the tendency of the mobility decrease of the interfacial

organic layer for these systems with increasing filler content. The glass

transition temperature was found to increase with increasing powder content.

It was also shown that the absolute values of the specific heat capacity of


composites are smaller than those of unfilled polymers not only in solid but

also in liquid states. The difference in solid state is thought to disappear in a

high enough temperature range owing an intensification of the

macromolecular thermal vibrations but however it was not clearly

detected [22]. These works, for the first time, show the possibility to

investigate the interface in polymer nanocomposites calorimetrically.

Giannelis et al. have investigated polymer silicate composites by

different methods as well as DSC [34]. It was shown that on a local scale,

intercalated polymers exhibit relaxation for a wide range of temperatures, with

a significant suppression (or even absence) of cooperative dynamics typically

associated with the glass transition. The glass transition temperature of

polystyrene filled with organically modified silicate (C18FH) was reported also

to disappear. Namely, the absence of any thermal transition for the

intercalated polymer in the conventional glass transition temperature region

was observed similar to [36]. For the polyimides with silicate obtained by

sol-gel method no glass transition was detected in data from DSC

measurements at 50 wt% nanofiller loading [27]. One has to mention that the

absence of the step in glass transition temperature range was observed only

for the polymer with relatively low molecular weight (5000 g/mol) while for the

systems with higher molecular weight the calorimetric relaxation strength was

only lowered. Taking these results into consideration, the polymers to be

chosen for the investigation in this work were synthesized with a wide

molecular weight distribution. The low molecular weight fraction is expected to

interact with nanoparticles easier than that of high molecular weight.

On the contrary to the mentioned above, the glass transition was

reported also not to be influenced et al. [22, 25, 27, 29, 32], to be shifted to

higher [22-28] or lower [29, 30] temperatures as well.

The interfacial interaction can be varied by the preparation method of

the polymer nanocomposites, by change of nanoparticles type, dimensions,

surface modification and polymer type as reported in literature. Reviewing the

investigations of the interfacial layer in polymer nanocomposites one may

conclude that this is still an open question. However it is generally reported

that in polymer nanocomposites the polymer layer on the nanoparticle surface

26 Chapter 2

is thought to be immobilized, e.g. the chain mobility is reduced, in spite of rare

situations as in [31, 138, 149] for instance or in the others given in the

introduction. But to the best of my knowledge, there is no evidence of DSC

confirmation of the interfacial immobilization of the polymer by nanofiller where

the nanoparticles surface is not organically treated.

2.2. Semicrystalline polymers

A similar situation, a polymer interacting with rigid particles, is present

in semicrystalline polymers. Semicrystalline polymers consist of crystallites

(lamellae) and an amorphous fraction which thickness is in the range of

ca. 10 nm. The polymer nanocomposites are usually filled with particles of

similar size. Therefore the interface between amorphous and crystalline

fractions in the semicrystalline polymers can be treated in the same way as for

polymer nanocomposites if an immobilized layer exists.

The quantification of the immobilized amorphous polymer by the

crystals, i.e. a rigid amorphous fraction (RAF), was introduced for

semicrystalline polymers [51, 157, 158], see Wunderlich for a recent

review [55]. Similar procedure may be performed for the polymer

nanocomposites as well. Consequently the amount of immobilized layer may

be available from the calorimetric measurements as described in [157]. The

understanding of its formation and devitrification both in semicrystalline

polymers and polymer nanocomposites can help to obtain materials with

controlled properties.

2.2.1. RAF in semicrystalline polymers

Semicrystalline polymers have frequently a negative contribution to the

heat capacity between glass transition and melting, linked to the RAF [55].

Because of the need to accommodate flexible polymer molecules of typically

1–100 μm length into micro- and nanophases, there is usually a strong

coupling between crystal and amorphous phases due to the frequent crossing

of the interface by the long molecules. In all polymers, this strong coupling

between the phases results in a broadening of the glass transition to higher

temperature, as seen for instance for PET [159, 160]. In many polymers this

coupling causes a separate glass transition for the RAF, as summarized


in [55]. An effect due to the RAF was first reported for several semicrystalline

polymers as a deficit in calorimetric relaxation strength (Δcp) at glass

transition [157, 161].

The heat capacity of the semicrystalline poly(oxymethylene)s between

the glass transition and the melting temperature, as shown in [157], indicated

much lower levels than expected from a two-phase crystallinity model as

shown in Fig. (2.6). The dashed line in Fig. (2.6) corresponds to liquid

poly(oxymethylene) heat capacity and the dotted line is a guess at the low

temperature continuation. The heavy line is the experimental data presented.

The dash-dotted lines correspond to the calculated data for the indicated

percentages of “rigid” phase.

Figure 2.6. Heat capacity of poly(oxymethylene) when fitted assuming 56%

crystallinity, 24% rigid amorphous, and 20% mobile amorphous

poly(oxymethylene). The dashed line correspond to liquid

poly(oxymethylene) heat capacity, the dotted line is a guess at

the low temperature continuation. The heavy line is the

experimental data. The dash-dotted lines represent calculated

data for the indicated percentages of “rigid” phase (crystalline

and rigid amorphous). Only the 80% curve fits the data [157]

28 Chapter 2

The authors showed that the straight lines of Fig. (2.6) are tangents to

the 100% crystalline samples [157]. A larger rigid fraction (0.8) than calculated

from crystallinity (0.56 for Fig. (2.6)) according to the two-phase model was

needed to match experimental data and calculation. The experimental data lie

significantly lower than the calculated data from the two-phase model. This

means that there is a part in rigid fraction which does not contribute to the step

in heat capacity at glass transition. And the same situation was found for the

other poly(oxymethylene)s investigated [157].

In addition, in [157] the data for chemically different samples felt on

slightly different curves. The only interpretation of those results could be that

the crystallinity model is not suitable for the description of heat capacities of

poly(oxymethylene) in this temperature range. To derive a possible structure

parameter for heat capacity the authors assumed that, based on the normal

beginning of the glass transition, a portion of the non-crystalline fraction is

gaining normal mobility at the glass transition. This part of the non-crystalline

fraction was called “mobile amorphous” and treated similar to the super cooled

liquid, with a heat capacity identical to the data extrapolated from the melt.

Figure 2.7. Subdivision of the heat capacity cp of semicrystalline

poly(oxymethylene) into “rigid amorphous” and “mobile

amorphous” at 265 K [157]

The remaining non-crystalline fraction of the sample, which was called

“rigid amorphous”, was assumed to depend on sample structure, and possibly


also on crystallization condition, see Fig. (2.7). One has also to mention that

the curves calculated using the crystallinity from the heat of fusion as a

calculated parameter (0.56) are far out of any reasonable experimental

uncertainty. The negative and positive heat capacity deviations for 38

semicrystalline poly(oxymethylene)s and poly(oxyethylene)s in the

temperature range between glass and melting transition have been clearly

delineated in [157]. The negative deviation was linked to an added fraction of

RAF, while the positive deviation was assigned to processes such as defect

formation or beginning of melting, i. e. gaining of mobility and possibly

disordering. The RAF in poly(oxymethylene) was found to be constant up to

the melting region, in contrast to polypropylene, where it is decreasing with

increasing temperature [157].

The concept of a rigid amorphous fraction can also be applied for other

relaxation strength measurements than heat capacity. Mechanical [162, 163]

and dielectric spectroscopy result in nearly the same RAF as determined from

heat capacity increments [52, 164]. From the dielectric data not only the

relaxation strength at the dynamic glass transition but also the relaxation

strength for the secondary (more local) relaxations is available. Dobbertin et

al. [52] report about calorimetric and dielectric measurements on the same

semicrystalline PET samples. The question arises if the β-relaxation

(connected with local movements) is similarly influenced by the crystals than

the dynamic glass transition? Such local movements are not possible in the

crystalline part but could be expected to occur in the whole non-crystalline

part. Fig. (2.8) compares the normalized dielectric intensities for the α- and

β-relaxation and the α-relaxation strength from calorimetric measurements.

This confirms the introduction of RAF in semicrystalline polymers, i.e. that the

deviations from the two-phase model for the α-relaxation are present in the

dielectric data too. In [52] the authors found that the secondary β-relaxation

follows the two–phase model as shown in Fig. (2.8). This means that a local

movement is possible in the RAF but not a cooperative segmental motion

(α-relaxation, glass transition).

30 Chapter 2

Figure 2.8. Normalized intensity for α (ε, cp) and β (ε) for differently

crystallized PET samples [52]

Obviously the length scale probed by the different measurements is

different and yields different outcomes regarding the existence of a RAF. The

results of further investigations from the dielectric relaxation and calorimetry

allowed authors to conclude that both independent measurements yield a

correlation length of some nm for the undisturbed glass transition. This allows

concluding that the RAF layer thickness should be most likely in the same

range in the semicrystalline polymers and possibly also polymer

nanocomposites.

2.2.2. Vitrification of RAF

The mentioned above reveals that the RAF existence in semicrystalline

polymers is already confirmed by different methods. Of interest is the question

when the RAF is formed. Fig. (2.9) represents the results from

quasiisothermal crystallization measurements of two polymers [165]. As seen

in Fig. (2.9) the measured heat capacity becomes smaller than the baseline

heat capacity according the two-phase model (curve d), indicating the

occurrence of significant RAF during the crystallization process. On the other

hand, the expected heat capacity, taking into account the RAF obtained at the

glass transition (line e) is in perfect agreement with the measured value at the

end of isothermal crystallization. There is no difference in the amount of RAF

at crystallization and the glass transition temperature; also Tg is more than


30 K below the crystallization temperature in the case of polycarbonate (PC).

Therefore, one can conclude that the total RAF of PC and

poly(3-hydroxybutyrate) PHB is vitrified (formed) during the isothermal

crystallization. No additional vitrification occurs on cooling from the

crystallization to the glass transition temperature.

(a) (b)

Figure 2.9. Time evolution of heat capacity during quasiisothermal

crystallization of (a) - PC at 456.8 K and (b) – of PHB at 296 K,

temperature amplitude 0.5 K and period 100 s, curve a. Curves b

and c correspond to solid and liquid heat capacities from the

ATHAS database [166], respectively. Curve d was estimated

from a two-phase model and curve e from a three-phase model.

The squares represent measurements at modulation periods

ranging (a) – 30 to 12 000 s and (b) – 240 to 1 200 s. Curve f

shows the exothermal effect in the total heat flow rate [165]

Cebe et al. have investigated the formation of the RAF for isotactic

polystyrene (iPS) [167]. The cold crystallization of iPS resulted in the

formation of an RAF, which increases with crystallization time and

temperature in a manner analogous to the development of the crystalline

fraction. Authors conclude that the RAF is formed at nearly the same time as

the crystalline phase and increases more rapidly after spherulite impingement.

Consequently the formation (vitrification) of RAF can be followed by

calorimetric methods. But there are cases when vitrification can not be

determined from heat capacity. For instance, the reversing melting occurring

32 Chapter 2

during the quasiisothermal crystallization of poly(ether ether ketone) (PEEK)

as discussed in [168].

Figure 2.10. Specific heat capacity of PEEK as a function of time from the

data shown in [168]. Curve a - cp value from the measured heat

flow rate, b - expected baseline heat capacity.

In Fig. (2.10) the measured complex heat capacity and expected

baseline heat capacity are shown. In case of PEEK complex heat capacity

increases during crystallization, while baseline heat capacity decreases. If one

wants to study crystallization by TMDSC measurement conditions must be

chosen to fulfill requirements of linearity and stationarity as discussed in [169,

170]. Changes in sample properties (e.g. degree of crystallinity) must be

negligible during one modulation period. But even if these conditions are

fulfilled melting and subsequent crystallization may occur during one period of

the temperature modulation and contribute to the measured heat capacity.

Finally, an excess heat capacity can be seen. Fig. (2.10) shows that the

measured heat capacity behaves different from the expected baseline heat

capacity with increasing crystallinity. The difference can be described as an

excess heat capacity which stays constant after the end of main

crystallization. It can be related to reversing melting during crystallization

[168].

2.2.3. Devitrification of RAF

In spite of the rare situations like described for PEEK the vitrification of

the RAF in semicrystalline polymers can be followed by isothermal


crystallization. The question arises, at which temperature the RAF devitrifies.

Quantitative DSC and TMDSC are the key macroscopic techniques which

allow the characterization of the intermediate phase by evaluation of the glass

transition, the quantitative evaluation of the amount of a RAF, and the

differentiation of various types of RAF via its separate Tg below, at, or above

Tm [55]. Since the main temperature range for characterization lies between Tg

and Tm, a range where the increase in heat capacity due to conformational

motion can compensate the decrease due to the RAF, and where its increase

due to the RAF glass transition may be competing with the beginning of

melting and reorganization of crystals and of reversible melting. Therefore to

detect when the RAF devitrifies calorimetrically is a very difficult task. Special

techniques like quasiisothermal TMDSC and frequency- and amplitude-

dependent measurements need to be tried to avoid the problems mentioned.

The problem of reversible melting is introduced in more detail in [160].

Quasi-isothermal TMDSC in the melting range should, according to [160],

have no contribution from melting and/or crystallization to the reversing heat

capacity. Fig. (2.11) shows, however, that this is not the case. A reversing

contribution to the heat capacity is present and depends on the crystallization

conditions. Although the contribution is much less than that of the total heat of

fusion, the reversing contribution is not negligible. The only interpretation of

this observation is that the polymer molecules that contribute to the reversing

heat capacity are still attached to crystals that melt at a higher temperature

and can serve as molecular nuclei. After the heating cycle a number of melted

polymer molecules, which are still attached to higher melting crystals can

recrystallize during the cooling cycle with negligible supercooling. Overall, this

process yields a reversible, apparent heat capacity contribution similar to

[168].

34 Chapter 2

(a) (b)

Figure 2.11. Reversing heat capacity by quasi-isothermal TMDSC (a) - on

cooling from the melt (filled circles) and (b) - on heating from the

quenched, amorphous sample. The thin lines indicate the

ATHAS database data [166] for the amorphous and crystalline

PET; the broken lines indicate the computed heat capacity for

(a) - 49% and (b) – 40% crystalline PET. The open circles are

melt-crystallized PET on the quasi-isothermal upon step-heating

as reference [160]

The RAF is the part of the non-crystalline PET that does not participate

in the measured Δcp at the glass transition but, on the other hand, does also

not contribute to the heat of fusion [157]. The figure shows that the reversing

heat capacities reach the expected equilibrium heat capacity of the

semicrystalline PET derived from the ATHAS database [166] at about 430 to

450 K. Unfortunately, this temperature is sufficiently close to the beginning of

melting that the actual crossover temperature may be somewhat higher due to

some low temperature reversible melting. And such problems limit the

possibilities of the RAF devitrification detection.

Wunderlich et al. however discussed the RAF determination for the

special case where the limitations mentioned above do not appear as a

disturbing factor. For the understanding of the mechanism of formation and

devitrification of the RAF the quasi-isothermal TMDSC of poly(oxy-2,6-


dimethyl-1,4-phenylene) (PPO) is described in [171]. In Fig. (2.12) the

measured, reversing heat capacity and the crystallinity of PPO are plotted

together.

Figure 2.12. Comparison of the measured heat capacity of semicrystalline

PPO with its change in crystallinity and RAF [171]

As the temperature increases the crystallinity and the RAF decrease,

but at different rates. melting is completed at about 510 K. Up to about 495 K

the crystallinity decreases very little, while the RAF loses almost 20% of its

value, which is in accordance with the assumption that the surrounding glass

must become mobile first, before melting can occur. Between 495 and 510 K

the decrease of both, the RAF and the crystallinity, is close to linear, with the

RAF losing three times as much solid fraction as the crystallinity. In this

temperature range the crystallinity is lost parallel to the loss of the RAF.

Cebe et al. offered a mechanism of RAF devitrification for iPS [172].

Taking into consideration the formation of RAF at crystallization

temperature (Tc), the authors pointed out that RAF is stable at temperatures

below Tc. [56]. Furthermore, heat capacity measurements above the melting

point suggest that only one phase exists at high temperature, i.e. 100% liquid

mobile amorphous fraction (MAF), i.e. not immobilized amorphous polymer.

Therefore, the RAF must be relaxed at some temperature between Tc and the

upper melting point. To provide further evidence for devitrification of RAF,

Fig. (2.13b) shows the temperature dependent heat capacity data in expanded

scaling, for Tc = 155 °C, and for predictions based on the three-phase model

36 Chapter 2

(dark solid curve). Also shown are the predictions based on a two-phase

model (light solid curve). In Fig. (2.13b) at temperatures below the annealing

peak, experimental heat capacity data matches the three-phase model

baseline. At temperature just above the annealing peak, the system

approaches to the two-phase model, in which only crystals and liquid (MAF)

exist. Thus, as temperature increases from below Ta to above Ta, the system

exhibits a transition from three-phase to two-phase. Such a transition turns the

RAF into an identical amount of MAF.

