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Rheological and thermo-mechanical properties
of PLA-based miscible blends and composites
Rheologische und thermomechanische
Eigenschaften von PLA-basierten mischbaren
Blends und Kompositen
Der Technischen Fakultät
der Friedrich-Alexander-Universität
Erlangen Nürnberg
zur
Erlangung des Doktorgrades Dr.-Ing.
vorgelegt von
Xiaoqiong Hao
aus Henan, China
Als Dissertation genehmigt
von der Technischen Fakultät
der Friedrich-Alexander-Universität Erlangen Nürnberg
Tag der mündlichen Prüfung: 08.03.2016
Vorsitzender des Promotionsorgans: Prof. Dr. Peter Greil
Gutachter: Prof. Dr.-Ing. habil. Dirk W. Schubert
Prof. Dr.-Ing. Volker Altstädt
Table of Contents
1. Introduction ......................................................................................................................................... 1
2. Literature review ................................................................................................................................. 3
2.1 Biopolymer and PLA ..................................................................................................................... 3
2.2 Current approach to improve PLA properties ............................................................................... 5
2.2.1 PLA based nanocomposites .................................................................................................... 5
2.2.2 PLA based blends ................................................................................................................... 7
2.3 Effect of particles on polymer blends ............................................................................................ 9
2.3.1 Nanoparticles in the miscible polymer blends ........................................................................ 9
2.3.2 Nanoparticles in the immiscible polymer blends ................................................................. 10
2.4 Rheological properties of polymer composites ........................................................................... 11
2.4.1 Dynamic mechanical experiment ......................................................................................... 12
2.4.2 Creep and creep recovery experiment .................................................................................. 14
2.5 Shape memory polymer ............................................................................................................... 16
2.5.1 Classification of SMPs ......................................................................................................... 17
2.5.2 Structure and mechanism of semi-crystalline SMPs ............................................................ 18
2.6 Entanglements and the tube model .............................................................................................. 21
2.7 Motivation ................................................................................................................................... 26
3. Material and sample preparation ....................................................................................................... 28
3.1 Matrix polymers .......................................................................................................................... 28
3.2 Filler materials ............................................................................................................................. 29
3.3 Sample preparation ...................................................................................................................... 29
3.3.1 Preparation samples for rheological measurements ............................................................. 30
3.3.2 Preparation of cast films for biaxial stretching ..................................................................... 32
3.3.2 Preparation of biaxially stretched films ................................................................................ 32
4. Characterization methods .................................................................................................................. 33
4.1 Analytical characterization .......................................................................................................... 33
4.1.1 Size exclusion chromatography (SEC) ................................................................................. 33
4.1.2 Thermogravimetric analysis (TGA) ..................................................................................... 33
4.1.3 Differential scanning calorimetry (DSC) ............................................................................. 33
4.2 Morphological characterization ................................................................................................... 34
4.3 Rheological characterization ....................................................................................................... 34
4.3.1 Dynamic mechanical thermal analysis (DMTA) .................................................................. 34
4.3.2 Oscillatory shear rheology .................................................................................................... 34
4.4 Shape memory characterization .................................................................................................. 36
II Table of contents
5. PLA/silica composites ....................................................................................................................... 39
5.1 Morphological characterization ................................................................................................... 39
5.2 Thermal behavior ........................................................................................................................ 41
5.3 Rheological investigation ............................................................................................................ 43
5.3.1. Linear viscoelastic region .................................................................................................... 43
5.3.2. Thermal stability .................................................................................................................. 44
5.3.3. Dynamic mechanical experiments ....................................................................................... 45
5.3.4. Creep-recovery experiments ................................................................................................ 48
5.3.5. Zero shear viscosity and steady-state compliance ............................................................... 54
5.4 A model to describe the interactions in PLA/silica composites .................................................. 55
5.4.1 Interaction between silica particles and PLA matrix ............................................................ 55
5.4.2 Interaction between silica particles ...................................................................................... 58
5.5 Conclusions ................................................................................................................................. 62
6. PLA/PMMA blends ........................................................................................................................... 64
6.1 PLA/PMMA 7N blends with different compositions .................................................................. 64
6.1.1 Thermo-mechanical properties of PLA/PMMA blends ....................................................... 65
6.1.2 Melt rheology of PLA/PMMA blends .................................................................................. 71
6.1.3 Interactions of PLA and PMMA via molecular entanglements............................................ 79
6.2 PLA/PMMA 50/50 blends with different molecular structures .................................................. 84
6.2.1 Molecular Characterization of PMMA 6N, 7N and 8N ....................................................... 84
6.2.2 Thermal behavior of PMMA and PLA/PMMA 50/50 blends .............................................. 85
6.2.3 Rheological properties of neat PMMA ................................................................................. 86
6.2.4 Interactions of PLA and PMMA via molecular entanglements in symmetrical PLA/PMMA
blends............................................................................................................................................. 88
6.3 Shape memory property of PLA/PMMA blends and the underlying mechanism ....................... 90
6.3.1 Influence of stretching parameters on the shape memory properties ................................... 91
6.3.2 Influence of molecular structure of PMMA and blend composition on the shape memory
properties ....................................................................................................................................... 98
6.3.3 The shape memory mechanism of PLA/PMMA blend system .......................................... 100
6.4 Conclusions ............................................................................................................................... 105
7. PLA/PMMA/silica nanocomposites ................................................................................................ 107
7.1 Morphological characterization ................................................................................................. 108
7.1.1 Dispersion of nanosilica in PLA and PMMA ..................................................................... 108
7.1.2 Dispersion of nanosilica in PLA/PMMA blends ................................................................ 108
7.2 Preferential adsorption on nanosilica by one of the components of PLA/PMMA blends ......... 111
7.3 Thermo-mechanical properties .................................................................................................. 114
7.3.1 DSC .................................................................................................................................... 114
III
7.3.2 Dynamic mechanical analysis (DMTA) ............................................................................. 115
7.4 Rheological properties of PLA/PMMA/silica nanocomposites ................................................ 116
7.4.1 Oscillatory strain sweep ..................................................................................................... 116
7.4.2 Oscillatory time sweep ....................................................................................................... 117
7.4.3 Oscillatory frequency sweep .............................................................................................. 118
7.4.4 Molecular entanglement ..................................................................................................... 122
7.4.5 Creep and recovery experiment .......................................................................................... 124
7.5 The influence of nanosilica on the shape memory properties of uniaxially stretched PLA/PMMA
blends .............................................................................................................................................. 125
7.6 The shape memory of biaxially stretched films......................................................................... 127
7.7 Conclusions ............................................................................................................................... 133
8. Summary and Outlook ..................................................................................................................... 135
9. Summary (in German) ..................................................................................................................... 139
10. Appendix ....................................................................................................................................... 143
10.1 Reproducibility of rheological measurements ......................................................................... 143
10.2 The melt density of PLA and PMMA at 200 °C ..................................................................... 144
10.3 Thermogravimetric analysis (TGA) of nanocomposites ......................................................... 145
10.4 The stress-strain curves of semi-crystalline polymers during cold or hot-deformation .......... 147
Abbreviations and symbols ................................................................................................................. 149
References ........................................................................................................................................... 152
Acknowledgements ............................................................................................................................. 166
List of Publication ............................................................................................................................... 168
1. Introduction
Nowadays, polymers have been widely used in our daily life and play more important roles in
many application fields due to the unique structure, low density, easy processability and
modification [Millich and Carraher (1977)]. Nonetheless, the inherent properties in their pure
state usually cannot meet the rapid growing demands. With the development of science and
technology, many modification methods have been developed to extend their range of
applications. The often used methods include blending, filling, surface and chemical
modification [Hamielec and Tobita (1992), Meister (2000)].
Blending polymers is the most versatile and economic method to obtain balanced properties
based on two or more polymers [Paul and Barlow (1979), Roland and Ngai (1991)].
Depending on the interactions between the components, the blends can be classified into fully
miscible blends, partially miscible blends and immiscible blends. The strong interaction
between the components in miscible blends will result in good physical properties. Another
usually used modification method is adding fillers into the polymer matrix which can
influence the mechanical, electrical, magnetically properties, especially for fillers in nanoscale.
Polylactide (PLA), as one of the most promising linear aliphatic thermoplastic polyester, has
attracted attentions not only in academia but also in industry [Garlotta (2001), Singh and Ray
(2007)]. It can be synthesized by ring opening polymerization of lactides which are typically
derived from renewable biomass via fermentation [Sawyer (2003)], and be considered as
“Green plastic”. Since PLA is biodegradable, biocompatible and nontoxic to the human body
and the environment [John et al. (2000), Lim et al. (2008), Shao et al. (2013)], it has been
widely used in biomedical, agricultural and industrial fields [Garlotta (2001), Auras et al.
(2004)]. Up to 2010, PLA has owned the second largest consumption volume in all the
bioplastics. Unfortunately, a few drawbacks such as the obvious brittleness, low melt strength,
2 1.Introduction
poor heat resistance limit the expansion and diversification of PLA’s application [Lim et al.
(2008)]. In this work, the properties of PLA were modified by adding silica of different
particle sizes, or alternatively, by blending with PMMA of various molar masses through melt
mixing approach. The influences of silica with various particle size and PMMA on the
thermo-mechanical and rheological properties of PLA were fully investigated.
In addition, the PLA/PMMA blend system is a typical miscible amorphous/semi-crystalline
polymer blend with shape memory potential. To our knowledge, the shape memory properties
of the miscible amorphous/semi-crystalline polymer blends have been extensively studied in
the previous work [Behl and Lendlein (2007)]. However, the underlying shape memory
mechanism needs further investigation. In this work, the entanglement network formed in
miscible PLA/PMMA blend is investigated by a rheological approach, and its influence on the
shape memory performance of PLA/PMMA blends is studied as well. Furthermore, we
systematically studied the influences of stretching parameters, blend composition and
component molar mass on the shape memory properties of PLA/PMMA blends. These results
provide a novel understanding into mechanism underlying shape memory properties for
miscible semi-crystalline/amorphous SMPs.
2. Literature review
2.1 Biopolymer and PLA
Biopolymers are polymers produced by living organisms, and therefore, are biodegradable
and recyclable in the environment [Nair and Laurencin (2007)]. They could be obtained from
renewable resources so it’s also called as “green” polymeric materials. In the past few decades,
extensive research focused on biopolymers of different groups, as well as on their copolymer
and blends. Biopolymers have been well established in the field of medicine [Hollander and
Hatton (2004)], and now they are more and more widely used in economical production [Lim
et al. (2008)]. Due to the strong demand to improve the relationship between environment and
industrial manufacture, biopolymers from renewable resources provide the best way to
maintain sustainable development [Sudesh and Iwata (2008)].
Polylactide (PLA) is one of the most commonly used synthetic aliphatic polyesters that have
the best cost efficiency. In the early research stage of PLA, it is mainly applied for tissue
scaffold or implant devices due to its limited molar mass and high cost [Lim et al. (2008)].
Nowadays, the PLA with high molar mass has been developed and widely used as food
packaging or other industrial products [Auras et al. (2003)].
Figure 2.1 Chemical structures of L-, meso- and D-lactides [Nampoothiri et al. (2010)].
4 2. Literature review
PLA can be synthesized from lactic acid, which is a natural product obtained from
microorganisms through fermentation [Sawyer (2003)]. Lactide, the cyclic dimer of lactic
acid, is formed by the condensation of two lactic acid molecules and gives rise to L-lactide
(LL-lactide), D-lactide (DL-lactide) and meso-lactide (LD-lactide) (as shown in Figure 2.1).
After a ring opening polymerization of lactide by using a catalyst under vacuum or inert
atmosphere, high molar mass PLA (𝑀𝑊 ≥ 100,000) is obtained. By controlling the reaction
time, temperature and catalyst, we can get PLA with different ratio and sequence of L- and D-
lactide units.
Figure 2.2 General structure of PLA [Jamshidi et al. (1988)].
The commercial grades of PLA are copolymer of poly(L-lactide) (PLLA) and poly(D,L-
lactide) (PDLLA), and the general structure of PLA is shown in Figure 2.2. Depending on the
composition of L- and DL-lactide, PLA can crystallize into three forms (, and ) [Auras et
al. (2004)]. PLA polymer with L-enantiomer above 90% tends to form crystallites and reduce
the optical purity [Jamshidi et al. (1988)].
Since PLA is compostable and derived from renewable sources, it is regarded as an ideal
material to fulfill current environmental concerns in view of environmental pollution and
excessive consumption of fossil resources. Generally, unoriented PLA is quite brittle, but
possesses good strength and stiffness. After processing by suitable methods, such as biaxial or
uniaxial stretching, the mechanical properties of oriented PLA are even better than
polystyrene (PS), high density polyethylene (HDPE), polypropylene (PP) [Auras et al. (2005)].
2.1 Biopolymer and PLA 5
However, some shortcomings of PLA hinder its application. Firstly, as a typical of aliphatic
polyester, the melt strength of PLA is relatively poor [Carlson et al. (1999)]. Moreover, the
thermal stability at typical processing conditions is a critical issue for rheological
measurements. The presence of ester groups was responsible for the decrease of PLA molar
mass at high temperature [Murariu et al. (2008)]. In addition, the brittleness and low heat
resistance of PLA also limit its application in industry. Therefore, some modified methods
were proposed to improve PLA’s performance.
2.2 Current approach to improve PLA properties
The primary objective of materials modification is to improve mechanical, thermal properties
and subsequently the flow properties during processing, which could be easily achieved by
blending with other polymers or fillers.
2.2.1 PLA based nanocomposites
Polymer nanocomposites refer to multiphase materials which are formed by polymer and
nanofiller (at least one dimension in the nanoscale, 100nm) [Hussain et al. (2006)]. Various
types of nanofillers have been considered as filler for PLA matrix in order to enhance its
thermo-mechanical and the theological properties, as well to provide functional properties
such as conductivity, fire-resistance and optical property, etc.
PLA/clay nanocomposites are commonly referred to PLA reinforced by aluminosilicate-based
nanofillers, which include layered silicates (montmorillonite), sepiolite and halloysite
nanotubes [Raquez et al. (2013)]. Layered silicate nanocomposites show prominent
thermomechanical, barrier and fire resistance properties even at low filler content [Ray and
Okamoto (2003)]. The improvements of these characteristics by the nanocomposites are
strongly related to the dispersion level of the layered silicates in the polymer matrix.
Therefore, the preparation method is the key factor to influence the final properties.
6 2. Literature review
PLA/cellulose nanocomposites are totally biomaterials as cellulose is the most abundant
biopolymer on earth. Cellulose is a natural polymer, a long chain of linked sugar molecules. It
is an important component of the primary cell wall of green plants, and the basic building
block for many textiles and papers [Atalla (1999)]. Nanocellulose substrates have attracted a
lot of attention in the nanocomposites field due to the excellent properties such as nanoscale
dimensions, low density, unique morphology and good mechanical strength, as well they are
renewable and biodegradable [Raquez et al. (2013)].
PLA/carbonaceous nanocomposites are developed for its exceptional properties in terms of
stiffness and conductivity. Carbon nanotubes (CNT) have gained main interest as nanofiller
due to its exceptional physical properties, high electrical and thermal conductivity [Spitalsky
et al. (2010)]. CNTs are an allotropic form of carbon like diamond, graphite or fullerenes, and
its distribution in the polymer matrix play an important role in its conductivity. CNTs are
usually dispersed in PLA via solvent-evaporation, in situ polymerization and melt blending,
Wu et al. [Wu et al. (2010)] investigated the influence of the aspect ratio of CNTs and the
formation of percolating network in PLA/CNTs nanocomposites on the rheological, electrical
conductivity and mechanical properties of PLA. The results demonstrated that the CNTs with
high or low aspect ratios displayed different structural characteristics in PLA matrix.
PLA/metal and PLA/metallic oxide nanocomposites have attracted interest due to their
distinct optical, magnetic, antibacterial, electrical and catalytic properties, etc. Nanosilver
compounds are usually used for hygienic and healing purposes [Tolaymat et al. (2010)].
PLA/silver nanocomposites are developed mainly for its antibacterial activity. In addition, the
thermal properties and tensile strength are also improved due to the improvement in the
kinetics of PLA crystallization induced by nanosilver.
PLA/silica nanocomposites have been widely studied due to the advantages of silica (SiO2)
such as low cost, great natural abundance, high thermal resistance and surface
2.2 Current approach to improve PLA properties 7
functionalization, etc. Polymer nanocomposites based on silica can be used for many
applications including optical devices, coatings and flame-retardant materials [Zou et al.
(2008)]. In the work of Yan et al. [Yan et al. (2007)], PLA/silica nanocomposites are
synthesized by a sol-gel process. The thermal stability of the samples is improved by silica
loading. In addition, the presence of even small amounts of silica greatly improved the tensile
strength of the samples. Li et al. [Li et al. (2012)] studied the rheology and biodegradation of
melt compounded PLA/silica nanocomposites. A percolated silica network was formed if the
silica loading reached up to 5 wt%. Moreover, the biodegradation rates were enhanced in the
PLA/silica nanocomposites compared to neat PLA.
As discussed above, nanofiller represent an interesting way to improve the properties of PLA
via PLA based nanocomposites. The dispersion of nanofillers in PLA is still a challenge to
achieve the desired performance.
2.2.2 PLA based blends
A blend of two polymers can be characterized as miscible or immiscible, depending on
whether the blend shows a single phase or phase separation. In miscible blends, the two
components form a homogeneous single phase, and the interactions between the component’s
molecules are favorable and result in good physical properties. Miscible blends typically
result in an average of properties of the components, e.g. rheology and appearance.
Immiscible blends usually possess poor physical properties, delaminate upon impact, and
often differ greatly in viscosity, stability or polarity, but may still have very useful properties
[Robeson (2007)].
According to the literature, it is found that just a few polymers are miscible with PLA. Sheth
et al. [Sheth et al. (1997)] found that poly(ethylene glycol) (PEG) with high molar mass
( 𝑀𝑤 = 20,000 g/mol) was miscible with PLA when PEG content was below 50%.
8 2. Literature review
Furthermore, PEG could plasticize PLA, resulting in higher elongations and lower modulus
values. A lower molar mass polyethylene oxide (PEO) (𝑀𝑤 = 300 − 1000 g/mol) is miscible
with PLA as well. However, PLA is immiscible with higher molar mass polypropylene oxide
(PPO).
Polyvinyl acetate (PVA) is also reported to be miscible with PLA [Gajria et al. (1996)] at all
blend ratio. Low level of PVA content (5-30%) can increase the tensile strength and
elongation of PLA probably due to some interactions taking place in that composition region.
A vast difference in the weight loss of pure PLA and 95/5 PLA/PVA blend was observed.
This difference in the thermal degradation was induced by the vast difference in the surface
tension of pure PLA films and the blends.
PLA appears to be miscible with poly(methyl methacrylate) (PMMA), many other acrylates
and copolymers of (meth)acrylates [Eguiburu et al. (1998), Zhang et al. (2003)]. In the blends
of crystalline PLLA and amorphous PDLA with PMMA and poly(methyl acrylate) (PMA)
prepared by solution/precipitation and solution-casting film methods, only one glass transition
temperature 𝑇𝑔 is observed. The values of 𝑇𝑔 follow the Gordon-Taylor theory for miscible
blend systems. In the PLLA/PMMA blend, the crystallization of PLLA is greatly restricted by
the presence of amorphous PMMA. However, for PLLA/PMA blend, the crystallization of
PLLA is largely favored by PMA.
Due to the significant differences in polarity between PLA and other polymers, PLA is
immiscible with polyolefin (polypropylene and polyethylene), styrenic resins (polystyrene
etc.), polycarbonate (PC) and Acrylonitrile butadiene styrene (ABS). Block copolymer
compatibilizers are developed to increase the compatibility between the components in the
blends.
2.2 Current approach to improve PLA properties 9
In this work, miscible PLA/PMMA blends were prepared by melt blending to produce shape
memory polymers. The influence of blend composition on the shape memory performance
will be studied in detail.
2.3 Effect of particles on polymer blends
In the production process of polymeric materials, composites are usually manufactured by
mixing different macromolecules or with solid “fillers” to improve the impact strength,
modulus, processability, conductivity, flammability or appearance [Meijer et al. (1988)]. The
addition of nanofillers can improve the performance significantly due to the small-size effect.
During the processing of polymer blends, the nanofillers can be used as compatibilizers to
improve the stability of the morphology by preventing domain coarsening. Therefore, it is of
great significance to study the how the nanofillers influence the overall phase behavior of the
polymer blends.
2.3.1 Nanoparticles in the miscible polymer blends
According to the previous work [Lipatov (2002), Lipatov et al. (2002), Huang et al. (2005),
Lipatov and Alekseeva (2007)], the addition of filler into a miscible blend can change the
temperature and kinetics of phase separation, or modify the shape of the phase diagram.
Lipatov and co-workers [Lipatov et al. (2002), Lipatov (2006)] found that for nanocomposites
based on blends containing chlorinated polyethylene and copolymers of ethylene with vinyl
acetate, the presence of fumed silica could influence the temperature of phase separation. It
can shift the temperature either up or down, depending on the amount of fillers. When the
filler content is around a value that can transit both components into the state of a border layer,
the phase separation temperature will be increased. At lower filler content, the redistribution
of the blend components may decrease the phase separation temperature. The possible
interpretation for this phenomenon is the change of the thermodynamics of interaction near
10 2. Literature review
the surface of the fillers which is induced by the selective absorption of one of the
components. In addition, a redistribution of the blend components according to their
molecular masses between the boundary region and the matrix happened.
In the work of Huang et al. [Huang et al. (2005)], it was found that the introduction of silica
particles increased the phase separation temperature of PMMA/poly(styrene-stat-acrylonitrile)
(SAN) blends, and the thermodynamic interaction parameter was correspondingly decreased.
The phase stability of PMMA/SAM mixtures was improved by the addition of silica. The
mechanism for this enhancement was related to the preferential adsorption and specific
interactions between fillers and one of the components (PMMA) of the blend.
Composto et al. [Chung et al. (2004), Chung et al. (2007)] found that the incorporation of
silica nanoparticles slowed down the phase separation process in PMMA/SAN blends. It was
found that the nanoparticles were observed to partition into the PMMA-rich phase and stratify
during phase separation. Similar phenomena were observed in PS/PVME
(polystyrene/polyvinyl methyl ether) blends filled with fumed silica [Gharachorlou and
Goharpey (2008)]. Silica nanoparticles acted as an obstacle to the coarsening of the
morphology and segregated in the PVME-rich phase during the phase separation process.
2.3.2 Nanoparticles in the immiscible polymer blends
Adding solid particles to immiscible polymer blends is usually used to improve the
mechanical properties or obtain high electric conductivity [Fenouillot et al. (2009)]. Moreover,
the morphology of the immiscible polymer blends is also significantly influenced.
In the immiscible blend, the distribution of nanoparticles is mainly determined by the physical
interactions between the surface of the nanoparticles and the polymer components, viscosity
ratio of polymers and the compounding conditions. Generally speaking, the particles will
2.3 Effect of particles on polymer blends 11
selectively locate in one of the phase of the immiscible polymer blend and its uneven
distribution can be predicted calculating the wetting coefficient ωa [Fenouillot et al. (2009)],
ωa =γp−A−γp−B
γAB (2.1)
where 𝛾𝑝−𝐴 , 𝛾𝑝−𝐵 and 𝛾𝐴𝐵 are the interfacial tensions between particles and polymer A,
particles and polymer B, polymer A and B, respectively. If 𝜔𝑎˃1 , the particles will be
preferentially distributed in polymer B; when −1˂𝜔𝑎˂1, the particles will be preferentially
distributed at the interface of polymer blend; when 𝜔𝑎˂ − 1 , the particles will be
preferentially distributed in polymer A.
Nanoparticles can also be used a compatibilizer for immiscible blends. One of the early
reports about nanoparticles utilized as compatibilizer is on carbon black dispersed in
elastomers [Callan et al. (1971)]. Elias [Elias et al. (2007)] found that the most pronouncing
compatibilizing effect can be achieved when the particles are present in the interface of two
polymers.
In this thesis, the effect of silica nanoparticles on the phase behavior of miscible PLA/PMMA
blends will be investigated.
2.4 Rheological properties of polymer composites
During the conversion process of PLA and its composites, the rheological properties exhibit a
significant effect on the melt flow of materials. Since the rheological performance of PLA are
highly dependent on molar mass, fillers, processing temperature and shear rate, all these
factors must be taken into account during the optimization of process parameters [Lim et al.
(2008)].
The viscoelastic behavior of polymer melts can be characterized by zero-shear viscosity and
recoverable compliance, which could be determined by dynamic mechanical and creep-
recovery experiments in the linear region of deformation. These characterizations give the
12 2. Literature review
relaxation time for final stress equilibration in the melts and get into insight the interactions
inside [Macosko and Larson (1994)].
2.4.1 Dynamic mechanical experiment
In industrial processes, polymers are usually processed as concentrated solution or melts.
Rheological properties are important in evaluating the processing evolution. The rheology of a
viscoelastic material is intermediate between that of a perfectly elastic solid and a purely
viscos fluid.
In dynamic mechanical measurement, a sinusoidal stress 𝜏(𝑡) is applied to the sample with
the fixed angular frequency 𝜔 and the stress amplitude value 𝜏0, and the response (𝑡) is a
sinusoidal deformation shifted according to /𝜔. The ratio between the viscous and elastic
behavior in polymer melt can be described by the tangent of the phase angle (𝑡𝑎𝑛).
𝑡𝑎𝑛 = 𝐺′′/𝐺′ (2.2)
The complex modulus 𝐺∗(𝜔) can be determined by the ratio of the shear stress amplitude and
the deformation amplitude. The complex viscosity |𝜂∗|(𝜔) can be determined from the
complex modulus from
|𝜂∗(𝜔)| = |𝐺′′(𝜔)|/𝜔 (2.3)
When 𝜔 is getting small, |𝜂∗(𝜔)| becomes independent of 𝜔 . In the terminal regime, the
constant viscosity value is called zero shear viscosity 𝜂0,
lim𝜔→0|𝐺′′(𝜔)|/𝜔 = 𝜂0 (2.4)
The material is said to act as linear viscoelastic when the study of viscoelastic materials is
under very small strain or deformation [Carreau et al. (1997)]. The linear steady-state elastic
compliance 𝐽𝑒0 can be obtained by
2.4 Rheological properties of polymer composites 13
lim𝜔→0|𝐺′′(𝜔)|/𝜔2 = 𝐽𝑒0𝜂0
2 (2.5)
It is well known, that the presence of nanoparticles in polymer matrices can not only change
the physical properties but also can influence the processability of polymer melts significantly
[Bhattacharya et al. (2008)]. Rheology is a powerful tool to investigate the mesoscopic and
microscopic structure of a polymer/filler system, and it can also be used to predict the
processing behavior of various composites [Han (2007)]. Dynamic-mechanical experiments
are constantly used to study the moduli and viscosity of polymer melts over the whole range
of frequencies. For polymer nanocomposites, the storage modulus 𝐺′ and loss modulus 𝐺′′
increase gradually with the increased filler fraction especially at low frequencies. Furthermore,
the moduli become more and more independent on frequency [Münstedt et al. (2010)]. When
the nanofillers fraction exceeds a certain level, the interactions between the particles are very
strong and lead to a solid-like behavior. Cassagnau [Cassagnau (2008)] studied the frequency
dependence of polystyrene/silica nanocomposites and found that percolated silica network
structures formed when the silica loading reaches up to 5 wt%. Wu et al. [Wu et al. (2006)]
found that the percolation threshold φ𝑐 of the PLA/CNTs (carbon nano-tubes)
nanocomposite was about 4 wt%, and the percolated network was very sensitive to both the
quiescent and the large amplitude oscillatory shear deformation. Wu et al. [Wu et al. (2007)]
pointed out that the viscoelastic properties are highly related to the fillers’ dispersion state and
the interactions between the fillers and the polymer matrix. A distinct increase of the moduli
and viscosity was observed in the low angular frequency range with the loading of fillers. In
the paper of Osman et al. [Osman et al. (2004)], this increase was attributed to hydrodynamic
effects caused by the presence of solid particles in the melt stream.
The oscillatory shear measurement can get the short time response of material melts while the
creep recovery test get the long time response [Münstedt et al. (2010), Triebel et al. (2011)].
14 2. Literature review
In order to study the elasticity of polymer melts, creep-recovery experiment up to steady state
was conducted.
