Encyclopedia of Polymer Blends, Volume 3 E-Book
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- Herausgeber: Wiley-VCH
- Kategorie: Fachliteratur
- Serie: Encyclopedia of Polymer Blends
- Sprache: Englisch
A complete and timely overview of the topic, this Encyclopedia imparts knowledge of fundamental principles and their applications for academicians, scientists and researchers, while informing engineers, industrialists and entrepreneurs of the current state of the technology and its utilization.
- The most comprehensive source on polymer blends available on the market
- Offers a complete and timely overview of the topic
- Each article presents up to date research & development on a topic and its basic principles and applications, integrates case studies, laboratory and pilot plant experiments, and gives due reference to published and patented literature
- Equips academics, scientists and researchers with knowledge of fundamentals principles and their applications, and informs the engineers, industrialists and entrepreneurs about the state of the art technology and its applications
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Veröffentlichungsjahr: 2016
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Contents
Cover
Related Titles
Title Page
Copyright
Preface
List of Contributors
Chapter 1: Glass-Transition Phenomena in Polymer Blends
1.1 Introduction
1.2 Phenomenology and Theories of the Glass Transition
1.3 Manipulating the Glass Transition
1.4 Experimental Means of Determination
1.5 Blend Morphology and Glass Transitions
1.6 Analyzing Glass Transitions in Single-Phase Systems
1.7 Case Studies
1.8 Concluding Remarks
Abbreviations
References
Chapter 2: Crystallization and Melting Behavior in Polymer Blends
2.1 Introduction
2.2 Miscibility of Polymer Blends
2.3 Miscible Blends
2.4 Immiscible Blends
2.5 Compatibilized Polymer Blends
2.6 Summary and Conclusions
2.7 Nomenclature
References
Chapter 3: Morphology and Structure of Crystalline/Crystalline Polymer Blends
3.1 Introduction
3.2 Systems with Small Melting Point Difference
3.3 Systems with Large Melting Point Difference
3.4 Concluding Remarks
Acknowledgment
References
Chapter 4: Rubber–Plastic Blends: Structure–Property Relationship
4.1 Introduction: Key Challenges
4.2 Rubber Toughening of Thermoplastics
4.3 Models for Rubber Toughening of Plastics
4.4 Characterization of Rubber–Plastic Blends
4.5 Experimental Rubber–Plastic Blends
4.6 Thermoplastic Vulcanizates
4.7 Blends Made during Polymerization
4.8 Conclusions
References
Chapter 5: Morphology of Rubber/Rubber Blends
5.1 Introduction
5.2 Characterization Techniques for Rubber Blends
5.3 Effect of Material Parameters and Processing on Structure and Morphology of Rubber Blends
5.4 Distribution of Fillers and Cure Balance in Rubber Blends
5.5 Morphology and Properties of Different Rubber Blends
5.6 Conclusions
References
Chapter 6: Phase Morphology and Properties of Ternary Polymer Blends
6.1 Introduction
6.2 Miscibility of Polymers in Ternary Polymer Blends
6.3 Formation of Phase Morphology
6.4 Properties of Ternary Polymer Blends
6.5 Conclusions and Future Development
References