Figure 2.13. Standard DSC scan ((a) – wide scaling, (b) – expanded scaling)

showing specific heat capacity vs. temperature at heating rate of

10 K/min for iPS cold-crystallized at 155 °C for 12 h. The dashed

line is the heat capacity of 100% liquid, while the dotted line is

the heat capacity of 100% solid obtained from the ATHAS

database [166]. In part (a) the solid line and in part (b) the dark

solid line represents the baseline heat capacity based on the

three-phase model, while the light solid line indicates the

baseline heat capacity based on the two-phase model [56, 172,

173]

Using Fourier-Transformation-Infrared spectroscopy (FTIR), wide angle

X-ray scattering (WAXS), and standard DSC scanning, the crystalline fraction

appears to be unaffected by the transition of RAF into MAF, at least within the

error limits of the crystallinity measurement [172]. It is possible that a tiny

amount of crystals, within the error limits, melts at Ta. However, as the authors

demonstrated that it is not possible for the entire endotherm area at Ta to arise

from crystal melting. Therefore, the authors assign the annealing peak in

Fig. (2.13a) as the devitrification of the rigid amorphous fraction, which


transforms RAF into equilibrium liquid without detectable melting of the

crystals. They assume that the relaxation of RAF occurs as a sigmoidal

change in the baseline heat capacity, accompanied by an excess enthalpy.

But the assumption of RAF devitrification by [172] was disproved later

by Minakov et al. using high-rate calorimetry [57].

Figure 2.14. Heat capacity of iPS sample crystallized at 140 °C for 12 h at

heating rate 10 K/min (dashed line) and 30,000 K/min (solid line)

[57]. Expected heat capacities [166] for the liquid, the crystalline

and the semicrystalline iPS according a two- and three-phase

model.

In order to check the hypothesis by Cebe et al. the authors compare in

Fig. (2.14) heat capacities at slow and fast heating, 10 K/min and

30 000 K/min, for the iPS sample crystallized at 140 °C. For both heating rates

above glass transition heat capacity follows the line expected from a three-

phase model taking into account crystalline, mobile amorphous and rigid

amorphous fractions, for details see e.g. [56, 165, 167]. For the low heating

rate after the first endothermic peak heat capacity coincides with that

expected according a two-phase model taking into account crystalline and

mobile amorphous fractions only as already shown by Cebe et al. [56]. If the

first endothermic peak is caused by an enthalpic relaxation of the RAF one

would expect to see a similar effect or at least some step in the heat capacity

38 Chapter 2

curve at temperatures around 160 °C for the fast heating too. But there is

nothing to see at fast heating. Heat capacity reaches the liquid line above the

single melting peak. This indicates that melting of crystals and relaxation of

the RAF occurs in the temperature range of the broad single melting peak,

most probably simultaneously. There is a solid fraction of about 0.55 as for the

slowly heated sample, which is indicated by the three-phase line in Fig. (2.11),

[57]. At fast heating one sees a significant shift of the glass transition to higher

temperatures. The Tg of the mobile amorphous fraction shifts from 100 °C at

10 K/min to about 115 °C at 30 000 K/min. Considering the same apparent

activation energy for the relaxation of the RAF, the beginning of the heat

capacity increase (peak or step) should be shifted to 160 °C. But on fast

heating nothing special happens around 160 °C. It is therefore unlikely that the

annealing peak is related to the nonreversing enthalpic relaxation of the RAF

only. As shown for PC and PHB [165] and for iPS [56] heat capacity changes

from the value expected from a three-phase model to that according a

two-phase model in the temperature range of the low temperature endotherm.

Combining these earlier observations with a continuous melting–

recrystallization–remelting model, which is supported by the results obtained

by Strobl et al. [174, 175] too and the fast heating experiments, one can

discuss the observations as follows. At low heating rates melting of the

crystals starts at the rising flank of the lowest temperature endotherm. Parallel

to crystal melting the RAF surrounding the just molten crystals relaxes. As

shown in [58, 176] the melt is than in a state (conformation) allowing very

rapid (within milliseconds) recrystallization. This recrystallization creates more

stable crystals but does not significantly change overall crystallinity. Assuming

a continuous melting–recrystallization–remelting the remaining amorphous

material in between the crystals may not be vitrified as in the case of slow

isothermal crystallization [56, 165]. If the amorphous material does not vitrify

heat capacity should be the same as expected from a two-phase model as

soon as the continuous melting–recrystallization–remelting starts and that

seems to be what is observed [57].

The mentioned above demonstrates that the question, at which

temperature the RAF devitrifies, is still under discussion. However Wunderlich


et al. discussed the relaxation of RAF in semicrystalline polymer (PPO) using

TMDSC which is a special case when the difficulties like beginning of the

melting, reorganization or reversing melting do not arise. In this work I tried to

find a solution by means of a model system – polymer nanocomposites. It is

hoped that absence of any transition of inorganic fraction in the range from Tg

up to the degradation temperature of the truly amorphous polymer will help to

avoid the difficulties occurring for semicrystalline polymers. For that one has

first to obtain the polymer nanocomposites exhibiting a RAF. Next the RAF

should be quantified in the same way as for semicrystalline polymers from

heat capacity data as shown in [157]. Then the devitrification of the RAF could

be investigated by increasing mobility of the polymer chains of the RAF by

increasing temperature or adding some plasticizer. For the two later points

heat capacity must be determined with adequate precision

2.3. Heat capacity determination

Heat capacity of polymeric materials can be measured by calorimetry.

The applications and interest in calorimetry in material science have grown

enormously during the last half of the 20th century and the beginning of the

21st. Different calorimetric methods are utilized to get information about the

thermal properties of the materials, such as adiabatic [177], AC [178, 179],

DSC [180-182] and TMDSC [183-191]. But the DSC and TMDSC are used in

this work due to the simplicity of use and the uncertainty in measurement

results of ca. 2% or even better [192].

Two basic types of differential scanning calorimeters must be

distinguished:

• Heat flux DSC

• Power compensation DSC.

They differ in the design and measuring principle. Common to all DSCs is a

differential method of measurement which is defined as follows: A method of

measurement in which the measured quantity (measurand) is compared with

a quantity of the same kind, of known value only slightly different from the

value of the measurand, and in which the difference between the two values is

measured [193].

40 Chapter 2

The characteristic feature of all DSC measuring systems is the

twin-type design and the direct in-difference connection of the two measuring

systems which are of the same kind.

The heat flux DSC belongs to the class of heat-exchanging

calorimeters [181]. In heat flux DSCs a defined exchange of the heat to be

measured with the environment takes place via a well-defined heat conduction

path with given thermal resistance. The primary measurement signal is a

temperature difference; it determines the intensity of the exchange and the

resulting heat flow rate (Φ) is proportional to it. In commercial heat flux DSCs,

the heat exchange path is realized in different ways, but always with the

measuring system being sufficiently dominating compared to the heat transfer

inside the sample.

The power compensation DSC belongs to the class of

heat-compensating calorimeters [181]. The heat to be measured is (almost

totally) compensated with electric energy, by increasing or decreasing an

adjustable Joule’s heat.

Figure 2.15. Power compensation DSC (Perkin Elmer Instruments). Set-up of

the measuring system. Sample measuring system with sample

crucible, microfurnace and lid, reference sample system

(analogous to sample), 1 heating wire, 2 resistance

thermometer. Both measuring systems, separated from each

other, are positioned in a surrounding (block) at constant

temperature.

Sample ReferencePlatinum Alloy

PRT Sensor

PlatinumResistance Heater

Heat Sink

SampleSample ReferenceReferencePlatinum Alloy

PRT Sensor

PlatinumResistance Heater

Heat Sink

(1)

(2)


Since the DSC used for this work is a power compensation DSC the

measuring system of it (Perkin Elmer DSC) is described in more details. The

measuring system (Fig. (2.15)) consists of two identical microfurnaces which

are mounted inside a thermostated aluminium block. The furnaces are made

of a platinum-iridium alloy, each of which contains a temperature sensor

(platinum resistance thermometer) and a heating resistor (made of platinum

wire).

There are several variants of measuring possibilities known using DSC.

Two widely used techniques were applied: heating or cooling of the sample

with a linear temperature program (linear scanning) and temperature

modulated DSC. Both of them will be described in this chapter. The output

signal from a DSC is the differential heat flow rate as a function of time. The

procedures required to evaluate the measured curve differ from one case to

another as shown in Chapters (2.3.1) and (2.3.2).

2.3.1. Linear scanning

The results from DSC measurements with linear temperature program could

be treated in different ways. As first the “2-curve” heat capacity determination

is presented. The use of normal, not hermetically sealed, DSC aluminium

pans (with a lid which rests on the sample and may be lightly closed by

crimping) always gives the heat capacity at constant pressure.

The procedure can be followed by Fig. (2.16). The temperature-time

curve during the experiment is shown by the red line, the response of the

calorimeter for empty pan (baseline) and sample are given as blue and green

lines respectively. The two-curve determination of heat capacity is performed

as follows.

1. Determination of the heat flow rate of the baseline Φ0(T), using empty

pans in the sample and reference ovens. The temperature program should

only be started when the isothermal heat flow rate at the starting temperature

Tst has been equilibrated for at least 1 minute. At the beginning and the end of

the temperature program isothermal segments are performed at temperatures

Tst and Tend, respectively. For the evaluation procedure all tree regions of the

42 Chapter 2

curve are needed: isothermals at start and end temperatures and scanning

region.

0 2 4 6 8 10

20

30

40

50

60

70

40

60

80

100

120

140

160

180

Isothermalat Tend

Hea

t flo

w ra

te in

mW

Time in min

Saphire

Sample

Baseline

Tempe

rature

progra

m

scanning mode

Isothermalat Tst

Endo

Φre

f−Φ0

Temperature in °C

ΦS−Φ

0

Figure 2.16. The two- and three-curve determination of the heat capacity; red

line – temperature program, green line – sample (pure PMMA)

measurement, blue line – baseline measurement, black line –

sapphire measurement, Tst and Tend – start and end

temperatures, respectively. Heating rate is 10 K/min, sample

mass 15 mg and sapphire mass 131 mg (PerkinElmer Pyris

Diamond DSC)

2. The sample of known mass is placed into the sample pan (or into a pan

of same type and mass as used for 1. on the sample side). Nothing should be

changed on the reference side. The same experimental procedure as for the

baseline measurement must be used for the sample measurement. A

correction for asymmetry of the measuring system is performed by subtracting

the empty scan in time domain from the sample measurement. Small

differences in start and end isotherms can be corrected by subtracting a

straight line bringing the end points of the isotherms to zero. After these

corrections heat capacity can be obtained from

cp S⋅mS⋅β = KΦ(T) ⋅(ΦS - Φ0) (2.1)

cp S, mS and ΦS are the specific heat capacity, mass and heat flow rate of the

sample, β is average heating rate, KΦ(T) is a temperature dependent

calibration factor and Φ0 is the heat flow rate of the empty pan measurement.


3. For the three-curve determination one more step is needed. KΦ(T)

could be neglected if the measurement of the calibration substance is also

performed under exact the same conditions. A calibration substance of known

mass and heat capacity cref (for this work sapphire was used) is placed into

the sample pan (or into a pan of same type and mass) while no other

parameter is changed. In analogy to Eq. (2.1) above one gets the following.

cp ref⋅mref⋅β = KΦ(T) ⋅(Φref - Φ0) (2.2)

And the specific heat capacity (at a given temperature) can be

calculated by a simple comparison of the heat flow rates into the sample and

the calibration substance as illustrated in Fig. (2.16).

refS

ref

ref

SS c

mm

c ⋅Φ−ΦΦ−Φ

=0

0 (2.3)

As mentioned above this method can be used only having stable start

and end temperature isotherms, which is not always the case for the polymer

nanocomposites. As the nanoparticles have large surface area, even after

drying at reduced pressure and temperatures above Tg some quantity of water

or solvent may remain.

The water is evaporated during the heating scan which causes

instability of the end isotherm as shown in Fig. (2.17). The red line

corresponds to the first heating scan of PS filled with SiO2 and the blue line to

the second. As it is seen from inset the isotherm at the end temperature for

the first scan not only differs from that of the second scan but also does not

reach a constant value.

Assuming that an immobilized polymer needs more energy e.g. higher

temperatures than Tg to relax, one expects the devitrification of RAF in the

range from Tg of the MAF up to degradation of the polymer. One method to

detect a possible devitrification of the RAF is therefore the heating of the

nanocomposite up to the degradation temperature of the polymer and check if

there is any additional transition.

44 Chapter 2

0 2 4 6 8 10 12 14

20

22

24

26

28

30

32

9 10 11 12 13 1426.1

26.2

26.3

26.4

26.5

40

60

80

100

120

140

160Isothermal

at Tend

1st heat 2nd heat

Hea

t flo

w ra

te in

mW

Time in min

ENDO

Isothermalat Tst

Temperature in °C

Figure 2.17. Heat flow rate as a function of time for PS / 24 wt% SiO2 system.

Heating rate is 10 K/min, sample mass is 18 mg (PerkinElmer

Pyris Diamond DSC); magenta line represents the temperature

program

In this case the end temperature isotherm cannot be stable because at

the temperatures close to the degradation of the polymer partial degradation

occurs. To avoid problems with unstable isotherms at the highest

temperatures the temperature interval of interest can be divided in small

temperature steps of some Kelvin, each followed by an isotherm. Then the

curve can be evaluated at least until the isotherms become unstable. This

method is called StepScan DSC which is a special version of temperature

modulated DSC [191].

2.3.2. StepScan DSC

TMDSC is an extension to conventional DSC which provides

information about the reversing and nonreversing characteristics of thermal

events. The additional information from TMDSC allows unique insights into the

structure and behaviour of materials. However the StepScan DSC (SSDSC), a

combination of short heating (or cooling) steps with isotherms of different

length determined by a stability criterion, is used in this work from which

similar information like in TMDSC is available. The detailed description of this

method [191] is given as next. A typical temperature-time-profile and the heat-


flow rate are shown in Fig. (2.18) for an initially amorphous PEEK sample. In

Fig. (2.19) the details of the data treatment are given.

1000 2000 3000 4000 5000-6-4-202468

101214161820

Hea

t flo

w ra

te in

mW

time in s

0246

250 300 350 400

112116

T in °C

2.53.03.5

exo

Heating rate (q

o ) in K m

in-1

Figure 2.18. StepScan DSC measurement of initially amorphous PEEK from

100 °C to 380 °C. The inset shows a part of the temperature

profile (step height δT = 2 K, heating rate q = 20 K min-1,

tiso max = 1 min, stability criterion = 0.02 mW absolute) and the

resulting heat-flow rate. The bottom part shows the mean

underlying heating rate which varies according to the length of

the isotherms. In the empty pan corrected heat-flow rate at about

1,000 s glass transition, at about 1,500 s cold crystallization and

around 4,500 s melting can be seen. (PerkinElmer Pyris

Diamond DSC)

In SSDSC heat capacity can be determined in several ways. As in

common DSC for each heating period the heat-flow rate displacement in

steady state is measured and heat capacity is obtained from Eq.(2.4).

op q

HFC = (2.4)

HF is the heat-flow rate necessary to heat the sample with the rate qo.

This evaluation requires steady state for each heating step. Therefore the

heating and the isothermal segment should not be too short, at least 20 s for a

power compensated PerkinElmer Instruments DSC. If the isothermal period is

too small it may happen that heat-flow rate does not go back to zero as

46 Chapter 2

expected. This was taken into account during the DSC measurements of

polymer nanocomposites performed to get the precise heat capacity data. The

isothermal segment was chosen to allow return to the equilibrium value for

each system under investigation.

Heat capacity can also be obtained from the ratio of the applied heat

and the resulting temperature step. In SSDSC the temperature step, δT, is

predefined and the heat-flow rate response, HF(t), is measured. Heat capacity

can be obtained from the area under the heat-flow rate peaks according

∫=St

p dttHFT

C0

)(1δ (2.5)

where ts is step time consisting of heating and isothermal time for each

individual step. By varying the step time the relevant time scale of a SSDSC

experiment can be varied. The details of the data treatment of a SSDSC

measurement are shown in Fig. (2.19).

140

150

160

170

Tem

pera

ture

in °C

1.5

2.0

2.5

3.0

Heating rate (q

o ) in K m

in-1

0.00.30.60.9 endoarea

Hea

t flo

wra

te in

mW

5000 5100 5200 5300 5400 5500 5600 5700

0123

cp from area

time in s

Spec

ific

heat

capa

city

in in

J/gK

Figure 2.19. StepScan DSC measurement of a PHB/PCL 50/50 blend in the

temperature region of PHB melting. δT = 1 K, q = 5 K min-1,

tiso max = 1 min, absolute criterion 0.001 mW. (PerkinElmer Pyris

Diamond DSC)

In StepScan DSC the length of the isotherms is not always predefined.

Depending on the setting of the equilibration criteria the next step in

temperature occurs as soon as the criterion is fulfilled. Consequently the


underlying heating rate may change depending on sample response. This can

be seen in Figs. (2.18) and (2.19) at melting. This allows a dramatically

reduction of measuring time in case of long equilibration times. As long as no

time dependent processes occur in the sample the instrument will only stay at

the isotherm for the time needed to reach steady state after the temperature

step. For the power compensating DSC this time is in the order of 0.5 min. If a

time dependent process yields an increasing heat-flow rate at longer times,

the length of the isotherm will be adapted accordingly. This way controlled rate

DSC experiments can be performed [194-196]. In Figs. (2.18) and (2.19) the

maximum time for the isotherms was set to 1 minute. Therefore the results

correspond to a time scale of 1 minute also most of the steps were actually

much shorter.