2.4.2 Creep and creep recovery experiment
The creep-recovery experiment is the most important rheological tool to measure the elasticity
of polymer melt. The process of creep and creep recovery experiment is shown in Figure 2.3.
In the creep test section, a constant shear stress 𝜏0 is applied to the sample at the creep time
𝑡𝑐𝑟 =0 s and the time-dependent deformation 𝛾(𝑡𝑐𝑟) is measured. The creep compliance
𝐽𝑐𝑟(𝑡𝑐𝑟) can be defined as
𝐽𝑐𝑟(𝑡𝑐𝑟) =𝛾(𝑡𝑐𝑟)
𝜏0= 𝐽0 + 𝜓(𝑡𝑐𝑟) +
𝑡𝑐𝑟
𝜂0 (2.6)
𝐽𝑐𝑟(𝑡𝑐𝑟) is independent of the creep stress 𝜏0 in the linear range of deformation. 𝐽0 is the
instantaneous elastic compliance, 𝜓(𝑡𝑐𝑟) is the creep function and 𝜂0 is the zero shear-rate
viscosity. Here 𝑡𝑐𝑟
𝜂0 is the irreversible viscous term and 𝜓(𝑡𝑐𝑟) is the viscoelastic part. If the
creep time 𝑡𝑐𝑟 is sufficiently long, 𝜓(𝑡𝑐𝑟) and 𝐽0 become negligible, and 𝜂0 can be determined
from experiments conducted at stresses in the linear regime as
lim𝑡𝑐𝑟→∞𝑡𝑐𝑟
𝐽𝑐𝑟(𝑡𝑐𝑟)= 𝜂0 (2.7)
At the time 𝑡𝑐𝑟 =𝑡0, the stress 𝜏0 is set to zero and then the recovery section begins. In this
stage, the recoverable compliance 𝐽𝑟(𝑡𝑟) was defined as
𝐽𝑟(𝑡0, 𝑡𝑟) =𝛾𝑟(𝑡0 ,𝑡𝑟)
𝜏0= 𝐽0 + 𝜓(𝑡0, 𝑡𝑟) (2.8)
𝛾𝑟(𝑡0, 𝑡𝑟) = 𝛾(𝑡0) − 𝛾(𝑡0, 𝑡𝑟) (2.9)
2.4 Rheological properties of polymer composites 15
so in the linear range of deformation, 𝐽𝑟(𝑡𝑟) is a measure of the elasticity of the material
[Triebel and Münstedt (2011)]. If the creep recovery time is long enough in the linear regime,
the steady-state recoverable compliance 𝐽𝑒0
can be obtained:
lim𝑡𝑜→∞𝑡𝑟→∞
𝐽𝑟 (𝑡0, 𝑡𝑟) = 𝐽𝑒0 (2.10)
Figure 2.3 Schematic diagram of a creep-recovery measurement.
Münstedt et al. [Münstedt et al. (2010)] reported that the elastic properties in the linear range
of deformation of a PMMA melt were more significantly influenced by the addition of fillers
than the viscos ones, and this effect was strongest in the steady-state which needs a long time
to reach. Unfortunately, the often used dynamic-mechanical experiments were not very
suitable for this situation due to its short time window. It was also found that the elastic
properties of nanocomposites at long experimental times were a very sensitive tool to get an
insight into the interaction between particles and molecular chains.
Triebel [Triebel et al. (2011)] investigated the influence of matrix polydispersity on the
elasticity of polystyrene (PS) melts filled with nanosilica. It was found that the linear steady-
state recoverable compliance of the broad PS 158K was about one order of magnitude higher
for the narrow aPS. The polydispersity could significantly influence the elasticity of the
16 2. Literature review
matrix. However, the incorporation of nanosilica at the same fraction leads to a more
pronounced increase of elasticity for the narrow sPS compared to the broad PS 158K.
2.5 Shape memory polymer
Blend made from PLA and PMMA is a typical miscible amorphous/semi-crystalline polymer
blends with shape memory potential, which has received increasing interests in recent years.
Understanding the shape memory mechanism would be beneficial for designing novel
polymers with desired shape memory properties.
Among the stimuli sensitive polymers, shape memory polymers (SMPs) have gained
extensive research interest on account of their intrinsic advantages such as low density, large
attainable strain, easy processing and low cost [Xie (2011)], which provide a great potential
for applications in packaging, sensors, switches, smart textiles, drug delivery and medical
devices [Liu et al. (2007), Mather et al. (2009), He et al. (2012), Luo and Mather (2013)].
Generally speaking, SMP can memorize its original shape and, therefore, will return to it from
a temporary shape when exposed to an appropriate stimulus [Fengkui Li et al. (1997)]. At
molecular level, SMP usually consist of two components: a soft switching phase and a hard
stationary phase [Lendlein et al. (2005), Xie (2010), You et al. (2012)]. The former is
responsible for the fixation of the materials’ temporary shape, while the latter determines the
original shape. The molecules in the switching phase are linked to each other by net-points,
which could be of either physical or chemical nature [You et al. (2012)]. When the
temperature surpasses the switching or transition temperature 𝑇𝑡𝑟𝑎𝑛𝑠, the polymer networks
exhibit “super-elasticity”, and the entropy is considered as the driving force for shape
recovery in SMP systems [Hu et al. (2012)].
2.5 Shape memory polymer 17
2.5.1 Classification of SMPs
The reported SMPs include electrically sensitive, light sensitive, pH sensitive, magnetic
sensitive, thermal or moisture induced materials, based on the external stimulus [Hu et al.
(2012)]. According to the nature of the net-points, SMPs can be derived into two main classes:
physical cross-links and chemical cross-links [Meng and Hu (2009), Chang (2012)]. For the
SMPs based on chemical cross-links, the rubbery elasticity derived from covalent cross-links
leads to a great shape memory property, which can be tuned by the degree of covalent
crosslinking [Jung et al. (2010)]. On the other hand, the crystalline or rigid amorphous phase
usually act as hard stationary phase for SMPs based on physical cross-links, and the transition
temperature 𝑇𝑡𝑟𝑎𝑛𝑠 is either a glass transition temperature 𝑇𝑔 or a melting temperature 𝑇𝑚 .
Numerous SMPs have been prepared by physical cross-linking, such as linear block
copolymers [Behl and Lendlein (2007)] or copolyesters [Booth et al. (2006)].
Extensive studies have been carried out to develop new materials with a shape memory
potential and to expand their application. Based on the structural principles of polymers,
SMPs with new macromolecular structures can be synthesized or designed [Behl et al. (2009),
Wu et al. (2013)]. In the work of Lendlein et al. [Lendlein et al. (2001)], an AB-polymer
network showing shapes memory property was first reported. Oligo(ɛ-caprolactone) were
incorporated covalently into the thermoset as a crystallizable triggering segments that could
fix the secondary shape by physical cross-links. This polymer system showed excellent shape-
memory properties with recovery rate above 99% after 3 cycles. Liu et al. [Liu et al. (2006)]
prepared SSPs based on poly(methyl methacrylate)-co-(N-vinyl-2-pyrrolidone) and linear
poly(ethyleneglycol) (PMMA-co-VP/PEG). The shape memory network was prepared by
radical copolymerization of MMA and VP in the presence of linear PEG based on hydrogen-
bonding. Thereafter, multiple types shape memory polymers containing AB-polymer
networks were introduced [Lendlein et al. (2005), Bellin et al. (2007), Li et al. (2014)].
18 2. Literature review
However, the methods of synthesis and modification of the designed networks are proved to
be complicated and inconvenient. To some extent, ease the processing has been an issue for
large-scale utilization of SMPs. Hence, polymer blending offers a simply way to tune the
shape memory property of conventional SSP and fabricate new SMP system. Miscible SMP
blend systems have aroused considerable scientific and practical interest since it possess
favorable phase interaction and good physical properties.
2.5.2 Structure and mechanism of semi-crystalline SMPs
The shape memory mechanism of SMP is significantly influenced by the materials’ molecular
structure [Hu et al. (2012)]. It’s well known that cross-linked polyethylene (PE) is one of the
first SMPs based on the semi-crystalline phase [Ota (1981)]. The crystalline phase with a
crystal melting temperature (𝑇𝑚) serves as switching phase, while the chemically cross-linked
PE network is used to memorize the permanent shape after deformation upon heating. For this
type SMP, the switching temperature is 𝑇𝑚.
Another type of semi-crystalline SMP has a switching temperature related to 𝑇𝑔 of the
amorphous phase. Compounding miscible amorphous/crystalline polymers is proved to be an
attractive method to prepare shape memory polymers that have a single phase and 𝑇𝑔. Mather
et al. [Campo and Mather (2005)] reported two types of miscible amorphous/semi-crystalline
polymer blends: amorphous polyvinyl(acetate) (PVAc)/semi-crystalline PLA and amorphous
PMMA/semi-crystalline polyvinylidene fluoride (PVDF). In this shape memory polymer type,
the crystalline phase served as physical cross-links and the amorphous phase between the
crystals acted as the switching phase. Furthermore, 𝑇𝑡𝑟𝑎𝑛𝑠 of the blends which can be tailored
by the composition is chosen as transition temperature [Liu et al. (2007)]. Furthermore, You
and co-workers [You et al. (2012), You et al. (2012)] investigated the shape memory
mechanism of another miscible amorphous/crystalline polymer blend composed of PVDF and
acrylic copolymer (ACP). It was found that the PVDF crystals worked as fixed phase and the
2.5 Shape memory polymer 19
amorphous phase between the PVDF crystals worked as the switching phase. In addition, the
tie molecules among the fixed PVDF crystalline phase significantly influenced the shape
memory properties. It was found that, the blend with 50 wt% PVDF possessed the best shape
memory properties.
Accordingly, the shape-memory mechanism of semi-crystalline polymers related to 𝑇𝑚 or 𝑇𝑔
are proposed and depicted schematically in Figure 2.4.
Figure 2.4 Schematic diagrams for the shape memory properties of semi-crystalline polymers
related to (a) 𝑇𝑚 and (b) 𝑇𝑔.
In order to fulfill the increasing requirements of developing new SMP with good
processibility and functional features, a new miscible SMP blend system based on semi-
crystalline PLA and amorphous PMMA was prepared by melt mixing. PLA is one of the most
promising biopolymers for its excellent performances such as biocompatibility,
biodegradability and nontoxic to the human body and environment [John et al. (2000)].
20 2. Literature review
Furthermore, PMMA is also a biocompatible polymer which has been used as bone cement
since about half a century [Goncalves et al. (2013)]. Nowadays it is widely applied as
ophthalmic implants [Lloyd et al. (2001), Panahi-Bazaz et al. (2009)]. Therefore, a binary
blend of these two polymers has a potential to be used in biomedical application. As reported
by Samuel et al. [Samuel et al. (2013), Samuel et al. (2014)], melt mixing of semi-crystalline
PLLA with amorphous PMMA results in a miscible blend which has excellent shape memory
property. The mechanism for PLA/PMMA blends is more complex because the crystallinity is
changed with composition. Except the PLA crystals, the molecular entanglements also play an
important role to influence the shape memory behavior. So, in this work we will analyze these
factors in detail.
The shape memory properties of the miscible amorphous/semi-crystalline polymer blends
have been extensively studied in the previous work [Behl and Lendlein (2007)]. To our
knowledge, the reported shape memory mechanisms for miscible semi-crystalline/amorphous
polymer blends are mostly related to the crystallite in the blends, and the crystallinity is
simply suggested to be the dominate factor that influence the shape memory performance
[Campo and Mather (2005), Liu et al. (2007), You et al. (2012)]. In general, the crystals are
considered to serve as the physical cross-links and 𝑇𝑔 of the amorphous phase works as the
critical temperature for triggering the shape recovery. However, a promising shape memory
performance still can be observed when the blends are amorphous, and the molecular
entanglements are believed to serve as physical cross-links [Samuel et al. (2014)]. The
researches have pointed out the switching phase and stationary phase of SMPs, but the
correlation between the shape memory performance and the degree of crystallinity or
molecular entanglement is still unclear. For the miscible amorphous/semi-crystalline polymer
blends, we believe the underlying shape memory mechanism needs further investigation.
2.6 Entanglements and the tube model 21
2.6 Entanglements and the tube model
As mentioned above, the molecular structure of miscible PLA/PMMA blend can significantly
influence the shape memory performance and the processing behavior the melts. Rheology
has been demonstrated to be a suitable tool to provide information about molecular structure
in the polymer melts [Dealy and Larson (2006)].
In concentrated solutions or melts with high molar mass, the flexible polymer chains are
invariably entangled with the neighbors and create topological entanglement junctions [De
Gennes (1979), Edwards (1986)]. That is, the motion of a chain is significantly impeded by
the topological constraints, and its ability to relax after deformation is imposed [Dealy and
Larson (2006)]. These constraints called entanglements are formed in molten state and can be
fixed when the temperature decreases so that the rubbery or glassy state is attained. The
entanglement junctions in the melt are randomly distributed and are constantly formed and
destroyed in the entanglement and disentanglement process [Wu (1989)]. As well known,
chain entanglements play an important role in controlling solid mechanical [Kausch (1987)],
melt rheological [Ferry (1980)], and adhesive properties of polymers [Wu (1982)].
The tube model has been considered as the most established framework to understanding
polymer dynamics in the melt state [De Gennes (1979), Doi and Edwards (1986)]. For the
monodisperse homopolymer melts with linear architecture, the dynamics of long chains are
governed by reptation behavior due to the presence of entanglements while that for short
chains follows Rouse behavior [Rouse Jr (1953), De Gennes (1979)]. For unentangled
polymer melts, the zero-shear viscosity 𝜂0 is found to be proportional to the molar mass 𝑀,
and this relationship is valid for linear polymers of low molar mass (Equation 2.11). When the
molar mass exceed a critical value 𝑀𝑐 , a new relationship is found between 𝑀 and 𝜂0
(Equation 2.2).
22 2. Literature review
𝑓𝑜𝑟 𝑀 < 𝑀𝑐: 𝜂0 ∝ 𝑀 (2.11)
𝑓𝑜𝑟 𝑀 > 𝑀𝑐: 𝜂0 ∝ 𝑀3.4 (2.12)
The major breakthrough in the theory of entangled polymers is Edwards’ tube model of
entangled polymer networks [Doi and Edwards (1986)]. In tube model, the entanglements of a
test chain with the surrounding chains (matrix) are effectively confined to a tube-like region.
The surrounding chains restrict the transverse motion of the test chain, while forces will move
it for the most part along the axis of the tube. Figure 2.5a shows the tube model of the test
chain (red) entangled with surrounding chains.
Figure 2.5 (a) Schematic diagrams for tube model formed by molecular entanglements in
polymers, (b) Entanglements between similar and dissimilar molecular chains in miscible
binary blend.
Cooper-White and Mackay [Cooper‐White and Mackay (1999)] studied the dynamic
viscoelastic behavior of PLLA melts with molar masses ranging from 2,000 to 360,000 g/mol.
It was found that the critical molar mass 𝑀𝑐 for PLLA is approximately 16,000 g/mol. Dorgan
and his group [Dorgan et al. (1999), Dorgan et al. (2000), Dorgan et al. (2005)] investigated
the linear and branched PLAs in the molten state, the results show that for PLA with a 98:2
ratio of L and D enantiomeric monomers, the entanglement molar mass is ca. 9,000 g/mol.
2.6 Entanglements and the tube model 23
Wu and Beckerbauer [Wu and Beckerbauer (1992)] found that the plateau modulus 𝐺𝑁0 and
entanglement molar mass 𝑀𝑒 of PMMA are strongly dependent on tacticity, while they are
independent of molar mass and its distribution. 𝐺𝑁0 and 𝑀𝑒 of PMMA vary from 2.410
5 Pa
and 16400 g/mol for pure isotactic chains to 5.3 105 Pa and 7800 g/mol for fully syndiotactic
chains. Moreover, log 𝐺𝑁0 and log 𝑀𝑒 vary linearly with diad fractions.
In a miscible binary blend formed by high molar mass polymers, there exist interactions
between similar and dissimilar molecular chains. Therefore, three kinds of entanglements are
formed in the blends. Figure 2.5b shows entanglements between similar and dissimilar
molecular chains in miscible binary blend. A double reputation model has been proposed [Des
Cloizeaux (1988)] and applied to miscible blends [Pathak et al. (2004)].
In the work of Wu [Wu (1987)], the effects of molecular structure and specific interchain
interactions on the entanglement network were investigated. The entanglement behavior
between dissimilar chains in a PVDF/PMMA blend is compared with that for similar chains.
The results reveal that, in these blends, the similar chains are more likely to entangle with
each other than dissimilar ones, resulting in a large reduction of zero shear viscosity in the
blends.
The molar mass where the zero shear viscosity changes from a linear dependence on 𝑀𝑤 to
𝑀𝑤3.4 is defined as the critical molar mass for entanglement, 𝑀𝑐 [Dealy and Larson (2006)].
The molar mass between entanglements 𝑀𝑒 , is the most fundamental material parameter to
study the molecular dynamics.
For an elastomer, the equilibrium shear modulus for infinitesimal deformation can be defined
based on the classical theory of rubber elasticity [Ferry (1980)],
𝐺𝑒 = 𝑣𝑅𝑇 (2.13)
24 2. Literature review
where 𝑣 is the number of moles of network strand per unit volume, 𝑅 is the gas constant and
𝑇 is the absolute temperature. The entangled melts can be considered to be a rubber in which
the crosslink network is replaced by the entanglement network. Therefore, Ferry defined the
plateau modulus of entangled melts by
𝐺𝑁0 = 𝑣𝑒𝑅𝑇 =
𝑅𝑇
𝑀𝑒 (Ferry definition) (2.14)
where is the melt density. The entanglement molar mass is given by
𝑀𝑒 = 𝑎
/𝑒 (2.15)
It is worth noting that the definition of 𝑀𝑒 by Equation 2.14 reflects all the relaxation that
occurs in response to the initial stress, except for the extremely short-time glassy modes. One
fifth of the initial stress relaxes before the entanglement network interrupts the process due to
the relatively fast Rouse modes allow re-equilibration of tension along the chains. Thus, the
observed plateau modulus in experiments is expected to be about 4/5 of the definition in
Equation 2.14. Graessley and Fetters [Ferry (1980), Graessley (1980), Fetters et al. (1994)]
modified the definition and gave a new correlation between 𝑀𝑒 and 𝐺𝑁0 by the following
equation,
𝐺𝑁0 = 4𝜌𝑅𝑇/5 𝑀𝑒 (Graessley-Fetters definition) (2.16)
where 𝑅 is the gas constant, 𝜌 is the density.
In order to determine the value of plateau modulus 𝐺𝑁0 , some semi-empirical methods have
been developed to extract 𝐺𝑁0 from the linear viscoelastic experiments [Plazek and Echeverrıa
(2000), Dealy and Larson (2006), Liu et al. (2006)], and a cross-check of all available
methods is the best way to achieve maximum accuracy.
2.6 Entanglements and the tube model 25
Generally, the most common method to calculate 𝐺𝑁0 is measuring the linear viscoelastic
properties by oscillatory shear experiments. Especially for monodisperse polymers with high
molar mass (𝑀𝑤) and narrow molar mass distribution (MWD), 𝐺𝑁0 can be determined by the
value of storage modulus 𝐺′ at the frequency 𝜔𝑚𝑖𝑛 where 𝐺′′ reaches a minimum [Ferry
(1980)]:
𝐺𝑁0 = 𝐺′(𝜔)𝐺′′→𝑚𝑖𝑛𝑖𝑚𝑢𝑚. (2.17)
We call this approach the ‘minimum’ (MIN). Another method named “integral method (INT)”
is derived from the Kronig-Kramers relation for 𝐺′and 𝐺′′ [Sanders and Ferry (1969), Ferry
(1980)], 𝐺𝑁0 can be calculated by numerical intergration over the terminal relaxation peak of
𝐺′′(𝜔):
𝐺𝑁0 =
2
∫ 𝐺′′(𝜔)𝑑𝑙𝑛𝜔
+
− (2.18)
Unfortunately, PLA/PMMA blend is a typical polydisperse system, and the molar mass
difference between the components makes it impossible to get 𝐺𝑁0 from above methods.
Therefore, these methods should be modified and extend to polydisperse polymers [Liu et al.
(2006)], and the modifications of the MIN method and INT method were developed as
follows [Wu (1985), Wu (1989), Wu and Beckerbauer (1992)]
𝐺𝑁0 = 𝐺′(𝜔)𝑡𝑎𝑛→𝑚𝑖𝑛𝑖𝑚𝑢𝑚. (2.19)
𝐺𝑁0 =
4
∫ 𝐺′′(𝜔)𝑑𝜔𝑙𝑛
𝜔𝑚𝑎𝑥
− (2.20)
A semi-quantitative method based on the terminal cross-point of 𝐺′and 𝐺′′ were developed by
Wu [Wu (1989)] and Nobile-Cocchini [Nobile and Cocchini (2001)]. When the value for
𝑀𝑤/𝑀𝑛 is less than 3, 𝐺𝑁0 can be calculated from the crossover modulus 𝐺𝑋 (𝐺𝑋 = 𝐺′ = 𝐺′′)
by the follow equation,
26 2. Literature review
log (𝐺𝑁
0
𝐺𝑋) = 0.38 +
2.63 log (𝑀𝑤/𝑀𝑛)
1+2.45 log (𝑀𝑤/𝑀𝑛) (2.21)
It is worth noting that the method base on crossover modulus to extract the value of 𝐺𝑁0 is a
semi-quantitative method. The plateau modulus obtained from Equation (2.21) is tentative due
to the approximation embedded in the relationships.
2.7 Motivation
Although PLA has modulus and tensile strength comparable to petroleum-based polymers,
some shortcomings such as brittleness, poor melt strength, and low heat and impact resistance
limit its application. Therefore, adding fillers, or blending with other polymers to form
composites is considered as an effective and simple way to extend and improve the properties
of PLA. The main conversion methods for PLA are usually based on melt processing. The
polymers are heated above their melting point, and shaped to the desired form with a cool
process. Therefore, understanding the thermal, crystallization and melt rheological behavior
of the composites is critical to optimize the materials properties. In the first part of this work,
we will study how silica with different particles size influences the thermo-mechanical and
rheological properties of PLA. What kinds of interactions exist in PLA/silica composites?
What’s the relationship between the particles size and the rheological percolation threshold of
PLA/silica composites?
PLA/PMMA blends with various compositions and molar masses of the PMMA used were
prepared by melt mixing. In this section, we will study how PMMA influence the thermo-
mechanical and rheological properties of PLA? How PLA and PMMA molecules interact
with each other in the blends? What’s the shape memory mechanism of PLA/PMMA blends?
What factors could influence the shape memory behaviors of PLA/PMMA blends?
In the third part of this work, nanosilica particles are added into PLA/PMMA 50/50 blends.
We will study how nanosilica particles disperse in the miscible PLA/PMMA 50/50 blends.
2.7 Goals of this work 27
What kinds of interaction exist in PLA/PMMA/silica mixtures? How nanosilica influence the
dynamic heterogeneity and shape memory behaviors of PLA/PMMA blends?
At the last part of this work, biaxially oriented films based on neat PLA, PLA/PMMA blends
and PLA/PMMA/silica nanocomposites were produced through simultaneous stretching at
different temperatures, strain rates and ratios. The influences of stretching parameters on the
shape memory performances of biaxially stretched films are analyzed in this work.
3. Material and sample preparation
3.1 Matrix polymers
In this work, the main matrix polymers are PLA and PMMA. PLA (Ingeo 4032D), containing
2% D-lactic acid and 98% L-lactic acid, was purchased from NatureWorks (USA). It is a
commercial grade designed for biaxially stretched films, having a density of 1.24 g/cm3 at
room temperature and 1.12 g/cm3 at 200 °Ϲ (see the appendix for the determination of melt
density). The weight average molar mass and number average molar mass are 𝑀𝑤= 210.2
kg/mol and 𝑀𝑛=123.5 kg/mol, respectively. Table 3.1 shows some material parameters of the
polymers.
Three different PMMA from Evonik Röhm GmbH (Germany) were chosen as blend partners
for PLA: Plexiglas® 6N (PMMA 6N), Plexiglas® 7N (PMMA 7N) and Plexiglas® 8N
(PMMA 8N). The glass transition temperatures 𝑇𝑔 of three PMMA increase from 6N to 8N.
From the 1H-NMR analysis (500 MHz, CDCl3, 25 °С) [Samuel et al. (2013)], the syndiotactic
sequences rr (42%) and atactic sequences mr (45%) are predominant in PMMA 6N. A similar
tacticity is found for PMMA 7N with the syndiotactic sequences rr (46%) and atactic
sequences mr (43%). For PMMA 8N, the syndiotactic sequences rr (49%) and atactic
sequences mr (42%) are predominant.
Table 3.1 Properties of the polymers used in this work.
Material Product
Density (g/cm3)
𝑴𝒘
(kg/mol) 𝑴𝒘/𝑴𝒏 𝑻𝒈(°С) Room
temperature 200 °Ϲ
PLA
PMMA
Ingeo 4032D
Plexiglas 8N
1.24
1.19
1.12
1.13
210
116
1.75
1.99
60
117
PMMA Plexiglas 7N 1.19 1.13 92 1.93 110
PMMA Plexiglas 6N 1.19 1.13 65 2.07 96
3.3 Sample preparation 29
3.2 Filler materials
Three different types of spherical silica particles were used in this study. For convenience, we
record these three particles as silica 300, silica OX50 and silica 63, respectively. They have
different average primary particle size 𝑑 and specific surface area (SSA, measured by BET-
method). The characters of these silica particles are given in Table 3.2 and Figure 3.1. The
silica particles are all spherical, and all silica particles were used as received without any
further treatments.
Figure 3.1 SEM images of silica particles (a) silica 300, (b) silica OX50, (c) silica 63.
Table 3.2 The characterizations of three kinds of silica particles.
Silica Supplier 𝒅 (𝐧𝐦) SSA (m2/g) Density (g/cm
3)
AEROSIL® 300 Evonik
Industries AG
7 300±30 2.2
AEROSIL®
OX50
Evonik
Industries AG
40 50±15 2.2
TIXOSIL 63 Rhodia Group 9000 55±5 2.2
3.3 Sample preparation
All the composites and blends used in this work are listed in Table 3.3. Before the melt
mixing, the polymers were dried at 80 °С under vacuum for more than 6 h to remove moisture
to prevent hydrolytic degradation of the PLA and PMMA. In addition, the silica particles
were dried in a vacuum oven at 80 °С for more than 24 h to remove moisture.
30 3. Material and sample preparation
Table 3.3 The PLA/silica composites, PLA/PMMA blends and PLA/PMMA/silica
nanocompsoites used in this work.
polymer matrix filler/blend Filler/blend
content
Mixing
temperature (°C)
PLA
silica 300 0, 1.1, 2.8, 5.8,
9.0 vol. % 180 silica OX 50
silica 63
PLA PMMA 7N
100/0, 90/10,
70/30, 50/50,
30/70,10/90,
0/100 (wt. /wt.)
200
PLA
PMMA 6N
PMMA 7N
PMMA 8N
50/50 (wt. /wt.) 200
PLA/PMMA 50/50
silica 300 0, 2, 5, 10 wt. % 200
PLA/PMMA 80/20 silica 300 0, 2 wt. % 200
3.3.1 Preparation samples for rheological measurements
In order to compare the difference of rheological properties of PLA induced by silica particles
of various sizes, an internal mixer (Haake polyDrive, Thermo Scientific, Germany) was used
to prepare the PLA/silica composites with 1.1, 2.8, 5.8 and 9.0 vol. % of silica (silica 300,
silica OX50, silica 63). The melt mixing was carried out at 180°C for 10 min with a rotational
speed of 100 rpm. For comparison, the neat PLA was subjected to the same treatment as the
composite. For rheological measurement, the samples are compression molded after extrusion
to 2 mm thick disk-shape plates with a diameter of 25 mm at 180 °C and 200 bars. Prior to the
blending and the measurements, all the samples were dried in a vacuum oven at least 12h at
80 °C to remove moisture.