Chapter 7: Morphology and Structure of Polymer Blends Containing Nanofillers
7.1 Introduction
7.2 Type of Nanofiller Used in Polymer Nanocomposite
7.3 Nanostructural Characterization
7.4 Partially Miscible Polymer Blends Containing Nanoparticle
7.5 Immiscible Polymer Blends Containing Nanoparticle
References
Index
End User License Agreement
List of Tables
Table 1.1
Table 1.2
Table 1.3
Table 1.4
Table 1.5
Table 1.6
Table 1.7
Table 1.8
Table 2.1
Table 2.2
Table 2.3
Table 2.4
Table 3.1
Table 3.2
Table 5.1
Table 5.2
Table 6.1
Table 7.1
Table 7.2
Table 7.3
Table 7.4
List of Illustrations
Figure 1.1
Figure 1.2
Figure 1.3
Figure 1.4
Figure 1.5
Figure 1.6
Figure 1.7
Figure 1.8
Figure 1.9
Figure 1.10
Figure 1.11
Figure 1.12
Figure 1.13
Figure 1.14
Figure 1.15
Figure 1.16
Figure 1.17
Figure 1.18
Figure 1.19
Figure 1.20
Figure 1.21
Figure 1.22
Figure 1.23
Figure 1.24
Figure 1.25
Figure 1.26
Figure 1.27
Figure 1.28
Figure 1.29
Figure 1.30
Figure 1.31
Figure 1.32
Figure 1.33
Figure 1.34
Figure 1.35
Figure 1.36
Figure 1.37
Figure 1.38
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Figure 1.40
Figure 1.41
Figure 1.42
Figure 1.43
Figure 1.44
Figure 1.45
Figure 1.46
Figure 1.47
Figure 1.48
Figure 1.49
Figure 2.1
Figure 2.2
Figure 2.3
Figure 2.4
Figure 2.5
Figure 2.6
Figure 2.7
Figure 2.8
Figure 2.9
Figure 2.10
Figure 2.11
Figure 2.12
Figure 2.13
Figure 2.14
Figure 2.15
Figure 2.16
Figure 2.17
Figure 2.18
Figure 3.1
Figure 3.2
Figure 3.3
Figure 3.4
Figure 3.5
Figure 3.6
Figure 3.7
Figure 3.8
Figure 3.9
Figure 3.10
Figure 3.11
Figure 3.12
Figure 3.13
Figure 3.14
Figure 3.15
Figure 3.16
Figure 3.17
Figure 3.18
Figure 3.19
Figure 3.20
Figure 3.21
Figure 3.22
Figure 3.23
Figure 3.24
Figure 3.25
Figure 3.26
Figure 3.27
Figure 3.28
Figure 3.29
Figure 3.30
Figure 3.31
Figure 5.1
Figure 5.2
Figure 5.3
Figure 5.4
Figure 5.5
Figure 5.6
Figure 5.7
Figure 5.8
Figure 5.9
Figure 5.10
Figure 5.11
Figure 5.12
Figure 5.13
Figure 5.14
Figure 5.15
Figure 5.16
Figure 6.1
Figure 6.2
Figure 6.3
Figure 6.4
Figure 6.5
Figure 6.6
Figure 6.7
Figure 6.8
Figure 6.9
Figure 6.10
Figure 6.11
Figure 6.12
Figure 6.13
Figure 6.14
Figure 6.15
Figure 6.16
Figure 6.17
Figure 6.18
Figure 6.19
Figure 6.20
Figure 6.21
Figure 6.22
Figure 6.23
Figure 6.24
Figure 6.25
Figure 6.26
Figure 6.27
Figure 6.28
Figure 6.29
Figure 6.30
Figure 6.31
Figure 6.32
Figure 6.33
Figure 6.34
Figure 6.35
Figure 6.36
Figure 6.37
Figure 6.38
Figure 6.39
Figure 6.40
Figure 7.1
Figure 7.2
Figure 7.3
Figure 7.4
Figure 7.5
Figure 7.6
Figure 7.7
Figure 7.8
Figure 7.9
Figure 7.10
Figure 7.11
Figure 7.12
Figure 7.13
Figure 7.14
Figure 7.15
Figure 7.16
Figure 7.17
Figure 7.18
Figure 7.19
Figure 7.20
Figure 7.21
Figure 7.22
Figure 7.23
Figure 7.24
Figure 7.25
Figure 7.26
Figure 7.27
Figure 7.28
Figure 7.29
Figure 7.30
Figure 7.31
Figure 7.32
Figure 7.33
Figure 7.34
Figure 7.35
Figure 7.36
Figure 7.37
Figure 7.38
Figure 7.39
Figure 7.40
Figure 7.41
Figure 7.42
Figure 7.43
Figure 7.44
Figure 7.45
Figure 7.46
Figure 7.47
Figure 7.48
Figure 7.49
Figure 7.50
Figure 7.51
Figure 7.52
Figure 7.53
Figure 7.54
Figure 7.55
Guide
Cover
Table of Contents
Begin Reading
Chapter 1
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Related Titles
Isayev, Avraam I. (ed.)