Because the data treatment performed in time domain in SSDSC is

straight forward and not based on Fourier analysis there is no need for

linearity or steady state neither during the heating nor during the isothermal

step. Therefore this method is very attractive for reasonable precise heat

capacity determination in relatively short time. In order to get the SSDSC

results more precise the baseline measurements were also performed and

subtracted during specific heat capacity determination.

It is also important to show that by SSDSC the uncertainty of the

absolute values of the instruments used for this work is acceptable and in

agreement with literature data. In Fig. (2.20) an example is shown. The black

line represents the specific heat capacity of sapphire from the literature. The

data for two sapphire samples with different masses are given as well. The

specific heat capacity of light (35 mg) and heavy (131 mg) sapphire samples

match perfectly together but not with the literature data. The uncertainty

makes about ± 1 % due to improper setting of the calibration factor KΦ(T).

Higher precision of the measurement can be reached if the curves are

corrected by one sapphire measurement as described above for the

three-curve heat capacity determination (Chapter (2.3.1)). The blue line in

Fig. (2.20) represents the corrected curve for the light sample considering the

heavy sample as calibration standard. It is seen that the precision of the data

corrected in this way is better than ± 1 %.

48 Chapter 2

80 90 100 110 120 130 140 150 160 1700.86

0.88

0.90

0.92

0.94

0.96

0.98

1.00

110 115 120

0.91

0.92

0.93

0.94

Literature data Sapphire (131 mg) Sapphire (35 mg) Sapphire corrected

Spec

ific

heat

cap

acity

in J/

k*g

Temperature in °C

Figure 2.20. Specific heat capacity of sapphire as a function of sample

temperature; (black) – literature, (red) and (green) – measured

data for heavy (131 mg) and light (35 mg) samples respectively,

(blue) – corrected data of light sample (Perkin Elmer Pyris 1

DSC); The inset shows a magnified interval of the curves where

the discrepancy between them is largest. δT = 3 K, q = 6 K min-1,

tiso max = 1 min, absolute criterion 0.0001 mW. (PerkinElmer Pyris

Diamond DSC)

Fig. (2.20) shows that the error in specific heat capacity determination

is in agreement with the data given in [197] or even below that of [192] for the

SSDSC measurements. Improvement of the heat capacity data from SSDSC

by applying the sapphire correction is marginal and uncertainties other than

that of heat capacity are more important for the nanocomposite

measurements. Therefore SSDSC without further sapphire correction has

been used as a reasonable precise measuring technique to obtain the specific

heat capacity of polymer nanocomposites prepared as described next. The cp

measurements have been carried out on a PerkinElmer PYRIS Diamond

DSC. The following measurement conditions were applied: 3 K step at heating

rate 6 K/min, 1 min isotherm (in case of high pressure pans 2 min),

temperature range 30 – 170 °C, absolute criterion was chosen equal to

0.0001 mW to have surely equilibrated system at the end of each step.

3. EXPERIMENTAL

3.1. Materials

For the investigation amorphous polymers with different functionality

have been chosen. The purpose was to obtain polymer nanocomposites with

different interfacial interaction strength from “no” up to strong interaction. One

of the polymers used is polystyrene - a polymer which appeared to dewet the

silicate surface [138] because of a lack of covalent or hydrogen bonds with

silica; therefore has no well-defined tendency to form any bond with

nanoparticles surface. The other two are poly(methyl methacrylate) and

poly(butyl methacrylate). Poly(methyl methacrylate) is often reported to exhibit

an interaction with the silicates.

styrene methyl methacrylate butyl methacrylate

Figure 3.1. Chemical structure of the monomers used

Poly(butyl methacrylate) has longer side group (butyl, as shown in

Fig. (3.1)) which exhibits more hydrophobic properties than methyl group of

poly(methyl methacrylate) and this may influence the interaction strength

between polymer matrix and nanoparticles. Considering hydrophility of silica a

weaker interaction is expected.

The polymerization was carried out for producing the PMMA filled with

silicon dioxide and Laponite RD nanocomposites. The reagents were

prepared for the polymerization as follows. The methyl methacrylate

(monomer) was distilled under reduced pressure; potassium persulfate

(initiator) - 99%, sodium dodecyl sulphate (surfactant for classical emulsion

C

C

O

O

H2C **

CH3

C H4 9

butyl

C H

C

C

OCH3

O

H2C **

3

methyl

HCH2C **

50 Chapter 3

polymerization) – 98.5% were used as received. All the chemicals mentioned

were received from Sigma Aldrich GmbH. The PMMA and PBMA for the

solution method were received from Scientific Polymer Products, Inc.

(http://www.scientificpolymer.com/catalog/description.asp?QproductCode=006

). The shear mixed samples were kindly provided by colleagues from the

Department of Polymer Structures, Leibniz Institute of Polymer Research,

Dresden. The PMMA synthesized by microemulsion polymerization was used

to get the PMMA/SiO2 nanocomposites and for the PMMA/Al2O3

nanocomposites the PMMA Oroglas, ARKEMA and Nanodur Al2O3

(d ≈ 36 nm) from Nanophase (www.nanophase.com). PS was kindly provided

by BASF.

The porous silicon dioxide nanopowder (spherical particles with

d ≈ 10 nm) – 99.5% with 530-690 m2/g specific surface area was used to

prepare the nanocomposites with PS, PMMA and PBMA. The specific surface

area given is much larger for the porous materials due to the fractal surface.

Assuming that the polymer cannot penetrate into the pores of nanosized silica

the spherical shape of nanoparticle should be considered. Taking this into

account the effective specific surface area of SiO2 nanoparticles (d ≈ 10 nm) is

estimated as 63 m2/g. Silica properties are related to the surface chemistry of

the samples. The hydroxy (OH) groups are generally bounded via the valence

bond with Si atoms on the silica surface (hydroxyl coverage), and in some

cases with Si atoms inside the particles of silica. In 1930’s studies of the

condensation processes of silicic acids (see [198] for review) showed that

hydroxyl (silanol) groups, ≡Si-OH, should be present on the surface of

silicates and silicas. Now numerous spectral and chemical data

unambiguously confirm the presence of the OH groups on such SiO2 surface.

Silanol groups are formed on the surface by two main processes [198]. First,

such groups are formed in the course of silica synthesis, e.g. during the

condensation polymerization of Si(OH)4 (Fig. (3.2a)). Here, the supersaturated

solution of the acid becomes converted into its polymeric form, which then

changes into spherical colloidal particles containing Si-OH groups on the

surface. Upon drying, the hydrogel yields xerogel, the final product, which

retains some or all of the silanol groups on its surface. Secondly, surface OH

Experimental 51

groups can form as a result of rehydroxylation of dehydroxylated silica when it

is treated with water or aqueous solutions. The surface silicon atoms tend to

have a complete tetrahedral configuration, and in an aqueous medium their

free valence becomes saturated with hydroxyl groups (Fig. (3.2b)).

Figure 3.2. The formation of silanol groups on the silica surface: (a)

Condensation polymerization; (b) Rehydroxylation [198]

The surface properties of amorphous silica, which is considered to be

an oxide adsorbent, in many cases depend on the presence of silanol groups.

At a sufficient concentration these groups make such a surface hydrophilic.

Figure 3.3. Types of silanol groups and siloxane bridges on the surface of

amorphous silica, and internal OH groups. Qn - terminology is

used in NMR, where n indicates the number of bridging bonds

(-O-Si) tied to the central Si atom: Q4, surface siloxanes; Q3,

single silanols; Q2, geminal silanols (silanediols) [198].

52 Chapter 3

The OH groups act as the centers of molecular adsorption during their

specific interaction with adsorbates capable of forming a hydrogen bond with

the OH groups, or, more generally, of undergoing donor–acceptor interaction.

Surface OH groups are subdivided as following (Fig. 3.3): (i) isolated free

(single silanols), ≡SiOH; (ii) geminal free (geminal silanols or silanediols),

=Si(OH)2; (iii) vicinal, or bridged, or OH groups bound through the hydrogen

bond (H-bonded single silanols, H-bonded geminals, and their Hbonded

combinations). On the SiO2 surface there also exist surface siloxane groups or

≡Si-O-Si≡ bridges with oxygen atoms on the surface. At last, there is

structurally bound water inside the silica skeleton and very fine

ultramicropores, d < 1 nm (d is the pore diameter), i.e. internal silanol groups.

The properties of the silica surface are very essential and may affect the

interfacial interaction as well as the heat capacity measurements as described

in Chapter (4).

Laponite RD, a synthetic hectorite clay Mg5.34Li0.66Si8O20(OH)4Na0.66

made up of nearly monodisperse, thin cylindrical platelets, with a crystalline

unit cell, rather similar to that of the natural montmorrilonite phyllosilicates

[199-201], was kindly provided by Southern Clay Products (Gonzales, Texas).

The platelets are of mean diameter d ≈ 30 nm and thickness l = 1 nm. The

specific surface area of the platelets is 320 m2/g. These clay particles are

composed of a central sheet of octahedrally coordinated magnesium ions

(with lithium ion substitution) between two tetrahedrally coordinated silica

sheets. Substitution of lithium for magnesium in the central sheet gives rise to

a net negative charge on the faces of the particles that is balanced by sodium

counterions. The counterions become unbound when Laponite RD is

dispersed in aqueous solution, leading to a charged colloidal suspension. The

edge charge of a Laponite RD particle is pH dependent [202]. At high pH, the

edge charge is negative, implying overall repulsive electrostatic interactions

between Laponite RD particles in solution.

Experimental 53

Figure 3.4. Laponite RD dispersion in water. A single disc of Laponite RD,

diameter 30 nm, thickness 1 nm,

A detailed phase diagram for Laponite RD suspensions at high pH as a

function of ionic strength and Laponite RD volume fraction has previously

been established [203]. At low ionic strength and high pH, increasing

Laponite RD volume fraction leads to a transition from a liquidlike phase to a

solidlike phase in which the system becomes jammed in a glassy state.

The polymer inorganic nanocomposites based on these polymers and

nanoparticles were prepared in different way as described next.

3.2. Preparation methods

Different preparation methods are utilized to obtain the polymer

nanocomposites expecting a variation of the interfacial interaction between

nanoparticles and polymer matrix. The preparation methods - solution mixing,

shear mixing (Leibniz Institute of Polymer Research, Dresden), classical

emulsion and non-surfactant microemulsion polymerization, were used during

this work.

3.2.1. Solution method

For this preparation method 2 g polymer was dissolved in 10 ml

chloroform at room temperature. The corresponding quantity of nanopowder

was dispersed in 15 ml chloroform by sonification. The sonification has been

performed by “Sonifier 250” (Branson Ultrasonics, USA) instrument for 30

minutes at output control position “4” and 40% duty cycle. Then the polymer

solution was added to the nanoparticles suspension in chloroform without

stopping the sonification. After that the mixture was sonificated under the

same conditions as for nanopowder suspension for 20 minutes. Then the

54 Chapter 3

mixture was heated up to the boiling temperature of chloroform, Tb = 66 °C

without sonification under stirring conditions. After evaporation of solvent the

composite obtained was heated up to Tg + 20 °C for 10 minutes. All the

samples were dried under reduced pressure (10-2 mbar) at 150°C for 8 hours

before the experiments.

3.2.2. Shear mixing

Two series of the samples of PMMA by microemulsion polymerization

filled with SiO2 (d = 10 nm) and PMMA, Arkema Oroglas™ VS-UVT Acrylic,

Injection Molding Grade, filled with Nanodur Al2O3 (www.nanophase.com,

d = 36 nm) have been prepared by shear mixing on a co-rotating twin screw

extruder ZE 25 (Berstorff, Germany). These series have been received from

the department of Polymer Structures, Leibniz Institute of Polymer Research,

Dresden.

3.2.3. Classical emulsion polymerization

The polymer nanocomposites by classical emulsion polymerization

were synthesized at the Department of Chemistry, State Engineering

University of Armenia. The polymerization was carried out at 70°C using

2 wt% (in respect to monomer) potassium persulfate as an initiator. First the

monomer (2 ml) was mixed with nanopowder (SiO2) and sonificated for 20 min

as described for solution method. Then to the reaction media 80 ml of 3 wt%

sodium dodecyl sulphate solution in bidistilled and deionized water was added

and heated up to 70°C. After, 10 ml of initiator aqueous solution was given to

the reaction media and polymerized under stirring conditions for 7 hours. The

monomer concentration in respect to water was about two volume percent in

order to keep the temperature of the reaction media constant. This is needed

because the polymerization is an exothermal process. Then the polymer latex

obtained was centrifuged with 6000 rpm. The sediment was mixed with 50 ml

of bidistilled and deionized water to wash away the surfactant. This was

repeated 3 times and the final sediment was dried under reduced pressure

10-2 mbar at 150°C for 8 hours.

Experimental 55

3.2.4. Microemulsion polymerization

Another type of emulsion polymerization which was performed in this

work is the non-surfactant microemulsion polymerization. Because the

obtained latexes have been characterized by REM which show that the latex

particles are in the range of maximum 0.5 μm, the polymerization is called

microemulsion. Microemulsion polymerization was carried out again at 70°C

using 2 wt% potassium persulfate as initiator. First the monomer (2 ml) was

mixed with nanopowder (SiO2) and sonificated for 20 min as described for

solution method. Then, without stopping the sonification, to the reaction media

80 ml bidistilled and deionized water was added and heated up to 70°C. After,

10 ml of initiator aqueous solution was given to the reaction media and

polymerized under sonification for 4 hours. No surfactant was used to avoid

unwanted interaction with the nanoparticles. For the high content polymer

nanocomposites’ preparation 15 ml chloroform was added to the monomer +

SiO2 suspension in order to avoid gelation and to allow better dispersion of the

nanoparticles. Later, the chloroform was just evaporated during the heating

before polymerization starts.

In case of Laponite RD nanofiller the synthesis was carried out starting

with the preparation of 70 ml suspension of nanopowder in water by

sonification at the same conditions as for SiO2 for 20 min. After, the monomer

was added and the reaction media was emulsified and heated up to the 70°C

under sonification. Then 10 ml aqueous solution of initiator was given to the

reaction media. Polymerization was carried out with continuous pulsing

sonification for 4 hours. Then the polymer latex obtained was centrifuged and

the sediment was dried as described above.

The obtained samples were then characterized by different methods.

3.3. Characterization

3.3.1. Gel permeation chromatography

The polymer nanocomposites obtained by microemulsion

polymerization have been characterized also by gel permeation

chromatography (GPC) in order to get the idea about the molecular weight

56 Chapter 3

and polydispersity of them. The GPC measurements were performed at the

Department of Polymer Structures, Leibniz Institute of Polymer Research,

Dresden.

The number average molecular weight is a way of determining the

molecular weight of a polymer. Polymer molecules, even ones of the same

type, come in different sizes (chain lengths, for linear polymers), so the

average molecular weight will depend on the method of averaging. The

number average molecular weight is the common, mean, average of the

molecular weights of the individual polymers. It is determined by measuring

the molecular weight of n polymer molecules, summing the weights, and

dividing by n as follows (Eq. (3.1)).

∑∑=

i i

ii in N

MNM (3.1),

where Ni is the number of molecules of molecular weight Mi. An alternative

measure of the molecular weight of a polymer is the weight average molecular

weight which calculated by Eq. (3.2).

∑∑=

i ii

ii iw MN

MNM

2

(3.2)

The polydispersity index (PDI), is a measure of the distribution of molecular

weights in a given polymer sample. The PDI calculated is the weight average

molecular weight divided by the number average molecular weight. It indicates

the distribution of individual molecular weights in a batch of polymers. The PDI

has a value always greater than 1, but as the polymer chains approach

uniform chain length, the PDI approaches unity (1) (http://en.wikipedia.org).

The data for the PMMA obtained by microemulsion polymerization is

given in Fig. (3.5). From the data represented in Table (1) come the following

results for this certain sample: Mn = 61 x 103 g/mol, Mw = 332 x 103 g/mol, PDI

Mw/Mn = 5.5.

Experimental 57

5 6 7 8

1. Messg. 2. Messg. Standard Mp = 138 500 g/mol

Elutionsvolumen [ml] Figure 3.5. The chromatogram of the PMMA by microemulsion

polymerization; blue and red lines correspond to 1st and 2nd

measurements and black one represents a standard

The filled samples prepared by the microemulsion polymerization have

not been characterized by GPC because of the nanofiller presence which

which may cause uncertainties in obtained data. It is assumed that the

molecular weights of the polymer in the composites are very much similar to

those of pure polymer.

The high PDI means that the polymer obtained has a very wide

molecular weight distribution as shown in Fig. (3.5) too.

Tabble 1. Gel permeation chromatographic data of pure PMMA by

microemulsion polymerization

Sample

Nr Mn in g/mol Mw in g/mol PDI

1A 63 x 103 337 x 103 5,349

1B 58 x 103 326 x 103 5,590

Average 61 x 103 332 x 103 5,5

58 Chapter 3

And as it was already mentioned in Chapter (2) the presence of the low

molecular weight polymer fraction is attractive for the interfacial interaction

between polymer matrix and the nanoparticles.

To characterize the latexes obtained by the microemulsion

polymerization and also to get information about the dispersion of the

nanoparticles in the polymer matrix electron microscopic methods were used.