PLA and PMMA (PMMA 7N) with different compositions (100/0, 90/10, 70/30, 50/50, 30/70,
10/90, 0/100 by weight) were melt mixed by means of an internal mixer (Haake polyDrive,
Thermo Scientific, Germany) at 200 °C for 10 min with a rotational speed of 100 rpm. Then,
the materials were hot pressed (T= 200 °C) into: sheets and disk-shape plates. The sheets with
a dimension of 85 mm × 85 mm × 0.3 mm were compression molded. After that, rectangular
3.3 Sample preparation 31
films of 25 mm × 5 mm were cut from the central part of the sheets for shape memory and
DMTA tests. The disk-shape plates with a diameter of 25 mm and 2 mm thickness were
prepared for rheological measurements at the same conditions as the sheets.
In order to study the influence of nanosilica on the dynamic rheology and molecular
entanglement of PLA/PMMA blends, PLA/PMMA (PMMA 7N) (weight ratio: 50/50) with
nanosilica (silica 300) contents of 0, 2, 5 and 10 wt% were melt mixed at 200 °C and 100 rpm
for 10 min with an internal mixer (Haake polyDrive, Thermo Scientific, Germany). In order to
compare the interaction between nanosilica with the components, neat PLA, neat PMMA,
PLA/10 wt% nanosilica and PMMA/10 wt% nanosilica were prepared at the same conditions.
The samples were then compression molded to two shapes at 200 °C. Rectangular films with
a dimension of 25 mm × 5 mm × 0.3 mm were prepared for differential scanning calorimetry
(DSC) test and dynamic mechanical thermal analysis (DMTA). Disk-shape plates (25 mm
diameter, 2 mm thickness) were made for dynamic rheological measurements. For
convenience, the processed neat PLA/PMMA blend and PLA/PMMA/silica nanocomposites
are, respectively, designated as P/P/Si x in the following discussion, and x represents the
silica weight content (wt %) in the nanocomposite.
In order to investigate the effect of the molar mass of PMMA on the molecular entanglements
and viscoelastic properties of symmetric PLA/PMMA blends, PLA/PMMA 50/50 blends
(PMMA 6N, PMMA 7N, PMMA 8N) were prepared by melt-blending at 200 °C and 100 rpm
for 10 min by an internal mixer (Haake PolyDrive, Thermo Scientific, Germany), respectively.
Neat PLA and PMMA are also prepared by the same way as reference samples. The samples
were then compression molded to 2 mm thick disk-shape plates with a diameter of 25 mm at
200 °C. Disk-shape plates of neat PLA and PMMA with 25 mm diameter and 1 mm thickness
were also prepared to make symmetrical bilayers based on PLA and PMMA.
32 3. Material and sample preparation
3.3.2 Preparation of cast films for biaxial stretching
Neat PLA, PLA/PMMA 80/20 blend and nanosilica (silica 300, 2wt%) filled PLA/PMMA
80/20 blend were first extruded through a slit die by a twin screw extruder at 200 °C followed
by cooling in a water bath unit, whose temperature was set to 20 °C. The samples were
granulated and then extruded through a wide flat die into cast film with a thickness of 0.3 mm.
The cast films were cut into 85 mm 85 mm specimens to be used in a Brückner laboratory
biaxial stretcher. The nanosilcia used here is silica 300, which is untreated before blending.
3.3.2 Preparation of biaxially stretched films
These cast films were then stretched in the partly molten state at temperatures of 80 and 90 °C
on a Brückner biaxial stretching device. In this work, the simultaneous stretching mode is
used and the strain rate changes from 20%/s to 100%/s. Before the stretching process, the
samples were pre-heated for 40 s to reach the desired temperature in the hot-air oven. Samples
were then stretched and then finally quickly cooled to room temperature. For each test, the
load and time were recorded and converted into stress vs. biaxial draw ratio curves. The
samples were then used for shape memory testing.
4. Characterization methods
4.1 Analytical characterization
4.1.1 Size exclusion chromatography (SEC)
The information of the molar mass 𝑀𝑤 and molar mass distribution 𝑀𝑤/𝑀𝑛 of PLA and
PMMA were obtained by size exclusion chromatography (GPCmax, Malvern). For PMMA,
the measurements were carried out at 25 °C with tetrahydrofuran (THF) and a constant flow
rate of 1 ml/min. In our study, the GPC calibration standard is PS. The lowest value of
polymer nominal 𝑀𝑃 is 580 g/mol, the highest value of 𝑀𝑃 value is 6870000 g/mol, and
𝑀𝑊/𝑀𝑛 is ca. 1.0.
4.1.2 Thermogravimetric analysis (TGA)
In order to determine the thermal stability of samples, thermogravimetric analysis (TGA 2950,
TA Instruments) were carried out under a nitrogen atmosphere. A constant heating rate of
10 °C/min is applied and the weight loss is recorded with a thermo-scale. The samples with an
initial weight of 20 mg to 30 mg were heated up to 500 °C.
4.1.3 Differential scanning calorimetry (DSC)
The thermal behavior is characterized using differential scanning calorimetry (DSC, TA
Q2000, TA Instruments, USA) under nitrogen. All samples were heated from ambient
temperature to 200 °C at a heating rate of 10 °C /min, kept there for 3 min to eliminate the
thermal history, and then cooled down to 20 °C at a rate of 10 °C/min. A second heating run
at the same conditions as the first one was applied in order to determine the glass transition
temperature 𝑇𝑔, taken at the inflection point of heat flow change. The degree of crystallinity
𝑋𝑐 of PLA and its semi-crystalline blends can be calculated by subtracting the cold
34 4. Characterization methods
crystallization enthalpy from the melting enthalpy, taking the concentration of PLA in the
blend (𝑤 ) into account (as shown in Equation 4.1). The melting enthalpy 𝐻𝑓𝑜 of 100%
crystalline PLA is 93 J/g.
𝑋𝑐 = 𝐻𝑓/(𝐻𝑓𝑜𝑤) (4.1)
4.2 Morphological characterization
Field Emission Scanning Electron Microscope (FE-SEM) (LEO 435 VP, Carl Zeiss
Microscopy, Germany) was used to investigate the morphology of silica particles and the
distribution of silica in the polymer matrix, as well the morphology of the fracture surface of
blends and composites. Before SEM observation, the molded specimens were fractured in
liquid nitrogen to get undeformed fracture surfaces and then coated with gold using Sputter
Coater S150B from Edwards.
4.3 Rheological characterization
4.3.1 Dynamic mechanical thermal analysis (DMTA)
Dynamic mechanical thermal analysis (DMTA) was conducted by means of DMTA IV
(Rheometric Scientific, USA) in the tensile mode. The dynamic storage and loss moduli were
determined as a function of temperature from 15 to 150 °C at a frequency of 1 Hz and a
heating rate of 2 °C/min. Herein, the glass transition temperature of the amorphous phase is
determined from the peak in the curve of loss modulus 𝐸′′ (also named -relaxation
temperature).
4.3.2 Oscillatory shear rheology
The rheological measurements were performed under a nitrogen atmosphere using ARG2
rheometer (TA-ARG2, TA Instruments, USA) in plate-plate geometry (25mm diameter, 2mm
4.3 Rheological characterization 35
gap). A new sample was used for each run, and a waiting time of 5 min was applied for each
test. The reproducibility for all rheological experiments in shear was better than ± 5% (as
shown in Appendix).
4.3.2.1 Dynamic mechanical experiment
In order to determine the linear viscoelastic region, dynamic strain sweep was performed first.
Figure 4.1 schematically illustrates the strain sweep in the linear region and nonlinear region.
The viscoelastic response can be quantified by two parameters: elastic storage modulus 𝐺′(𝜔)
and the viscous loss modulus 𝐺′′(𝜔). In the linear regime, the strain amplitude is sufficiently
small and both viscoelastic moduli are independent of strain amplitude. A strain level γ = 1%
was chosen for the linear rheological measurements.
Figure 4.1 Schematic illustration of the strain sweep test at a fixed frequency.
The oscillatory time sweeps at low frequencies were performed to ensure the long time
rheological tests were carried out in linear range, angular frequency of 0.1 rad/s and strain of
1% were adopted. 𝐺′ in the terminal region is very sensitive to detect molecular changes,
corresponding to the thermal stability of the sample. A maximum change in the storage
modulus 𝐺′ of 5 % from the initial value was used to estimate the range of thermal stability.
36 4. Characterization methods
Oscillatory frequency sweeps ranging from 0.1 (or 0.05) to 500 rad/s were performed at
different temperatures for neat PLA and PLA composites. The storage modulus 𝐺′, the loss
modulus 𝐺′′ and complex viscosity |𝜂∗| were recorded as functions of angular frequency.
4.3.2.2 Creep and creep recovery experiment
In order to ensure the rheological tests were performed in the linear regime, the stress
dependence of the creep compliance and recoverable compliance were performed on neat
PLA and its composites in a stress range from 20 to 500 Pa at a temperature of 180 °C or
200 °C.
The creep compliance 𝐽𝑐𝑟(𝑡𝑐𝑟) of PLA and the composites can reach the constant double
logarithmic slope of one in ca. 1000 s, which indicated that after this time the creep
compliance reached the terminal flow zone. In consideration of the thermal stability of PLA, a
creep time 𝑡0=2000 s was chosen. The duration of the recovery section did last twice the time
of the creep section, and the experiments also proved that the recovery time of 𝑡𝑟 = 4000 s
was sufficient to reach the steady state for all the samples.
4.4 Shape memory characterization
The shape memory characterization experiments were carried out by means of DMTA IV
(Rheometric Scientific, USA) in the film tension mode. Rectangular films with the same
dimensions as for dynamic mechanical measurements were used. The whole process of the
shape memory test is shown in Figure 4.2. The sample was first uniaxial stretched to an
extensional strain of = 100% with a strain rate of 0.02 s-1
at the sample’s 𝑇𝑔 . The
deformation was frozen in by quenching the sample to ambient temperature by liquid nitrogen.
The shape recovery potential was quantified by a thermal shrinkage test. The stretched films
were put into an oil bath at a temperature of the sample’s 𝑇𝑔 +10 °C for 120 s, and then
4.4 Shape memory characterization 37
rapidly cooled down to room temperature. The time of 120s is sufficient for complete
recovery.
Two quantities were determined from this test. The shape fixing ratio 𝑅𝑓
𝑅𝑓 =𝐿2−𝐿0
𝐿1−𝐿0 × 100% (4.2)
quantifies the amount of deformation that was lost during the cooling and is, therefore, not
available for the shape recovery process, while the shape recovery ratio 𝑅𝑟
𝑅𝑟 =𝐿2−𝐿3
𝐿2−𝐿0 × 100% (4.3)
is a measure of the effectiveness to recover the initial shape.
Figure 4.2 Schematic diagram for shape memory test of PLA/PMMA blend.
In this work, neat PLA, PLA/PMMA 80/20 blend and PLA/PMMA 80/20 blend filled with
2wt% nanosilica were chosen to produce biaxially stretched films. The cast films were first
cut into 85 mm 85 mm specimens. Same stretching temperature (80 °C/90 °C) were chosen
in this work to draw films at a strain rate of 20%/s up to different stretch ratios .
38 4. Characterization methods
Figure 4.3 Schematic diagram for shape memory test of biaxially stretched films.
In order to measure the shape memory properties of biaxially stretched films, the cast films
were treated under a specific temperature-deformation program as shown in Figure 4.3. The
recovery temperature is 10 °C higher than the stretching temperature. The shape fixing ratio
𝑅𝑓 and shape recovery ratio 𝑅𝑟 can be determined according to Equations (4.2 and 4.3).
5. PLA/silica composites
As discussed in Chapter 2, adding particles into polymers is a commonly used way to improve
the polymers’ mechanical properties and processibility due to the interactions between
particles and polymer matrix. It is well established that the mechanical properties of polymers,
such as stiffness, strength and toughness, are significantly influenced by particles size,
particle/matrix interface adhesion and particle loading [Fu et al. (2008)]. Nofar et al. [Nofar et
al. (2013)] reported the effects of various additives with different sizes on the crystallization
kinetics of PLA, and Katsikis et al. [Katsikis et al. (2007)] studied the thermal stability of
PMMA/silica nano- and microcomposites by dynamic-mechanical experiments. However,
there is little publication focus on the correlation between particle size and rheological
behavior in PLA/silica system. In this work, three types of silica particles from micro- to
nano- were chosen to fill PLA, and the viscous and elastic properties of PLA with different
silica concentrations were fully studied by dynamic-mechanical experiments and creep-
recovery experiments. In addition, a model was proposed to describe the interactions between
polymer matrix and particles. The correlation between particles size and rheological behavior
of PLA/silica composites was fully analyzed.
5.1 Morphological characterization
The morphologies of silica particles and the fracture surfaces of PLA/silica composites with
silica loading of 2.8 vol. % are shown in Figure 5.1. It can be seen that the geometries of three
types of silica particles are similar to spherical shape, and the particle size increases from
silica 300 to silica 63. As for PLA/silica composites, silica particles were detected as white
dots. As shown in Figure 5.1a’and b’, some silica agglomerates with diameters of hundred
nanometers can be observed in the nanocomposites, but the overall dispersion of the
nanosilica is quite good in the matrix. The agglomerates of nanosilica possess a significant
40 5 PLA/silica composites
porous structure, which means the polymer melt can infiltrate the agglomerates and interact
with the surface of nanoparticles [Münstedt et al. (2010)]. Furthermore, note that the
microsilica particles also disperse well in the polymer matrix (as shown in Figure 5.1c).
However, the interaction between microsilica and matrix is relatively weak, and some
particles nearly split away from polymer matrix at the fracture surface.
Figure 5.1 SEM images of silica particles and the fracture surfaces of composites with silica
loading 2.8 vol. %. (a) silica 300 (a’) PLA/silica 300 (b) silica OX50 (a’) PLA/silica OX50
(b’) silica 63 (c’) PLA/silica 63.
5.2 Thermal behavior 41
5.2 Thermal behavior
The thermograms of PLA/silica composites with various silica concentrations obtained by the
second heating scan are shown in Figure 5.2. The neat PLA exhibits a glass transition
temperature 𝑇𝑔 around 60 °C, while the cold crystallization peak 𝑇𝑐 is centered at 112.3 °C.
With the addition of silica particles, there is no change of 𝑇𝑔 can be observed, while the cold
crystallization peak and the melting peak are significantly influenced.
Figure 5.2 DSC curves of neat PLA and PLA/silica composites at heating rate of 10 °C/min.
(a) PLA/silica 300 (b) PLA/silica OX50 (c) PLA/silica 63.
42 5 PLA/silica composites
Two melting peaks corresponding to α and α’ crystals [Pan et al. (2007)] can be clearly
observed between 160 to 170 °C for neat PLA. For neat PLA, this bimodal melting peak is
due to the melting of the imperfect α crystals and the formation of more perfect α’ crystals at
high temperature [Zhang et al. (2014)]. The presence of silica particles reduced the melting
peak of the α crystal, which indicates the perfect degree of crystallization for PLA is
improved. For all the composites, the cold crystallization enthalpy 𝐻𝑐𝑐 is reduced with the
addition of silica, and 𝑇𝑐 shifts to lower temperatures, of the range of 106.0-112.3 °C,
depending on the silica concentration and the particle size. The crystallinity (𝑋𝑐) of PLA can
be calculated by subtracting the cold crystallization enthalpy from the melting enthalpy,
taking the content of PLA in the mixtures (𝑤) into account (as shown in Equation 4.1). The
melting enthalpy of 100% crystalline PLA is taken as 93 J/g [Zhang et al. (2003)]. The values
of 𝑇𝑔, 𝑇𝑐 and 𝑋𝑐 of PLA in PLA/silica composites obtained from the second heating scan are
listed in Table 5.1.
Table 5.1 The values of 𝑇𝑔, 𝑇𝑐, 𝐻𝑐𝑐 and 𝑋𝑐 of PLA in PLA/silica composites in the second
heating scan.
silica concentration
(vol. %) 𝑻𝒈 (°C) 𝑻𝒄 (°C) 𝑯𝒄𝒄 (J/g) 𝑿𝒄(%)
PLA 0 60.0 112.3 26.2 6.5
PLA/silica
300
1.1 60.0 106.0 23.4 10.4
2.8 60.0 106.0 22.9 11.1
5.8 60.0 106.0 19.5 13.9
PLA/silica
OX50
1.1 60.0 107.2 24.9 7.9
2.8 60.0 107.2 22.5 8.5
5.8 60.0 107.2 21.4 10.7
PLA/silica 63
1.1 60.0 111.5 26.2 6.6
2.8 60.0 107.8 21.8 8.7
5.8 60.0 107.2 20.8 10.1
Obviously, the crystallinity of PLA is improved by the addition of silica particles, and 𝑋𝑐
increases with silica loading. The impact of silica 300 on 𝑋𝑐 is more prominent compared to
silica OX50 and silica 63, which is due to the effect of particle size. Therefore, it can be
5.2 Thermal behavior 43
concluded that in PLA/silica composites, silica particles work as nucleating agents and
improve the crystallinity of PLA. At the same filler concentration, the nanosilica with smaller
particle size provides much more nucleating agents, and therefore shows a more remarkable
impact on the crystallization of PLA.
5.3 Rheological investigation
5.3.1. Linear viscoelastic region
The linear viscoelastic region of the PLA matrix and the PLA/silica composites with the
highest silica concentration is determined by a dynamic strain sweep at 180 °C with a
frequency of 1 Hz (ω=6.28 rad/s). Figure 5.3 shows the dependence of storage modulus 𝐺′ on
the strain amplitude γ from 0.01% to 100% for neat PLA and PLA/9 vol. % silica composites
(the maximum silica concentration in this work). Apparently, the linear viscoelastic region of
the composites was reduced by the addition of silica, and the reduction was more pronounced
for silica 300 which has the smallest particle size. Therefore, all following oscillatory shear
rheological measurements were conducted at γ = 1%.
Figure 5.3 Dynamic strain sweeps for neat PLA and PLA/9 vol.% silica composites at 180 °C
with a frequency of 1 Hz (ω=6.28 rad/s).
44 5 PLA/silica composites
5.3.2. Thermal stability
When long measuring time and high temperatures are applied, a precondition for reliable
rheological measurement is the thermal stability of the material. Particularly, PLA is easier to
degrade in comparison with petroleum-derived polymers when exposed to elevated
temperatures for a long time.
A very common way to evaluate the thermal stability of samples in molten state is to measure
the storage modulus 𝐺′ as a function of time at a defined temperature in the terminal region
(𝜔=0.05 rad/s) [Katsikis et al. (2007), Münstedt et al. (2010)]. Figure 5.4 shows the thermal
stability tests of PLA without and with 2.8 vol. % silica at the processing temperature 180 °C.
Samples are regarded to be stable as long as the deviation of 𝐺′(𝑡)/𝐺0′ is smaller than 5% (𝐺0
′
is the storage modulus at 𝑡=0). In this experiment, a moderate silica concentration (2.8 vol. %)
was chosen to investigate the thermal stability of composites due to the increase of 𝐺′(𝑡)/𝐺0′
induced by the particle diffusion at higher silica concentration.
Figure 5.4 Relative change of storage modulus as a function of the residence time at 180 °C
for neat PLA and PLA mixed with 2.8 vol. % silica.
As can be seen from Figure 5.4, the stable time for neat PLA is about 5000 s at 180 °C and a
slight increase (approximately 1000 s) can be found after mixing with 2.8 vol. % microsilica.
5.3 Rheological investigation 45
At the same temperature, the stable time for PLA/nanosilica mixtures significantly increase to
ca. 8000 s. In addition, it is found that the thermal stability of PLA/silica 300 is better than
PLA/silica OX50 nanocomposite, which may arise from the smaller particles size and larger
SSA of silica 300. The results indicate that the addition of silica particles could improve the
thermal stability of PLA, and the particles with smaller size will lead to a greater
enhancement. The results of thermal stability measurements also provide a time limitations
for the long time rheological measurement.
5.3.3. Dynamic mechanical experiments
The presence of various silica particles has different influence on the linear viscoelasticity of
PLA matrix. Figure 5.5 shows the frequency dependence of storage modulus 𝐺′ (Figure 5.5a,
b and c) and complex viscosity |𝜂∗| (Figure 5.5a’, b’ and c’) with various silica types and
silica concentrations at 180 °C. At low frequencies, the neat PLA shows the typical terminal
behavior with the scaling of 𝐺′ ∝ 𝜔2, which is consistent with the linear viscoelastic theory.
For PLA/silica 300 nanocomposites (Figure 5.5a), 𝐺′ at the low-frequency region increases
with the addition of nanosilica prominently and reaches an approximately frequency-
independent plateau at the high concentration (above 2.8 vol. %). At the high concentration
level, PLA/silica 300 nanocomposites exhibit solid-like response in the low frequency region,
indicating the formation of a percolated silica network [Wu et al. (2006)]. It is noteworthy that
the terminal behavior of PLA disappears gradually with increased nanosilica loading, and no
zero shear viscosity could be detected after the formation of interparticle network (as shown
in Figure 5.5a’). In Figure 5.5b and b’, the moduli and viscosity of PLA/silica OX50
nanocomposites exhibit similar response to the increased silica concentration. However, the
impact induced by silica OX50 is less pronounced due to its relatively larger particle size (40
nm). The nanocomposites do not display solid-like response even at the maximal silica
concentration in this work.
46 5 PLA/silica composites
Figure 5.5 Storage modulus (𝐺′) and complex viscosity |𝜂∗| as function of angular frequency
ω for PLA/silica composite with various silica concentrations. (a, a’) PLA/silica 300 (b, b’)
PLA/silica OX50 (c, c’) PLA/silica 63.
The effect of microsilica (silica 63) on the rheological properties is shown in Figure 5.5c and
c’. It is found that the influence induced by microsilica on 𝐺′ and |𝜂∗| is extremely weak in
comparison with nanosilica. With the increased microsilica loading, 𝐺′ varies negligibly in the
5.3 Rheological investigation 47
high frequency region, and just a slight increase is found at the low frequency range. The
terminal region also disappeared with the addition of high content of microsilica. Moreover,
the viscosity of the microcomposite shows slight dependence on the addition of silica due to
the weak interaction between microsilica and PLA matrix.
As discussed above, 𝐺′ in the low frequency region is more sensitive to silica loading. In
order to investigate the effects of different silica particles on linear rheological properties of
PLA, storage moduli of the composites at a fixed frequency of 0.1 rad/s as a function of silica
concentration are plotted in Figure 5.6. The results show that silica 300 lead to a dramatic
increase in the rheological properties at high silica concentrations, and the inflection point can
be considered as the rheological percolation threshold φ𝑐. No inflection point can be observed
for PLA/silica OX50 and PLA/silica 63 system in the whole concentration range.
According the results shown in Figure 5.5 and 5.6, φ𝑐 for PLA/silica 300 nanocomposites is
between 2.8 vol. % and 5.8 vol. %. While φ𝑐 for PLA/silica OX50 and PLA/silica 63 are both
above 9 vol. %. Moreover, the effect of silica OX50 is less effective in comparison with silica
300 due to its relatively larger particle size. The microsilica shows the least impact on the
rheological properties.
Figure 5.6 Storage modulus (𝐺′) of PLA reinforced with various silica particles at a fixed
frequency of 0.1 rad/s as a function of silica concentration: (a) PLA/silica composites below
the rheological percolation threshold φ𝑐, (b) PLA/silica 300 in the whole concentration range.
48 5 PLA/silica composites
5.3.4. Creep-recovery experiments
In order to investigate the melt elasticity of composites and reliably get insight into the
interactions between particles and polymer matrix, creep-recovery experiment has been
proved to be a good choice [Lamnawar et al. (2011), Triebel and Münstedt (2011)].
5.3.4.1. Stress dependence of creep-recoverable compliance
It has been demonstrated that the interaction between polymer matrix and filler is responsible
for the increased retardation time that determines the recoverable compliance [Lamnawar et al.
(2011)]. The applied creep stress could significantly influence the number of adherent
molecules. In order to apply the creep-recovery experiment in the linear region, the following
recoverable compliance should be independent on the creep stress. Therefore, the dependence
of the recoverable compliance on creep stress was evaluated on neat PLA and PLA/silica
composites in a stress range from 20 to 500 Pa at 180 °C. As can be seen from Figure 5.7a, a
double logarithmic slope of 1 is reached in the terminal flow zone for all the creep curves, and
the creep curves and recoverable curves of neat PLA at different stresses overlap each other
perfectly, i.e., the experiments are all performed in the linear range of deformation. The
recoverable compliance of neat PLA is found to be a constant value of Je0 =1.7×10
-5 Pa
-1 at
180 °C, and PLA shows a linear behavior in the whole stress range applied.
5.3 Rheological investigation 49
Figure 5.7 Stress dependence of the creep compliance and recoverable compliance for (a)
neat PLA, (b) PLA/silica 300, (c) PLA/silica OX50 and (d) PLA/silica63 composites with a
filler fraction of 2.8 vol. % at 180 °C.
Similar measurements were performed on PLA/silica composites with a filler concentration of
2.8 vol. %. As shown in Figure 5.7b, c, and d, the creep compliances don’t show any
noticeable dependence on the creep stress, but the recoverable compliances exhibit
remarkable stress dependence. The recoverable compliance curves applied at 20 Pa and 50 Pa
are almost overlap, indicating the creep-recovery experiments are performed at linear region
of deformation. Starting at a stress of 200 Pa, the steady-state recoverable compliances
decrease with increased creep stress. Therefore, a pronounced nonlinearity occurs when the
creep stress is larger than 200 Pa for the filled samples. This result may be induced by the
50 5 PLA/silica composites
disintegration of attached molecules from the particles surfaces [Münstedt et al. (2010)].
Based on these results, a stress of 50 Pa can be chosen for the creep experiments.
5.3.4.2. Creep time dependence of creep-recoverable compliance
Figure 5.8 Creep-recovery compliance as a function of recovery time of PLA/2.8 vol. %
silica 300 composite following various creep times 𝑡𝑐𝑟.
In order to ensure that a steady state of the preceding creep behavior was fully reached, the
recovery compliances of PLA/2.8 vol. % silica 300 composite following different preceding
creep time 𝑡𝑐𝑟 (2000 s, 3000 s, 4000 s) at 180 °C are displayed in Figure 5.8. For a creep time
of 2000 s, the plateau in the recoverable compliance is found to be 0.0073 Pa-1
. An increase of
the creep time (up to 3000 s) results in a plateau of 0.0074 Pa-1
at the same experimental
condition. When the creep time is 4000 s, the plateau is around 0.0078 Pa-1
. The difference
between these plateaus is quite small, indicating that an increase of the creep time doesn’t
change the recoverable compliance. Hence, the creep time 𝑡𝑐𝑟=2000 s is long enough to reach
the steady state.
5.3 Rheological investigation 51
5.3.4.3. Temperature dependence of creep-recoverable compliance
The temperature dependence of creep-recovery compliance for neat PLA and PLA/2.8 vol. %
silica 300 composite are shown in Figure 5.9. Both the creep curves of neat PLA and the
composite are increasing with the increased temperature due to the lower viscosity at higher
temperatures. It evident that the linear steady state recoverable compliance of the neat PLA is
independent of temperature and attains a value of Je0 =1.7×10
-5 Pa
-1. However, for PLA/silica
composites, the steady state recoverable compliance shows a weak temperature dependence.
As shown in Figure 5.9b, the steady state recoverable compliance obtained at lower
temperatures is smaller than that determined at higher temperatures.
Figure 5.9 Temperature dependence of the creep compliance and recoverable compliance for
neat PLA and PLA/silica composites with a filler fraction of 2.8 vol. % at the temperature
range of 180-210 °C.
5.3.4.4. Concentration dependence of creep-recoverable compliance
The creep and recoverable compliance in dependence on the silica concentration are shown in
Figure 5.10. For PLA/silica 300 and PLA/silica OX50 nanocomposites, the creep compliances
decrease observably with the addition of nanosilica when the filler concentration is above 1.1
vol. %, corresponding to the increase of viscosity of the composites. A double logarithmic
52 5 PLA/silica composites
slope of 1, indicating the steady state of creep compliance, is reached in the terminal flow
zone for the creep curves when the filler’s concentration is below φ𝑐. On the other side, when
the nanosilica concentration is above φ𝑐, such as PLA/5.8 vol. % silica 300 (as shown in
Figure 5.10a), there is no steady state reached with the increase of creep time, while the creep
curve approaches closely to a plateau at very long creep time. According to Ferry [Ferry
(1980)], this effect can be explained due to the formation of network phase.
Figure 5.10 Creep-recovery experiments on PLA filled with various volume fractions of
silica (a) silica 63 (b) silica OX50 (c) silica 300.