Encyclopedia of Polymer Blends
Volume 1: Fundamentals
2010
PRINT ISBN: 978-3-527-31929-9
Isayev, Avraam I. (ed.)
Encyclopedia of Polymer Blends
Volume 2: Processing
2011
PRINT ISBN: 978-3-527-31930-5
Elias, H.
Macromolecules
Volume 1: Chemical Structures and Syntheses
2005
Print ISBN: 978-3-527-31172-9
Elias, H.
Macromolecules
Volume 2: Industrial Polymers and Syntheses
2007
Print ISBN: 978-3-527-31173-6
Elias, H.
Macromolecules
Volume 3: Physical Structures and Properties
2007
Print ISBN: 978-3-527-31174-3
Elias, H.
Macromolecules
Volume 4: Applications of Polymers
2008
Print ISBN: 978-3-527-31175-0
Edited by Avraam I. Isayev
Encyclopedia of Polymer Blends
Volume 3: Structure
All books published by Wiley-VCH are carefully produced. Nevertheless, authors, editors, and publisher do not warrant the information contained in these books, including this book, to be free of errors. Readers are advised to keep in mind that statements, data, illustrations, procedural details or other items may inadvertently be inaccurate.
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Print ISBN: 978-3-527-31931-2ePDF ISBN: 978-3-527-65399-7ePub ISBN: 978-3-527-65398-0Mobi ISBN: 978-3-527-65397-3oBook ISBN: 978-3-527-65396-6
Preface
Encyclopedia of polymer blends will include scientific publications in various areas of blends. Polymer blends are mixtures of two or more polymers and/or copolymers. Polymer blending is used to develop new materials with synergistic properties that are not achievable with individual components without having to synthesize and scale up new macromolecules. Along with a classical description of polymer blends, articles in the encyclopedia will describe recently proposed theories and concepts that may not be accepted yet but reflect future development. Each article provides current points of view on the subject matter. These up-to-date reviews are very helpful for understanding the present status of science and technology related to polymer blends.
The encyclopedia will be the source of existing knowledge related to polymer blends and will consist of five volumes. Volume 1 describes the fundamentals including the basic principles of polymer blending, thermodynamics, miscible, immiscible, and compatible blends, kinetics, and composition and temperature dependence of phase separation. Volume 2 provides the principles, equipment, and machinery for polymer blend processing. Volume 3 deals with the structure of blended materials that governs their properties. Volume 4 describes various properties of polymer blends. Volume 5 discusses the blended materials and their industrial, automotive, aerospace, and other high technology applications. Individual articles in the encyclopedia describe the topics with historical perspective, the state-of-the-art science and technology and its future.
This encyclopedia is intended for use by academicians, scientists, engineers, researchers, and graduate students working on polymers and their blends.
Volume 3 is devoted to the structure of blended materials that governs their properties and consists of seven chapters. These chapters cover glass transition phenomena, crystallization and melting behavior, structure–property relationship, morphology and structure of polymer blends, and blends containing various nanofillers. Existing theoretical approaches to describe morphology and structure of blends are extensively discussed. The importance of flow, rheology of components, and rheological aspects of blends is emphasized. These aspects are detailed below and build on each other.
Chapter 1 addresses a number of topics, including general phenomenology, theories, and metrology of the glass transition of polymer blends. The theoretical foundations and practical examples from the analysis of experimental data for miscible systems including binary polymer blends, oligomer/polymer mixtures, and copolymers are critically reviewed. The applicability ranges of important theoretical, semi-empirical, or purely phenomenological mixing rules used for describing the compositional dependence of the glass transition are explored. Examples demonstrating the physical meaning of model parameters are given. A number of case studies involving hydrogen-bonding binary polymer blends and ternary polymer systems are presented. The chapter ends by summarizing general rules that relate the results of glass-transition studies with structural characterizations and miscibility evaluations of polymer blends.
Chapter 2
List of Contributors
Sudhin Datta
ExxonMobil Chemical Co.