3.3.2. Electron Microscopy

The raster electron microscopic (REM) characterization for some

latexes of the synthesized polymer nanocomposites was performed by “DSM

960A” TEM, Carl Zeiss at the “Center of Electron Microscopy”, University of

Rostock.

(a) (b) Figure 3.6. Raster electron microscopic images of PMMA with 27 wt%

Laponite RD (DSM 960A, Carl Zeiss)

The REM images with different magnification are given in Fig. (3.6) for

PMMA relatively highly filled with Laponite RD nanoparticles. Here one can

identify the latex spherical particles of nanocomposite in the range of 150 nm

in diameter. Due to the latex particle size the polymerization carried in this

work is called microemulsion. Similar pictures are received also for all other

systems prepared by microemulsion polymerization using sonification of the

reaction meadia.

2 μm 5 μm

Experimental 59

The transition electron microscopic (TEM) characterization for some

polymer nanocomposites obtained with intermediate filler content was

performed by “EM 902A” TEM, Carl Zeiss to show the degree of

deagglomeration of nanoparticles and how they are dispersed in the polymer

matrix.

The prepared nanocomposite samples have been pressed at 150°C

under 2 bar excess pressure. Then a thin film of about 10 μm has been cut

from the sample and was embedded into the epoxy resin “Araldit” by Fluka,

Switzerland. The resin was cured at 58 °C in a thermostat for 2-3 days. The

ultrathin cuts (50-100 nm) were obtained on ultramicrotome “Ultrotom III” by

LKB, Sweden and then fixed on the cupper grid and contrasted by uranyl

acetate and plumbum citrate.

Figure 3.7. TEM images of PMMA with 4 wt% SiO2 nanocomposite obtained

by microemulsion polymerization (EM 902 A, Carl Zeiss)

But even by contrasting the samples it was not possible to obtain a

clear picture where one may distinguish the core/shell morphology of

nanoparticles as it is reported for poly(styrene – methylmethacrylate)/SiO2

[99]. The possible reason for that could be the modification of nanoparticles by

oleic acid used for the modification of the nanoparticles in [99] which may be

also considered as additional contrasting. In this work SiO2 nanoparticles have

not been modified in order to obtain two-component systems and

consequently to investigate the interaction between only polymer and

nanopartilces assuming that the SiO2 nanopowder received from Sigma

200 nm50 nm

60 Chapter 3

Aldrich has no surface treatment or some adsorbed substances. Fig. (3.7)

shows that nanoparticles are agglomerated but even in such condition they

still have a great specific surface area.

In recent decade clay nanoparticles have attracted increased attention

due to enhanced functional properties of polymer clay nanocomposites.

Assuming that huge surface area of clay nanoparticles may also result for

composites in a RAF, the PMMA filled with Laponite RD nanocomposites have

been also prepared.

(a) (b)

Figure 3.8. TEM images of PMMA with 11 wt% (a) and 27 wt% (b)

Laponite RD nanocomposite (EM 902 A, Carl Zeiss)

In Fig. (3.8) the TEM images for different Laponite RD loadings of

PMMA nanocomposites are given. One can recognize that at medium filler

contents (11 wt%, Fig. (3.6a)) the clay platelets cover only the surface of

polymer latex particles and with increasing filler content (27 wt%, Fig. (3.8b))

Laponite RD nanoparticles start to penetrate also into the latex particle.

Fig. (3.8) shows that the filler is evenly distributed in the polymer matrix but

these images do not clearly answer the question if the clay nanoparticles are

fully exfoliated in these systems or not.

3.3.3. Small angle X-ray scattering

The SAXS experiments have been performed to clarify if the

Laponite RD clay nanoparticles are exfoliated or not. The measurements have

200 nm 200 nm

Experimental 61

been performed at the Department of Polymer Structures, Leibniz Institute of

Polymer Research, Dresden. The Laponite RD nanopowder has been used as

it was received and also dispersed in water and then dried as described for

sample preparation. The composites have been pressed at 150 °C and dried

as for sample preparation. The analyzing SAXS device was a KRATKY

compact camera (AntonPaar Graz, Austria).

0.1 0.2 0.3 0.4 0.5 0.6 0.7 0.8 0.9 1.0 1.10.1

1

25 10 12345

Nano-powder without drying Nano-powder dried PMMA/11 wt% Laponite RD (28) PMMA/27 wt% Laponite RD (30)I*

s /cp

s*nm

-1

s /nm-1

d(nm) (~1.2 ) (~1.2 ) (~ 2.0) (~ 2.0)

d /nm

Figure 3.9. Common comparison of Lorentz-corrected SAXS-curves

(Ism(s)*s vs. s) for Laponite RD with (red line) and without (blue

line) drying and for PMMA filled with 11 wt% (green line) and

27 wt% (magenta line) Laponite RD

The data shown in Fig. (3.9) were obtained without any treatment like

absorption correction, background correction or desmearing procedures.

Calculation of BRAGG values d from the positions of the scattering maxima

(layer reflections) was performed as follows.

2 d sin Θ = n λ (3.3)

Here 1/d = s is the scattering vector and λ is wave length (Cu-Kα

radiation ≈ 0.154 nm). Fig. (3.9) presents the outcome of the SAXS

experiments performed.

Fig. (3.9) demonstrates that in polymer clay nanocomposites obtained

the nanoparticles are not exfoliated. The comparison of red (nanopowder

without drying) and blue (dried nanopowder) lines allows making a conclusion

that there is no influence of drying on the clay structure. For both of them the

62 Chapter 3

distance between platelets made up about 1.2 nm. Observing the data for

nanocomposites, one sees that even at low clay loadings still not an exfoliated

system has been obtained. The nanoplatelets with about 2 nm interlayer

distance are present. Consequently the intercalated polymer-nanocomposites

have been obtained by microemulsion polymerization of PMMA with Laponite

RD.

3.3.4. Thermogravimetry

To get the content of the nanofiller in the polymer nanocomposites

thermogravimetric measurements for all samples have been performed on a

Labsys, Setaram, instrument at 2 K/min heating rate in the temperature range

from 30 to 650 °C under air. Thermal degradation of PMMA filled with SiO2

and Laponite RD is presented below in Figs. (3.10).

100 200 300 400 500 600-100

-90

-80

-70

-60

-50

-40

-30

-20

-10

0 PMMA pure (1) 4 wt% SiO2 (2) 15 wt% SiO2 (3) 22 wt% SiO2 (4) 30 wt% SiO2 (5) 47 wt% SiO2 (6) 66 wt% SiO2 (7) 73 wt% SiO2 (8) SiO2 pure

Mas

s los

s in

%

Temperature in °C

(a)100 200 300 400 500 600

-100

-90

-80

-70

-60

-50

-40

-30

-20

-10

0

PMMA pure (27) 11 wt% Laponite RD (28) 14 wt% Laponite RD (29) 27 wt% Laponite RD (30) 42 wt% Laponite RD (31) 59 wt% Laponite RD (32) Laponite RD pure

Mas

s los

s in

%

Temperature in °C

(b)

100 200 300 400 500 600-100

-90

-80

-70

-60

-50

-40

-30

-20

-10

0

Mas

s los

s in

%

Temperature in °C

PS pure (42) 9 wt% SiO2 (43) 22 wt% SiO2 (44) 46 wt% SiO2 (46)

(c)

Figure 3.10. Thermal degradation of PMMA SiO2 (a) and Laponite RD (b)

nanocomposites synthesized by microemulsion polymerization

and PS SiO2 (c) nanocomposites synthesized by solution

method (Setaram Labsys, TGA/DSC)

Experimental 63

Thermal behavior of the samples based on PMMA and SiO2 is very

similar for each series. Therefore the data only one of them (with the largest

number of samples) - PMMA with SiO2 by microemulsion polymerization is

shown. Fig. (3.10) shows that polymer nanocomposites obtained in this work

while being filled with SiO2 degrade in one step and those with Laponite RD

obviously in two steps. This can be explained by the modification of clay

nanoparticles which also may interact with polymer matrix in a different way

compared to SiO2. Another observation is that the degradation behavior of

composites is not much influenced with increasing filler content. The shift in

the temperature, when degradation starts, makes up maximum 50 K.

100 200 300 400 500 600 700

-80

-70

-60

-50

-40

-30

-20

-10

0

10

500 550 600 650 700-78

-77

-76

-75

-74

1st run 2nd run 3rd run

Mas

s los

s in

wt%

Temperature in °C

1%+

Figure 3.11. Independent measurements of PMMA with 25 wt% SiO2 (15)

nanocomposite using Setaram Labsys, TGA/DSC instrument

The precision of thermogravimetric measurements is certainly of

importance because deviations in filler content may result in misleading

information of RAF existence. The PMMA filled with 25 wt% of SiO2 (15)

nanocomposite obtained by solution method has been independently

measured three times at the same conditions to check the uncertainty of the

measurements. The inset in Fig. (3.11) shows that the deviation is about ±1%.

Information about the polymer-filler system, its preparation method,

nanoparticles size and the filler content which was received by

thermogravimetric measurements, one can find in Table (2).

64 Chapter 3

Table 2. Nanofiller contents for all samples prepared

N Preparation method Polymer Nanofiller type

Nanofiller dimensions

Nanofiller content, wt%

1 0 2 4 3 15 4 22 5 30 6 47 7 66 8

Microemulsion Polymerization PMMA SiO2 D = 10 nm

73 9 0

10 9 11 35 12

Classical emulsion polymerization

PMMA SiO2 D = 10nm

53 13 0 14 10 15 25 16 40 17

Solution method using PMMA (1)

PMMA SiO2 D = 10 nm

48 18 0 19 9 20 28 21 35 22

Solution method using PMMA from Scientific

Polymer Products PMMA SiO2 D = 10 nm

48 23 0 24 5 25 10 26

Shear mixing using PMMA (1)

PMMA SiO2 D = 10 nm

20 27 0 28 11.4 29 14 30 27 31 42 32

Microemulsion polymerization

PMMA Laponite RD d = 1 nm,

D = 30 nm

59 33 0 34 5 35 10 36 15 37

Shearmixing using PMMA Oroglas,

ARKEMA PMMA Al2O3 D = 30 nm

20 38 0 39 12 40 20 41

Solution method using PBMA from Scientific

Polymer Products PBMA SiO2 D = 10 nm

37 42 0 43 9 44 24 45

Solution method using PS 168N, BASF

PS SiO2 D = 10 nm

41

Experimental 65

The characterization of the polymer nanocomposites obtained by

different preparation methods was performed by the available techniques.

After that the existence of a possible RAF was checked.

3.4. RAF determination

Having the polymer nanocomposites prepared and characterized the

question of RAF existence has to be answered. This information can be

available applying the method described by Wunderlich [157] for

semicrystalline polymers. The calorimetric relaxation strength at the glass

transition Δcp can be considered as a tool for determination of the RAF in

semicrystalline polymers and the immobilized fraction in polymer

nanocomposites too.

Fig. (3.12) represents this method for semicrystalline polycarbonate.

The green line is corresponding to the liquid state of polycarbonate. In the

work that polymer has 23% [165] crystallinity which means that magenta line

should represent specific heat capacity above glass transition for such

crystallinity. But as one can see from Fig. (3.12) this is not the case. The

specific heat capacity and consequently Δcp from the calorimetric

measurement is much smaller than expected from 23% crystallinity.

380 390 400 410 420 430 440 450 460

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

e f

d

a

cpb (χcrystal= 0.23)

cpb (χsolid(Tg) = 0.49)

b - c p solid

c - cp liquid

Spec

ific

heat

cap

acity

in J/

K*g

Temperature in K Figure 3.12. Amorphous and semicrystalline polycarbonates specific heat

capacity as a function of temperature; introduction of RAF, see

text [165]

66 Chapter 3

This can be explained by the formation of RAF which does not

contribute to the step height at glass transition. The estimation of RAF is

shown in the following steps. The mobile amorphous fraction can be estimated

having the step in specific heat capacity of semicrystalline polymer Δcp sc from

the measurement (black line in Fig. (3.12)) and the Δcp a for the amorphous

polymer.

51.0)( =Δ

Δ=

ap

scpgma c

cTχ (3.4)

The RAF can be estimated according Eq. (3.5).

26.0)()(1)( =−−= gcgmagra TTT χχχ (3.5)

The same consideration will be applied to detect the existence of an

immobilized polymer fraction (RAF) in polymer nanocomposites. To apply this

method to the samples obtained one needs reasonable precise heat capacity

data, which are available from StepScan DSC measurements as discussed

above.

3.5. Annealing experiments

In this chapter another method to check the RAF existence is described

which is independent on drawing tangents on the heat capacity curves for the

calorimetric relaxation strength determination. This will help to prove the

results obtained from the Δcp determination.

There is a time-dependence of the supercooled polymers properties,

whereby their physical behavior changes as a function of annealing time at

constant temperature. The annealing of polymers can be understood in terms

of their amorphous structure by reference to a typical schematic enthalpy-

temperature diagram, presented in Fig. (3.13). On cooling from an equilibrium

liquid, the enthalpy departs from equilibrium (for simplicity indicated here as a

linear temperature dependence of the enthalpy) and forms a glass at a critical

temperature called the glass transition temperature, Tg, which depends on

cooling rate. The glassy state is characterized by an excess of enthalpy and

consequently there will be a thermodynamic driving force to reduce the

enthalpy towards equilibrium if the annealing temperature Tannealing is held

Experimental 67

constant after cooling through Tg. This reduction in specific enthalpy for pure

amorphous polymers and nanofilled systems should differ considering an

immobilized fraction. If there is a fraction of amorphous polymer in the

composite which does not contribute to the glass transition, then the enthalpy

relaxation of it will have a deficit in comparison to that of the pure polymer.

This assumption is also valid considering the following

∞Δ=Δ∗Δ HTcp (3.6)

Here ΔH∞ is the difference between starting H0 and equilibrium H∞ enthalpy of

the sample. Enthalpy relaxation ΔH(t) Fig. (3.13) can be presented as

anHHH −=Δ 0 (3.7)

Eq. (3.6) shows that ΔH is somehow related to Δcp but the experiments are in

totally different time scales.

Han

H0

Tannealing

liquid

qTg(q)

T

H

glass

annealing

H∞

Figure 3.13. Schematic enthalpy - temperature diagram showing the change

in enthalpy that occurs on cooling at rate q from the equilibrium

liquid (red curve), and the definition of the rate dependent glass

transition temperature Tg(q) (blue curve). Annealing at

temperature Tannealing reduces the enthalpy from H0 to Han

towards an equilibrium value H∞. The black curve describes the

heating after annealing.

Taking into account the arguments above the annealing experiments

have been performed to check the RAF existence independent from Δcp

68 Chapter 3

determination uncertainties. The temperature-time profile of the annealing

measurements performed is presented below in Fig. (3.14a). First, the sample

is heated up to the maximum temperature much above Tg which is in the case

of PMMA 170 °C. At that temperature any kind of thermal history or

mechanical or thermal stresses are excluded. Then the sample is cooled

down at 10 K/min to the current annealing temperature and annealed for

10 hours. Firstly the annealing time was chosen one hour but as the error bar

for enthalpy relaxation determination procedure was larger than an effect

which could be seen, the time of annealing was extended to ten hours. After

cooling to the minimum temperature the sample is heated up to the maximum

temperature. For excess cp determination, the sample was cooled and without

any annealing heated up again at exactly the same conditions. This heating

curve was used as a baseline for the two-curve heat capacity determination to

get the excess heat capacity data. An example of data treatment for pure

PMMA is given in Fig. (3.14b).

0 20 600 620 640 660 680

40

60

80

100

120

140

160

Tem

pera

ture

in °C

Time in min

tanneal= 600 min

70 80 90 100 110 120 130 140 1508

9

10

11

12

13

14

15

-0.2

0.0

0.2

0.4

0.6

0.8

1.0

PMMA annealed at 105°C Baseline

Temperature in °C

Hea

t flo

w in

mW

, End

o up

Excess specific heat capacity in J/K*g

sample

(a) (b)

Figure 3.14. Heat flow rate and excess specific heat capacity as a function of

sample temperature of PMMA annealed at 105°C for 10 h; red

line – first heating after annealing, blue line – baseline without

annealing, magenta line – excess heat capacity

The red line (Fig. (3.14a,b)) represents the first heating scan of the

sample after annealing. Then the excess specific heat capacity of the sample

was determined in accordance with two-curve determination using the second

heating scan (blue line) as a baseline. The magenta line in Fig. (3.14b) shows

Experimental 69

the outcome of the data treatment described. This was performed for different

annealing temperatures to get the dependence of the enthalpy relaxation on

the annealing temperature for the pure polymer and filled systems.

Quantitative comparison of such data allows drawing the conclusion about the

RAF existence in polymer nanocomposites. Namely, the area under the

excess specific heat capacity is determined by integration between 70 and

140 °C. The obtained data are normalized to the polymer mass to allow a

direct comparison between the pure polymer and the nanocomposites. The

excess heat capacity data for pure PMMA, PS and PBMA and its

nanocomposites filled with SiO2 prepared by solution method are given in

Appendix (A2).

70 Chapter 3

4. RESULTS

4.1. DSC measurements

In Chapter (2) several examples of RAF in semicrystalline polymers and

the discussion about the problems of detection of RAF devitrification are given.