As for the following recovery stage, when the filler’s concentration is below φ𝑐 , the
recoverable compliance shows a significant increase with the addition of silica and then a
decrease after a critical concentration value. Therefore, it can be concluded that the highest
value of 𝐽𝑒0 of polymer/nanosilica system occurs at the filler concentration around φ𝑐. When
5.3 Rheological investigation 53
the filler’s concentration is above φ𝑐, the difference between the curves of creep compliance
and recoverable compliance is extremely small, indicating the solid-like behavior of the
samples.
For polymer/microsilica system as shown in Figure 5.10c, the curves of creep compliance
display little change with the increase of silica fraction, but the recoverable compliance
significantly increases. The magnitude of the increment is amazing even at low filler
concentration.
5.3.4.5. A comparison between complex viscosity and creep compliance
Figure 5.11 Complex viscosity |𝜂∗| as a function of ω in comparison with 𝑡/𝐽(𝑡𝑐𝑟) in
dependence on creep time 𝑡 at 180 °C.
In Figure 5.11, the magnitude of the complex viscosity as a function of the angular frequency
is compared with 𝑡/𝐽(𝑡𝑐𝑟) as a function of creep time 𝑡 for PLA filled with silica particles.
The same filler fraction of 2.8 vol. %, which is below φ𝑐 for all PLA/silica systems was
chosen. According to Equation (2.7), the quantity of 𝑡/𝐽(𝑡𝑐𝑟) at long enough time is equal
54 5 PLA/silica composites
to 𝜂0 , and this correlation is proved in Figure 5.11. These results also demonstrate the
accuracy and consistency of these two experimental methods [Münstedt et al. (2010)].
5.3.5. Zero shear viscosity and steady-state compliance
The zero shear viscosity is significantly influenced by the addition of fillers [Bhattacharya et
al. (2008)]. As shown in Figure 5.12, a linear relation between 𝑙𝑜𝑔𝜂0 and silica concentration
is found when the filler concentration is below φ𝑐,
𝑙𝑜𝑔𝜂0 = 𝑎 + 𝑏 (5.1)
The slopes 𝑎 for PLA/silica 300, PLA/silica OX50 and PLA/silica 63 composites are 10.6, 5.1
and 0.7, respectively. It can be concluded that the particles with larger size could induce a
slow growth rate of zero shear viscosity.
Figure 5.12 Zero shear viscosity 𝜂0 as a function of filler fraction for PLA/silica composites.
The linear steady state elastic compliance 𝐽𝑒0 is also presented as a function of filler
concentration in Figure 5.13. 𝐽𝑒0 for neat PLA is around 1.710
-5 Pa
-1. When the filler
concentration is below φ𝑐, the elastic properties of the composites increase with the addition
of silica more rapidly for the PLA/nanosilica system in comparison with the PLA/microsilica
system, and the growth rate is strongly depend on the particles size. At the same filler faction
5.3 Rheological investigation 55
of 2.8 vol. %, an increase by a factor of 10 was found for PLA/silica 63, while the increase
factor of PLA/silica 300 and PLA/silica OX50 is 200 and 60, respectively. Compared to the
results shown in Figure 5.12, it is found that 𝐽𝑒0 is distinctly stronger influenced by the
addition of particles than 𝜂0.
Figure 5.13 Steady-state compliance 𝐽𝑒0 as a function of filler fraction for PLA/silica
composites.
5.4 A model to describe the interactions in PLA/silica composites
Understanding the interactions in polymer/particle systems is a key element to reveal the
structure-property correlations for polymer composites. In particular, the flow characteristics
of the composites during processing at molten state are significantly influenced by the
interactions in the materials. In PLA/silica system, the main interactions included particle-
polymer interaction and particle-particle interaction [Asakura and Oosawa (1958), Fu et al.
(2008)]. A theoretical network model is verified in the following section.
5.4.1 Interaction between silica particles and PLA matrix
With the incorporation of silica into the PLA matrix, the interactions between particles and
polymer matrix can significantly influence the thermo-mechanical and rheological properties
of the composite [Sarvestani and Picu (2004), Bhattacharya et al. (2008)]. PLA molecules are
56 5 PLA/silica composites
assumed to attach on the surface of silica particles by physical adhesion forces, which will
hinder the mobility of those PLA molecules and lead to longer retardation times. In
PLA/silica composites, the primary adhesion forces between silica particles and PLA matrix
are the Van-der-Waals forces, which can be enhanced by increasing the particle contact area
[Hays et al. (1996), Quesnel et al. (2002)]. The particle surface area for silica 300, silica
OX50 and silica 63 can be easily calculated as all these silica particles are spherical and
monodisperse. Based on the specific surface area of each particle listed in Table 3.2, the
particle surface areas in m2 per cm
3 of PLA/silica composites can be calculated as a function
of the volume fraction (as shown in Figure 5.14).
Figure 5.14 Specific particles surface area in m2 per cm
3 for PLA/silica composites as a
function of silica concentration.
It is evident that the increase of surface area as a function of concentration for silica 300 is
much stronger than that of silica OX50 and silica 63. Although silica OX50 is in nano-scale,
but its SSA is relatively small. The specific particles surface area for silica OX50 is similar
with that of silica 63, indicating the Van-der-Waals forces in both PLA/silica composites are
similar with each other. However, the viscosity and elasticity for PLA/silica OX50
nanocomposites are much larger than that of PLA/silica 63 microcomposites at the same filler
5.4 A model to describe the interactions 57
concentration (as shown in Figure 5.12 and 5.13). This result demonstrates that the particle
size play a crucial role in the polymer/filler composites.
The existence of hydrogen-bonding interactions between the Si-OH and the C=O of PLA
chains has been reported in the previous work [Wen et al. (2009), Li et al. (2012)] . As shown
in Figure 5.15, silica particles possess Si-OH groups on the surfaces, and PLA has C=O
groups in the molecular chains. Therefore, hydrogen bonding interactions between PLA and
silica may exist, and this assumption is proved by Fourier transform infrared (FTIR) analysis
[Wen et al. (2009)]. All silica particles have a similar chemical structure, but the surface area
per volume unit for microsilica is much lower compared to nanosilica. For the equal filler
fraction, the total specific surface area of nanosilica is relatively larger. Hence, nanosilica can
adsorb more polymer molecules and result in greater interactions in comparison with
microsilica. Although nanosilica will form some loose agglomerates in the polymer matrix,
these agglomerates possess a significant porous structure, which means the polymer melt can
infiltrate the agglomerates and still interact with the surface of nanoparticles [Münstedt et al.
(2010), Triebel (2011)].
Figure 5.15 Schematic illustrations of the hydrogen-bonding interactions between the PLA
and silica particles.
58 5 PLA/silica composites
Therefore, it can be concluded that the interactions between particles and polymer molecules
are mainly determined by the particle size and specific surface area in PLA/silica composites,
and this interaction can significantly hinder the mobility of the attached PLA molecules. The
formation of hydrogen-bonding interactions between the PLA and silica particles increased
the thermal stability of PLA, which can be used to interpret the better stability of PLA/silica
composites in comparison with unfilled PLA.
5.4.2 Interaction between silica particles
Generally speaking, the interaction between small particles is mainly determined by the
attractive forces like Van-der-Waals forces and repulsing forces like electrostatic forces
[Hamaker (1937)]. However, the attractive forces between nanoparticles get very pronounced
due to the large specific surface area, which makes the repulsing forces can be ignored. In this
case, the attractive forces will lead to the formation of agglomerates [Rhodes (1981), Zhang et
al. (2004)]. Moreover, from the chemical point of view, reactions of hydroxyl groups on the
surface of silica particles under the elimination of a water molecule are possible when the
distance between particles is very small [Bakar et al. (2009), Kourki and Famili (2012)]. This
chemical reaction also contributes to the formation of nanosilica agglomerates.
In polymer/nanoparticles mixture, the interactions between particles can be ignored at low
filler contents, which are below the rheological percolation threshold φ𝑐. However, a particle
network will be formed when the filler content exceeds φ𝑐. A theoretical network model was
proposed by Sarvestani [Sarvestani and Picu (2004)], they suggested that the polymeric chains
would bridge neighboring fillers and form a nanoparticle network when the wall-to-wall
distance between fillers was small than the bridging segment. For the nanocomposites with
well dispersed non-agglomerated nanoparticles, the bridging segment is related to the radius
of gyration 𝑅𝑔 of polymer chains.
5.4 A model to describe the interactions 59
In order to study the formation of the silica network, a comparison between the mean particle
distance 𝐷 and the radius of gyration 𝑅𝑔 is carried out for PLA/silica composites with
different particle sizes and particle concentrations.
The inter-particle distance in case of an ideal distribution of spherical primary particles with
different diameters 𝑑 and filler concentrations can be calculated from the Equation (5.2),
according to Refs. [Münstedt et al. (2010), Xu et al. (2010)]. The results were showed in
Figure 5.16.
𝐷 = (√4π
3
3− 2) ∗ 𝑑 (5.2)
The radius of gyration 𝑅𝑔 of PLA molecules can be calculated by the method given by Flory
[Flory and Volkenstein (1969)]:
𝑅𝑔2 =< 𝑠2 >=
1
6< 𝑟2 > (5.3)
< 𝑟2 >= 𝑏02 × 𝑧 × 𝐶 × 𝑃 (5.4)
𝑟 is the chains’ end-to-end distance. Here z is the number of bonds of a monomer unit in the
main chain being 3 for PLA. 𝑏0 is the distance between neighboring atoms per repeat unit in
the main chain, which is made up by C-O and C-C bonds. So the average length of the
skeletal bond 𝑏0 = 2.9 Å. 𝑃 is the degree of polymerization of PLA, which can be calculated
from 𝑀𝑛/𝑀𝑔 . The characteristic ratio 𝐶 is ca. 6.7 for PLA. Consequently, for the PLA studied
𝑅𝑔 =< 𝑠2 >0.5≈ 21 𝑛 𝑚 (5.5)
As shown in Figure 5.16, the inter-particle distance D (ideal dispersion condition) is
calculated and plotted as function of particle fraction. In case of a volume concentration of 2.8
vol. %, particles of 7 nm in diameter would be 23 nm apart from each other, for 40 nm this
60 5 PLA/silica composites
quantity is 132 nm, and for 9 m microsilica is ca. 29.7 m. A comparison between 2𝑅𝑔 and
the inter-particle distance is often used to determine the existence of particle-particle
interactions via attached polymer molecules [Sarvestani and Picu (2004), Münstedt et al.
(2010)]. However, for PLA/silica nanocomposites, the nanosilica agglomerates are hard to
break during melt compounding. Although there are some local particle-particle interactions
in the nanocomposites due to the formation of silica aggregates, a continuous silica network
structure has not formed even the inter-particle distance is smaller than 2𝑅𝑔. According to the
rheological experiments, this concentration (2.8 vol. %) is below the rheological percolation
threshold of these three PLA/silica composite systems, i.e., there is no continuous silica
network structure formed. In comparison with particle-PLA interactions, the particle-particle
interactions under this concentration can be neglected.
Figure 5.16 The modified relationship to determine the rheological percolation threshold φ𝑐.
When the volume concentration increases up to 5.8 vol. %, the ideal inter-particle distance for
silica 300 is ca. 14 nm, which is much less than 2𝑅𝑔 of PLA molecules. As discussed above,
PLA/5.8 vol. % silica 300 nanocomposite exhibits a solid-like behavior at low frequencies,
indicating the formation of a continuous silica network, and the rheological properties are
mainly determined by particle-particle interactions and particle-PLA molecule interactions.
5.4 A model to describe the interactions 61
For PLA/silica OX50 system, although the nanoscale particles dispersed well in the matrix,
the inter-particles distance is larger than 2𝑅𝑔 of PLA molecules even at the highest filler
concentration (9 vol.%) in this work. Therefore, it is unlikely to form a continuous silica
network in this composite. That is to say, the interactions between particles via attached
molecules are impossible, and the rheological properties are related to particles-polymer
molecule interactions, only. In addition, for PLA/microsilica system, the particle-particle
interaction is negligible because of the extremely large particle size. Therefore, just silica
particle-PLA molecules interactions are responsible for the rheological behavior of
PLA/microsilica composites.
Figure 5.17 Schematic representation of the interactions between nanoparticles under ideal
and real distribution.
As shown in Figure 5.17, the bridging segment is around 2𝑅𝑔 for the nanocomposites with
well dispersed non-agglomerated nanoparticles [Sarvestani and Picu (2004)]. However, for
PLA/silica nanocomposites, the nanosilica agglomerates are hard to break during
compounding. Considering the impact of nanosilica agglomerates on the real distribution of
silica nanoparticles, using the comparison between 2𝑅𝑔 and 𝐷 to determine the formation of
continuous silica network is inaccurate. According to the rheological percolation threshold of
PLA/silica300 nanocomposite system (between 2.8 vol. % and 5.8 vol. %.), a modified
62 5 PLA/silica composites
relationship between 𝑅𝑔 and 𝐷 was concluded as shown in Figure 5.16. When inter-particle
distance is larger than 2𝑅𝑔, the interaction between the particles via attached molecules is
insignificant. Therefore, the rheological behavior is merely related to particle-polymer
interactions. In case of the inter-particle distance being smaller than 𝑅𝑔 of PLA molecules, a
continuous silica network will be formed. The interactions between particles are so strong that
the composites act as solid-like rheological behavior at low frequency. The particle-particle
interactions dominate the rheological properties.
A critical situation is that 𝐷 locating between 𝑅𝑔 and 2𝑅𝑔, the interactions between particles
and the rheological properties of nanocomposites are mainly determined by the dispersion of
fillers [Damm et al. (2008), Triebel et al. (2010)]. This model additionally confirms that the
particle size has a significant effect on the rheological properties of PLA/silica composites.
What calls for special attention is the limitation of this model for polymer filled by
microparticles. When the volume of microparticle is dominant in the composites, although
there is no silica network structure formed in this case, a solid-like flow behavior will occur
because of the replacement of flexible polymer chains by rigid particles.
5.5 Conclusions
The thermo-mechanical and rheological properties of PLA/silica composites with various
particle sizes and concentrations were investigated. It was found that silica particles worked
as nucleating agents in the composites and significantly improved the crystallinity of PLA.
The nanosilica shows stronger impact on crystallization compared with microsilica. In
addition, the influence of silica particle size and concentration on the rheological properties of
PLA/silica composites were investigated by dynamical mechanical and creep-recovery
experiments. The results demonstrate that the thermal stability of PLA is improved by the
addition of silica, especially for the silica particles with smaller size. In addition, the
5.4 A model to describe the interactions 63
rheological properties of PLA/silica composites show a strong dependence on the particle size
and concentration. The presence of nanosilica can significantly enhance the moduli and
viscosity of the composites, which is almost not influenced by the microsilica. When the filler
concentration is below the rheological percolation threshold φ𝑐, a linear relationship between
𝑙𝑜𝑔𝜂0 and filler concentration is found. Creep-recovery experiments reveal that the melt
elasticity of the composite is very sensitive to the addition of particles. Even for
PLA/microsilica system, the steady-state compliance 𝐽𝑒0 could be increased dramatically upon
the addition of silica. An enhancement of 𝐽𝑒0 by a factor 10 for PLA reinforced by microsilica
at a filler fraction of 2.8 vol. % is observed. Under the same silica concentration, the increase
factors for PLA/nanosilica OX50 (d=40 nm) and PLA/nanosilica 300 (d=7 nm)
nanocomposites are ca. 60 and 200, respectively. The particles with smaller size could induce
a greater enhancement on the melt elasticity. A model describing the interactions in
PLA/silica system was proposed on the basis of correlation between the radius of gyration of
polymer matrix 𝑅𝑔 and the mean distance between particles 𝐷. The model suggests that when
the mean distance between particles exceeds 2𝑅𝑔of polymer matrix, the rheological properties
are mainly determined by the particles-polymer interactions. Whereas the particle-particle and
particle-polymer interactions are responsible for the rheological properties as the mean inter-
particle distance is below the radius of gyration of polymer matrix.
6. PLA/PMMA blends
Blending polymers is a convenient way to develop polymeric materials with desirable
properties. Suitable blend partners could not only improve the physical properties, but also
create some novel features that are not available in the pure components. As is well known,
PMMA is one of the few polymers that are compatible with PLA. The blends show better
transparency, better mechanical behavior and processability in comparison with neat PLA.
Furthermore, PLA/PMMA blends have a potential to be used in biomedical application or
solar concentrators. Hence, it is of great significance to study the thermo-mechanical and
rheological properties of PLA/PMMA blends.
Accordingly, in this chapter, PLA/PMMA blends with various compositions and molar mass
of PMMA were prepared by melt mixing. The effects of the crystallites and chains
entanglement on the shape memory performances were investigated via thermo-mechanical
and rheological measurements. In addition, the influences of stretch parameters on the shape
properties of PLA/PMMA blends were fully studied.
6.1 PLA/PMMA 7N blends with different compositions
In this section, PMMA 7N is chosen to be blended with PLA with various compositions
(100/0, 90/10, 70/30, 50/50, 30/70, 10/90, 0/100 by weight) by melt mixing. The impact of
PMMA content on the thermo-mechanical and rheological properties of PLA/PMMA blends
will be investigated.
6.1 PLA/PMMA 7N blends with different compositions 65
6.1.1 Thermo-mechanical properties of PLA/PMMA blends
6.1.1.1 Differential scanning calorimetry (DSC)
The miscibility of PLA/PMMA blends in the solid state and the crystallization properties were
firstly investigated by DSC measurements. As shown in Figure 6.1, the neat PLA film
displays a glass transition temperature around 60 °C, followed by a cold crystallization peak
at ca. 112 °C. The neat PMMA film is completely amorphous and its glass transition
temperature is determined as 110 °C. The PLA/PMMA blends show a quite broad transition
characterized by only one single 𝑇𝑔 which locates between 𝑇𝑔 of the individual components
and increases with increased PMMA content (as shown in Table 6.1). It is worth noting that
the temperature span Δ𝑇𝑔 of the glass transition of the blends is much broader than that of the
neat components, which is generally acknowledged as a result of local heterogeneity and
concentration fluctuation in the segmental length scale [Wetton et al. (1978), Roland and Ngai
(1993), Shi et al. (2013)].
Figure 6.1 DSC thermograms in the second heating scans of PLA/PMMA blends.
66 6 PLA/PMMA blends
Comparison to the Predictions of the Lodge−McLeish Model
Figure 6.2 Variation of 𝑇𝑔 with blend composition for PLA/PMMA blends. The red and black
lines represent the fit to Equation (6.3) for PLA and PMMA, respectively. The pink line is 𝑇𝑔
calculated by the Fox’s equation. The solid dots (●) represent 𝑇𝑔 of PLA/PMMA blends
determined by DSC.
In a miscible polymer blend system, the glass transition temperatures of the blends can be
theoretically calculated by the Fox’s equation. As shown in Figure 6.2, 𝑇𝑔 determined by DSC
(blue dots) are different with that calculated by the Fox’s equation (pink dotted line). This
difference is attributed to the dynamic heterogeneity in the molecular length.
In PLA/PMMA blends, the local dynamics of the components’ chains may exhibit different
dependences on temperature and overall composition. In order to take into account the
dynamic heterogeneities in this miscible blend system, a concept of “self-concentration”
which is caused by the intermolecular connectivity is introduced by Lodge and McLeish
[Lodge and McLeish (2000)]. According to the Lodge-McLeish (LML) model, the effective
local concentration 𝑒𝑓𝑓
is given by
6.1 PLA/PMMA 7N blends with different compositions 67
𝑒𝑓𝑓
= 𝑠
+ (1 − 𝑠) (6.1)
where 𝑠 is the “self-concentration” of monomers, and is the average bulk composition. The
determination of 𝑠 is based on the volume actually occupied by a Kuhn length’s worth of
monomers, divided by 𝑉 = 𝑙𝑘3
𝑠
=𝐶𝑀0
𝑛𝑏𝑁𝑎𝑉𝑉 (6.2)
where 𝐶 is a polymer’s specific constant the characteristic ratio, 𝑀0 is the repeating unit
molar mass, 𝑛𝑏 is the number of backbone bonds per repeat unit, is the density, 𝑁𝑎𝑉 is the
Avogadro constant, and volume 𝑉 is assumed to be the volume occupied by the Kuhn length
(𝑙𝑘)’s worth of monomers, expressed by 𝑉 = 𝑙𝑘3.
In a binary polymer blend, a consequence of “self-concentration” is that, A monomers
experience local environments that tend to a rich in A, and the same status for B monomers.
The self-concentrations of components A (PLA) and B (PMMA) are taken as 0.37 and 0.25,
according to the calculation based on Equation (6.2), respectively. In a miscible blend, the
component with higher 𝑇𝑔 exhibits a lower 𝑠 and vice versa. The dynamics of the average
blend composition can be better represented by the dynamics of component with high 𝑇𝑔, as
proposed in a previous study [Lodge and McLeish (2000)].
The effective glass transition temperature 𝑇𝑔,𝑒𝑓𝑓 for each component can be calculated by the
modified Fox’s equation,
1
𝑇𝑔,𝑒𝑓𝑓,𝐴=
𝑒𝑓𝑓,𝐴
𝑇𝑔,𝐴+
1−𝑒𝑓𝑓,𝐴
𝑇𝑔,𝐵 (6.3)
The effective glass transition temperature for the components can be calculated and compared
with the calorimetric 𝑇𝑔 of the blends in Figure 6.2. According the concept proposed by
68 6 PLA/PMMA blends
Lodge and McLeish [Lodge and McLeish (2000)], the prediction of the LML model should
meet the physics that the lower 𝑇𝑔 component (PLA) tend to have segmental dynamic in the
blends closer to its own bulk, weakly dependent on concentration. On the other hand, the
dynamics of the higher 𝑇𝑔 component (PMMA) would be more representative of the average
blend composition [Lodge and McLeish (2000)]. This effect could be demonstrated at high
PMMA content region (>50%). However, when PMMA content is below 50%, the
calorimetric 𝑇𝑔 of the blends are more close to 𝑇𝑔,𝑒𝑓𝑓 of PLA. The plausible reason for this
phenomenon may be due to the restriction of chain mobility in the vicinity of the crystallites
[Ngai and Roland (1993)].
Crystallization of PLA/PMMA blends.
From the DSC results, the two melting peaks of the neat PLA corresponding to α crystal and α’
crystal could be clearly observed at 163 °C and 169 °C, respectively [Pan et al. (2007)].
However, the tendency of PLA to crystallize is obviously suppressed by the incorporation of
PMMA. The crystallinity 𝑋𝑐 of the blend is listed in Table 6.1. It is worth noting that, if the
PMMA content is below 50%, a weak melting peak can be found at 165 °C in the blends.
Consequently, the degree of crystallinity 𝑋𝑐 decreased with the increasing PMMA content.
For PMMA content above 50%, no melting peak was found. Hence, the PLA in these blends
is completely amorphous. As the blends have shown to be miscible, the various polymer
chains (PLA and PMMA) in the amorphous phase are randomly distributed and interact with
each other [Graessley (1965)]. The segmental mobility of PLA chains in the amorphous may
be significantly affected by the interactions between PMMA and vice versa. The interactions
between PMMA and PLA will be discussed in more detail later in this section.
A single 𝑇𝑔 is obtained for all the PLA/PMMA blends, indicating the intimate mixing between
PLA and PMMA chains. There are two elements can explain the suppression of PLA
6.1 PLA/PMMA 7N blends with different compositions 69
crystallinity with the addition of PMMA: i) the incorporation of PMMA into PLA matrix
dilutes the PLA chains in the blends, ii) the molecular entanglements between PLA and
PMMA chains reduced the mobility of PLA chains in the matrix, which can prevent the
alignment of PLA chains. Moreover, the work of Samuel [Samuel et al. (2013)] demonstrated
the absence of any other reactions between PMMA and PLA phase. Therefore, the crystalline
phase of PLA was suppressed when PMMA was introduced.
Table 6.1 Glass transition temperature 𝑇𝑔 , associated broadness Δ 𝑇g and crystallinity 𝑋𝑐 of
PLA/PMMA cast films.
DSC DMTA
PLA/PMMA 𝑻𝒈a
(°C) Δ𝑻𝐠b
(°C) 𝑿𝒄c (%) 𝑻𝐠
d (°C) Δ𝑻𝐠
e (°C)
100/0 60 14 10.8 59 19
90/10 62 17 6.3 61 26
70/30 67 28 3.4 65 36
50/50 75 32 0.1 73 45
30/70 85 36 0 82 50
10/90 100 29 0 98 36
0/100 110 18 0 109 30 a
Evaluated by DSC at the inflection point (second heating scan). b
Evaluated by DSC as 𝑇𝑔−𝑒𝑛𝑑 – 𝑇𝑔−𝑜𝑛𝑠𝑒𝑡 . c
Evaluated by DSC (first heating scan). d
Evaluated by DMTA. e
Evaluated by DMTA as half bandwidth of the 𝑇g peak.
6.1.1.2 Dynamic mechanical analysis (DMTA)
Beside DSC, DMTA is another common method to characterize the glass transition. As
shown in Figure 6.3, a single -relaxation corresponding to the glass-rubber transition of
polymers was observed. It is evident that the -relaxations shifts to higher temperature with
increasing PMMA content, and the broadness of glass transition region of the blends is also
enhanced compared to the neat components. For neat PLA, the -relaxation temperature
determined by DMTA is about 2 °C smaller than its glass transition temperature evaluated by
DSC, but the broadness of -relaxation process is much larger than that measured by DSC.
70 6 PLA/PMMA blends
As listed in Table 6.1, the broadness of glass transition determined by DSC and DMTA
methods exhibit similar tendencies with the variation of PMMA content. These differences
between the values of 𝑇g and Δ𝑇g are induced by the different measuring principles of DSC
and DMTA [Rahman et al. (2007)].
Figure 6.3 The dynamic loss moduli 𝐸′′ as a function of temperature determined at a
frequency of 1 Hz and a heating rate of 2 °C/min.
According to Jordan et al. [Jordan et al. (2014)], two transition regimes can be identified in
the loss modulus curves of neat PMMA. The low temperature process (only partially revealed
in the temperature regime measured) is associated with the β-relaxation, while the one at
higher temperature is associated with the glass transition (-relaxation). As illustrated in
Figure 6.3, the high temperature flank of the β-relaxation at low temperature increases
gradually with PMMA addition. Especially for the blends with PMMA contents above 50%,
the high temperature shoulder of the β-relaxation peak at low temperature range is quite
obvious.
6.1 PLA/PMMA 7N blends with different compositions 71
6.1.2 Melt rheology of PLA/PMMA blends
6.1.2.1 Linear rheological region
In order to determine the linear viscoelastic limits of PLA/PMMA blends, the dynamic strain
sweep measurements were carried out at 200°C and an angular frequency of 6.28 rad/s. As
shown in Figure 6.4, the storage moduli 𝐺′ of various blends increase obviously with the
presence of PMMA due to the enhanced rheological behavior. The linear viscoelastic limits of
neat PLA is around the strain of 50%, and the incorporation of PMMA into PLA matrix
reduced the linear viscoelastic region. Interestingly, the linear regions of PLA/PMMA blends
are smaller than that of the neat components, and PLA/PMMA 50/50 blend shows the
smallest linear limit of 4%. This effect may be related to the self-concentration in miscible
PLA/PMMA blends.
Therefore, the linear viscoelastic properties of PLA/PMMA blends were determined at a
strain of 1% in this work.
Figure 6.4 The storage modulus 𝐺′ of various PLA/PMMA blends as a function of strain
amplitude at 200°C.
72 6 PLA/PMMA blends
6.1.2.2 Thermal stability
Dynamic time sweeps were performed to measure the thermal stability of PLA/PMMA blends
in the linear range of deformation. The degradation of PLA and PMMA during the rheological
measurements should be considered because polymers are quite unstable in the molten state.
Dynamic time sweeps were performed to measure the thermal stability of PLA/PMMA blends
in the linear range of deformation. The thermal stability is evaluated using the storage
modulus 𝐺′ as a function of time in the terminal region (0.05 rad/s) at 200 °C. Samples are
regarded to be stable as long as 𝐺′ doesn’t diverge from the original 𝐺′ by more than 5%. As
shown in Figure 6.5, the time of thermal stability for neat PLA at 200 °C is around 1000 s. At
the same temperature, the addition of PMMA distinctly increases the stabling time of the
blends. For PLA/PMMA 50/50 blend sample, no change of 𝐺′within the stability limit even
after 8000 s is observed.