Namely the main difficulty occurs because of the overlapping of melting of the

crystalline fraction, reorganization or reversing melting with the RAF relaxation

processes. That is why it is not possible to clearly recognize which process

begins first – melting of crystals or relaxation of the RAF. Such information is

needed to understand the exact mechanism of the RAF devitrification. In this

work it was tried to simplify the task – to exclude the melting of the crystalline

fraction. The crystalline lamellae have been “replaced” by inorganic

nanoparticles of the similar dimensions which are dispersed in a truly

amorphous polymer matrix, i.e. polymer nanocomposites have been used for

investigations. It is assumed that the RAF is formed also in such systems. In

this work the PMMA, PBMA and PS filled with spherical SiO2 nanoparticles

with ca 10 nm diameter and Laponite RD clay nanoparticles with 1 nm

thickness have been used. Different polymers and nanoparticles have been

chosen to investigate the influence of the polymer and nanofiller structure on

the formation of RAF.

First the existence of RAF in the polymer nanocomposites has to be

checked. Wunderlich [157] has introduced a method of RAF determination for

semicrystalline polymers which is described in Chapter (3). According to that

the step height at the glass transition Δcp can be used for the determination of

the RAF in semicrystalline polymers. This method was first applied to detect a

possible immobilized fraction in polymer nanocomposites.

Next, an example is presented how to apply that method to polymer

nanocomposites. In the Fig. (3.1) the green line corresponds to the heat

capacity of the polymer in the liquid state and the blue line to that of the solid

state of PMMA. The Δcp determination is also graphically shown by vertical

double arrow at the glass transition temperature for the pure polymer.

72 Chapter 4

60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3

cp 3phase

c p 2phase c p solid

PMMA pure (1) PMMA+47% SiO2 (6) 2 phases PMMA+47% SiO

2 (6) 3 phases

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

cp liquid

Tg

Δcp

Figure 4.1. Specific heat capacity of PMMA with 47 wt% SiO2

nanocomposite; straight lines are for solid and liquid states for

the pure polymer (green) and polymer nanocomposite according

to two- (magenta) and three-phase (black) model

Assuming that there is 47 wt% SiO2 for instance in the composite, the

magenta lines correspond to the simple mixing rule of the PMMA and SiO2,

e.g. two phase model of the nanocomposite. If the specific heat capacity of the

sample coincides with the black line which shows smaller step at glass

transition, then the deficit in the Δcp can be explained by the existence of RAF,

e.g. three-phase model of the polymer nanocomposites. According to this

model the nanocomposites consist of nanofiller, mobile amorphous fraction

(MAF) and RAF.

It has to be checked if there is RAF in polymer nanocomposites. In case

there is RAF in composite samples obtained one needs the data from the DSC

measurements for them all. Here below in Fig. (4.2, a-d) the specific heat

capacity as a function of sample temperature for four different

“polymer+nanoparticle” systems is given. The normalized data to the polymer

mass for all the other systems are presented in Appendix (A1).

Results 73

30 40 50 60 70 80 90 100 110 120 130 140 1500.70.80.91.01.11.21.31.41.51.61.71.81.92.02.12.22.3

ATHAS data PS pure (42) 9 wt% SiO2 (43) 24 wt% SiO2 (44) 46 wt% SiO2 (45)

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

(a)SiO2

-40 -20 0 20 40 60 80 100

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PBMA pure (38) 12 wt% SiO

2 (39)

20 wt% SiO2 (40) 37 wt% SiO

2 (41)

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

SiO2 (b)

30 40 50 60 70 80 90 100 110 120 130 140 150 160 1700.70.80.91.01.11.21.31.41.51.61.71.81.92.02.12.22.3

ATHAS data PMMA pure (18) 9 wt% SiO

2 (19)

28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

SiO2 (c)30 40 50 60 70 80 90 100 110 120 130 140 150 160 170

0.70.80.91.01.11.21.31.41.51.61.71.81.92.02.12.22.3 ATHAS data

PMMA pure (27) PMMA+11 wt% Laponite RD (28) PMMA+14 wt% Laponite RD (29) PMMA+27 wt% Laponite RD (30) PMMA+42 wt% Laponite RD (31) PMMA+59 wt% Laponite RD (32)

Step

Scan

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

SiO2 (d)

Figure 4.2. Measured specific heat capacity determined in respect to sample

mass for (a) – PS, (b) – PBMA, (c) – PMMA filled with spherical

SiO2 particles of 10 nm diameter prepared by solution method

and (d) - PMMA filled with Laponite RD nanocomposites

synthesized by microemulsion polymerization

But one has to mention that the specific heat capacity data has been

calculated in respect to the sample mass. This means that during the specific

heat capacity calculation the heat flow rate obtained from the measurements

has been divided by the mass of “polymer + nanofiller”. The specific heat

capacity data for polymer nanocomposites are shifted to lower values with

increasing filler content. The lowering is expected considering the additivity of

heat capacity because SiO2 specific heat capacity is lower than that of

polymers used. But such a shift is not observed for semicrystalline polymers

where the specific heat capacity of the crystalline, the rigid amorphous and not

immobilized fractions are very similar below the glass transition. As for

74 Chapter 4

semicrystalline polymers a slight increase in glass transition temperature was

observed for all nanocomposites after carefully drying at reduced pressure.

0 10 20 30 40 50 60 70-1

0

1

2

3

4

5

6

7

8

PMMA/SiO2 (1-8) PMMA/Laponite (27-32) PS/SiO2 (42-45)

T g -

T g pu

re in

K

Filler content in wt% Figure 4.3. Glass transition temperature of the different nanocomposites as

function of filler content. Half step temperature from StepScan

DSC measuremnts. Tg pure PMMA = 111 °C, Tg pure PS = 99 °C

The glass transition temperature of the nanocomposites was

determined as the half step temperature from the StepScan DSC

measurements. The values are slightly different compared to scan

measurements at 10 K/min because of the different time scale of the

experiment. Glass transition temperature and width of the step in heat capacity

were only little affected by the addition of nanofillers. This is different from

semicrystalline polymers where always a significant increase in glass transition

temperature and a broadening of the transition interval is seen, e.g. [159].

One has also to mention that the cp data for pure polymers should

coincide with ATHAS database data [166] within the uncertainty of DSC

measurements (±1 %) and ATHAS data bank (±5 %). But as it can be seen

from Fig. (4.2b) the PBMA specific heat capacity data are shifted to lower

absolute values in comparison to that of ATHAS database for about 7-8%

which can just be explained by the uncertainties given. For the other pure

polymers the measured curves are in good agreement with the ATHAS

database values.

Results 75

30 40 50 60 70 80 90 100 110 120 130 140 150 160 1701.0

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PMMA pure (18) PMMA+35% SiO2 (21) PMMA+48% SiO2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C

Δcp

Figure 4.4. The determination of the calorimetric relaxation strength at glass

transition for the PMMA SiO2 nanocomposites

The specific heat capacity data given are needed to estimate the

calorimetric relaxation strength at glass transition in order to check if there is a

RAF [157] in polymer nanocomposites obtained or not. As it was already

mentioned one has to draw the tangents outside the glass transition region.

Then the step height at Tg has to be estimated as it is shown in Fig. (4.4)

which is Δcp. The determination of Δcp for PMMA pure and filled with 35 wt%

and 48 wt% SiO2 nanocomposites is presented in Fig. (4.4).

In Fig. (4.4) the tangents drawn in solid and liquid states of specific heat

capacity curves for filled systems appear to have different slopes from that of

the ATHAS database [166] and pure polymer. Moreover the slope change is

larger with increasing filler content. This is discussed next in Chapter (4.2).

4.2. Specific heat capacity correction

For the comparison of the results from calorimetric measurements the

exclusion of the SiO2 contribution to the total heat capacity of the polymer

nanocomposite is needed. Considering the additivity of the heat capacity the

SiO2 contribution should be subtracted which is described below step by step.

In Fig. (4.5) the specific heat capacity dependence on sample temperature for

PMMA nanocomposites is presented as it was measured. As there is a

number of series for different preparation methods, only the data for one of

76 Chapter 4

them are given here as an example and the corrected data to the polymer

mass in Appendix (A1).

30 40 50 60 70 80 90 100 110 120 130 140 150 160 170

0.8

1.0

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PMMA pure (18) 9 wt% SiO2 (19) 28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22) SiO2 lit. data SiO2 dried up to 550°C

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C Figure 4.5. Specific heat capacity of polymer fraction as function of

temperature for PMMA SiO2 nanocomposites prepared by

solution method (the measured data)

In Fig. (4.1) the outcome of DSC measurements for PMMA SiO2

nanocomposites prepared by solution method is presented. To check the data

precision the specific heat capacity of PMMA from ATHAS database is also

given [166]. From the graph it is seen that the data for pure polymer does not

coincide with that of ATHAS database. But the difference is within the range of

the precision in absolute values of specific heat capacity which is for these

measurements ± 2 %, see Chapter (3). Moreover in solid state the pure

polymer and ATHAS database curves are parallel which means that the

measurements themselves gave reliable results. The discrepancy between

measured specific heat capacity of pure PMMA and that of ATHAS database

above Tg in liquid state is not known.

The SiO2 specific heat capacity is presented by dotted lines as well. The

dark yellow line corresponds to the literature data [204] for the bulk material

and the navy one is the measured specific heat capacity of SiO2 nanopowder

after drying at 550 °C in DSC in dry nitrogen. It is clearly seen that the

measured data has higher absolute values than that available from the

literature. It can be explained by the contribution of bounded water on the

Results 77

nanoparticles surface due to their great surface area. This indicates the

difficulties with precise cp measurements for systems with large surface areas.

But the nanocomposites were prepared by solution method which is described

in Chapter (3.2) and then dried under reduced pressure at 150 °C which is well

above the Tg of PMMA. Therefore for the subtraction the more reliable data

from the literature was taken because the surface of the nanoparticles is most

probably covered by polymer and the specific heat capacity should not include

a large contribution from water but this is not known in detail.

First step is to determine heat capacity (J/K) of each sample by

multiplication of the specific heat capacity data (J/K*gsample) by the mass of the

sample. Out of this step one gets the heat capacity data as it is shown in Fig.

(4.6).

30 40 50 60 70 80 90 100 110 120 130 140 150 160 1700

5

10

15

20

25

30

35

40

PMMA pure (18) 9 wt% SiO2 (19) 28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22)

Hea

t Cap

acity

in m

J/K

Temperature in °C

SiO2

Figure 4.6. Heat capacity of PMMA SiO2 nanocomposites prepared by

solution method (the heat capacity for the SiO2 fraction of each

sample is also given)

The contribution of the SiO2 fraction of each sample is shown in Fig. (4.6) in

the similar color as used for the corresponding sample. The heat capacities for

the polymer in each polymer nanocomposite after subtraction are shown in

Fig. (4.7).

78 Chapter 4

30 40 50 60 70 80 90 100 110 120 130 140 150 160 1700

5

10

15

20

25

30

35

40

PMMA pure (18) 9 wt% SiO2 (19) 28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22)

Hea

t Cap

acity

in m

J/K

Temperature in °C Figure 4.7. Heat capacity of the polymer of PMMA SiO2 nanocomposites

prepared by solution method

The heat capacity data shown Fig. (4.7) could be compared with each

other only as specific heat capacity data after the individual division by the

polymer masses for each sample.

30 40 50 60 70 80 90 100 110 120 130 140 150 160 1701.0

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PMMA pure (18) 9 wt% SiO

2 (19)

28 wt% SiO2 (20)

35 wt% SiO2 (21) 48 wt% SiO

2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure 4.8. Specific heat capacity of the polymer fraction as function of


solution method

From the division we get the specific heat capacity in J/K*gpolymer. It is

expected that the curves coinciding in the solid state with ATHAS database

data or at least with the pure PMMA due to the same degrees of freedom,

which is not the case as shown in Fig. (4.8). What can be seen is that the

Results 79

specific heat capacity data for the filled systems lie higher than that of the pure

polymer and ATHAS database. This means that by the subtraction described

above the contribution of the SiO2 is overcompensated. But it is known from

Chapter (3.3) that the thermogravimetric measurements have approximately

± 1% uncertainty. Taking this into account and assuming that the determined

filler content is not correct, the filler content was varied within the uncertainty.

This way the heat capacity data for any filler loading could be shifted to that of

the pure polymer by changing the filler content within 1-2% and performing the

subtraction steps as described. Consequently one can rely on the results

obtained from the DSC measurements only within the uncertainty of heat

capacity and filler content.

There may be another reason for the discrepancy too. The specific heat

capacity of the SiO2 nanopowder is not known precisely. Fig. (4.9) represents

the outcome of the subtraction procedure using the measured SiO2 data

(dotted line). Comparing this graph with that shown in Fig. (4.5) one can see

that there is nearly no difference in absolute values before and after

subtraction. And the final values differ from the ATHAS database values by up

to 25% and cannot be explained by uncertainty of neither thermogravimetric

nor DSC measurements. This indicates that the cp measurement of the SiO2

nanopowder is superimposed by adsorbed water heat capacity and water

desorption.

As it was already mentioned, in solid state the polymer cp data for each

sample should coincide with that of the pure polymer due to the similar degree

of freedom. And the deviations in cp are the result of the influence of

measurement precision and adsorbed water. In order to allow a direct

comparison of the specific heat capacity, the curves (Fig. (4.8)) were finally

shifted to the ATHAS data bank value at 60 °C. The result is shown in

Fig. (4.10).

80 Chapter 4

30 40 50 60 70 80 90 100 110 120 130 140 150 160 170

0.8

1.0

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PMMA pure (18) 9 wt% SiO2 (19) 28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C

SiO2

Figure 4.9. Specific heat capacity of polymer fraction as function of


solution method (the measured cp of SiO2 nanopowder was used

for the subtraction)

30 40 50 60 70 80 90 100 110 120 130 140 150 160 170

1.2

1.4

1.6

1.8

2.0

2.2 ATHAS data PMMA pure (18) 9 wt% SiO2 (19) 28 wt% SiO2 (20) 35 wt% SiO2 (21) 48 wt% SiO2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure 4.10. Specific heat capacity as a function of sample temperature for

PMMA SiO2 nanocomposites prepared by solution method

(corrected to the polymer mass by SiO2 contribution subtraction

and vertically shifted to ATHAS data [166] at 60 °C)

In Fig. (4.10) it is now clearly visible that the relaxation strength at glass

transition for different filler loadings is not the same. The higher the filler

content, the lower is the step height at glass transition even the data are

corrected for the filler content and presented in respect to polymer mass. This

Results 81

deficit in Δcp means that there is some RAF. Therefore the results obtained so

far can be explained by introduction of a RAF in the PMMA SiO2 and

Laponite RD nanocomposites (see Appendix (A1)).

As the cp of SiO2 has linear temperature dependence in the temperature

range of interest, the cp data of the samples may be also corrected to that of

the polymer by shift and rotation of the originally measured curves. Again

assuming that in solid state the polymer cp should be the same for all samples

the originally measured data was simply recalculated to that of the polymer

fraction for each loading and then by shifting and rotating fitted to the ATHAS

database data in the solid state in the temperature range of 30 – 80 °C. It is

interesting to check if there is significant difference between the polymer cp

obtained by the SiO2 contribution subtraction and by such simple manipulation.

30 40 50 60 70 80 90 100 110 120 130 140 150 160 170

1.2

1.4

1.6

1.8

2.0

2.2

1.4

1.6

1.8

2.0

2.2

2.4

2.6

PMMA/48 wt% SiO2

SiO2 subtracted PMMA/48 wt% SiO2

not corrected

ATHAS data PMMA pure (18) 9 wt% SiO

2 (19)

28 wt% SiO2 (20)

35 wt% SiO2 (21)

48 wt% SiO2 (22)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure 4.11. Specific heat capacity as a function of sample temperature

PMMA SiO2 nanocomposites prepared by solution method

(recalculated to the polymer mass and fitted to ATHAS data [166]

in solid state)

For the direct comparison the specific heat capacity of PMMA / 48 wt%

SiO2 was corrected to the polymer mass as shown in Fig. (4.10) (black dotted

line, Fig. (4.11)). It is seen that there is nearly no discrepancy between the

data corrected by the subtraction of bulk SiO2 cp (black dotted line) and that of

just shifted and rotated (magenta dotted line). Uncertainty has been estimated

82 Chapter 4

as 1 - 2%. As one can see even the shape of glass transition, as well as the

slopes in liquid state are similar in both cases (the dotted lines).

The data for PS (Fig. (4.12)) as well as all the other samples

(Appendix (A1)) were also corrected in the similar way. The PS filled with SiO2

nanocomposites do not exhibit RAF as it is seen from the Fig. (4.12). The data

for each nanofiller loading coincides with that of ATHAS database after

recalculation to the polymer mass. The only difficulty to draw such a

conclusion comes from the specific heat capacity curve for highest filler

content (cyan, Fig. (4.12)). The step in polymer cp at glass transition for that

sample is smaller that those for the other samples. This may be explained by

the influence of the large quantity of adsorbed water on the free (in case of no

interaction) surface of nanoparticles but not known exactly.