Figure 6.5 The storage modulus 𝐺′ of PLA/PMMA blends obtained in time sweep at 200 °C.
6.1.2.3 Dynamic mechanical experiments
As shown in Figure 6.6a, the complex viscosities |𝜂| at low frequencies increase with the
addition of PMMA, and a plateau of |𝜂| which is equal to zero shear viscosity 𝜂0 could be
reached in the low frequency range for all the blends. PMMA shows larger viscosity
6.1 PLA/PMMA 7N blends with different compositions 73
compared with PLA at the same temperature, and the viscosities of PLA/PMMA blends are
located between the neat components.
A Cole-Cole curve is useful to obtain some viscoelastic properties of polymers, removing the
effect of frequency. Figure 6.6b presents the 𝜂′′ as a function of 𝜂′ for neat components and
the blends. The characteristic relaxation time 0 can be determined at angular frequency
corresponding to the maximum 𝜂′′ [Dealy and Larson (2006)]. The Cole-Cole plot of each
PLA/PMMA blend displays a smooth and single semicircle like the neat components,
indicating the excellent miscibility of the PLA/PMMA blend melts.
Figure 6.6 (a) The complex viscosity |𝜂| of PLA/PMMA blends as a function of angular
frequency and (b) Cole-Cole plots for PLA/PMMA blends obtained by oscillatory frequency
sweep at 200 °C.
It is worth mentioning the viscosities of the PLA/PMMA blends increase with the addition of
PMMA. According to Haley et al. [Haley and Lodge (2004)], if |𝜂| of the blend can be
described by,
𝜂𝑏𝑙𝑒𝑛𝑑 = 𝑤1|𝜂1
| + 𝑤2|𝜂2| (6.4)
where 𝑤 is the weight fraction of the component, the blend will be considered to be
unentangled (i.e. Rouse like). However, as shown in Figure 6.7, the correlation between the
74 6 PLA/PMMA blends
complex viscosity and PMMA content is not well described by the linear model of Equation
6.4, but rather represented by the exponential form of the complex viscosity. This result
suggests the existence of an entanglement network in the PLA/PMMA blends [Liu et al.
(2002)]. Therefore, the additivity of 𝑙𝑜𝑔𝜂0 not only demonstrates the intimate mixing between
PLA and PMMA chains, but also indicates the existence of molecular entanglement in
PLA/PMMA blends.
Figure 6.7 (a) zero shear viscosity 𝜂0 at 200 °C versus blend composition (b) 𝑙𝑜𝑔𝜂0 at 200 °C
versus blend composition.
In Figure 6.8, the plot of the phase angle as a function of the absolute value of the complex
modulus |𝐺∗| (van Gurp-Palmen (vGP) plot [Van Gurp and Palmen (1998)], is used
additionally to evaluate the miscibility of PLA/PMMA blends. For neat PLA, the values
increased with the reduction of |𝐺∗| towards 90°. For PLA/PMMA blends, the curves of the
blends lay between the curves of the neat components, a slight deviation from the curves of
neat PLA was observed with the increased PMMA content.
6.1 PLA/PMMA 7N blends with different compositions 75
Figure 6.8 (a) Van Gurp-Palmen plots of phase angle versus complex modulus |𝐺∗| of
various PLA/PMMA blends at 200 °C, and (b) Van Gurp-Palmen plots of the PLA/PMMA
50/50 blend at various temperatures between 180 °C and 210 °C.
The influence of temperature on the Van Gurp-Palmen plots of PLA/PMMA 50/50 blend as
an example is shown in Figure 6.8b. The curves of different temperatures perfectly
superimpose, indicating that the blend, as well as all other blends investigated, is thermo-
rheologically simple. This fact verifies the miscibility of two polymers in the molten state
[Ferry (1980)], which is based on the fact that the various relaxation times of
thermorheological simple polymers exhibit a similar temperature dependence [Van Gurp and
Palmen (1998)].
6.1.2.4 Time-temperature superposition (TTS)
Time-temperature superposition (TTS) principle is frequently used to determine the phase
separation temperature of polymer blends and estimate the miscibility of blend in molten state
[Jeon et al. (2000)]. Furthermore, the construction of master curves obtained by TTS principle
can extend the time or frequency range in rheological measurements. Figure 6.9 and Figure
6.10, respectively, show the master curves of the neat polymers and their blends at a reference
76 6 PLA/PMMA blends
temperature 𝑇𝑟𝑒𝑓 =200 °C obtained by horizontal shifting along the angular frequency axis
using the shift factor 𝑎𝑇.
Figure 6.9 Master curves of 𝐺′ , 𝐺′′ , |𝜂∗| and 𝑡𝑎𝑛𝛿 for neat PLA and neat PMMA at a
reference temperature of 200 °C. The rheological data were measured at various temperatures
between 180 and 210°C (□■180 °C, △▲190 °C, ○●200 °C, ▽▼210 °C). 𝛼𝑇 is the shift
factor for constructing the master curves.
It is evident that all master curves follow the TTS principle well in the whole frequency range,
indicating the miscibility of PLA/PMMA blends at the experimental temperature range.
6.1 PLA/PMMA 7N blends with different compositions 77
Figure 6.10 Master curves of 𝐺′, 𝐺′′, |𝜂∗| and 𝑡𝑎𝑛𝛿 for PLA/PMMA blends (70/30, 50/50,
30/70, 10/90) at a reference temperature of 200 °C. The rheological data were measured at
various temperatures between 180 and 210°C (□■180 °C, △▲190 °C, ○●200 °C, ▽
▼210 °C). 𝛼𝑇 is the shift factor for constructing the master curves.
To quantify the temperature dependence of the viscoelastic properties for all blends, the
values of 𝑎𝑇 are presented in an Arrhenius plot in Figure 6.11. All the curves can be
described by the Arrhenius equation as follows [Fesko and Tschoegl (1971)],
𝑙𝑜𝑔𝑎𝑇 (𝑇) =𝐸𝑎
2.303𝑅(
1
𝑇−
1
𝑇𝑟𝑒𝑓) (6.5)
where 𝑅 is the universal gas constant, and 𝐸𝑎 is the activation energy of flow. The 𝐸𝑎 values
for neat PLA is 80 kJ/mol and for neat PMMA is 182 kJ/mol, which are consistent with the
reported values in the literature [Agrawal et al. (1997), Holland and Hay (2002)]. Moreover,
the 𝐸𝑎 values of PLA/PMMA blends linearly depend on the composition as expected for a
miscible system.
78 6 PLA/PMMA blends
Figure 6.11 (a) Arrhenius plot of time-temperature superposition shift factor 𝛼𝑇 for various
PLA/PMMA blends, and (b) Activation energy 𝐸𝛼 versus PMMA content.
6.1.2.6 Creep-recovery experiments
Figure 6.12 The creep compliance and recoverable compliance of various PLA/PMMA
blends at 200 °C.
The creep and recoverable compliance of PLA/PMMA blends with different compositions are
plotted in Figure 6.12. The creep compliance is decreasing with increased PMMA content,
corresponding to the increased viscosity of the blends. All the experiments were performed in
the linear range of deformation, and a linear steady-state recoverable compliance 𝐽𝑒𝑜 is reached
within the chosen recovery time 4000 s. 𝐽𝑒𝑜 for neat PLA is ca. 1.7×10
-5 Pa
-1, and 𝐽𝑒
𝑜 for neat
6.1 PLA/PMMA 7N blends with different compositions 79
PMMA is around 2.9 ×10-5
Pa-1
. The 𝐽𝑒𝑜 values of the blends are located between the neat
components. The differences between these 𝐽𝑒𝑜 values are extremely small, indicating the
intimate mixing between PLA and PMMA chains. The melt elasticity of miscible blend is not
influenced by the composition.
Figure 6.13 Complex viscosity |𝜂∗| as a function of ω in comparison with 𝑡/𝐽(𝑡𝑐𝑟) in
dependence on creep time 𝑡 at 200 °C.
In Figure 6.13, the complex viscosity |𝜂∗| as a function of ω is compared with 𝑡/𝐽(𝑡𝑐𝑟) in
dependence on creep time 𝑡 for PLA/PMMA blends. The zero shear viscosity obtained by
oscillatory shear experiments is equal to that determined by creep compliance. A steady-state
is reached for all the blends, and the creep time 𝑡 to reach this steady state is increasing with
the addition of PMMA.
6.1.3 Interactions of PLA and PMMA via molecular entanglements.
Entanglements between polymer molecules are essential in understanding mechanical [Wu
(1982), Kausch (1987)] and rheological [Ferry (1980)] properties of polymers. The average
molar mass between adjacent temporary entanglement points, 𝑀𝑒 , is one of the most
fundamental material parameters to investigate the structure of molten polymers [Ferry
80 6 PLA/PMMA blends
(1980)]. According to Graessley and Fetters [Graessley (1980), Fetters et al. (1994)], the
entanglement molar mass 𝑀𝑒 can be deduced from the plateau modulus 𝐺𝑁0 ,
𝑀𝑒 = 4𝜌𝑅𝑇/5𝐺𝑁0 (6.6)
where 𝜌 denotes the density, 𝑅 the gas constant and 𝑇 the absolute temperature.
For a miscible binary blend, the subscripts 1 and 2 are used to describe the components PLA
and PMMA, respectively. There are three types of possible entanglements: 1-1, 2-2 and 1-2
(assumed to be equal to 2-1). Herein, 𝑀𝑒1 and 𝑀𝑒2 are defined as the average entanglement
molar mass of the neat components (similar chains), 𝑀𝑒12 is the average entanglement molar
mass of a hypothetical pure component formed by PMMA chains entangled with PLA
molecules (dissimilar chains). Hence, we can deduce the entanglement state in similar and
dissimilar chain by comparing the value of 𝑀𝑒1, 𝑀𝑒2 and 𝑀𝑒12.
Wu [Wu (1987)] showed that the density of contacts between two similar chains is
proportional to the square of volume fraction 𝜙12 or 𝜙2
2 while that between two dissimilar
chains has shown to be proportional to 2𝜙1𝜙2.The total average entanglement density can be
defined as [Wu (1987)]
𝜌
𝑀𝑒=
𝜙12𝜌1
𝑀𝑒1+
𝜙22𝜌2
𝑀𝑒2+
2(𝜙1𝜙2)(𝜌1𝜌2)
𝑀𝑒12 (6.7)
By substituting Equation (6.6) into Equation (6.7), 𝑀𝑒12 is obtained by
𝑀𝑒12 = 2𝜙1𝜙2𝑅𝑇(𝜌1𝜌2)1
2/(𝐺𝑁0 − 𝜙1
2𝐺𝑁10 − 𝜙2
2𝐺𝑁20 ) (6.8)
6.1 PLA/PMMA 7N blends with different compositions 81
Figure 6.14 Frequency sweep of PLA/PMMA 50/50 and PLA/70%PMMA blends at 140 °C.
As shown in Figure 6.9 and 6.10, 𝐺𝑁0 for neat PMMA and PLA/90%PMMA can be
determined as 𝐺′ at the frequency where 𝑡𝑎𝑛𝛿 reaches a minimum (MIN method) [Liu et al.
(2006)]. As shown in Figure 6.14, the minimum 𝑡𝑎𝑛𝛿 for PLA/PMMA 50/50 and
PLA/70%PMMA blends can be obtained at 140 °C, and the corresponding 𝐺𝑁0 is determined
on the basis of Equation (2.19). However, for neat PLA and PLA/PMMA blends (PMMA
content<50%), the MIN method is not viable to determine 𝐺𝑁0 as the low experimental
temperature will result in the crystallization of PLA. So this method cannot be used to
determine 𝐺𝑁0 for all the samples. As alternative, Wu [Wu (1989)] and Nobile-Cocchini
[Nobile and Cocchini (2001)] proposed a semi-quantitative method based on the crossover
modulus 𝐺𝑋 ( 𝐺𝑋 = 𝐺′ = 𝐺′′ ). Wu [Wu (1989)] found the relationship between 𝐺𝑁0 and
crossover modulus 𝐺𝑋 if 𝑀𝑤/𝑀𝑛 < 3 by the following equation,
log (𝐺𝑁
0
𝐺𝑋) = 0.38 +
2.63 log (𝑀𝑤/𝑀𝑛)
1+2.45 log (𝑀𝑤/𝑀𝑛) (6.9)
For miscible blends, 𝑀𝑤 and 𝑀𝑛 are averaged values and can be calculated by the following
expression [Struglinski and Graessley (1985)]
𝑀𝑤 = 𝑤1𝑀𝑤1 + 𝑤2𝑀𝑤2 (6.10)
82 6 PLA/PMMA blends
1/𝑀𝑛 = 𝑤1/𝑀𝑛1 + 𝑤2/𝑀𝑛2 (6.11)
where 𝑤1 and 𝑤2 are the weight fractions of PLA and PMMA, respectively.
The entanglement molar mass 𝑀𝑒 and 𝑀𝑒12 of the various PLA/PMMA blends determined by
the “crossover modulus” and “MIN” methods are listed in Table 6.2.
Table 6.2 Average entanglement molar mass 𝑀𝑒, 𝑀𝑒12 and average entanglement density 𝑒
of PLA/PMMA blends.
PLA/PMMA
blend
𝑴𝒆 a (kg/mol)
𝑴𝒆 b
(kg/mol)
𝒆 a
(×10-4
mol/cm3)
𝑴𝒆𝟏𝟐 a (kg/mol)
100/0 3.8 ± 0.1 ---- 2.9 ± 0.1 ----
90/10 3.4 ± 0.1 ---- 3.3 ± 0.1 2.8± 0.1
70/30 3.1 ± 0.2 ---- 3.6 ± 0.2 2.9 ± 0.2
50/50 3.2 ± 0.1 3.5 ± 0.1 3.5 ± 0.1 3.0 ± 0.1
30/70 3.7 ± 0.2 4.7 ± 0.2 3.0 ± 0.2 3.1 ± 0.2
10/90 4.3 ± 0.1 6.2 ± 0.1 2.6 ± 0.1 2.1 ± 0.1
0/100 6.0 ± 0.1 7.1 ± 0.1 1.9 ± 0.1 ---- a
Estimated by using “crossover modulus-based method”. b
Estimated by using “MIN method”.
The 𝑀𝑒 value of pure PMMA obtained from “MIN method” agrees well with that reported in
literature [Wu (1987), Zhang et al. (2015)]. However, there is no available 𝑀𝑒 for pure PLA
that can be determined by “MIN method” even if the experimental temperature is decreased to
its crystallization temperature range. The 𝑀𝑒 values of neat PMMA and the PLA/PMMA
blends (PMMA content ≥50%) obtained from the “MIN method” is somewhat higher than
those determined by “crossover modulus-based method”, and this variation should be ascribed
to the semi-quantitative nature of the latter method. It is worth noting that the 𝑀𝑒 values
determined by both methods exhibit the same trend versus PMMA content.
6.1 PLA/PMMA 7N blends with different compositions 83
Figure 6.15 Average entanglement molar mass 𝑀𝑒 , 𝑀𝑒12 and entanglement density 𝑒
PLA/PMMA blends as a function of blend composition.
As shown in Figure 6.15, it can be observed that the values of 𝑀𝑒12 are in general smaller
than 𝑀𝑒 of neat components and the blends, indicating that the dissimilar chains are more
likely to entangle with each other than similar ones in the PLA/PMMA blends [Wu (1987)].
This result can partially explain the larger entanglement density for the blends compared to
the neat components. When PMMA content is below 50%, the increase of 𝑀𝑒12 with the
addition PMMA demonstrates the strong tendency of PLA chains to self-association rather
than inter-association with PMMA chains [Es‐Haghi et al. (2007)], corresponding to the self-
concentration model in miscible blends. The 𝑒 values of the blends are increasing with
PMMA content up to 50% and then decreasing with PMMA content.
The results about the degree of entanglement can be used to explain why PLA/PMMA blends
have improved shape memory properties compared to the neat components. Whereafter, shape
memory measurements will be carried out to confirm the impact of molecular entanglement
on the shape memory performance in the section 6.3.
6.2 PLA/PMMA 50/50 blends with different molecular structures
In this section, PMMA (PMMA 6N, PMMA 7N, PMMA 8N) with different molar masses are
used to blend with PLA. The effect of PMMA’s molecular structure on the thermo-
mechanical and rheological properties of PLA/PMMA 50/50 blends will be investigated. This
study is relevant in understanding the effects of molecular structure and specific inter-chain
interaction on molecular entanglement.
6.2.1 Molecular Characterization of PMMA 6N, 7N and 8N
The molar mass distributions of the PMMA 6N, 7N and 8N samples are plotted in Figure 6.16.
Obviously, the average molar mass 𝑀𝑤 of PMMA is increasing from 6N to 8N, and these
PMMA samples have similar and broad molar mass distribution in terms of the degree of
dispersion, 𝑀𝑤/𝑀𝑛. The 𝑀𝑤 of these PMMA are all above the critical molar mass 𝑀𝑐 (ca.
3104 g/mol) for entanglement coupling [Tadano et al. (2014)]. Therefore, the zero-shear
viscosity 𝜂0 of PMMA should scale in the following manner with the molar mass 𝑀𝑤,
𝜂0 = 𝐾𝑀𝑤3.4 (6.12)
In general, this relationship holds for linear polymers except for extremely broad blends, and
it could be check by the oscillatory shear experiments.
In addition, the radius of gyration 𝑅𝑔 of the PLA and PMMA chains could be calculated based
on Equation (5.3 and 5.4). The characteristic ratio 𝐶 is ca. 8.2 for PMMA [Wool (1995)]. As
listed in Table 6.3, the radius of gyration of PLA is ca. 21.6 nm, which is much larger than
that of PMMA. The radii of PMMA are increasing from 6N to 8N.
6.2 PLA/PMMA 50/50 blends with different molecular structures 85
Figure 6.16 Molar mass distribution of the PMMA grades under investigation.
6.2.2 Thermal behavior of PMMA and PLA/PMMA 50/50 blends
The thermal behavior of PLA, PMMA and the various PLA/PMMA 50/50 blends were
measured by DSC, as shown in Figure 6.17. The glass transition temperature 𝑇𝑔 of neat PLA
is around 60 °C, and 𝑇𝑔 for neat PMMA 6N, 7N and 8N are ca. 96 °C, 110 °C and 117 °C,
respectively. The differences in PMMA’s 𝑇𝑔 are related to the tacticity and molar mass of
PMMA. This effect will be discussed later in detail.
Figure 6.17 DSC thermograms in the second heating scan for (a) neat PLA and PMMA, (b)
PLA/PMMA 50/50 blends.
86 6 PLA/PMMA blends
A single broad 𝑇𝑔 for PLA/PMMA 50/50 blends were observed, confirming the miscibility of
the blends. 𝑇𝑔 for PLA/PMMA 50/50 blends are 72 °C, 75 °C and 77 °C, corresponding to
PMMA 6N, 7N, 8N, respectively. Furthermore, the broadness of 𝑇𝑔 of the PLA/PMMA 50/50
blends are increasing from PMMA 6N to 8N. According to the literature [Wetton et al. (1978),
Roland and Ngai (1993)], the broadening of 𝑇𝑔 is a result of concentration fluctuation and
local heterogeneity in the blend system.
In addition, it is worth mentioning that these three PLA/PMMA 50/50 blends are almost
amorphous, and no obvious melting peak could be observed in the DSC heating curves.
6.2.3 Rheological properties of neat PMMA
The zero shear viscosity 𝜂0 of PMMA could be determined by oscillatory shear experiments
and then plotted as a function of molar mass 𝑀𝑤, as shown in Figure 6.18a. A well-known
power law dependence with an exponent for the PMMA products of 3.62 is observed for the
correlation between 𝜂0 and 𝑀𝑤. Figure 6.18b presents 𝜂′′ as a function of 𝜂′ for neat PMMA.
The characteristic relaxation time 0 can be determined at angular frequency corresponding to
the maximum 𝜂′′ [Dealy and Larson (2006)], as listed in Table 6.3.
Figure 6.18 (a) Zero shear viscosity at 180 °C as a function of molar mass and (b) Cole-Cole
plots determined at 180 °C for PMMA 6N, PMMA 7N and PMMA 8N.
6.2 PLA/PMMA 50/50 blends with different molecular structures 87
The master curves of neat PMMA obtained by time-temperature superposition (TTS)
principle at a temperature range of 170-210 °C are shown in Figure 6.19. The temperature
dependence of the shift factors followed the Arrhenius equation (Equation (6.5)) and the
corresponding values of the activation energy 𝐸𝑎 are given in Table 6.3. It is evident that 𝐸𝑎
for neat PMMA is increasing from 6N to 8N. Furthermore, a plateau zone of 𝐺′ curves of
PMMA can be perceivable at the high frequency range. The plateau modulus 𝐺𝑁0 can be
determined at the reference temperature 180 °C based on the “MIN” method, as listed in
Table 6.3.
Figure 6.19 Master curves of 𝐺′, 𝐺′′, 𝜂∗ and 𝑡𝑎𝑛𝛿 for PMMA at a reference temperature of
180 °C. The rheological data were measured at 180-210°C. (□■170°C, ○●180°C, △▲190°C,
▽▼200°C, ◇◆210°C). 𝛼𝑇 is the shift factor for constructing the master curves.
88 6 PLA/PMMA blends
As shown in Table 6.3, the entanglement molar mass 𝑀𝑒 of PMMA are decreasing from 6N to
8N, and the average entanglement densities of PMMA are increasing from 6N to 8N.
According to previous studies [Wu and Beckerbauer (1992), Huang et al. (2011)], the
molecular entanglement behavior of PMMA is mainly influenced by chain tacticity, and the
pure syndiotactic PMMA shows the least 𝑀𝑒 in comparison with other PMMA. In this work,
the proportions of syndiotactic sequence increase from PMMA 6N to 8N (as shown in section
3.1), corresponding to the decreased 𝑀𝑒.
Table 6.3 Characterization of the neat components.
PLA PMMA
6N 7N 8N
Glass transition temperature 𝑇𝑔
(°C)
60 96 110 117
Radius of gyration 𝑅𝑔 (nm) 21.6 4.5 5.5 6.1
Zero shear-rate viscosity at
180 °C (×104 Pas)
0.5 8.6 28.3 56.1
Relaxation time 𝜏0 at 180 °C (s) 0.006 0.2 0.7 1.6
Plateau modulus 𝐺𝑁0 at 180 °C
(×105 Pa)
7.6 3.6 5.4 5.6
Activation energy 𝐸𝑎 (kJ/mol) 80.1 165.8 182.0 189.2
Entanglement molar mass 𝑀𝑒
(kg/mol)
4.2 7.2 6.1 5.8
Average entanglement density 𝑒
(×10-4
mol/cm3)
2.9 1.6 1.9 2.0
6.2.4 Interactions of PLA and PMMA via molecular entanglements in symmetrical
PLA/PMMA blends
The molecular entanglement in miscible blends is significantly influenced by the molecular
structure of the components. As discussed above, the dissimilar chains are more likely to
entangle with each other than the similar ones in PLA/PMMA blend.
As shown in Figure 6.20, TTS works well in the temperature range from 140-200 °C for all
PLA/PMMA 50/50 blends, indicating the miscibility of the blends in the experimental
6.2 PLA/PMMA 50/50 blends with different molecular structures 89
temperature range. In addition, 𝑡𝑎𝑛𝑚𝑖𝑛 is reached in the high frequency range and the
plateau modulus 𝐺𝑁0 for the blends can be obtained by “MIN” method, as listed in Table 6.4.
Figure 6.20 Master curves of 𝐺′ , 𝐺′′ , |𝜂∗| and 𝑡𝑎𝑛𝛿 for PLA/PMMA 50/50 blends at a
reference temperature of 180 °C. The rheological data were measured at various temperatures
between 140 and 200 °C (□■140 °C, △▲160 °C, ○●180 °C, ▽▼200 °C).
Table 6.4 Plateau modulus 𝐺𝑁0 (180 °C), entanglement molar mass 𝑀𝑒 and entanglement
density 𝑒 of PLA/PMMA 50/50 blends.
PLA/PMMA
50/50
𝑮𝑵𝟎
(×105 Pa
)
𝑴𝒆 (kg/mol)
𝒆
(×10-4
mol/cm3)
6N 9.0±0.2 3.7±0.2 2.9±0.2
7N 9.6±0.1 3.5±0.1 3.2±0.1
8N 9.9±0.1 3.4±0.1 3.4±0.1
90 6 PLA/PMMA blends
As can be seen from the Table 6.4, the average entanglement molar mass 𝑀𝑒 of PLA/PMMA
50/50 blends is decreasing slightly from 6N to 8N, while the entanglement densities 𝑒 is
increasing from 6N to 8N. In general, the entanglement molar mass of polymers is mainly
determined by the molecular structures, such as chain length, molar mass, nature of chain
substituents and spatial arrangement [Chalmers and Meier (2008)].
As discussed above, PMMA 8N has the largest molar mass and chain length, and the
conformational behavior of PMMA is studied in section 6.1. PLA/PMMA 8N 50/50 has the
largest entanglement density due to the larger chain length and higher proportion of
syndiotactic sequences of PMMA 8N.
Therefore, it can be concluded that the molecular structure of the component can significantly
influence the entanglement work in the miscible blends. The polymer with larger chain length,
molar mass and higher proportion of syndiotactic sequence is much easier to entangle with
other chains, leading to a higher entanglement density.
6.3 Shape memory property of PLA/PMMA blends and the underlying
mechanism
PLA/PMMA blends are typical SMP formed by a semi-crystalline polymer and an amorphous
polymer. The shape memory properties of PLA/PMMA blend films can be quantified by
Equations (4.2 and 4.3) based on Figure 4.2. In this section, the influence of stretching
temperature, strain rate, molar mass and blend composition on the shape fixing ratio 𝑅𝑓 and
shape recovery ratio 𝑅𝑟 were systematically studied. The maximum strain of DMTA, used a
deformation of 100%, is applied to all samples.
6.3 Shape memory property of PLA/PMMA blends 91
6.3.1 Influence of stretching parameters on the shape memory properties
6.3.1.1 Stretching temperature
The glass transition temperature 𝑇𝑔 is around 60 °C for neat PLA and 75 °C for PLA/PMMA
7N 50/50 blends. In order to study the influence of stretching temperature 𝑇𝑠 on the shape
memory performance, three stretching temperatures 𝑇𝑔, 𝑇𝑔+10 °C, and 𝑇𝑔+25 °C were chosen
to draw the films at a constant strain rate of 0.02/s for 50 s. The temperature for shape
recovery measurements (𝑇𝑟) is 10 °C above the stretching temperatures. For each sample, at
least three samples are measured.
Figure 6.21 Stress-strain curves of (a) neat PLA and (b) PLA/PMMA7N 50/50 blend films
stretched at various temperatures at a strain rate of 0.02 /s.
The stress-strain curves of films based on neat PLA and PLA/PMMA 7N 50/50 blend
determined at various temperatures are shown in Figure 6.21. When the PLA films were
stretched at 𝑇𝑔 (60 °C), a linear stress-strain relationship corresponding to elastic deformation
could be observed up to about 3% deformation, where an inconspicuous yield point was
obtained. After a slight decrease of the stretching stress at the yield point, a stable stress was
reached with the increase of deformation. The constant stress portion of the curve is the
elastic-plastic region, which is induced by the orientation of the free chains in the amorphous
92 6 PLA/PMMA blends
phase around 𝑇𝑔. In addition, no strain hardening can be observed up to the maximum strain
of 100%. Similar stress-strain curve is obtained when the PLA films are stretched at 𝑇𝑔+10 °C
(70 °C). However, the modulus and the constant stretching stress were reduced with increased
stretching temperature.
When the stretching temperature of PLA film is much higher than its 𝑇𝑔 (85 °C), the samples
react like a rubber. A highly nonlinear stress-strain response followed by a constant flow
stress could be observed. During this nonlinear stress-strain region, crystalline grains induced
by chains orientation and cold crystallization could be formed for semi-crystalline polymers
(as shown in Table 6.5). Afterwards, typical equilibrium is reached between the
recrystallization and stress relaxation induced by high temperature.
The stress-strain curves of films made of PLA/PMMA 7N 50/50 blend show s similar
response to an increase in stretching temperature. As shown in Figure 6.21b, when the sample
is stretched at 𝑇𝑔 (75 °C), a stable flow stress is reached after the yield point at about 3%
deformation. With the increase of stretching temperature, the linear stress-strain relationship
transforms into a nonlinear stress-strain response before a constant tress is arrived. The value
of the flow stress is reduced, and an obvious stress relaxation behavior can be observed when
the stretching temperature is above 100 °C.
After stretching and quenching to the room temperature, the length the sample was recorded
and then a shape recovery process was carried out in the oil bath. The shape fixing ratio 𝑅𝑓
and shape recovery ratio 𝑅𝑟 tested at different temperatures are listed in Table 6.5.
As shown in table 6.5, the shape fixing ratios 𝑅𝑓 for all samples are above 95%, and 𝑅𝑓 is
increasing with increased stretching temperature 𝑇𝑠. When 𝑇𝑠 is much higher than 𝑇𝑔, 𝑅𝑓 is
close to 100%. On the other side, the shape recovery ratio 𝑅𝑟 shows a different dependence on
increased 𝑇𝑠 . A higher 𝑇𝑠 could result in a lower 𝑅𝑟 , and the maximum value of 𝑅𝑟 is
6.3 Shape memory property of PLA/PMMA blends 93
obtained at 𝑇𝑠 = 𝑇𝑔 . Therefore, it can be concluded that the shape memory properties of
PLA/PMMA blends show obvious “temperature memory effect”, i.e., the shape fixing ratios
could be improved by increased stretching temperature, while the shape recovery ratios are
reduced by increased stretching temperature.
Table 6.5 The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of PLA and PLA/PMMA 7N
50/50 blend films stretched at various temperatures, and the crystallinity 𝑋𝑐 of cast films after
stretching.
𝑻𝒔 (°C) 𝑹𝒇 (%) 𝑹𝒓 (%) 𝑿𝒄 (%) 𝑻𝒓 (°C)
PLA
60 95.1 0.8 88.6 0.9 15.5 70
70 98.5 1.5 86.3 0.8 16.2 80
85 99.2 1.0 73.9 1.3 20.3 95
PLA/PMMA7N
50/50
75 96.2 1.3 98.9 0.7 0.2 85
85 98.8 1.6 97.5 1.4 0.5 95
100 99.3 1.1 91.8 1.2 3.8 110
6.3.1.2 Strain rate
In this section, three strain rates 0.02/s, 0.005/s, 0.002/s were chosen to investigate their
influences on the shape memory performance of films made of PLA/PMMA 7N 50/50 blend
at various stretching temperatures. The stress-strain curves obtained at various strain rates are
shown in Figure 6.22.
94 6 PLA/PMMA blends
Figure 6.22 Stress-strain curves of films made of PLA/PMMA7N 50/50 blend stretched at (a)
75 °C and (b) 100 °C at different strain rates.
As illustrated in Figure 6.22a, a stable stress level was reached after the yield point at ca. 3%
deformation when the films were stretched at 𝑇𝑔 (75°C). The stress-strain curves obtained at
different strain rates coincide. Thus, the flow stress is not influenced by the strain rate at 𝑇𝑔.
On the other side, when the samples are stretched at higher temperatures (100 °C), the stress-
strain curves are significantly influenced by the strain rate. After a nonlinear stress-strain
response, a dynamic equilibrium state is reached for all samples. It is found that lower strain
rates leads to a lower stretching stress due to stress relaxation processes.
The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 tested at various temperatures and strain
rates are calculated and listed in Table 6.6. When the samples are stretched at 𝑇𝑔, lower strain
rates result in a slightly higher shape fixing ratio, while the strain rate shows no impact on the
shape recovery ratio. On the contrary, the strain rate shows little impact on the shape recovery
ratio when the samples are stretched at higher temperatures. Smaller strain rates leads to a
significantly reduced shape recovery ratio.
Table 6.6 The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of PLA/PMMA 7N 50/50
blend films stretched at various strain rates.
Strain rate 𝑻𝒔 (°C) 𝑹𝒇 (%) 𝑹𝒓 (%) 𝑿𝒄 (%) 𝑻𝒓 (°C)
0.02/s 75 96.2 1.3 98.9 0.7 0.2 85
0.005/s 75 96.6 0.9 98.5 1.1 0.4 85
0.002/s 75 97.1 1.8 98.3 0.6 0.9 85
0.2/s 100 98.3 1.1 91.8 2.2 0.8 110
0.005/s 100 98.8 0.9 90.4 1.4 1.9 110
0.002/s 100 99.6 0.8 82.9 1.7 2.6 110
When the sample is stretched at low strain rate, the relaxation of “soft domains” is enhanced,
which could increase the shape fixing ratio. Moreover, a slight increase of crystallinity can be
6.3 Shape memory property of PLA/PMMA blends 95
observed after long time of stretching. These two factors are responsible for the reduced shape
recovery ratio.
Therefore, it could be concluded that a lower strain rate could slightly increase the shape
fixing ratio of PLA/PMMA blends, while decrease the shape recover ratio. In order to reach a
higher shape recovery ratio in the shape memory test, the larger strain rate (0.02/s) was
chosen for further experiments.
6.3.1.3 Discussion
For neat PLA films, the PLA crystallites mainly serve as physical cross-links to keep the
permanent shape of the samples and the amorphous phase between the crystallites works as
switching phase. When PLA films are stretched at lower temperatures (around 𝑇𝑔), there are
two reasons would lead to the reduced recovery ratios: (1) a slight increase of the crystallinity
after stretching; (2) chain slippage at the interfacial regions between the crystalline and
amorphous phase. The chain slippage and crystallization could change the proportions of
switching phase and stationary phase. Accordingly, this effect would lead to a plastic
irreversible modification during the stretching process and significantly reduce the shape
recovery ratio.
As shown in Table 6.5, the neat PLA film exhibits a crystallinity of about 10.8% before
stretching and 15.5% after stretching. This increase of crystallinity may be induced by a strain
(or local orientation)-induced crystallization of PLA, which has been demonstrated by
previous studies [Huang et al. (2011), Saeidlou et al. (2012)]. When the films are stretched
above the glass transition temperature, the active chains in the amorphous phase rearrange
along the deformation direction. Therefore, the stretching process would result in an increase
in crystallinity, and a schematic is given in Figure 6.23.
96 6 PLA/PMMA blends
Figure 6.23 Schematic illustration of microstructure changes of PLA/PMMA blend films
under strain deformation and crystallization.
The cold crystallization of PLA starts around 85 °C (as shown in Figure 6.1). Therefore, when
PLA films are stretched at higher temperature (≥ 85 °C), in addition to the strain-induced
crystallization, cold crystallization is also possible to increase the crystallinity of the samples.
As shown in Figure 6.21a, a stress relaxation behavior can be observed when PLA films are
stretched at 85 °C. This relaxation behavior is a partial relaxation of “soft domains”, which
may reduce the orientation of switching phase. Moreover, this relaxation behavior will lead to
a better shape fixing ratio but reduce the shape recovery ratio.
Therefore, the low recovery ratios of PLA films at high temperatures could be attributed to (1)
partial modification of “stationary phase” induced by strain-induced crystallization and cold
crystallization, (2) plastic irreversible deformation induced by chain slippage between the
crystalline and amorphous chains, (3) relaxation of “soft domains” [Ratna and Karger-Kocsis
(2008), Samuel et al. (2014)].
For semi-crystalline PLA/PMMA blends (PLA rich), a similar mechanism can be used to
explain the influence of temperature on the shape memory properties as for neat PLA.
6.3 Shape memory property of PLA/PMMA blends 97
Consequently, the stretching at higher temperatures shows a positive effect on the shape
fixing ratio, but exhibits a significantly and negative impact on the shape recovery ratio on
neat PLA films and semi-crystalline PLA/PMMA blend films.
For PLA/PMMA 50/50 blend films, the temperature shows similar impact on the shape
memory performances as for neat PLA, but with different mechanisms. The PLA/PMMA
50/50 blend is almost amorphous with an extremely low crystallinity (0.1%). The molecular
entanglements mainly serve as physical cross-links to keep the original shape. When the films
are stretched at low temperatures (around 𝑇𝑔 ), the stress-strain curves shows no stress
relaxation behavior during stretching process. In addition, there is little change of the
crystallinity of the films, indicating the good stability of “hard domains” during stretching.
Therefore, a good shape recovery ratio is obtained for PLA/PMMA 50/50 blend stretched at
𝑇𝑔.
With the increase of stretching temperature, stress-relaxation behavior could be observed
during stretching process. The stress-relaxation behavior will not only reduce the orientation
of switching phase, but also induce disentanglement during stretching [Ratna and Karger-
Kocsis (2008), Wang and Li (2015)]. Moreover, the DSC measurements proved the increase
of crystallinity which will induce a plastic irreversible deformation (as shown in Table 6.5).
Consequently, it can be concluded that the shape fixing ratio of PLA/PMMA 50/50 blend
films is improved with increased stretching temperature, while the shape recovery ratio is
reduced by the increased stretching temperature.
98 6 PLA/PMMA blends
6.3.2 Influence of molecular structure of PMMA and blend composition on the shape
memory properties
6.3.2.1 PMMA’s molecular structure
The influences of PMMA’s molecular structure on the thermo-mechanical and rheological
properties of PLA/PMMA 50/50 blends have been discussed in section 6.2. In this section, its
effect on the shape memory performances of PLA/PMMA blends will be discussed based on
the 50/50 blend.
Figure 6.24 Stress-strain curves of films made of PMMA/50%PMMA6N, 7N and 8N blends
stretched at a strain rate of 0.02 /s at their respective 𝑇𝑔.
As shown in Figure 6.24, the stress-strain curves of symmetric PLA/PMMA blends with
PMMA of various molecular structures display little difference when they are stretched at
same condition at their respective 𝑇𝑔. The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of
three blends were calculated and listed in Table 6.7. It is evident that the blends show similar
𝑅𝑓 and 𝑅𝑟 when they are stretched at their respective 𝑇𝑔. According to the results listed in
Table 6.4, the entanglement density is increasing from 6N to 8N, but the increment is so small
that little difference could be observed in their shape memory performance. Consequently, it
can be concluded that the molar mass of PMMA shows a negligible influence on the shape
6.3 Shape memory property of PLA/PMMA blends 99
memory properties on PLA/PMMA blends due to the small difference between entanglement
densities. Therefore, PMMA 7N was chosen for further experiments.
Table 6.7 The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of films made of
PLA/PMMA 50/50 blends stretched at their respective 𝑇𝑔.
PLA/PMMA 50/50 𝑻𝒈 (°C) 𝑹𝒇 𝑹𝒓 𝑻𝒓 (°C)
6N 73 96.1 1.7 98.2 2.1 83
7N 75 96.2 1.3 98.9 0.7 85
8N 77 95.9 2.2 98.7 1.2 87
6.3.2.2 PMMA’s content
For miscible amorphous/crystalline polymer blends used as shape memory polymers, 𝑇𝑔 of
the amorphous phase is the critical temperature for triggering the shape recovery [Mather et al.
(2009)]. According to our study in the section 6.3.1 and the investigations of Samuel et. al
[Samuel et al. (2014)], the shape memory performance of PLA/PMMA blends presents a
significant “temperature memory effect”. Therefore, the sample’s 𝑇𝑔 is chosen as switching
temperature to compare the influence of blend composition on the shape memory properties.
Figure 6.25 Stress-strain curves of PLA/PMMA blend films stretched at their respective 𝑇𝑔
with a strain rate of 0.02/s.
100 6 PLA/PMMA blends
As shown in Figure 6.25, for each PLA/PMMA blend film, a stable flow stress is reached
after the yield point at ca. 3% deformation when it is stretched at their respective 𝑇𝑔. The
stretching stresses of the blends are quite similar and range from 2.0106 to 3.010
6 Pa.
Furthermore, no stress relaxation or strain hardening behavior could be observed in the steady
state stress levels when the samples are stretched at their respective 𝑇𝑔.
The shape fixing ratios 𝑅𝑓, shape recovery ratios 𝑅𝑟 and crystallinity 𝑋𝑐 after stretching as a
function of PMMA content are listed in Table 6.8. It is evident that all the samples show
shape fixing ratios around 95%, independent of the blend composition. However, the shape
recovery ratios of PLA/PMMA blends exhibit strong composition dependence. The highest
shape recovery ratio was observed for blends containing 30-70% PMMA, and the
corresponding 𝑅𝑟 values for these samples are all above 95%.
Table 6.8 The shape fixing ratio 𝑅𝑓 , shape recovery ratios 𝑅𝑟 and crystallinity 𝑋𝑐 after
stretching of PLA/PMMA blend films stretched at their respective 𝑇𝑔.
PLA/PMMA 𝑇𝑔 (°C) 𝑅𝑓 (%) 𝑅𝑟 (%) 𝑋𝑐 (%) 𝑇𝑟 (°C)
100/0 60 95.1 0.8 88.6 1.1 15.5 70
90/10 62 95.8 1.2 95.1 0.8 8.2 72
70/30 67 95.9 1.5 98.2 1.2 4.9 77
50/50 75 96.2 1.3 98.9 0.7 0.2 85
30/70 85 96.1 1.0 96.9 0.8 0 95
10/90 100 96.6 1.4 86.9 0.5 0 110
0/100 110 96.2 1.3 81.1 1.0 0 120
6.3.3 The shape memory mechanism of PLA/PMMA blend system
As discussed above, the shape memory properties of PLA/PMMA blends show distinct
composition dependence. Therefore, the corresponding mechanism can be divided into two
categories according to the blend composition: semi-crystalline blends and amorphous blends.
For the semi-crystalline blends, the PLA crystallites and molecular entanglement play the role
6.3 Shape memory property of PLA/PMMA blends 101
of physical cross-links. On the other side, for amorphous PLA/PMMA blends, just molecular
entanglements serve as physical cross-links.
Figure 6.26 Shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of PLA/PMMA blends as a
function of PMMA content, and the shape memory mechanism at different PMMA contents.
As shown in Figure 6.26, the neat PLA film shows a moderate shape recovery ratio. The PLA
crystallites are considered as the stationary phase for neat PLA, and the active chains between
crystallites play a role of the switching phase. During the stretching process, these disordered
active chains would rearrange along the deformation direction (as shown in Figure 6.23),
leading to an increase of crystallinity. Consequently, the strain-induced crystallization of neat
PLA results in an irreversible modification during the stretching process and significantly
reduces the shape recovery ratio. On the other hand, molecular slippage at the interfacial
regions between the crystalline and amorphous phase will occur upon long-term stress [You et
al. (2012)]. This factor could also reduce the shape recovery ratio of neat PLA.
For the blends with PMMA content less than 50% (PLA-rich), the crystallization of PLA is
strongly hindered by the presence of PMMA (as shown in Figure 6.27). This effect can be
explained by two factors: (1) the incorporation of PMMA into PLA matrix dilutes the PLA
chains in the blends which disrupt the ordered arrangement of PLA chains, (2) the molecular
102 6 PLA/PMMA blends
entanglements between PLA and PMMA chains reduced the mobility of PLA chains in the
matrix, thus limiting the capability of PLA chains to adopt an optimal alignment. Moreover,
the work of Samuel [Samuel et al. (2013)] demonstrates the absence of any other reactions
between PMMA and PLA phase. Therefore, the crystalline phase of PLA was suppressed
when PMMA was introduced.
Figure 6.27 The crystallization 𝑋𝑐 of the films made of PLA/PMMA blends before and after
stretching at their respective 𝑇𝑔.
In addition, the introduction of PMMA into PLA creates new physical cross-links which are
formed by the interactions between PLA and PMMA chains. It could be noticed that the shape
recovery ratios seem to increase with the amount of PMMA up 50%. In this case, in addition
to the PLA crystallites acting as physical cross-links to keep the original shape, the role of
molecular entanglements in the amorphous phase become more important [Ratna and Karger-
Kocsis (2008)] and contribute to the good shape memory properties of the blends. In Figure
6.28a, the shape recovery ratio 𝑅𝑟 shows exponential dependence on both the crystallinity 𝑋𝑐
and the entanglement density 𝑒. However, 𝑅𝑟 decrease with 𝑋𝑐 but increase almost linearly
with 𝑒.
6.3 Shape memory property of PLA/PMMA blends 103
During the stretching process at 𝑇𝑔, it is evident that the crystallinity of the blends is increased
due to the strain-induced crystallization (as shown in Figure 6.27) [Huang et al. (2011),
Saeidlou et al. (2012), Yin et al. (2015)]. Moreover, the disentanglement of long linear
polymer chains during the crystallization process has also been demonstrated [Luo and
Sommer (2012)]. Although PLA/30%PMMA blend and PLA/PMMA 50/50 blend have a
similar entanglement density in the molten state (as shown in Figure 6.28b), the
disentanglement of polymer chains due to the crystallization process could lead to a lower
entanglement density in PLA/30%PMMA blend in comparison with PLA/PMMA 50/50 blend.
Consequently, PLA/PMMA 50/50 blend exhibits better shape recovery ratio than
PLA/30%PMMA blend.
Figure 6.28 The dependence of shape recovery ratio 𝑅𝑟 on (a) crystallinity 𝑋𝑐 (b)
entanglement density 𝑒 for the blends with PMMA contents above 50%.
Concerning the stability of the stationary phase, for semi-crystalline blends, the molecular
entanglements between PLA and PMMA chains seem to have a positive impact on the shape
recovery ratio, while the crystallinity shows a negative impact on it.
104 6 PLA/PMMA blends
Figure 6.29 Shape recovery ratio 𝑅𝑟 as a function of entanglement density 𝑒 for the blends
with PMMA content above 50%.
When PMMA dominates in the blends (PMMA-rich), PLA crystallization is completely
suppressed by the presence of PMMA. The PLA/PMMA blends (PMMA content >50%) are
amorphous and only the entanglement network is responsible for maintaining the original
shape. As shown in Figure 6.29, for the amorphous PLA/PMMA blends in which PMMA
accounts for more than 50% of mass, the shape memory potential is arising from the
entanglement network and strongly dependent on the entanglement density.
According to the previous investigations, the shape memory mechanism of miscible semi-
crystalline/amorphous blends can be concluded from PLA/PMMA blend system. When the
blends are semi-crystalline, the crystallites and molecular entanglement act as physical cross-
links to keep the permanent shape and recovery to it. However, a negative influence will be
induced by crystallites due to a strain-induced or temperature-induced crystallization during
the deformation above the glass transition temperature. For the amorphous blends, the
entanglement network serves as physical cross-link and plays a crucial role to influence the
shape memory performances. In summary, we propose that the entanglement network in
miscible semi-crystalline/amorphous blend is the main factor to maintain the shape recovery
property, which could be compromised by the increased crystallinity. The results of our study
6.3 Shape memory property of PLA/PMMA blends 105
provide a novel understanding of the mechanisms underlying the shape memory properties for
miscible semi-crystalline/amorphous SMPs.
6.4 Conclusions
The thermo-mechanical and rheological properties of miscible PLA/PMMA blends with
different blend compositions and PMMA molar masses were investigated in this work. The
crystallinity of the blends was found to decrease upon mixing with PMMA up to 50% where a
transition from the semi-crystalline to amorphous phase occurs. A smaller entanglement
molar mass for 𝑀𝑒,𝑃𝐿𝐴−𝑃𝑀𝑀𝐴 compared to 𝑀𝑒 of neat components or blends was derived from
the oscillatory shear rheological measurements on the PLA/PMMA blends, suggesting that
the dissimilar chains are more likely to entangle with each other than the similar ones. For the
semi-crystalline blends with PMMA contents less than 50%, an increase in entanglement
density is induced by the addition of PMMA which leads to a decreased crystallinity and an
enhanced shape recovery ratio for the blends. This result demonstrates that the molecular
entanglement has a positive influence on the shape memory performance, whereas crystallites
exhibit an opposite impact which may arise from the strain-induced crystallization and the
molecular slippage between the crystalline and amorphous chains network occurs upon long-
term stress. In the case of an amorphous blends in which PMMA accounts for more than 50%,
an positive dependence was derived between the entanglement network and shape recovery
ratio, indicating that the entanglement network is solely contributing to the shape recovery
potential. In summary, for miscible amorphous/semi-crystalline SMPs, the entanglement
network serves as a critical element in determining the shape recovery capability, while the
effect of crystallites may be an impediment via strain-induced crystallization.
In addition, the influences of stretching temperature, strain rate and component molar mass on
the shape memory performances of PLA/PMMA blends were systematically studied. It was
106 6 PLA/PMMA blends
found that the higher stretching temperature or lower strain rate would result in a larger shape
fixing ratio, but a lower shape recovery ratio. The molar mass of the component could
influence the critical temperature for triggering the shape recovery, but show little influence
on the shape memory properties when they the samples are stretched at their respective 𝑇𝑔.
7. PLA/PMMA/silica nanocomposites
In production processes, the commonly used polymeric materials are usually manufactured by
mixing different macromolecules or incorporating solid “filler” to improve the impact
strength, modulus, processability, conductively, flammability or appearance, etc. [Meijer et al.
(1988)] Polymer nanocomposites have gained extensive interests since the incorporation of
well dispersed nanoparticles can significantly improve the mechanical, optical or barrier
properties. In particular, for polymer blends, the addition of filler can result in either an
increase or a decrease of the temperature of phase separation, change of the interaction
parameter between two components, a modification of the shape of the phase diagram or a
change in the kinetics of phase separation [Lipatov (2002), Lipatov et al. (2002), Ginzburg
(2005), Huang et al. (2005), Lipatov and Alekseeva (2007)]. In recent years, it has been
recognized that small fractions of nanoscale fillers can significantly affect the rheological and
mechanical properties of polymers [Krishnamoorti et al. (1996), Krishnamoorti and Giannelis
(1997), Hoffmann et al. (2000), Lamnawar et al. (2011)]. Moreover, the presence of
nanoparticles can significantly affect the thermodynamic phase behavior of polymer blends [Ji
et al. (2000), Lee et al. (2006)].
In the present work, the intermolecular cooperativity and segmental dynamics in
PLA/PMMA/nanosilica mixtures were studied. The effect of nanosilica (silica 300, 𝑑=7 nm)
on the phase behavior, molecular entanglement, thermo-mechanical and rheological properties
PLA/PMMA 7N 50/50 blends was systematically investigated. For convenience, the unfilled
and filled PLA/PMMA 7N 50/50 blends are designed as P/P/Si x in the following discussion.
Here x represents the silica content (wt %) in the nanocomposite.
108 7 PLA/PMMA/silica nanocomposites
7.1 Morphological characterization
7.1.1 Dispersion of nanosilica in PLA and PMMA
It is well known that the dispersion of nanofillers in the polymer matrix plays an important
role in affecting the physical properties of polymer. A homogeneous dispersion of nanofillers,
together with an optimized interaction between nanofiller and polymer matrix, will effectively
improve the thermal mechanical and rheological properties of the polymer matrix [Li et al.
(2012)]. The dispersion of nanosilica (10 wt%) in PLA and PMMA matrix was firstly
investigated by FE-SEM. As shown in Figure 7.1, the nanoparticles dispersed evenly in both
the PLA and PMMA matrix and exhibited agglomerates with particle size less than 50 nm.
The dispersion level of nanosilica in PLA matrix is highly similar with that in PMMA matrix.
Figure 7.1 SEM micrographs of the fractured surfaces of PLA/10% silica and PLA/10%
silica nanocomposites.
7.1.2 Dispersion of nanosilica in PLA/PMMA blends
In order to elucidate the effect of nanosilica on the phase morphology, the PLA/PMMA 50/50
blend is chosen. The fracture surfaces of samples with and without nanosilica were observed
by FE-SEM. Clearly, as shown in Figure 7.2a, the binary PLA/PMMA blend shows typical
single phase morphology, indicating the good miscibility between PLA and PMMA. For the
7.1 Morphological characterization 109
nanocomposites (Figure 7.2b, c and d), nanosilica particles were detected as white dots. It is
evident that nanosilica particles can form aggregates that are uniformly dispersed in the blend
matrix with average size less than 100 nm, even at low silica contents (2 wt%, Figure 7.2b). A
specific “crater” structure is formed around the nanosilica aggregates which appears to serve
as the central core. With the increase of nanosilica loading, an increase in both the size and
number of the nanosilica aggregates are observed, whereas the dimension of the “crater”
structures is remarkably reduced (Figure 7.2c, d). At higher silica contents (> 5 wt%), the
size of the aggregates can even exceed 300 nm. These observations agree well with the notion
that the higher nanofiller contents, the larger the aggregate forms due to the strong
interactions among the nanoparticles [Li et al. (2012)].
Figure 7.2 SEM micrographs of the fractured surfaces of (a) P/P/Si 0, (b) P/P/Si 2, (c) P/P/Si
5, (d) P/P/Si 10.
110 7 PLA/PMMA/silica nanocomposites
An interesting comparison can be made as shown in Figure 7.3. Although the same silica
content of 10 wt% in all the nanocomposites, the dispersions of nanosilica in neat PLA or
PMMA matrix differs significantly from that in polymer blend matrix. In addition, the
nanosilica aggregates formed in the neat polymer matrix are much smaller than that formed in
blend matrix. Interestingly, the specific “crater” structure is solely observed in the fractured
surfaces of the nanocomposites based on PLA/PMMA blends.
Figure 7.3 The dispersion of nanosilica in PLA, PMMA and PLA/PMMA blends with the
same filler concentration of 10 wt%.
The influences of nanoparticles on the overall phase behavior of miscible blend have
extensively been studied [Lipatov (2002), Lipatov et al. (2002), Huang et al. (2005), Lipatov
(2006), Chung et al. (2007), Lipatov and Alekseeva (2007)]. One of the explanations is
7.1 Morphological characterization 111
related to the specific interactions and preferential adsorption on the fillers by one of the
components of the blend [Huang et al. (2005)]. This preferential adsorption will lead to the
unique morphological characterization in the nanocomposites based on miscible blend.
7.2 Preferential adsorption on nanosilica by one of the components of
PLA/PMMA blends
The interactions between nanosilica and the pure components were first investigated by
oscillatory shear rheology. In general, the enhancement of the complex viscosity |𝜂∗| of
nanocomposites mainly depends on the processing method, the dispersion of nanofillers and
the interactions between the nanofillers and the polymer matrix [Bar-Yam (1997)]. As
depicted in Figure 7.1, a similar dispersion of nanosilica (10 wt%) is observed in PLA matrix
and PMMA matrix which were prepared following the same manufacturing procedure.
Therefore, the different enhancements on the linear viscoelastic behaviors are mainly arisen
from the interactions between the nanosilica and the polymer matrix.
Figure 7.4 Normalized complex viscosity |𝜂∗| for the PLA and PMMA nanocomposite
samples containing 10 wt% of nanosilica at 200 °C.
Figure 7.4 illustrates the complex viscosity |𝜂∗| of PLA/10 wt% silica and PMMA/10 wt%
silica nanocomposites relative to |𝜂∗| of their corresponding neat polymer matrix at the
112 7 PLA/PMMA/silica nanocomposites
temperature equal to their compounding temperature (200 °C). Normalized complex viscosity
(|𝜂𝑁𝑎𝑛𝑜𝑐𝑜𝑚𝑝𝑜𝑠𝑖𝑡𝑒∗ |/|𝜂𝑀𝑎𝑡𝑟𝑖𝑥
∗ |) curves of different polymers with equal loadings of nanofillers
can provide insight about the affinity of the nanofiller towards each polymer [Abbasi Moud et
al. (2015), Rostami et al. (2015)]. From the results presented in Figure 7.4, it is found that the
normalized complex viscosity for PLA is clearly much higher than that of PMMA at the
whole frequency range, demonstrating a stronger affinity of nanosilica particles towards the
PLA molecular chains as compared to PMMA chains.
Form the molecular perspective, both PLA and PMMA have C = O groups that can interact
with the silanol groups on the surface of silica to form hydrogen bond [Wen et al. (2009)] (as
shown in Figure 7.5). Thus the weaker binding affinity of silica towards PMMA can be
ascribed to the steric hindrance introduced by the CH3 − C = O groups in PMMA.
Consequently, PLA molecules have stronger affinity with nanosilica in comparison with
PMMA molecules.
Figure 7.5 Schematic illustrations of the hydrogen-bonding interactions between (a) PLA and
silica particles, (b) PMMA and silica particles.