30 40 50 60 70 80 90 100 110 120 130 140 150

1,2

1,4

1,6

1,8

2,0

2,2 ATHAS data PS pure (42) PS+9 wt% SiO2 (43) PS+24 wt% SiO2 (44) PS+46 wt% SiO2 (45)

Spec

ific

Hea

t Cap

acity

in J/

K*g

sam

ple

Temperature in °C Figure 4.12. Specific heat capacity as a function of sample temperature PS

SiO2 nanocomposites prepared by solution method (recalculated

to the polymer mass and fitted to ATHAS data [166] in solid

state)

In Chapter (4.2) it was mentioned that large scatter in Δcp data for

PBMA filled with SiO2 nanocomposites is a result of wide glass transition

temperature interval which makes it difficult to draw the tangents to the liquid

and solid states of the polymer cp data. This might be solved by the extension

of the measurement temperature range but was not possible in solid state due

to the instrument limitation (-50 °C) and polymer degradation in liquid state.

Results 83

Anyway for this system the cp data of nanocomposites were also divided by

polymer fraction, shifted and rotated to that of the pure PBMA to observe at

least qualitatively the thermal behavior of the polymer fraction in

nanocomposites at glass transition.

-40 -20 0 20 40 60 80 1000.8

1.0

1.2

1.4

1.6

1.8

2.0

1.2

1.4

1.6

1.8

2.0

2.2

2.4 ATHAS data PBMA pure (38) 12 wt% SiO2 (39) 20 wt% SiO2 (40) 37 wt% SiO2 (41)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C

Δcp

Figure 4.13. Specific heat capacity as a function of sample temperature

PBMA SiO2 nanocomposites prepared by solution method

(recalculated to the polymer mass and fitted to ATHAS data [166]

in solid state)

In Fig. (4.13) it is clearly seen that the cp data for filled systems and

pure PBMA are very much identical after the correction described. This means

that there is not a significant amount of RAF in PBMA filled with SiO2

nanocomposites. Such a disagreement between this observation and the Δcp

values may be explained by the example of Δcp determination varying the

slope of the tangents drawn which is presented in Fig. (4.13) with respect to

the right scale. The magenta tangents are drawn following the slopes in solid

and liquid states of the measured cp data for PBMA filled with 37 wt% of SiO2.

The navy tangents are parallel to those of ATHAS database data. It is obvious

that by small variation of the tangent slope determined Δcp data might have up

to 30% uncertainty. Therefore the method of RAF recognition in polymer

nanocomposites based on Δcp determination is not applicable for all systems

and one needs to check the RAF existence in polymer nanocomposites by

another independent experiment described in Chapter (4.4).

84 Chapter 4

But the purpose of this work is to quantify the immobilized polymer

fraction in the polymer nanocomposites showing RAF which is, according to

[157] available from the measured cp data.

4.3. RAF determination

Following the idea of RAF determination by [25, 36] Δcp and Δcp pure are

the calorimetric relaxation strength at glass transition of the nanocomposite

and pure polymer, respectively. An immobilized or rigid amorphous fraction

can be determined from heat capacity according Eq. (3.5) replacing the

crystalline by the nanoparticle fraction [25, 36].

RAF = 1 – filler content - Δcp/ Δcp pure (4.1)

In Fig. (4.14) the Δcp data for PMMA filled with SiO2 nanoparticles

prepared by microemulsion polymerization are given. The normalization of the

data to the polymer mass is also needed to be able to compare the values for

different polymer nanocomposites in one graph directly. Depending on the

preparation method Δcp of the pure polymer for each series may differ.

In Fig. (4.14) according Eq. (4.1) the diagonal represents the case when

no RAF is present. In other words a two phase system (filler + polymer) is

present, which is expected if there is no interfacial immobilization. The steeper

red line is a guide for eye to show the decrease of the Δcp data for PMMA SiO2

nanocomposites by microemulsion polymerization. The upper arrow at 53 wt%

filler corresponds to the filler fraction. The lowest indicates the mobile

amorphous fraction (MAF) contributing to the calorimetric relaxation strength at

glass transition.

Results 85

0 10 20 30 40 50 60 70 80 90 1000,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

MAF

SiO2

Microemulsion polymerization (1-8)

Δcp

sam

ple /

Δcp p

ure

Filler content in wt%

RAF

Figure 4.14. Normalized calorimetric relaxation strength as a function of

nanofiller content for PMMA filled with SiO2 nanocomposites

prepared by microemulsion polymerization; the vertical magenta

double arrow indicates the amount of RAF at 53% filler (green

double arrow) and the blue one corresponds to MAF

The difference between the measured values and the diagonal (middle

arrow) represents the immobilized (rigid) fraction (RAF) which can be

calculated according Eq. (4.1). It corresponds to the mobile amorphous

fraction contributing to the relaxation strength at Tg.

In Fig. (4.15a) the data for all preparation methods are given. The points

for each preparation method lie nearly on the same red line as for

microemulsion polymerization. One can see that unexpectedly the interaction

strength between PMMA matrix and nanofiller surface, which should define the

amount of RAF, does not depend much on the preparation technique. As next

the results for PMMA filled with Laponite RD and aluminum oxide (Al2O3, ca 30

nm in diameter) nanocomposites are presented in Fig. (4.15b).

86 Chapter 4

0 10 20 30 40 50 60 70 80 90 1000,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0 Microemulsion p-n (1-8) Sol-n m-d with polymer (13-17) S-n m-d with standard (18-22) Classical emulsion pol-n (9-12) Shearmixing (23-26)

Δcp

sam

ple /

Δc p p

ure


RAF

(a)0 10 20 30 40 50 60 70 80 90 100

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

PMMA/Laponite RD (27-32) PMMA/Al2O3 (33-37)

Δc p

sam

ple /

Δc p p

ure


RAF

(b)


(a) nanofiller content for PMMA filled with SiO2 nanocomposites

prepared by microemulsion polymerization (squares), solution

method using synthesized PMMA (circles), solution method with

PMMA from Scientific Polymer Products (triangles), classical

emulsion polymerization (stars), shear mixing (diamonds); the

vertical double arrow indicates the amount of RAF at 53% filler

(b) PMMA with Laponite RD (green squares) and Al2O3 (cyan

squares) nanocomposites

Here the green line is a guide for eye showing the decrease of the

relaxation strength data for PMMA Laponite RD nanocomposites and is

steeper than that of PMMA SiO2 samples. On the contrary to that the diagonal

fits to the points for PMMA Al2O3 nanocomposites. From the description of the

materials used (Chapter 3) it is known that Laponite RD nanoparticles have

platelet-like form with 1 nm thickness. This means that they have larger

effective surface area than spherical SiO2 nanoparticles with 10 nm diameter.

Therefore it is expected that at the same nanofiller content exfoliated or

intercalated Laponite RD particles immobilize larger quantity of the polymer

than SiO2 what can be recognized in Fig. (4.15b). Even the glass transition

disappears (Δcp = 0) at about 50 wt% for the Laponite RD filler (Fig. 4.2d), e.g.

the whole polymer is immobilized by the nanoparticles. The Δcp data of PMMA

Al2O3 nanocomposites shows that there is no RAF detected. The calorimetric

relaxation strength data for PBMA and PS filled with SiO2 is presented in

Fig. (4.16).

Results 87

0 10 20 30 40 50 60 70 80 90 1000,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

Solution method PBMA / SiO2

(38-41)Δc

p sa

mpl

e / Δc

p pur

e


(a)0 10 20 30 40 50 60 70 80 90 100

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0

Solution method PS168N / SiO2

(42-45)

Δc p

sam

ple /

Δcp p

ure


(b)


nanofiller content for (a) - PBMA and (b) – PS with SiO2

nanocomposites

Fig. (4.16a) demonstrates the extremely large uncertainty of the data for

PBMA SiO2. The magenta line might show the RAF existence following the

points but taking into consideration the error bar no conclusion is drawn for

PBMA based systems. The discrepancy in data can be explained by the shape

of glass transition for PBMA. In Fig. (4.2b) one can see that the glass transition

has larger temperature interval than for PMMA and PS. This causes difficulties

to draw the tangents during Δcp determination. Fig. (4.16b) shows the Δcp data

for PS SiO2 system. The points lie close to the diagonal, which according

Eq. (4.1) represents the case when no RAF is present. In other words a two

phase system (filler + polymer) is present, which is expected if there is no

interfacial immobilization.

Having Δcp data for all systems observed the RAF content has been

estimated for all the systems prepared by means of Eq. (4.1). The results

obtained and normalized to the whole polymer fraction in composites are given

in Fig. (4.17) as function of the filler content. The red and green lines present

the data for PMMA SiO2 and Laponite RD nanocomposites respectively.

88 Chapter 4

0 10 20 30 40 50 60 70 80

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

1,0 PMMA/SiO2 (1-8) PMMA/Laponite RD (27-32) PS/SiO

2 (42-45)

PBMA/SiO2 (38-41)

RA

F / P

olym

er fr

actio

n

Filler content in wt% Figure 4.17. Normalized RAF as a function of filler content for (red) – PMMA

with SiO2, (green) - PMMA with Laponite RD, (blue) – PS with

SiO2 and (magenta) – PBMA with SiO2 nanocomposites

The blue and magenta lines correspond to the data for PS and PBMA

SiO2 nanocomposites respectively. As it can be seen from Fig. (4.8) the RAF

seems to be linearly dependent on the nanofiller content due to most likely

agglomeration of the nanoparticles. For Laponite RD filled PMMA

nanocomposites the RAF even saturated at highest filler concentration

(59 wt%).

The situation for the PS and especially PBMA composite samples is not

clear due to not well defined tangent construction. In these cases the

determination of Δcp for RAF estimation is uncertain. Molecular mobility at and

below the glass transition can be tested by annealing experiments too, see

[205] for a review. If a fraction of the amorphous polymer in the

nanocomposites is immobilized it is expected that enthalpy relaxation below

glass transition is reduced too [206, 207]. To check if this provides more

definite results regarding RAF the following annealing experiments were

performed.

4.4. Annealing experiments

As follows from the previous chapter at this step the results obtained

from Δcp estimation are checked by another method, independent on tangent

construction – annealing experiments. If a fraction of the amorphous polymer

Results 89

in the nanocomposites is immobilized it is expected that enthalpy relaxation

below glass transition is reduced too [206, 207]. To check if this provides more

definite results regarding RAF the following annealing experiments were

performed. Following the idea described in Chapter (3.5) it is shown that in this

work during the excess specific heat capacity determination from the

annealing experiments not an empty pan was used as a baseline but the

second heating scan without annealing of the same sample. This allows

comparing the behavior of only the polymer and any influence of adsorbed

water etc may be neglected.

As it is described in Chapter (3.5) the samples were heated well above

the Tg to erase the previous thermal history. Then after cooling at 10 K/min to

the annealing temperature they were annealed for 10 hours. For the detection

of the enthalpy change due to the annealing, the composite samples were

cooled and reheated again.

50 60 70 80 90 100 110 120 130 140-0.2

-0.1

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

Ta = 60°C

Ta = 80°C

Ta = 90°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = 95°C

Figure 4.18. Excess specific heat capacity after annealing at different

temperatures versus temperature for pure PS (42); the annealing

time is 10 h, sample mass 17 mg

The typical annealing peak is seen in the first heating, which is not

present in the next heating without annealing (Fig. (3.14)). The difference of

both curves yields excess heat capacity as shown in Fig. (4.18) for PS. The

data for PMMA and PBMA filled with SiO2 nanocomposites are presented in

Appendix (A2). Integration from 70 to 140 °C finally gives the specific enthalpy

90 Chapter 4

change ΔH during annealing. To allow a direct comparison the specific

enthalpy change was normalized to the polymer fraction of the

nanocomposites. Because of the small shift in glass transition temperature

results are plotted versus Tannealing - Tg. Here the peak position shifts from

higher to lower temperatures with decreasing annealing temperature.

-100 -80 -60 -40 -20 0 20-0,4

-0,3

-0,2

-0,1

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

PMMA pure (1) PMMA+15% SiO2 (3)

Enth

alpy

in J/

g poly

mer

Tanneal-Tg in K

(a)

-70 -60 -50 -40 -30 -20 -10 0 10

-0,2

-0,1

0,0

0,1

0,2

0,3

0,4

0,5

0,6

0,7

0,8

0,9

PBMA pure (38) PBMA+20% SiO2 (40)

Enth

alpy

in J/

g poly

mer

Tanneal-Tg in K

(b)

-35 -30 -25 -20 -15 -10 -5 0-0,2

0,0

0,2

0,4

0,6

0,8

1,0

1,2 PS pure (42) PS+24% SiO2 (44)

Enth

alpy

in J/

g poly

mer

Tanneal.-Tg in K

(c)

Figure 4.19. Enthalpy change (J/gpolymer) during annealing for 1 h as function

of the annealing temperature for (a) – PMMA, (b) – PBMA, (c) -

PS and its SiO2 nanocomposites; Tg PMMA = 111 °C,

Tg composite = 113 °C; Tg PBMA = 30 °C, Tg composite = 31 °C;

Tg PS = 99 °C, Tg composite = 103 °C

The area under the peak first increases and, after some maximum,

decreases with increasing annealing temperature. This is better seen in

Fig. (4.19) where the enthalpy change (J/gpolymer) during annealing for 1 h as

function of the annealing temperature is shown. The maximum for the samples

is in the glass transition region, just 5 – 10 K below Tg, as expected. The error

bars in Fig. (4.19) indicate that the effect observed is smaller or in the same

Results 91

range as the uncertainty of such measurements. Therefore the annealing time

has been extended to 10 h in order to increase the area under the peak in

excess cp data.

The enthalpies determined from the annealing peaks (Tannealing = 10 h)

for pure PS and its nanocomposite with 24 wt% of SiO2 are represented in the

Fig. (4.20). For this system no difference in the enthalpy relaxation for the pure

polymer and the nanocomposite is detected as shown in Fig. (4.20).

-120-110-100-90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20 30 40 50

-0,2

0,0

0,2

0,4

0,6

0,8

1,0

1,2

1,4 PS pure (42) 24 wt% SiO2 (44)

Enth

alpy

in J/

g poly

mer

Tanneal-Tg in K Figure 4.20. Enthalpy change (J/gpolymer) during annealing for 10 h as function

of the annealing temperature for PS and PS SiO2

nanocomposite; Tg PS = 99 °C, Tg composite = 103 °C

92 Chapter 4

-80 -60 -40 -20 0 20 40

-0,2

0,0

0,2

0,4

0,6

0,8

1,0

1,2

1,4

PBMA pure (38) PBMA+20% SiO2 (40)

Ent

halp

y in

J/g po

lym

er


of the annealing temperature for PBMA and PBMA with SiO2

nanocomposite; Tg PMMA = 30 °C, Tg composite = 31 °C

As expected from the dewetting properties of PS on silica surface pure PS and

the polymer fraction in the PS with 24 wt% SiO2 behave in the same way - the

data for them both coincide within the error limit of ±0.1 J/g.

Such a discrepancy may appear as a result of the error coming from the

thermogravimetric measurements, see Chapter (3.3). Fig. (4.21) demonstrates

the same situation for pure PBMA and the composite with 20 wt% SiO2.

-120-110-100 -90 -80 -70 -60 -50 -40 -30 -20 -10 0 10 20 30

-0,2

0,0

0,2

0,4

0,6

0,8

1,0

1,2

1,4 PMMA pure (18) PMMA+35% SiO2 (21)

Enth

alpy

in J/

g poly

mer


of the annealing temperature for PMMA and PMMA with SiO2

nanocomposite; Tg PMMA = 111 °C, Tg composite = 113 °C

Results 93

On the contrary, the enthalpy change for polymer fraction of PMMA with

35 wt% SiO2 depends on the normalized annealing temperature in different

way than those for PS and PBMA. In Fig. (4.22) the points for the filled system

lie significantly lower than those for pure polymer. This difference cannot be

explained by the error because the effect is larger than the uncertainty for the

enthalpy change which is ± 0.1 J/g. This means that not the whole polymer

fraction in the composite contributes to the enthalpy change, e.g. the missing

part is immobilized. Therefore one can conclude that there is RAF in PMMA

SiO2 nanocomposites, which is not the case for PS and PBMA filled with SiO2

systems.

-120 -100 -80 -60 -40 -20 0-0,3

-0,2

-0,1

0,0

0,1

0,2

0,3

0,4

0,5

0,6 PMMA (18) - (21) PS (42) - (44) PBMA (38) - (40)

ΔHpo

lym

er p

ure- Δ

Hpo

lym

er c

ompo

site

in J/

g poly

mer

Tanneal-Tg in K Figure 4.23. Enthalpy change difference versus normalized annealing

temperature for PMMA, PS and PBMA SiO2 nanocomposites

The mentioned above demonstrated by the data presented in

Fig. (4.23). Assuming that the deficit in the enthalpy change shown in

Fig. (4.22) is due to the RAF formation, the data of the polymer fraction in

nanocomposite has been subtracted from that of the pure polymer to show the

difference between systems with and without interfacial immobilization. One

expects that the outcome lies on the zero line (black horizontal line) within the

error bar if there is no RAF in the composite. This can be observed for the PS

and PBMA filled systems within the error equal to ± 0.15 J/g. Fig. (4.23) shows

that the values estimated by the subtraction for PMMA with 35 wt% SiO2 lie on

94 Chapter 4

the line corresponding to 0.37 J/g on enthalpy change difference scale which

is larger than the discrepancy of the data (± 0.15 J/g).