The SEM picture of P/P/Si 2 at high magnification (105) is presented in Figure 7.6. It’s
clearly to observe that the nanosilica particles form small agglomerates in the polymer matrix.
These agglomerates interact with nearby PLA or PMMA molecules to form the “crater”
7.2 Preferential adsorption on nanosilica 113
structure, as highlighted in Fig. 4 (left panel). A similar scenario is also observed in a
previous work in which the aggregates formed by nanoscale particles leads to the
redistribution of other components like poly(vinylidene fluoride) (PVDF)/PMMA blends
filled by modified multiwalled carbon nanotubes (MWNTs) [Sharma et al. (2014)]. Given the
fact the stronger binding affinity of nanosilica towards PLA, it can be expected that the PLA
molecules are more likely to adsorb on the surface of nanosilica in comparison with PMMA.
We propose a model describing the composition of the “crater” structure induced by the
addition of nanosilica particles, as seen in Figure 7.6 (right panel). A selective accumulation
of PLA molecules surround the surface of silica aggregates, while the PMMA molecules are
mainly distributed at the peripheric of the “crater” structure.
Figure 7.6 SEM micrographs of the fractured surfaces of P/P/Si 2 nanocomposites. The
specific “crater” structure is highlighted by the red circle (left panel). A schematic diagram
illustrating the composition of the “crater” structure in which a preferential adsorption of PLA
molecules on the surface of nanosilica aggregates is suggested.
114 7 PLA/PMMA/silica nanocomposites
7.3 Thermo-mechanical properties
7.3.1 DSC
The influence of nanosilica on the thermo-mechanical behavior of PLA/PMMA blends was
determined by DSC. As shown in Figure 7.7a, the DSC curves of the second heating scan
exhibit single broad glass transition temperature 𝑇𝑔, confirming the miscibility of
nanocomposites in the solid state. It is worth noting that 𝑇𝑔 is increasing with the addition of
nanosilica, which shows no effect on the 𝑇𝑔 of PLA/silica or PMMA/silica mixtures.
Lodge−McLeish Model: “self-concentration”
Figure 7.7 (a) DSC second heating scans of P/P/Si x, and (b) variation of 𝑇𝑔 with silica
concentration for PLA/PMMA blends.
According the previous studies [Wetton et al. (1978), Roland and Ngai (1993), Shi et al.
(2013)], the broadening of 𝑇𝑔 for PLA/PMMA/nanosilica attributes to the concentration
fluctuations and dynamic heterogeneity in the miscible blend. A concept of “self-
concentration” was proposed by Lodge and McLeish [Lodge and McLeish (2000)], and the
effective glass transition temperature 𝑇𝑔,𝑒𝑓𝑓 for each component could be calculated according
to Equations (6.1-6.3).
7.3 Thermo-mechanical properties 115
In a miscible blend, the dynamics of higher 𝑇𝑔 component in the mixture is more
representative of the average blend composition, while the dynamics of lower 𝑇𝑔 component
is similar to its pure component [Lodge and McLeish (2000)]. That is to say, the calorimetric
𝑇𝑔 of PLA/PMMA 50/50 blend should be in accordance with the 𝑇𝑔,𝑒𝑓𝑓 of PMMA at 50%
content. Surprisingly, a large discrepancy between the calorimetric 𝑇𝑔 of unfilled
PLA/PMMA blend and the effective glass transition temperature of PMMA is observed in
Figure 7.7b. With the increase of nanosilica concentration, this discrepancy between
calorimetric 𝑇𝑔 and 𝑇𝑔,𝑒𝑓𝑓,𝑃𝑀𝑀𝐴 is reduced, suggesting the presence of nanosilica increases the
concentration fluctuations and dynamic heterogeneity of PLA/PMMA blends.
Crystallization of PLA/PMMA/nanosilica mixtures
Interestingly, the presence of nanosilica slightly increases the crystallization tendency of PLA,
and a small but significant melting peak can be observed around 165 °C for P/P/Si 10
(crystallinity approach 0.4%, see Figure 7.7a). For P/P/Si 0, the crystallinity index of PLA is
extremely weak (approach zero) due to the intimate mixing of PLA and PMMA chains
[Samuel et al. (2014)]. A possible interpretation for the increase of the crystallinity of PLA is
the preferential interactions between nanosilica and PLA molecules, which would lead to the
local accumulation of PLA around the surface of nanosilica aggregates. These aggregates
seem to act as potential nucleation agent which is proposed to increase the crystallinity of
PLA [Nofar et al. (2013)].
7.3.2 Dynamic mechanical analysis (DMTA)
Figure 7.8 shows the storage modulus 𝐸′ and loss modulus 𝐸′′ as a function of temperature
for unfilled PLA/PMMA blend and nanocomposites containing 2, 5 and 10 wt% of nanosilica.
The moduli at temperatures below 𝑇𝑔 are increasing with the addition of nanosilica, indicating
the presence of nanosilica improves the stiffness of matrix. A well-defined relaxation peak
116 7 PLA/PMMA/silica nanocomposites
(i.e., transition) was displayed around 71 °C for unfilled PLA/PMMA blend. The peak of
loss modulus (𝑇) also moves to higher temperatures, corresponding to the result of DSC tests.
In addition, the broadness of transition is also increased with the addition of nanosilica.
Figure 7.8 The loss and storage modulus as function of temperature for PLA/PMMA 50/50
blends with different nanosilica contents.
7.4 Rheological properties of PLA/PMMA/silica nanocomposites
7.4.1 Oscillatory strain sweep
The linear viscoelastic region of the samples can be determined by the dynamic amplitude
measurements. Figure 7.9 exhibits the strain dependence of the dynamic storage modulus 𝐺′
of the pure PLA/PMMA blend and its nanocomposites at 200 °C at an angular frequency of
10 rad/s. Obviously, 𝐺′ at low strain range increases monotonously with increased silica
content, but the linear viscoelastic region of the nanocomposite was reduced by the presence
of nanosilica. The linear viscoelastic region for the unfilled PLA/PMMA blend is around 30%,
while that for P/P/Si 10 is just 1.5%. Therefore, the following dynamic rheological
measurements were performed at = 1%.
7.4 Rheological properties of PLA/PMMA/silica nanocomposites 117
Figure 7.9 The storage modulus 𝐺′ for pure PLA/PMMA blend and PLA/PMMA/silica
nanocomposites obtained in strain sweeps at 200 °C and 10 rad/s.
7.4.2 Oscillatory time sweep
Oscillatory time sweeps are very important for polymer/particle mixtures, which may undergo
macro- or microstructural rearrangements with time. These rearrangements may be induced
by the polymer degradation or particle diffusion in the composites.
For unfilled PLA/PMMA blends, thermal stability can be accessed via the change of storage
modulus 𝐺′ as a function of time. As shown in Figure 7.10, it is evident that the thermal stable
time for P/P/Si 0 is about 8000 s at 200 °C. With the incorporation of 2 wt% nanosilica, the
thermal stability exhibits a slight increase and stabling time shifts to 9000 s. This result
indicates the presence of nanosilica improved the thermal stability of PLA/PMMA blends.
118 7 PLA/PMMA/silica nanocomposites
Figure 7.10 Relative change of storage modulus as a function of the residence time at 200 °C
for P/P/Si x.
For PLA/PMMA blend containing high contents of nanosilica (>5 wt%), a distinct increase of
the normalized 𝐺′ can be observed in the short time range. This increase is induced by the
structure build-up in the nanocomposites at molten state [Eslami et al. (2009)]. 𝐺′ begins to
decrease when the experimental time exceed 10000 s, indicating that possibly the begin of
polymer degradation.
7.4.3 Oscillatory frequency sweep
Time-temperature superposition (TTS)
Time-temperature superposition (TTS) principle can be used to determine the phase
separation temperature of polymer blends [Jeon et al. (2000)], and can also extend the
frequency or time range [Van Gurp and Palmen (1998)]. It has been proved that for
temperatures in the homogeneous region, TTS principle worked well. While TTS principle
failed in the temperature range of phase separation [Nesarikar (1995), Kapnistos et al. (1996)].
Figure 7.11 shows the master curves of the P/P/Si x at the reference temperature of 200 °C
obtained by TTS.
7.4 Rheological properties of PLA/PMMA/silica nanocomposites 119
Figure 7.11 Master curves of 𝐺′, 𝐺′′, |𝜂∗| and 𝑡𝑎𝑛𝛿 for (a) P/P/Si 0, (b) P/P/Si 2, (c) P/P/Si 5,
(d) P/P/Si 10, at a reference temperature of 200 °C. The rheological data were measured at
190-220 °C (□■190 °C, △▲200 °C, ○●210 °C, ▽▼220 °C).
As shown in Figure 7.11a, the unfilled PLA/PMMA blend follows the known rheological
relations G′~𝜔2 and G′′~𝜔1 in the terminal region. The TTS principle holds satisfactorily for
𝐺′ , 𝐺′′ , 𝜂∗ and 𝑡𝑎𝑛𝛿 when the temperature is not higher than 210 °C, while TTS breaks
drown for 𝐺′ and 𝑡𝑎𝑛𝛿 at 220 °C. An obvious plateau is present in the low frequency region.
This result indicates that unfilled PLA/PMMA blend is a homogeneous mixture when the
temperature is below 210 °C, and phase separated when the temperature is above 210 °C.
Therefore, the phase separation temperature for the unfilled PLA/PMMA blend is between
210 and 220 °C.
120 7 PLA/PMMA/silica nanocomposites
The incorporation of nanosilica changed the slope of 𝐺′ in the terminal region, and its value
decreases with increased silica content. For P/P/Si 2 nanocomposite, a slope equal to 2 can
still be observed in the terminal region (Figure 7.11b), and the TTS principle applied well for
the temperature below 220 °C. A weak plateau of 𝐺′ can be observed at 220 °C in the low
frequency region, and the deviation of 𝑡𝑎𝑛𝛿 at low frequency region is also obvious. It can be
deduced that the phase separation temperature for P/P/Si 2 nanocomposite is around 220 °C,
which is slightly higher than that of unfilled PLA/PMMA blend. The slope of 𝐺′ at the
terminal region for P/P/Si 5 and P/P/Si 10 are both less than 2 (Figure 7.11c, d), while the
TTS principle worked well for both nanocomposites in the temperature range of the
experiments, indicating both of their phase separation temperatures are above 220 °C.
Han plots
In order to investigate the phase separation rheologically, an alternative method proposed by
Han is adopted. The Han plot (log𝐺′ vs. log𝐺′′) eliminates the effects of frequency, and it is
proved to be more sensitive to the phase separation induced by temperature [Nesarikar (1995),
Kapnistos et al. (1996)]. As shown in Figure 7.12a and 7.12b, the log𝐺′-log𝐺′′ curves of
P/P/Si 0 and P/P/Si 2 could be superposed on the same master curve at the temperatures from
180 to 210 °C, respectively. But the heterogeneous phase curves at 220 °C clearly deviate
from their own master curve. This result demonstrates their phase separation temperatures are
both located between 210 and 220 °C, agreeing with the conclusions obtained by TTS
principle.
7.4 Rheological properties of PLA/PMMA/silica nanocomposites 121
Figure 7.12 Master curves of Han plot (i.e. log𝐺′ vs. log𝐺′′), (a) P/P/Si 0, (b) P/P/Si 2, (c)
P/P/Si 5, (d) P/P/Si 10.
For P/P/Si 5 and P/P/Si 10 nanocomposite (Figure 7.12c, d), the homogeneous phase curves
coincide over the whole temperature range (180 to 220 °C) of the experiments, indicating that
their phase separation temperatures are above 220 °C. Consequently, the result from Han plot
is well consistent with that from TTS.
According to the analysis results from TTS and Han plot, it can be concluded that the
incorporation of nanosilica increases the phase separation temperature, suggesting that the
phase stability of PLA/PMMA blends is enhanced by the presence of nanosilica. Similar
effect induced by silica particles on the phase separation temperature was also observed by
other researchers [Lipatov et al. (2002), Huang et al. (2005)].
122 7 PLA/PMMA/silica nanocomposites
7.4.4 Molecular entanglement
As discussed above, the presence of nanosilica in PLA/PMMA blend can change its phase
behavior. It has been demonstrated that the preferential adsorption between nanosilica and one
of the components in PLA/PMMA blends will lead to the redistribution of the components in
the matrix bulk. According to previous work [Wu (1987), Zhang et al. (2015)] the molecular
entanglement in miscible blend is significantly impacted by the blend composition. Therefore,
the molecular entanglement in PLA/PMMA blend will be influenced by the incorporation of
nanosilica.
Figure 7.13 Storage modulus 𝐺′ and damping 𝑡𝑎𝑛 as a function of angular frequency 𝜔 for
P/P/Si x at 140 °C.
Figure 7.12 shows the master curves of PLA/PMMA/silica nanocomposites at a reference
temperature of 200 °C. The plateau modulus 𝐺𝑁0 can be determined by two methods: (I) 𝐺𝑁
0
equals to 𝐺′ where 𝑡𝑎𝑛 reaches a minimum using Equation (4.19); (II) 𝐺𝑁0 can be calculated
by the crossover modulus using Equation (4.21). As shown in Figure 7.13, 𝑡𝑎𝑛𝑚𝑖𝑛 of P/P/Si x
can be obtained at 140 °C. The entanglement molar mass and entanglement density of P/P/Si
x determined by both methods are listed in Table 7.1.
7.4 Rheological properties of PLA/PMMA/silica nanocomposites 123
Table 7.1 Plateau modulus 𝐺𝑁0 , entanglement molar mass 𝑀𝑒 and entanglement density 𝑒
of
P/P/Si x.
𝑮𝑵𝟎 a
(×105 Pa
)
𝒆 a
(×10-4
mol/cm3)
𝑮𝑵𝟎 b
(×105 Pa
)
𝒆 b
(×10-4
mol/cm3)
P/P/Si 0 10.4 ± 0.2 3.3 ± 0.2 8.8 ± 0.1 3.2 ± 0.1
P/P/Si 2 11.3 ± 0.1 3.5 ± 0.1 9.4 ± 0.1 3.4 ± 0.1
P/P/Si 5 12.9 ± 0.1 4.1 ± 0.1 10.4 ± 0.1 3.8 ± 0.1
P/P/Si 10 16.7 ± 0.2 5.2 ± 0.2 14.2 ± 0.1 5.1 ± 0.1 a
Estimated by using “crossover modulus-based methods” of Equation (4.21) at 200°C. b
Estimated by using “MIN methods” of Equation (4.19) at 140°C.
The relationship between entanglement density 𝑒 and nanosilica content are plotted in Figure
7.14. The values of 𝑒 can be calculated via Equation (4.19) and Equation (4.21). It is found
that the 𝑒 values calculated by “crossover modulus method” are a little higher than that
obtained by “MIN method” due to the semi-quantitative nature of the former method.
However, the results show the same variation tendency with the addition of nanosilica. It is
evident that 𝑒 exhibits a linear increase with silica content.
Figure 7.14 Entanglement density 𝑒 versus weight fraction of nanosilica for P/P/Si x.
The increase of entanglement density in PLA/PMMA/nanosilica mixture could be attributed
to (i) the redistribution of components around the surface of silica and the bulk matrix, (ii)
nanosilica particles also work as entanglement netpoints in the nanocomposites.
124 7 PLA/PMMA/silica nanocomposites
According to the discussion in section 6.1 and 6.2, the entanglement network in miscible
blend is significantly influenced by the molecular structure of the components and blend
composition. As stated previously, the incorporation of nanosilica into PLA/PMMA blends
can change the composition around the surface of silica and in the blend matrix due to
selective interactions (or preferential adsorption) between the surface of nanosilica and PLA
molecules. The original composition in blend matrix is 50/50. However, in the nanosilica
filled PLA/PMMA blends, the “crater” structure around silica agglomerate is rich in PLA
molecules, while the bulk matrix among the “crater” structures is rich in PMMA molecules.
As shown in Figure 7.2, the increase of nanosilica content increased the amount of “crater”
structure, i.e., the redistribution of molecular chains is exacerbated due to the addition of
nanosilica. Our previous investigations also demonstrated the dependence of molecular
entanglement on the blend composition in PLA/PMMA blends (See section 6.1). Therefore,
the entanglement network in PLA/PMMA blends is significantly influenced by the addition of
nanosilica. In addition, nanosilica particles with large specific surface area could strongly
interact with polymer chains. These nanoscale rigid particles could also play the role of
netpoints similar to the molecular entanglements. Therefore, nanosilica increased apparently
the entanglement density in PLA/PMMA/nanosilica mixture.
7.4.5 Creep and recovery experiment
In order to investigate the influence of nanosilica on the melt elasticity of PLA/PMMA blends,
the creep and recovery experiments of P/P/Si x were carried out at 200 °C. As shown in
Figure 7.15a, the creep compliance is decreasing with the addition of nanosilica,
corresponding to the increase of viscosity. For P/P/Si 0, P/P/Si 2 and P/P/Si 5, a slope of 1 in
the double logarithmic plot is reached with the chosen creep time of 2000 s. However, the
creep compliance of P/P/Si 10 approaches to a constant value with the increase of creep time.
This constant is named as elastic equilibrium compliance 𝐽𝑒 , which is induced by the
7.4 Rheological properties of PLA/PMMA/silica nanocomposites 125
formation of silica network in the matrix. The melt elasticity of PLA/PMMA/nanosilica
composites is increasing with silica concentration up to 10 wt%. As shown in Figure 7.15b,
the zero-shear viscosity of P/P/Si 10 is non-existent due to the formation of silica network.
Therefore, it can be deduced that the rheological percolation threshold φ𝑐 of
PLA/PMMA/nanosilica is around 10 wt% silica content.
Figure 7.15 (a) Creep and recovery curves of nanosilica for P/P/Si x performed at 200 °C,
and (b) complex viscosity |𝜂∗| as a function of ω in comparison with 𝑡/𝐽(𝑡𝑐𝑟) in dependence
on creep time 𝑡 at 200 °C.
7.5 The influence of nanosilica on the shape memory properties of
uniaxially stretched PLA/PMMA blends
The method described in section 6.3 is used to study the impact of nanosilica on the shape
memory performances of the symmetric PLA/PMMA blends modified by nanosilica. The
glass transition temperature of each sample is chosen as stretching temperature to avoid the
influence induced by temperature. The stress-strain curves of PLA/PMMA/nanosilica
composites determined at their respective 𝑇𝑔 are shown in Figure 7.16. For each sample, a
constant stress level was reached after the yield point at about 3% deformation, and no
126 7 PLA/PMMA/silica nanocomposites
relaxation could be observed during the stretching process. In addition, the stretching stress is
decreasing with the addition of nanosilica.
Figure 7.16 Stress-strain curves of PLA/PMMA/silica nanocomposites stretched at their
respective 𝑇𝑔 with a strain rate of 0.02/s.
The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 are calculated and listed in Table 7.2. It
was found that the incorporation of nanosilica has no influence on the shape fixing ratio of the
composites. According to the investigations in section 6.3, the shape fixing ratio is mainly
determined by the stretching temperature and strain rate. Moreover, a decrease of shape
recovery ratio is observed with the incorporation of nanosilica. The PLA/PMMA/silica
nanocomposites with 10 wt% nanosilica exhibit the lowest shape recovery ratio. Although the
incorporation of nanosilica increased apparently the entanglement density of PLA/PMMA
blend, the crystallinity was also increased due to the selectively dispersion of PLA chains
around silica surface. The strain induced crystallinity and relatively high 𝑇𝑔 (close to the cold
crystallization temperature) will significantly increase the crystallinity after stretching.
Therefore, the incorporation of nanosilica would reduce the shape recovery ratio of
PLA/PMMA 50/50 blends due to the increased crystallinity.
7.5 Shape memory properties of PLA/PMMA/nanosilica mixtures 127
Table 7.2 The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of PLA/PMMA 50/506N, 7N,
8N blend films stretched at their respective 𝑇𝑔.
𝑻𝒔 (°C) 𝑿𝒄a(%) 𝑹𝒇 𝑹𝒓 𝑿𝒄
b (%) 𝑻𝒓 (°C)
P/P/Si 0 75 0.1 96.2 1.3 98.6 0.7 0.2 85
P/P/Si 2 77 0.1 96.7 2.5 98.1 1.2 0.5 87
P/P/Si 5 79 0.2 96.5 0.4 97.4 1.5 1.1 89
P/P/Si 10 80 0.5 96.9 1.6 95.7 2.5 1.8 90 a
Evaluated by DSC before stretching. b
Evaluated by DSC after stretching.
7.6 The shape memory of biaxially stretched films
In order to produce shape memory films for industrial application, a biaxial stretching
equipment is used to produce films. In general, biaxially stretched films based on PLA are
usually used for packaging application, and PLA4032D is a grade designed for realization of
films. As discussed above, the addition of PMMA into PLA could increase its shape memory
properties, but still reduce the film formation ability. Hence, a PLA/PMMA 7N 80/20 blend
with low PMMA content is chosen to produce biaxially stretched films. For convenience,
PLA/PMMA 7N 80/20 blends can be abbreviated to P/20P. In addition, PLA/PMMA 7N
80/20 blends filled by 2 wt% silica 300 is designed as P/20P/Si 2.
In this section, neat PLA, P/20P and P/20P/Si 2 were chosen to produce biaxially stretched
films. The glass transition temperatures of these three samples are 60, 63, 64 °C, respectively.
When the cast films are stretched around 𝑇𝑔, only small stretch ratio (≤ 22) could be reached
without film tearing. In consideration of the small difference between samples’ 𝑇𝑔 , same
stretching temperature (𝑇𝑠= 80/90 °C) was chosen in this work to draw films at a strain rate of
20%/s up to different stretch ratios .
In order to measure the shape memory properties of biaxially stretched films, the cast films
were treated under a specific temperature-deformation program shown in Figure 7.19. The
128 7 PLA/PMMA/silica nanocomposites
recovery temperature is 10 °C higher than the stretching temperature 𝑇𝑠. The shape fixing
ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 can be determined based on Equations (4.11 and 4.12), as
shown in Table 7.3.
Figure 7.19 Schematic diagram for shape memory test of biaxially stretched films.
Figure 7.20 shows the nominal stress-strain curves of the cast films made of PLA, P/20P and
P/20P/Si 2 simultaneously stretched at 80 °C with a strain rate of 20%/s. The small deviation
between the curves of machine direction (MD) and transverse direction (TD) stresses is
induced by the mechanical artifact, and this deviation is not related to the pre-orientation of
cast films [Capt et al. (2001)]. As expected, a highly non-linear stress-strain response was
observed when the biaxially stretched films were in the rubbery state. It is evident that the
neat PLA has the lowest modulus, flow stress, yield stress and strain hardening stress. With
the incorporation of PMMA and nanosilica, the modulus and stresses increase, P/20P/Si 2
7.6 The shape memory of biaxially stretched films 129
nanocomposites exhibit the highest values. Moreover, pronounced stain hardening behavior
could be observed after 150% deformation (around stretch ratio 2.52.5) at 80 °C, and the
strain hardening modulus is enhanced by the addition of PMMA and nanosilica.
Figure 7.20 Stress-strain curves for the cast films made of PLA, P/20P and P/20P/Si 2
stretched up to different ratios at 80 °C with a strain rate of 20%/s.
When the samples are stretched at 80 °C, 3.53.5 is the maximum stretch ratio that could be
reached without the sample tearing. As the deformation temperature increase, the modulus
and stresses decreases, consistent with the previous reports [Menary et al. (2012)]. As shown
in Figure 7.21, the maximum stretch ratio for films stretched at 90 °C increased to 44, and
the strain hardening behavior occurred at higher strain (around stretch ratio 3.53.5),
indicating the relaxation behavior at higher temperature. Interestingly, PLA displayed the
largest strain hardening modulus while P/20P showed the smallest value, corresponding to the
crystallinity of the samples after stretching (as shown in Table 7.3). This result demonstrated
that the strain hardening behavior occurring at the high temperatures was related to the
recrystallization of the samples during stretching.
130 7 PLA/PMMA/silica nanocomposites
Figure 7.21 Stress-strain curves for the cast films stretched up to the maximum stretch ratio at
80 and 90 °C with a strain rate of 20%/s.
As discussed in section 6.31, a higher stretching temperature results in a better shape fixing
ratio. The stretching temperatures used in this work (80/90 °C) are much higher than 𝑇𝑔, and
the stretched films were cooling to room temperature under the fixation of stretcher’s clamps.
Therefore, a high shape fixing ratio around 99% is obtained for all samples.
As shown in Table 7.3, the shape recovery ratios of three films show different dependencies
on the stretching temperature and stretch ratio. When the cast films are stretched at 80 °C, the
stretch ratio exhibits more significant impact on the shape recovery ratio of neat PLA than for
P/20P and P/20P/Si 2 due to the higher crystallinity of neat PLA. For neat PLA films, the
increase of stretch ratio could dramatically reduce the shape recovery ratio. While for PLA,
P/20P and P/20P/Si 2, the impact of stretch ratio on the shape recovery ratio is relatively small
in comparison with neat PLA. Moreover, P/20P exhibits higher 𝑅𝑟 than P/20P/Si 2, implying
that the presence of nanosilica reduces the shape memory property of PLA/PMMA blends.
This result is in good agreement with the conclusions obtained at section 7.5.
7.6 The shape memory of biaxially stretched films 131
Table 7.3 The shape fixing ratio 𝑅𝑓 and shape recovery ratio 𝑅𝑟 of the cast films made of
PLA, P/20P and P/20P/Si 2 stretched at 80 or 90 °C with a strain rate of 20%/s up to different
ratios (), and the crystallinity 𝑋𝐶 of the cast films before and after stretching.
𝑿𝑪a 𝑻𝒔 (°C) 𝑹𝒇 (%) 𝑹𝒓 (%) 𝑿𝑪
b
PLA 10.8%
80
22 99 84.1 0.3 15.8%
33 99 52.5 0.1 24.5%
3.53.5 99 42.4 0.4 30.1%
90
22 99 78.7 0.3 17.4%
33 99 40.1 0.4 26.5%
3.53.5 99 36.6 0.2 33.3%
44 99 33.5 0.3 35.6%
P/20P
2.3%
80
22 99 98.1 0.1 3.4%
33 99 95.6 0.2 4.9%
3.53.5 99 94.4 0.2 6.8%
90
22 99 85.2 0.2 4.1%
33 99 46.2 0.3 6.2%
3.53.5 99 41.5 0.2 7.5%
44 99 38.8 0.4 8.2%
P/20P/Si
2
3.5%
80
22 99 97.0 0.1 4.5%
33 99 89.7 0.3 5.8%
3.53.5 99 80.3 0.1 7.9%
90
22 99 79.2 0.2 5.2%
33 99 43.3 0.3 8.1%
3.53.5 99 38.3 0.5 8.9%
44 99 36.6 0.2 9.3% a crystallinity of the cast films before stretching.
b crystallinity of the cast films after stretching.
When the films are stretched at 90 °C, the impact of stretch ratio on 𝑅𝑟 is more significant for
all the films compared with lower temperatures. A higher stretch ratio (>22) could reduce 𝑅𝑟
notably. This effect is attributed to the increased crystallinity (as shown in Table 7.3) and
longer time for stress relaxation (as shown in Figure 7.21) when films are stretched up to
higher ratios.
Therefore, in order to get films with good shape recovery, the stretching temperature should
be fixed around 𝑇𝑔, and the deformation should be lower than 100%. The higher temperatures
or larger stretch ratios will more significantly increase the crystallinity of the films and
132 7 PLA/PMMA/silica nanocomposites
improve their stress relaxation behavior during stretching process, but eventually reduce the
shape recovery ratio.
7.7 Conclusions
PLA/PMMA 50/50 blend filled by nanosilica in various concentrations were prepared via
melt blending. The results of the thermo-mechanical analysis demonstrated the miscibility of
the nanocomposites in the solid state, and the broadening of 𝑇𝑔 revealed that the incorporation
of nanosilica increased the concentration fluctuations and dynamic heterogeneity in
PLA/PMMA blends. A distinct “crater” structure at the cross section of
PLA/PMMA/nanosilica mixtures was observed by SEM, and the dimension of the “crater”
structure was reduced markedly with the increased nanosilica content. The preferential
interactions of nanosilica and the components of the blend were predicted by oscillatory shear
experiments. The results revealed that nanosilica has stronger interactions with PLA
molecules in comparison to PMMA molecules. Therefore, PLA molecules were selectively
adsorbed on the surface of nanosilica in the PLA/PMMA/nanosilica mixtures. According to
the analysis from TTS and Han plot, it could be concluded that the incorporation of nanosilica
increased the phase separation temperature in comparison with the unfilled PLA/PMMA
blend, and also enhanced its phase stability. The entanglement network formed in
PLA/PMMA blend was also influenced by the presence of nanosilica. The values of 𝐺𝑋 and
𝐺𝑁0 of the nanocomposites measured at 200 °C exhibited similar linear growth with the
addition of nanosilica. Moreover, the increase of silica content resulted in a linear decrease of
𝑀𝑒 and a linear increase of 𝑒. A possible mechanism for this result was proposed based on
the preferential adsorption of PLA molecules on the surface of nanosilica, and this
preferential adsorption resulted in the change of blend compositions in the “crater” structures
and the bulk matrix among the “crater” structures.