Consequently the polymer nanocomposites showing RAF have been

prepared. Having such model system, the devitrification of RAF is tried to be

detected as next.

4.5. Devitrification of RAF at high temperature

4.5.1. StepScan DSC

The specific heat capacity data from StepScan DSC for the composites

obtained has been regularly performed up to 170°C. In the temperature

interval between Tg and 170°C no RAF devitrification was observed. One of

the simplest methods to detect when the RAF relaxes is to heat the polymer

nanocomposite up to the degradation temperature of the polymer. The higher

the temperature is, the more energy is transferred to the polymer, therefore it

is expected that the chance of RAF devitrification increases.

60 80 100 120 140 160 180 200 220 2401,3

1,4

1,5

1,6

1,7

1,8

1,9

2,0

2,1

2,2

2,3

2,4

2,5 ATHAS data PMMA pure (1) PMMA/47 wt% SiO2 (6)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure 4.24. Specific heat capacity (J/K*gpolymer) of pure PMMA (red) and filled

with 47 wt% SiO2 nanocomposite (blue) measured in StepScan

mode up to the degradation of the polymer. δT = 3 K,

q = 6 K min-1, tiso max = 1 min, absolute criterion 0.0001 mW,

sample mass is 15 mg (PerkinElmer Pyris Diamond DSC).

In Fig. (4.24) the specific heat capacity of PMMA filled with 47 wt% SiO2

nanocomposite as a function of sample temperature is shown. As it is

Results 95

mentioned in Chapter (1) the RAF devitrification is expected to appear as a

second glass transition but as it is clear from the Fig. (4.24), there is no

steplike transition observed up to 230°C when the polymer starts to degrade.

4.5.2. High rate DSC

Influence of degradation on the heat capacity determination can be

reduced by using high heating rates. Using high rates up to 400 K/min the

polymer does not degrade even up to nearly 350 °C. And as the RAF

relaxation is expected to appear at the temperatures higher than conventional

Tg, it was supposed that this would help to detect the devitrification of RAF.

50 100 150 200 250 300 350

1.4

1.6

1.8

2.0

2.2

2.4

2.6 ATHAS data PMMA pure (1) PMMA+47 wt% SiO2 (6)

Spec

ific

heat

cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure 4.25. Hyper DSC [208] measurements at 400 K/min heating rate, mass

of each sample is 0.9 mg

But as it is seen from the Fig. (4.25) there is no any second glass

transition up to nearly 350 °C. This means that the interaction between

polymer matrix and nanoparticles is so strong that heating even up to such

high temperatures is not enough to devitrify RAF.

4.6. Plasticization experiments

As the high rate DSC was not enough to detect the RAF devitrification

calorimetrically, another method – plasticization of the polymer was used. The

plasticizer was added to lower the Tg hoping that the devitrification of the

immobilized polymer fraction will be also lowered. The idea is following. The

polymers we used degrade in the temperature range of 200-250 °C at low

96 Chapter 4

heating rates. This limits the measurement interval to between the Tg of MAF

and degradation temperature of the polymer. What one can do is to shift the Tg

of the MAF to lower temperatures which will enlarge that interval. So the

chance to devitrify the RAF is higher than without plasticizer addition. Of

course, this will work only in case if RAF is plasticized as well. For these

experiments the chloroform as plasticizing agent was used.

In Fig. (4.26) the specific heat capacity data from the pure PMMA and

PMMA with 47 wt% of SiO2 nanocomposites plasticized by different amounts

of chloroform is given.

-40 -20 0 20 40 60 80 100 120 140 160

0.6

0.8

1.0

1.2

1.4

1.6

1.8

2.0

2.2

2.4

2.6

2.8

3.0

32 wt% Chloroform

PMMA pure (1) PMMA with 47% SiO2 (6) ATHAS data

Spec

ific

Hea

t Cap

acity

in J/

K*g

dry

sam

ple

Temperature in °C

dry samples

32 wt% Chloroform

Figure 4.26. Specific heat capacity of the plasticized samples as a function of

sample temperature for PMMA pure (green curves) and

PMMA / 47 wt% SiO2, mass of dry samples is 15 mg (pure

PMMA) and 20 mg (PMMA / 47 wt% SiO2), the maximum

concentration of chloroform is 46 wt% in respect to polymer

The weight percentage is given in relation to the polymer mass. As one

can see the Tg of both pure PMMA and composite at 32 wt% chloroform is

nearly the same. This may mean that the organic solvent plasticizes the whole

polymer in composite, even immobilized fraction. But one has to be careful

while comparing different samples.

As it was mentioned above the Tg of RAF will be lowered only in case if

the solvent penetrates into the RAF. This can be checked by using the Fox

equation (Eq. (4.2)) [209] for two-component systems.

Results 97

solvg

solv

polg

pol

g TTTωω

+=1 (4.2)

Here Tg, Tg pol and Tg solv are the glass transition temperatures of

plasticized system, polymer and solvent respectively, ωpol and ωsolv are weight

fractions of the polymer and solvent respectively. For comparison of the data

obtained so called “calibration curve” is shown in Fig. (4.27). In other words, it

is the Tg of the pure polymer as a function of solvent concentration which

follows Fox equation. The same dependence was evaluated from the Fig.

(4.26) for the filled system and plotted with the calibration curve together. If the

solvent lowers the Tg of the polymer in both samples in the same way, this will

mean that it plasticizes the RAF also. And vice versa, if the data for the

composite do not follow the calibration curve, it will indicate that RAF is not

plasticized. Tg for pure polymer and the composite is not the same, therefore in

Fig. (4.27) the difference between dry sample Tg and that of the plasticized is

plotted versus solvent concentration.

0,0 0,1 0,2 0,3 0,4 0,5 0,6 0,7 0,8 0,9 1,0-200

-180

-160

-140

-120

-100

-80

-60

-40

-20

0 PMMA pure (1) PMMA with 47% SiO2 (6)

with respect to whole polymer with respect to MAF

Fox equation fit

T g pl

astic

ized

-Tg

dry in

K

Solvent content in mass fraction Figure 4.27. Normalized glass transition depending on the plasticizer

concentration, Tg dry PMMA = 111 °C, Tg dry composite = 118 °C

From this graph one can see that in comparison to the pure polymer,

the polymer fraction of the composite is completely plasticized. While solvent

concentration is estimated in respect to the whole polymer mass the data for

the filled system are lying on the same line as the data from the pure PMMA

within the error limit. This means that chloroform penetrates into either MAF or

98 Chapter 4

RAF in polymer nanocomposites obtained. But the devitrification of the RAF is

not observed from the data presented (Fig. (4.27)). There are two possible

explanations for that. First is that the RAF devitrification is spread over the

large temperature interval from the MAF Tg up to the degradation of the

polymer and by the calorimetric methods used during this work it is not

possible to detect it. The second is that the devitrification of RAF in polymer

nanocomposites does not occur before degradation of polymer.

5. DISCUSSION

Interaction at the filler polymer interface is considered to be important

for the behavior of polymer nanocomposites. Following the work of Lipatov

and Privalko [25, 36] the fraction of immobilized polymer (RAF) was quantified

for PMMA, PBMA and PS SiO2 and PMMA Laponite RD nanocomposites. For

the polymer fraction of the nanocomposites Eq. (4.1) can be rewritten for the

immobilized fraction of the polymer RAFpolymer [25, 36]

RAFpolymer = 1 – Δcp polymer/ Δcp pure (5.1)

From Figs. (4.11, 4.12) and Fig. (A1.5) the step height in specific heat

capacity Δcp polymer was determined as usual at Tg and normalized by the step

in heat capacity for the pure polymer. The data for different polymer-

nanoparticle systems are presented except those of PBMA based

nanocomposites because of the difficulty determining Δcp as disscused in

Chapter (4). No certain conclusion can be drawn for these systems due to

very broad glass transition. The Δcp data are not consistent; therefore the

specific heat capacity has to be measured from much lower temperatures to

lower the error in tangent construction for Δcp estimation.

Uncertainty of the step height for PMMA and PS based samples is

again mainly due to uncertainties in the slope of the tangents needed for the

determination. But in the normalized representation (Figs. (4.11, 4.12) and

Fig. (A1.5) the tangents should be parallel independent on filler content. Even

the situation is improved in Fig. (5.1) compared to Figs. (4.14-4.17)

determination of the tangents is still highly subjective. Nevertheless RAFpolymer

for PMMA and PS nanocomposites from normalized relaxation strength, was

determined from the data presented in and are shown in Fig. (5.1). According

Eq. (5.1) the rigid amorphous fraction of the polymer in the nanocomposite

(RAFpolymer ) could be obtained. The result for the 66 m% SiO2 PMMA

nanocomposite is indicated by the vertical arrows in Fig. (5.1). According

Eq. (5.1) a value of unity in Fig. (5.1) represents the case when no RAF is

present (two phase system; filler + polymer). The curved solid lines result from

a model assuming a constant ratio between RAF and filler content (equivalent

to the straight lines in Figs. (4.14-4.17)).

100 Chapter 5

0 10 20 30 40 50 60 70 80 90 1000.0

0.2

0.4

0.6

0.8

1.0

PS SiO2 (42-45) PMMA SiO2 (1-8) PMMA Laponite RD (27-32)

Δcp

poly

mer /

Δcp

pure

pol

ymer


RAFpolymer

MAFpolymer

Figure 5.1. Calorimetric relaxation strength of the polymer fraction as a

function of nanofiller content. Symbols: – PS with spherical

SiO2 nanoparticles; – PMMA with spherical SiO2

nanoparticles; – Laponite RD clay nanoparticles; synthesized

by in-situ microemulsion polymerization. The straight lines

through the measured points are guides to the eyes only. The

vertical double arrows indicate the amount of RAF and MAF at

70 m% filler for PMMA and spherical particles.

The ratio equals 0.1, 0.4 and 1 for PS, PMMA SiO2 and PMMA

Laponite RD nanocomposites, respectively. Assuming a decrease of the RAF

proportional to the polymer fraction yields the straight lines in Fig. (5.1). Even

the determination of Δcp polymer from Figs. (4.11, 4.12) and Fig. (A1.5) is a bit

more objective than Δcp sample from the measured data it does not give much

better results. In both cases a RAF is obviously detected for the PMMA

nanocomposites. For the PS nanocomposite the result is not definite.

Therefore an independent determination of the RAF was needed to allow

definite conclusions.

The occurrence of a RAF was confirmed for PMMA nanocomposites by

enthalpy relaxation studies too. For the PS and PBMA SiO2 nanocomposites

studied the result from heat capacity is not well defined but no RAF was

detected from enthalpy relaxation.

Discussion 101

Figure 5.2. Sketch of spherical (a, b) and layered (c) nanoparticles covered

by a layer of immobilized polymer (RAF). Total deagglomeration

of the particles is assumed in (a).

Existence of a RAF was probed by calorimetric experiments detecting

contributions from liquid like degrees of freedom to heat capacity and enthalpy

relaxation. In both experiments cooperative motions on a length scale of about

2 nm are probed [48, 210]. This is a much longer distance than the interaction

depth force range for a polymer molecule at the filler surface. The question

arises what the thickness of the immobilized layer around a nanofillers particle

is. From geometric consideration assuming spherical particles of 10 nm

diameter for the SiO2 filler or platelets with 1 nm thickness for the Laponite RD

filler and a density of 1 g/cm3 for the polymer and 2.4 g/cm3 for SiO2 a layer

thickness ranging from 2 nm to 1 nm follows from the data shown in

Figs. (4.14-4.17) and Fig. (5.1).

The relative amount of RAF (RAF/Filler) is significantly larger for the

PMMA Laponite RD nanocomposite compared to the PMMA SiO2

nanocomposite, Figs. (4.14-4.17) and Fig. (5.1). Despite this the thickness of

the RAF layer around the nanoparticles at low filler concentration, when

agglomeration is not dominant, is nearly the same – about 2.5 nm (Fig. (5.3)).

The detailed description of the RAF layer thickness estimation is given in

Appendix (A3). A similar value (2 nm) was found for the RAF layer at the fold

surface of semicrystalline PET [159] and 1.5 nm for a filled SBR 1500 rubber

[211]. The thickness of the immobilized layer is in all cases much thicker than

the range of the forces due to the interaction of the polymer with the

nanoparticle, which are in the order of several Å. Following the idea of the

Mobile

polymer

Rigid

amorphous

Nano-

particle

dRAF ca. 2 nm(b) (c) (a)

102 Chapter 5

importance of cooperatively rearranging regions (CRR) [48, 210, 212] for the

liquid like motions near the glass transition this observation can be explained.

0 10 20 30 40 50 60 70 800,6

0,8

1,0

1,2

1,4

1,6

1,8

2,0

2,2

2,4

2,6

2,8

3,0 PMMA/SiO2 (1-8) PMMA/Laponite RD (27-32)

d RAF in

nm

Filler content in mass% Figure 5.3. Thickness of the immobilized layer around the nanoparticles

(RAF) as a function of filler content. The line is a guide to the

eyes only.

Immobilizing a part of a polymer molecule at the interface affects the

movement of all neighboring segments within a CRR. Therefore the whole

CRR can not contribute to the liquid like motions and consequently not to the

increment of heat capacity at the glass transition. If one assumes that the

molecules are bounded only on one side of the CRR, at the interface, at the

opposite side at a distance of about 2 nm from the interface there is no local

immobilization of the polymer chains anymore and the “next” CRR behaves as

in a bulk liquid. Therefore no significant broadening of the glass transition is

observed for the nanocomposites as it would be expected for a gradual

change of mobility between the interface and the liquid polymer.

If immobilization was due to anchoring polymer molecules at the

nanoparticle surface, this would have significant influence on the way of

devitrification of the immobilized layer. For semicrystalline polymers it is often

argued that devitrification of the RAF occurs gradually above the common

glass transition [55]. Assuming a local immobilization of the polymer

molecules at the nanoparticles surface as the reason for the formation of a

rather thick RAF layer in turn requires a disappearance of the anchoring at the

Discussion 103

nanoparticle surface. No gradual increase in heat capacity (broad second

glass transition) is expected as long as the anchoring persists. Only removing

the anchors will allow the immobilized layer to relax, to devitrify. In order to

check the behavior of the RAF layer and to detect a possible second glass

transition, heat capacity measurements up to the degradation of the polymer

were performed. StepScan DSC was used to obtain precise heat capacity

data up to the beginning of degradation. Measurements were performed up to

degradation temperature but because of the isotherm after each 3 K

temperature step heat capacity could be obtained until the heat flow rate

during the isotherm was not stable anymore.

RAF devitrification is expected to appear as a second glass transition

but there is no steplike or gradual transition towards the liquid heat capacity

observed up to 230°C where the polymer starts to degrade (Fig. (4.24)).

Influence of degradation on the heat capacity determination can be further

reduced by using high heating rates. To shift beginning of degradation to

higher temperatures Hyper DSC measurements at 400 K/min heating rate

were performed [208]. At 400 K/min heating rate the polymer does not

degrade up to about 330 °C. But there is again no second glass transition

visible below 330 °C (Fig. (4.25)). This means that the interaction between the

polymer matrix and the nanoparticles is so strong that heating even up to such

high temperatures is not enough to remove the anchors and to allow

relaxation and devitrification of the RAF.

Plasticization experiments have been performed also to allow detection

of RAF devitrification. The idea behind is to lower the glass transition of the

polymer hoping that the RAF Tg is shifted to lower temperatures as well. The

data obtained could have two possible explanations. First is that the RAF

devitrification is spread over the large temperature interval from the MAF Tg up

to the degradation of the polymer and by the calorimetric methods used during

this work it is not possible to detect it. And the second is that the devitrification

of RAF in polymer nanocomposites does not occur before degradation of

polymer.

Interaction between PMMA and SiO2 at the interface of the

nanoparticles is expected to be weaker than a covalent bond, which is present

104 Chapter 5

in semicrystalline polymers if a polymer chain goes from a rigid crystal

lamellae through the interface and the immobilized layer to the mobile

amorphous polymer and eventually back into the same or another lamellae. If

the non-covalent bond between the inorganic nanoparticle and the PMMA

does not allow devitrification before degradation of the polymer occurs it is

very unlikely that in semicrystalline polymers the RAF devitrifies as long as the

polymer chains are covalently anchored to the rigid polymer crystals. Most

likely the polymer crystals must melt before the RAF can relax and devitrify.

This was demonstrated [57] for semicrystalline iPS by applying ultra fast

scanning rates to suppress reorganization of the crystals.

6. SUMMARY

The existence of an immobilized fraction in PMMA SiO2

nanocomposites was shown on the basis of heat capacity measurements at

the glass transition of the polymer. The results were verified by enthalpy

relaxation experiments below the glass transition. The immobilized layer is

about 2 nm thick at low filler content, if no agglomeration is present. At higher

filler content agglomeration becomes important and the layer thickness can

not be determined correctly.

The immobilized fraction in nanocomposites can not only be

determined from heat capacity as it is common for the rigid amorphous

fraction in semicrystalline polymers. The thickness of the layer is also similar

to that found in semicrystalline polymers and independent from the shape of

the nanoparticles.

Nanocomposites offer a unique opportunity to study the devitrification of

the immobilized fraction (RAF) without interference of melting of crystals as in

semicrystalline polymers. It was found that the interaction between the SiO2

nanoparticles and the PMMA is so strong that no devitrification occurs before

degradation of the polymer. No gradual increase of heat capacity or a

broadening of the glass transition was found by SSDSC up to the degradation

of the polymer and by high rate DSC and even by lowering the glass transition

of MAF by plasticization. The cooperatively rearranging regions (CRR) are

either immobilized or mobile. No intermediate states are found.

The results obtained for the polymer nanocomposites support the view

that the reason for the restricted mobility must disappear before the RAF can

devitrify. For semicrystalline polymers this means that rigid crystals must melt

before the RAF can relax. Only for semicrystalline polymers with significant

chain mobility inside the crystals RAF may devitrify before melting.

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APPENDIX

A1. Specific heat capacity data corrected

40 50 60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3 ATHAS data PMMA pure (9) 9 wt% SiO2 (10) 35 wt% SiO2 (11) 53 wt% SiO2 (12)

Step

Scan

Hea

t Cap

acity

in J/

K*g

poly

mer

Temperature in °C Figure A1.1. Specific heat capacity as a function of sample temperature for

PMMA SiO2 nanocomposites prepared by classical emulsion

polymerization (recalculated to the polymer mass and fitted to

ATHAS data in solid state)

40 50 60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3 ATHAS data PMMA pure (1) 4 wt% SiO2 (2) 15 wt% SiO2 (3) 22 wt% SiO2 (4) 30 wt% SiO2 (5) 47 wt% SiO2 (6) 66 wt% SiO2 (7) 73 wt% SiO2 (8)

Spec

ific

Hea

t Cap

acity

in J/

K*g

poly

mer


PMMA SiO2 nanocomposites prepared by microemulsion



A2

40 50 60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3 ATHAS data PMMA pure (13) 10 wt% SiO2 (14) 25 wt% SiO

2 (15)

40 wt% SiO2 (16) 48 wt% SiO2 (17)

Step

Scan

Hea

t Cap

acity

in J/

K*g

poly

mer


PMMA SiO2 nanocomposites prepared by solution method using

PMMA synthesized by microemulsion polymerization

(recalculated to the polymer mass and fitted to ATHAS data in

solid state)

40 50 60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3

ATHAS data PMMA pure (33) 5 wt% Al

2O

3 (34)

10 wt% Al2O3 (35) 15 wt% Al

2O

3 (36)

20 wt% Al2O3 (37)

Step

Scan

hea

t cap

acity

in J/

K*g

poly

mer


PMMA Al2O3 nanocomposites prepared by shearmixing

(recalculated to the polymer mass and fitted to ATHAS data in

solid state)

A3

40 50 60 70 80 90 100 110 120 130 140 150 160 1701.3

1.4

1.5

1.6

1.7

1.8

1.9

2.0

2.1

2.2

2.3 ATHAS data PMMA pure (27) 11 wt% Laponite RD (28) 14 wt% Laponite RD (29) 27 wt% Laponite RD (30) 42 wt% Laponite RD (31) 59 wt% Laponite RD (32)

Step

Scan

Hea

t Cap

acity

in J/

K*g

poly

mer


PMMA Laponite RD nanocomposites prepared by microemulsion



A2. The calorimetric data from annealing experiments

-20 0 20 40 60-0.04

-0.02

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.16

0.18

0.20

Ta = 25°C

Ta = 20°CTa = 10°C

Ta = 0°C

Ta = -20°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = -40°C

Ta = 25°C

Figure A2.1. Excess specific heat capacity as a function of temperature for

PBMA pure (38); the annealing temperatures are given for most

of the curves assigned with the same colour as for the curve,

annealing time is 10 h

A4

-20 0 20 40 60-0.04

-0.02

0.00

0.02

0.04

0.06

0.08

0.10

0.12

0.14

0.16

0.18

0.20

Ta = 25°C

Ta = 20°CTa = 10°C

Ta = 0°C

Ta = -20°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = -40°C

Ta = 25°C


PBMA filled with 20 wt% SiO2 (40); the annealing temperatures

are given for most of the curves assigned with the same colour

as for the curve, annealing time is 10 h

60 80 100 120 140

0.0

0.1

0.2

0.3

0.4

0.5

Ta = 110°C

Ta = 90°C

Ta = 70°CTa = 50°CTa = 30°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = 100°C


pure PMMA (18); the annealing temperatures are given for most

of the curves assigned with the same colour as for the curve,

annealing time is 10 h

A5

60 80 100 120

0.0

0.1

0.2

0.3

0.4

0.5

Ta = 90°CTa = 70°C

Ta = 110°C

Ta = 105°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = 100°C


PMMA filled with 35 wt% SiO2 (21); the annealing temperatures

are given for most of the curves assigned with the same colour

as for the curve, annealing time is 10 h

60 80 100 120 140-0.2

-0.1

0.0

0.1

0.2

0.3

0.4

0.5

0.6

0.7

0.8

0.9

1.0

Ta = 100°CTa = 80°C

Ta = 95°C

Exce

ss sp

ecifi

c he

at c

apac

ity in

J/K

*gsa

mpl

e

Temperature in °C

Ta = 90°C

Figure A2.5. Excess specific heat capacity as a function of temperature for PS

filled with 24 wt% SiO2 (44); the annealing temperatures are

given for most of the curves assigned with the same colour as for

the curve, annealing time is 10 h

A3. RAF layer thickness estimation

For the estimation full deagglomeration of nanoparticles is assumed. The

dimensions of nanoparticles are known. It is also assumed that in 1g sample

A6

(polymer nanocomposite) there are polymerpurep

samplepMAF c

cm

Δ

Δ= MAF, mRAF

immobilized polymer and mnp nanofiller. And according to Eq. (4.1) the mass

of immobilized polymer mRAF in 1 g sample is calculated (Eq. A3.1).

npMAFRAF mmm −−= 1 (A3.1)

From the mass fractions the volume fractions could be found by dividing by

their densities. To determine the volume VRAF of the immobilized polymer it is

assumed that at glass transition density of the polymer is not changed

extremely and equals the density ρpolymer at 25°C which is known from the

technical data of the product (ρRAF = ρMAF = ρpolymer 1.2 g/cm3 for PMMA, ρnp

= 2.4 g/cm3 for SiO2).

polymer

RAFRAF

mVρ

= (A3.2)

For the whole volume fraction of nanoparticles in 1 g nanocomposite sample

the volume of all nanoparticles VΣnp equals

np

npnp

mV

ρ=∑ . (A3.3)

To find the number of the nanoparticles nnp in 1 g sample the volume of a

single nanoparticle Vnp is needed. Considering the SiO2 nanoparticles as

spheres, its volume is estimated as

3

34

npnp RV π= , (A3.4)

where Rnp is a radius of a nanoparticle (5 nm). The number of nanoparticles is

approximated by

np

npnp V

Vn ∑= (A3.5)

The volume of immobilized polymer, VRAF per np, which covers one nanoparticle

is:

np

RAFnpperRAF n

VV = (A3.6)

A7

Figure A3.1. Schematic representation of (a) spherical SiO2 and (b) Laponite

RD nanoparticles covered by immobilized polymer

Assuming a coating of the nanoparticle by the RAF layer the total radius of the

rigid particle (nanoparticle + RAF) can be obtained from

3

4)(3

πnpnpperRAF

RAFnp

VVR

+=+ (A3.7)

Taking into account Fig. (A3.1a) and Eq. (A3.7) for Rnp+RAF determination, the

RAF layer thickness dRAF on the spherical nanoparticles can be finally

estimated as

npRAFnpRAF RRd −= + (A3.8)

dMAF

Rnp+RAF

Immobilized polymer

Nanoparticle

(a)

dRAF

dnp

d

(b)

Rnp

dRAF

A8

For the Laponite RD filled systems the assumption that platelets are exfoliated

and covered from both sides by polymer is made (Fig. (A3.1b)). The volume of

the “sandwich” model VS of Laponite RD filled systems is equal to:

)2( MAFRAFnpSS dddAV ++= (A3.9)

where AS is the area of the platelet, dnp, dRAF and dMAF are the thicknesses of

the nanoparticle, immobilized and mobile polymer layer respectively. On the

other hand VS is also available from (Eq. (A3.10)) through volume fractions of

MAF (ϕMAF), RAF (ϕRAF) or nanoparticles (ϕnp).

np

Snp

RAF

SRAF

MAF

SMAFS

AdAdAdV

ϕϕϕ===

2 (A3.10)

To obtain the volume fractions the weight fractions, which are available from

the Δcp data (mMAF, mRAF and mnp), are divided by the densities (Eq. (A3.11)).

np

npnp

RAF

RAFRAF

MAF

MAFMAF

mmmρ

ϕρ

ϕρ

ϕ === ; (A3.11)

The density of PMMA at 25 °C is used as ρpolymer for MAF and RAF, and

nanofiller density ρnp is taken equal to that of SiO2. Consequently the RAF

layer thickness for Laponite RD based nanocomposites is estimated by

Eq. (A3.12).

2SRAF

RAFV

dϕ

= (A3.12)

where VS is determined from the last term in A3.10 with dnp = 1 nm. The

results for PMMA / SiO2 and PMMA / Laponite RD nanocomposites are given

in Fig. (5.3).

Liste der Veröffentlichung

Publikationen

1. Davtyan, S.P., Tonoyan, A.O., Sargsyan, A.G., Schick, C., Tataryan,

A.A. (2007). “Physical-mechanical, Thermophysical and

Superconducting Properties of Polymer-Ceramic Nanocomposites.” J.

of Materials Processing Technology, submitted

2. Tonoyan, A.O., Poghosyan M.G., Sargsyan, A.G., Schick, C., Davtyan,

S.P. (2006) “I. Crystallization kinetics under nonisothermal

polymerization conditions” Izvestija NAS RA and SEUA, V. 59, N 2, p.

193

3. Davtyan, S.P., Tonoyan, A.O., Sargsyan, A.G., Schick, C. (2007). “II.

Crystallization kinetics under nonisothermal polymerization conditions.”

Advanced Materials and technologies, Proceedings of the International

Conference, Tbilisi 10-11 May, 2006; Nova Science Publishers, Inc.,

Ney York, accepted

4. Sargsyan, A.G., Tonoyan, A.O., Davtyan, S.P., Schick, C. (2007). “The

amount of immobilized polymer in PMMA SiO2 nanocomposites

determined from calorimetric data.” European Polymer Journal, in press

Tagungsbeiträge

1. Manukyan, L.S., Tonoyan, A.O., Sargsyan, A.G., Davtyan, S.P.

“Research of Stationary Area Frontal Polymerization of the

Metallomonomers.” (talk) Enikolopyan readings, International Scientific

Conference of SEUA, October 7 – 9 (2003), Yerevan, Armenia; a member

of International Organization Committee, Youth scientific committee.

2. Hayrapetyan, S.M., Sargsyan, A.G., Tonoyan, A.O., Davtyan, S.P.

“Intercalation in Superconducting Polymer Nanocomposites.” (talk)

Enikolopyan readings, International Scientific Conference of SEUA, March

10 – 12 (2004), Yerevan, Armenia; a member of International

Organization Committee, Youth scientific committee.

3. Sargsyan, A.S., Thomas, Se., Wurm, A., Thomas, Sa., Schick, C.

European "Rigid Amorphous Fraction of Polymer Nanocomposites and

Semicrystalline Polymers." (poster) Conference Calorimetry and Thermal

Analysis for Environment, ECCTAE 2005, September 6 - 11 (2005);

Zakopane, Poland.

4. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. „Rigid

Amorphous Fraction in Polymer Nano-Composites“ (poster) DPG

Frühjarstagung - CMD21, March 26 – 31 (2006), Dresden, Germany.

5. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Relaxation of

rigid amorphous phase in polymers.” 9th Lähnwitzseminar on Calorimetry,

May 29 - June 1 (2006), Rostock, Germany.

6. Sargsyan, A.S., Wurm, A., Tonoyan, A.O., Davtyan, S.P., Schick, C.

“Influence of Plasticizer on Glass Transition of Systems Showing a Rigid

Amorphous Fraction.” (poster) Thermo international 2006, July 30 -

August 4 (2006), Boulder, Colorado.

7. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Rigid

Amorphous Fraction in Polymer Nano-Composite.” (poster) Thermo

international 2006, July 30 - August 4 (2006), Boulder, Colorado.

8. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Thermal

Characterization of PMMA SiO2 Nano-Composites Prepared by Different

Methods.” (talk) North American Thermal Analysis Society (NATAS) 34th

Annual Conference, August 6 - 9 (2006), Bowling Green, Kentucky, USA,

Perkin Elmer Student Award winner.

9. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Rigid

Amorphous Fraction in Polymer Nano-Composites.” (poster) North

American Thermal Analysis Society (NATAS) 34th Annual Conference,

August 6 - 9 (2006), Bowling Green, Kentucky, USA.

10. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “When does the

rigid Amorphous Fraction in Polymer Nanocomposites Devitrify?” (poster)

North American Thermal Analysis Society (NATAS) 34th Annual

Conference, August 6 - 9 (2006), Bowling Green, Kentucky, USA.

11. Sargsyan, A.S., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Calorimetric

Investigation of PMMA SiO2 Nano-Composites Prepared by Different

Methods.” (poster) The 9th European Symposium on Thermal Analysis

and Calorimetry, August 27 - 31 (2006), Krakow, Poland.

12. Sargsyan, A.G., Tonoyan, A.O., Davtyan, S.P., Schick, C. “Thermal

Characterization of PMMA SiO2 Nano-Composites Prepared by Different

Methods.” (talk) Enikolopyan readings, International Scientific Conference

of SEUA, October 4 – 6 (2006), Yerevan, Armenia; a member of

International Organization Committee, Youth scientific committee.

13. Sargsyan, A.G., Tonoyan, A.O., Davtyan, S.P., Schick, C. „Rigid

Amorphous Fraction in Polymer Nano-Composites“ (poster) Enikolopyan

readings, International Scientific Conference of SEUA, October 4 – 6

(2006), Yerevan, Armenia; a member of International Organization

Committee, Youth scientific committee.

14. Sargsyan, A.G., Tonoyan, A.O., Davtyan, S.P., Schick, C. „Quantification

of the Immobilized Fraction in Polymer Nano-Composites“ (poster) 71th

Annual Meeting of the Deutsche Physikalische Gesellschaft - spring

meeting of the Division Condensed Matter, March 26 – 30 (2007),

Regensburg, Germany.

CURRICULUM VITAE

1. Full name: Sargsyan Albert

family first

2. Date and place of birth: June 24th 1980; Yerevan, Armenia

3. Present address: Max-Planck Str. 3A, 18059 Rostock,

Germany

4. Affiliation, title and degree: University of Rostock, Institute of

Physics

PhD student, Master of Science in Chemistry

5. Short scientific biography:

2004 Research engineer “Chemical Technologies”, PhD course

at State Engineering University of Armenia

2002 M. Sci. in Organic Chemistry

Thesis: Liquid phase non-catalytic oxidation of

halogenvinylic compounds by molecular oxygen

State Engineering University of Armenia

2000 Diploma in Chemistry

State Engineering University of Armenia

6. Employment:

2004- University of Rostock, Inst. of Physics PhD Student, Polymer Physics 2000-2004 Centre of Expertise at the Ministry of Justice of Armenia Expert of the laboratory of Investigations by Physical- Chemical methods 2001-2002 Centre of Investigations of Molecule Structure, Institute of

Fine Organic Chemistry, Yerevan, Armenia Operator at the Laboratory of Mass spectroscopy

7. Field of specialization:

Polymer chemistry and technology, calorimetry, polymer

nanocomposites

Aknowledgements

I want to thank my DAAD “Sandwich Program” co-supervisors Prof. Dr.

Anahit Tonoyan, Yerevan, and Prof. Dr. Christoph Schick, Rostock, very much

for giving me the opportunity to carry out this work. The very essential and

fruitful discussions helped me to get a deeper insight into many current

problems of calorimetry and polymer nanocomposites. I would like to thank

Prof. Dr. Sevan Davtyan, Yerevan, for valuable discussions and ideas.

I am very grateful to all my colleges (former and present) for their

support and friendship.

PD Dr. Doris Pospiech and her colleagues from the Leibniz Institute of

Polymer Research, Dresden, I acknowledge for the important contribution to

the characterization and processing of polymer nanocomposite samples.

I am very thankful to International PhD Program – “IPP made in

Germany” (University of Rostock), for the possibility to present my work at

several international conferences. Financial support of through a stipend from

DAAD is gratefully acknowledged.

This work would not be possible without the serious support and help

from my family.

Erklärung

Ich versichere hiermit an Eides statt, dass ich die vorliegende Arbeit

selbstständig angefertigt und ohne fremde Hilfe verfasst habe, keine außer

den von mir angegebenen Hilfsmitteln und Quellen dazu verwendet habe und

die den benutzten Werken inhaltlich und wörtlich entnommenen Stellen als

solche kenntlich gemacht habe.

Rostock,

Albert Sargsyan

quantification of the immobilized fraction in polymer ... of the immobilized fraction in polymer...

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