The influence of nanosilica on the shape memory property of symmetric PLA/PMMA blend
was also investigated. It was found that the addition of nanosilica reduced the shape recovery
ratio due to the increased crystallinity. Moreover, the shape memory performances of
134 8. Summary and Outlook
biaxially stretched films made of neat PLA, PLA, P/20P and P/20P/Si 2 were also
investigated. The larger stretch ratios could significantly reduce the shape recovery ratio,
especially at higher temperatures. The reduced shape recovery ratio is attributed to the longer
time to crystallization and stress relaxation when the films are stretched up higher ratios.
135
8. Summary and Outlook
Summary
Polylactide (PLA) is one of the most promising biopolymers that are made from renewable
sources. It is nontoxic to the human body and the environment and, therefore, highly
biocompatible. Although these properties make PLA to be widely used as food packaging or
biomedical materials, etc. some drawbacks, such as low heat resistance, poor thermal stability
and mechanical resistance hinder the application of PLA.
Mixing of inorganic particle into polymer material is widely used to tailor polymers physical
properties and processibility, in which the size and concentration of the former play a critical
role in polymer processing. In this study, rigid spherical silica particle with varying average
primary particle size (7 nm, 40 nm, and 9 μm) were used to reinforce polylactide (PLA) by
melt compounding.
The dependence of rheological properties of PLA/silica composites on the silica size and
concentration were examined by dynamical mechanical and creep-recovery experiments. Our
results demonstrate that mixing of silica particles into PLA matrix could increase the thermal
stability of PLA. Oscillatory shear tests in the linear viscoelastic regime revealed a strong
concentration-dependent behavior for the storage and loss moduli, and the complex viscosity
of the PLA/silica composites by the addition of nanosilica, while these properties were
slightly affected by the addition of microsilica at low frequencies. A linear relationship
between the silica concentration and the logarithm of zero shear viscosity (log0) of the
composites is found when the concentration is below the rheological percolation threshold,
whereas the growth rate is inversely influenced by the silica particle size. Creep-recovery
experiments indicated that the elastic properties of PLA were more sensitive to the addition of
silica than the viscous properties. Even for microsilica, a remarkable enhancement of the
elastic properties was found at low silica concentration. A model based on the radius of
136 8. Summary and Outlook
gyration of polymer matrix 𝑅𝑔 and the mean distance between particles 𝐷 is proposed to
describe the interactions between polymer matrix and particles. When 𝐷 is larger than 2𝑅𝑔,
the particles-polymer interactions are suggested to be responsible for the rheological
properties. On the other side, when 𝐷 is smaller than 𝑅𝑔, a silica network will be formed and
the rheological properties of the composites are dominated by the interactions of particle-
particle and particle-polymer.
The second object of this work is about the thermo-mechanical and rheological properties of
PLA/PMMA blends. The PLA/PMMA blend system is a typical miscible semi-crystalline/
amorphous polymer blend with shape memory potential and has received increasing interest
in recent years. With the incorporation of PMMA into PLA, a broad glass transition and
increased glass transition temperature 𝑇𝑔 fitting partially with the Lodge-McLeish model were
observed by differential scanning calorimetry (DSC) measurements. The broadening of glass
transition is attributed to the local nanoscale heterogeneities in the miscible blend system
which is related to the “self-concentration” of the components. The degree of molecular
entanglement of the blends was derived based on the oscillatory rheological measurements,
showing that the dissimilar chains are more likely to entangle with each other than the similar
ones, and the entanglement density 𝑒 is enhanced with increased PMMA content up to 50%
where a 100% recovery of the initial shape is yielded. The influences of stretching
temperature, strain rate, blend composition and molar mass on the shape memory
performance of PLA/PMMA blends were also investigated. It was found that the shape
memory properties of PLA/PMMA blends present a significant “temperature memory effect”.
The mechanism underlying the shape memory property is further validated by performing the
shape memory test on the PLA/PMMA blend films at their respective 𝑇𝑔. It is evident that for
the semi-crystalline blends (PLA rich), PLA crystallites and molecular entanglement provide
physical cross-links for shape recording, while a negative effect of crystallinity on the shape
137
recovery ratio is obtained due to the strain-induced crystallization and chain slippage between
the crystalline and amorphous chains. On the other hand, for the amorphous blends (PMMA
rich), the shape recovery ratio shows a strong exponential dependence on 𝑒 , and the
entanglement network is regarded as the most important factor in the shape memory
performance.
Nanosilica filled PLA/PMMA 50/50 blends with various silica concentrations (2, 5, 10 wt%)
were prepared by melt mixing. In this section, the effects of silica concentration and the
preferential adsorption between silica and the components on the thermo-mechanical and
rheological properties of PLA/PMMA blends were systematically investigated. The results of
DSC indicate the miscibility of PLA/PMMA/silica nanocomposites in the solid state, and the
incorporation of nanosilica does not only increase the glass transition temperature but also
extend the broadness of glass transition. Local nanoscale heterogeneities in the miscible
blends are proposed due to the self-concentration of the components, and the presence of
nanosilica increased the concentration fluctuation and dynamic heterogeneity of the
PLA/PMMA blends. A distinct “crater” structure can be observed in the fractured surfaces of
the blend nanocomposites. According to the results of SEM and rheological measurements, it
was proposed that PLA molecules were selectively adsorbed on the surface of nanosilica, and
the preferential adsorption of PLA changed the blend composition in the “crater” structure
and the bulk matrix. The phase separation temperature of the filled PLA/PMMA blends was
also increased by the incorporation of nanosilica, implying a potential role of nanosilica in
improving the phase stability of PLA/PMMA blends. In addition, nanosilica could apparently
increase the entanglement density and decrease the entanglement molecular weight of the
blends in a concentration-dependent manner. Finally, the influence of silica nanoparticles on
the shape memory performances of PLA/PMMA blends was investigated. Biaxailly stretched
138 8. Summary and Outlook
films were also produced to investigate the influences of stretch ratio and temperature on the
shape memory properties.
Outlook
For the purpose of application of shape memory films, it would be interesting to study the
influences of stretching mode, temperature, speed, ratio and blend composition on the shape
memory properties of PLA/PMMA films.
The diffusion process at the polymer/polymer interface of bilayer system based on PLA and
PMMA with varying molar masses will be investigated by small-amplitude oscillatory shear
measurements. We will study the inter-diffusion kinetics as well as the development of
interphase for symmetrical and asymmetrical bilayers based on PLA and PMMA in the
further work.
9. Summary (in German)
Polylactid (PLA) ist eines der vielversprechendsten Biopolymere, welche aus erneuerbaren
Quellen hergestellt werden. Es ist nicht toxisch für den menschlichen Körper und die Umwelt
und daher in hohem Maße biokompatibel. Obwohl PLA aufgrund dieser Vorteile weithin für
Lebensmittelverpackungen oder biomedizinische Materialien usw. verwendet werden könnte,
behindern einige Nachteile wie geringe Wärmebeständigkeit, geringe thermische Stabilität
und mechanische Festigkeit die Anwendung von PLA.
Mischungen von anorganischen Partikeln mit Polymermaterialien werden am häufigsten
verwendet, um ein prozessierbares Polymer mit den gewünschten physikalischen
Eigenschaften maßzuschneidern, wobei die Größe und die Konzentration der Partikel eine
kritische Rolle bei der Polymerverarbeitung spielen. In dieser Arbeit wurden starre,
sphärische Silicapartikel mit unterschiedlicher Größe (durchschnittliche Durchmesser von 7
nm, 40 nm, und 9 μm) verwendet, um PLA durch Schmelzmischen zu verstärken. Die
Abhängigkeit der rheologischen Eigenschaften der PLA/Silica- Komposite von Größe und
Konzentration der Silica ist durch dynamisch-mechanische Experimente und Kriech-Erhol-
Versuche ermittelt worden. Es wurde festgestellt, dass die Zugabe von Silica-Partikeln zur
PLA-Matrix die thermische Stabilität des PLA verbessert und dass kleinere Silikapartikel zu
höherer Stabilität führen. Oszillatorische Scherversuche im linear-viskoelastischen Bereich
zeigten eine deutliche Konzentrationsabhängigkeit des Speichermoduls, des Verlustmoduls
und der komplexen Viskosität der PLA/Nanosilica-Komposite im Vergleich zu
PLA/Mikrosilica Komposite.
Der logarithmisch des Nullscherviskosität aller Verbundstoffe stieg linear mit der Zugabe von
Silica an, solange wenn die Konzentration unterhalb der rheologischen Perkolationsschwelle
lag und die Wachstumsrate war umgekehrt proportional zur Partikelgröße. Kriech-Erhol-
Versuche zeigten, dass hinsichtlich der Zugabe von Silica die elastischen Eigenschaften des
140 9. Summary (in German)
PLA empfindlicher sind als die viskosen Eigenschaften. Sogar für Mikrosilica wurde eine
bemerkenswerter Verbesserung der elastischen Eigenschaften bei niedrigen Silica-
Konzentration gefunden. Es wurde ein Modell entwickelt, um die Interaktion zwischen
Polymermatrix und Teilchen zu beschreiben. Wenn der durchschnittliche Abstand zwischen
den Teilchen weit oberhalb des GyrationsdurchmesserS der Polymermatrix liegt, werden die
rheologischen Eigenschaften nur von Teilchen-Polymer-Wechselwirkungen beeinflusst.
Wenn der durchschnittlichen Abstand zwischen den Teilchen hingegen unterhalb des
Gyrationsradius der Polymermatrix liegt, wird eine Siliciumdioxid-Netzwerkstruktur gebildet
und die Teilchen-Teilchen- und Teilchen-Polymer-Wechselwirkungen scheinen die
rheologischen Eigenschaften zu bestimmen.
Das zweite Thema dieser Arbeit le handelt die thermo-mechanischen und rheologischen
Eigenschaften der PLA/PMMA-Mischung. Das PLA/PMMA-Blendsystem ist eine typische
mischbare teilbkristalline/amorph Polymermischung mit Formgedächtnispotenzial und hat in
den letzten Jahren zunehmendes Interesse erlangt. Bei der Verwendung von PMMA in PLA
wurden ein breiter Glasübergang und eine erhöhte Glasübergangstemperatur 𝑇𝑔 mit Hilfe des
Lodge-McLeish Modells und DSC Messungen beobachtet. Die Verbreiterung des
Glasübergangs wird durch lokale, nanoskalige Heterogenitäten in der Mischung verursacht,
die mit der Selbst-Konzentration der Komponenten zusammenhängen. Der Grad der
molekularen Verschlaufung der Mischungen wurde auf Basis der oszillatorischen
rheologischen Messungen bestimmt und zeigte, dass ungleiche Ketten einfacher als gleiche
Ketten miteinander verschlingen und dass die Verschlaufungsdichte 𝑒 mit erhöhter PMMA-
Konzentration (bis zu 50%) zunahm, wobei eine 99% Rückstellung auf die ursprüngliche
Form erhalten wurde. Die Einflüsse der Recktemperatur, Reckgeschwindigkeit,
Gblendzusammensetzung und des Molekulargewichts auf die Form-Gedächtnisleistung von
PLA/PMMA-Blends wurden ebenfalls untersucht. Es wurde festgestellt, dass die
141
Formgedächtniseigenschaften von PLA/PMMA Blends einen erheblichen "Temperatur-
Gedächtnis-Effekt" aufweisen.
Der Formgedächtnismechanismus wird weiterhin durch die Durchführung von
Formgedächtnistests an PLA/PMMA-Blend-Folien bei den jeweiligen 𝑇𝑔 validiert. Es ist
offensichtlich, dass für die teilkristallinen Mischungen (PLA-reich) PLA-Kristallite und
molekulare Verschlaufungen physikalische Vernetzungen für das Formgedächtnis bieten,
während eine negative Wirkung auf das Formgedächtnis auf Grund der Kristallinität durch die
dehnungsinduzierte Kristallisation und das Abgleiten von Keffen zwischen kristallinen und
amorphen Bereichen zu Stande kommt. Im Gegensatz dazu zeigt das
Formrückgewinnungsverhältnis für die amorphen Blends (PMMA-reich) eine stark positiv
lineare Anslieg mit 𝑒 und die Verschlaufung des Netzwerks wird als einer der wichtigsten
Faktoren im Bezug auf die Formgedächtnisleistung angesehen.
Nanosilica-gefülltes PLA/PMMA (50/50) mit verschiedenen Silica-konzentrationen (2, 5, 10
wt%) wurden durch Schmelzmischen hergestellt. In diesem Abschnitt wurden die
Einflussfaktoren, wie beispielsweise Silica-Konzentration, die Verteilung der Nanosilica, die
spezifischen Wechselwirkungen und bevorzugte Adsorption zwischen Nano-Siliciumdioxid
und der Komponente, auf die thermo-mechanischen und rheologischen Eigenschaften von
PLA/PMMA-Blends systematisch untersucht. DSC-Versuche zeigen die Mischbarkeit von
PLA/PMMA/Silica-Nanokompositen im festen Zustand, und die Verwendung von Nanosilica
erhöht nicht nur die Glasübergangstemperatur, sondern erweitert auch die Breite des
Glasübergangs. Aufgrund der Selbst-Konzentration der Komponenten können lokale
nanoskalige Heterogenitäten in den mischbaren Blends vermutet werden, und die
Anwesenheit von Nanosilica erhöht die Konzentrationsschwankungen und die dynamische
Heterogenität in den PLA/PMMA-Blends. Eine deutliche "Krater"-Struktur ist in den
Bruchflächen der Blend-Nanokomposite zu beobachten. Im Zusammenhang mit den
142 9. Summary (in German)
Ergebnissen der SEM und rheologischen Messungen wurde vorhergesagt, dass PLA Moleküle
selektiv auf der Oberfläche von Nano Silica adsorbiert werden, und die bevorzugte
Adsorption von PLA veränderte die Mischungszusammensetzung in der "Krater"-Struktur
und der reinen Matrix. Darüber hinaus wurde die Phasentrennungstemperatur der gefüllten
PLA/PMMA-Blends im Vergleich zu den ungefüllten Mischungen verbessert , und es zeigte
sich, dass die Zugabe von Nanosilica die Phasenstabilität der PLA/PMMA-Blends verbessert.
Entsprechend der auf rheologischen Messungen basierenden Berechnungen wurde aufgrund
der Umverteilung von PLA und PMMA-Molekülen in der Masse und rund um die Silica-
Oberfläche ein linearer Anstieg der Verschlaufungsdichte mit steigendem Silica-Gehalt
erreicht. Schließlich ist der Einfluß der Silica-Nanopartikel auf die Formgedächtnisleistung
von PLA/PMMA-Blends untersucht worden. Biaxial gestreckte Folien wurden ebenfalls
hergestellt, um die Einflüsse des Streckungsverhältnisses auf die
Formgedächtniseigenschaften zu untersuchen.
10. Appendix
10.1 Reproducibility of rheological measurements
For rheological measurement, it is important to check the reproducibility of rheological
measurements for neat and filled polymers. Frequency sweeps and creep-recovery
experiments for neat PLA and PLA with 2.8 vol. % nanosilica are used to check the
reproducibility of the rheological measurement in the linear region deformation.
Figure 10.1 The oscillatory frequency sweep curves of neat PLA and PLA/2.8 vol. % silica
OX50 nanocomposites measured at 180 °C.
For the rheological measurements in this work, at least three samples are used for each test.
As shown in Figure 10.1, the data of the storage modulus, loss modulus and complex
viscosity exhibit excellent reproducibility of neat PLA and PLA/silica nanocomposites. The
deviations between the results of two individual samples are less than 10%, indicating the
high reproducibility of oscillatory frequency sweeps. In addition, the creep-recovery
experiments also show excellent reproducibility, as shown in Figure 10.2.
144 10 Appendix
Figure 10.2 The creep-recovery curves of neat PLA and PLA/2.8 vol. % silica OX50
nanocomposites measured at 180 °C.
10.2 The melt density of PLA and PMMA at 200 °C
In order to calculate the molecular entanglement of PLA/PMMA blends in the molten state,
the melt density of PLA and PMMA should be measured. In this work, the melt flow rate
(MVR) at per10 min was measured by melt flow indexer (MI-4, Göttfert, Germany), and then
the samples were weighted. The melt density can be calculated by the ratio of MVR and the
mean weight per 10 min.
Table 10.1 The melt flow index of PLA and PMMA measured at 200 °C.
Waiting
time
(min)
MVR
(cm3/10min)
Mean
weight
(g/10min)
Density
(g/cm3)
PLA
200 °C, 2.16 Kg 5 4.06 4.54 1.12
200 °C, 2.16 Kg 10 4.83 5.46 1.13
200 °C, 2.16 Kg 20 4.64 5.20 1.12
PMMA
200 °C, 3.8 Kg 5 2.17 2.45 1.13
200 °C, 3.8 Kg 10 1.99 2.25 1.13
200 °C, 3.8 Kg 20 2.05 2.32 1.13
145
10.3 Thermogravimetric analysis (TGA) of nanocomposites
The thermal decomposition of neat PLA and silica particles was investigated as a function of
temperature under N2 atmosphere. As shown in Figure 10.3a, the onset degradation
temperature for neat PLA is around 300 °C, and that for silica particles are far above 400 °C.
Therefore, the chosen processing temperature (180 °C) for is below the thermal degradation
temperature of all the compositions. On the other hand, there is no change of weight can be
observed below 100 °C, suggesting that the drying process has removed the moisture of the
samples.
Figure 10.3 TGA scans of dried PLA and silica particles. The samples were heated at
5 °C/min from 20 °C to 500 °C under N2 atmosphere.
A comparative TGA of neat PLA and its composites with 2.8 vol. % silica is shown in Figure
10.3b. It is evident that the onset degradation temperature of the composites is larger than that
of neat PLA, indicating the PLA/silica composites have better thermal stability in comparison
with neat PLA. This enhancement may result from the interactions between organic and
inorganic phases [Fu and Qutubuddin (2001)]. The size of silica particle also exhibits impact
on the thermal stability, and silica with smaller particle size results in a larger enhancement of
the onset degradation temperature.
146 10 Appendix
In addition, the real silica concentration after sample preparation was also investigated by
TGA. As shown in Figure 10.3b, the real concentrations of the composites are close to the
desired silica contents. As the deviations between the real and the desired silica concentration
are very small, the values of the desired concentrations are used in this thesis.
Figure 10.4 TGA curves of PLA/PMMA 50/50 blends with different nanosilica contents at
the heating rate of 10 °C/min, recorded under nitrogen atmosphere.
The thermal stability properties of P/P/Si x nanocomposites was investigated as a function of
temperature, as shown in Figure 10.4. The unfilled PLA/PMMA blend shows the initial
degradation at ca. 290 °C, and this blend almost fully decomposed at 500 °C. In contrast to
unfilled PLA/PMMA blend, the initial degradation temperature of P/P/Si 2 is around 298 °C,
while that for P/P/Si 10 is at ca. 304 °C. All the mixtures displayed the residue of degradation
at 500 °C, and difference of residue is attributed to the different silica contents. Therefore, it
can be concluded that the presence of nanosilica improve the thermal stability of PLA/PMMA
blends.
147
10.4 The stress-strain curves of semi-crystalline polymers during cold or hot-
deformation
According to our previous study, the typical stress-strain curves of semi-crystalline polymers
during cold or hot-deformation can be conclude as shown in Figure 10.5. When the samples
are stretched below 𝑇𝑔, a very linear stress-strain relationship up to a well-defined yield point
can be observed. This linear portion of the curve corresponds to the elastic region and its
slope is the elasticity modulus [Besson et al. (2009)]. After the yield point, the stress
decreases slightly and then a strain hardening behavior occurs as deformation continues.
However, different stress-strain curves can be obtained with the increase of deformation
temperature [Guo-Zheng (2013)].
Figure 10.5 Stress-strain curves of semi-crystalline polymers during cold or hot-deformation
within relatively small strain.
When the films are stretched around 𝑇𝑔, a constant flow stress could be reached after the
elastic region deformation. In this case, the free chains in the amorphous phase work as
switching phase. For semi-crystalline polymer, strain-induced crystallization will result in a
plastic deformation in this region. Therefore, the strain in this region is a combination of
elastic and plastic deformation. There is no strain-hardening occurred due to the relatively
small deformation.
148 10 Appendix
The third type is when the samples are stretched at temperatures far above 𝑇𝑔. A constant flow
stress could be observed after a nonlinear stress-strain region. Crystalline grains are formed in
the nonlinear region, and the following constant stress is a dynamic equilibrium between the
recrystallization and stress relaxation induced by high temperature.
The stress-strain curves of the stretched films could be used to predict the following shape
memory performances. The stress relaxation behavior during the stretching can improve the
shape fixing ratio but reduce the shape recovery ratio. Moreover, the recrystallization and
disentanglement behavior during the stretching process would induce an irreversible plastic
deformation, which could significantly reduce the shape recovery ratio.
Abbreviations and symbols
PLA
PMMA
SiO2
Silica 300
Silica OX50
Silica 63
P/P/Si
P/20P
PP
PE
PVDF
POM
SEM
DSC
TGA
FTIR
SEC
DMTA
BET
THF
wt%
vol. %
TTS
Polylactide
Poly(methyl methacrylate)
Silica
AEROSIL® 300
AEROSIL® OX50
TIXOSIL 63
PLA50%/PMMA50%/silica 300
PLA80%/PMMA20%
Polypropylene
Polyethylene
Polyvinylidenfluorid
Polyoxymethylene
Scanning Electron Microscopy
Differential scanning calorimetry
Thermogravimetric analysis
Fourier transform infrared spectroscopy
Size exclusion chromatography
Dynamic mechanical thermal analysis
Brunauer–Emmett–Teller
Tetrahydrofurane
Weight fraction in %
Volume fraction in %
Time-temperature superposition
150 Abbreviations and symbols
Molar mass
Weight average molar mass
Number average molar mass
Entanglement average molar mass
Critical molar mass
Polydispersity
Activation energy
Glass transition temperature
Cold crystallization peak
Melting temperature
-transition temperature
Density
Density of amorphous phase
Radius of gyration
Average primary particle size
Inter-particle distance
Specific surface area
Crystallinity
Melting enthalpy
Content in the blend
Storage modulus
Loss modulus
Complex modulus
Phase angle
Tangent of the phase angle
Complex viscosity
Zero shear viscosity
Angular frequency
𝑀
𝑀𝑤
𝑀𝑛
𝑀𝑒
𝑀𝑐
𝑀𝑤/𝑀𝑛
𝐸𝑎
𝑇𝑔
𝑇𝑐
𝑇𝑚
𝑇
𝑎
𝑅𝑔
𝑑
𝑋𝑐
𝐻𝑓
𝑤
𝐺′
𝐺′′
𝐺∗(𝜔)
𝑡𝑎𝑛
𝜂∗(𝜔)
𝜂0
𝜔
D
SSA
151
𝑡𝑐𝑟
𝛾(𝑡𝑐𝑟)
𝐽𝑐𝑟(𝑡𝑐𝑟)
𝜓(𝑡𝑐𝑟)
𝐽0
𝜏0
𝐽𝑟(𝑡𝑟)
𝐽𝑒0
𝑅𝑓
𝑅𝑟
𝐺𝑁0
𝑅
𝑇
𝑒
𝑀𝑒
𝐺𝑋
𝐽𝑁0
γ
φ𝑐
Creep time
Time-dependent deformation
Creep compliance
Creep function
Instantaneous elastic compliance
Creep stress
Recoverable compliance
Steady-state recoverable compliance
Shape fixing ratio
Shape recovery ratio
Plateau modulus
Gas constant
Absolute temperature
Entanglement density
Entanglement molar mass
Crossover modulus
Plateau compliance
Strain
Rheological percolation threshold
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Acknowledgements
This thesis resulted from my work as a PhD student at the Institute for Polymer Materials of
the Material Science department at the Friedrich-Alexander-University Erlangen-Nuremberg.
During my study in Erlangen, I have had wonderful time with my colleagues and friends.
Here I would like to thank everybody who has supported my work in the past four years.
First and foremost, I want to thank China Scholarship Council (CSC) for financial support of
my PhD study.
I would like express my sincere appreciation to my supervisors Prof. Schubert for giving me
this opportunity to pursue my PhD in LSP group. He always supported me very well in my
research with patience and encouragement. In the last few months of my PhD study, he also
provided financial support for my work.
I gratefully acknowledge my teacher Dr. Joachim Kaschta for his help and support. He is my
second supervisor of my PhD study. He provided me extensive advice and time for discussion
throughout the whole time working on this thesis. The every step of my work has gotten his
valuable support. I also learned a lot of fundamental knowledge from this project under the
guidance of Dr. Kaschta. Without his help I would not able to finish this project. I also want
to acknowledge Prof. Helmut Münstedt. His work about rheological experiments gives me a
lot of inspiration.
My further gratitude goes to my colleagues Xianhu Liu, Yamin Pan and Hao Wang. We
studied at the same college (Zhengzhou University) before we came to Germany. When I
have problems in my work, they always discuss with me and help me out. Moreover, Xianhu
and Yamin also revised this thesis.
167
Mathias Bechert and Peter Kunzelmann are very nice to me and give me lots of help. I
appreciate Mathias Bechert and Jie Xu for revising my summary in German.
I would like to thank my office colleagues Yaping Ding, Atheer Alaa Abdulhussein, Jie Xu
and Volker Achenbach for our good collaboration in a nice and friendly atmosphere.
I am grateful to Dr. Zdeněk Starý for teaching me to do the heat shrinkage test. Magdalena
Papp is very kind to me and she is thanked for teaching me to use SEM. Inge Herzer is
thanked for DSC and GPC test. Jennifer Reiser taught me to use FTIR and also helped to do
TGA test. Susanne Michler is thanked for ordering the materials for this project and helping
me for rheological and DMTA measurements. Karl and Marko Heyder are thanked for
teaching me to use the extruder and biaxial stretching machines.
I also want to thank my colleagues Michael Härth, Bastian Walter, Franz Lanyi, Florian Küng,
Andreas Ziegmann and Jonas Daenicke. It is great honor to work with them.
Finally, I want to express my deepest gratitude to my family. My parents, my younger sister
and brother give me constant encouragement and support for my study. My husband Jing Han
always accompany with me at my hard time. His love, patience, encouragement and
companionship helped me to complete this thesis successfully.
I love you, LSP. I love you, Germany. It is a very precious experience for me to study in this
great country. I will miss the life and people here forever.
168 Acknowledgements
List of Publication
1. Xiaoqiong Hao, Joachim Kaschta, Xianhu Liu, Dirk W. Schubert. Entanglement network
formed in miscible PLA/PMMA blends and its role in rheological and thermo-mechanical
properties of the blends. Polymer, Volume 80, 2 December 2015, Pages 38-45
2. Xiaoqiong Hao, Joachim Kaschta, Xianhu Liu, Yamin Pan, Dirk W. Schubert,
Intermolecular cooperativity and entanglement network in a miscible PLA/PMMA blend in
the presence of nano-silica. Polymer, Volume 82, 15 January 2016, Pages 57-65
3. Xiaoqiong Hao, Joachim Kaschta, Dirk W. Schubert. Viscous and elastic properties of
polylactide melts filled with silica particles: effect of particle size and concentration.
Composites Part B Engineering, Volume 89, 15 March 2016, Pages 44–53
4. Xianhu Liu, Kun Dai, Xiaoqiong Hao, Guoqiang Zheng, Chuntai Liu. Crystalline structure
of injection molded β-isotactic Polypropylene: analysis of the oriented shear zone, ACS
Applied Materials & Interfaces, 2013, 52 (34), 11996–12002.