GCR-Product Fundamentals
5200 Bayway Drive
Baytown, TX 77520
USA
Avraam I. Isayev
The University of Akron
Department of Polymer Engineering
250 South Forge Street
Akron, OH 444325-0301
USA
Saleh A. Jabarin
University of Toledo
Polymer Institute
Department of Chemical and Environmental Engineering
2801 W. Bancroft Street
Toledo, OH 43606-3390
USA
Ioannis M. Kalogeras
University of Athens
Faculty of Physics
Department of Solid State Physics
15784 Zografos, Athens
Greece
V.N. Kuleznev
Lomonosov State University of Fine Chemical Technology
Prospekt Vernadskogo 86
119571 Moscow
Russia
Tian Liang
The University of Akron
Department of Polymer Engineering
250 South Forge Street
Akron, OH 444325-0301
USA
Elizabeth A. Lofgren
University of Toledo
Polymer Institute
Department of Chemical and Environmental Engineering
2801 W. Bancroft Street
Toledo, OH 43606-3390
USA
Kazem Majdzadeh-Ardakani
University of Toledo
Polymer Institute
Department of Chemical and Environmental Engineering
2801 W. Bancroft Street
Toledo, OH 43606-3390
USA
Yu. P. Miroshnikov
Lomonosov State University of Fine Chemical Technology
Prospekt Vernadskogo 86
119571 Moscow
Russia
Hossein Nazockdast
Amirkabir University of Technology
Department of Polymer Engineering
424 Hafez Ave.
158754413 Tehran
Iran
Zhaobin Qiu
Beijing University of Chemical Technology
State Key Laboratory of Chemical Resource Engineering
15 North Third Ring Road East
Chaoyang District
100029 Beijing
China
Shouke Yan
Beijing University of Chemical Technology
State Key Laboratory of Chemical Resource Engineering
15 North Third Ring Road East
Chaoyang District
100029 Beijing
China
1Glass-Transition Phenomena in Polymer Blends
2Crystallization and Melting Behavior in Polymer Blends
Saleh A. Jabarin, Kazem Majdzadeh-Ardakani, and Elizabeth A. Lofgren
Polymer Institute and Department of Chemical and Environmental Engineering, University of Toledo, Toledo, Ohio 43606-3390
2.1 Introduction
The history of polymer blends can be traced back more than a century to when the first blends of trans- and cis-1,4-polyisoprene (natural rubber with gutta-percha) were prepared in 1846 [1]. Preparing commercial blends of poly(phenylene ether) (PPE) and styrenics in 1965 led to the modern era of polymer blending. Since then, the manufacturing of polymer blends with the desired performances, by combining the unique properties of available polymers, has made such blends commercially important. The crystallization and melting characteristics of polymers and polymer blends determine many of their processing conditions and ultimate physical properties. Early reviews of polymer blends discuss the thermodynamics of mixing, phase separation, preparation, transport phenomena, chemical interactions, physical properties, and some commercial applications [1–5].
The study of polymer blends has been a very important scientific topic during the recent decades, and their applications are still growing because of the lower cost of blending processes compared to that of making new polymers. Many polymer characteristics, such as mechanical and barrier properties, chemical resistance, thermal stability, flame retardancy, and others, can be improved by blending different polymers. Volumes 4 and 5 of this encyclopedia series discuss this topic in more detail.
Important characteristics of polymer blends include the miscibility or immiscibility of their components. The properties such as crystallization and melting behavior of the blends are strongly influenced by the miscibility of the polymers. Most polymer pairs are thermodynamically immiscible; however, miscibility can be achievable with selected polymer combinations. The properties of miscible polymer blends are generally influenced by the chemical composition of each polymer. The fundamentals of miscible, immiscible, and compatible blends are included in Volume 1 of this encyclopedia series.
Many investigations have been conducted to characterize and document various preparation processes and the resultant properties of polymer blends [6–11]. Several books and review articles have also been published in this area of research [2,12–16]; however, fewer publications [12,17] comprehensively discuss the melting behavior and crystallization of polymer blends.
The following review comprises recent developments and progress related specifically to the crystallization and melting characteristics of polymer blends. It contains some basic aspects of crystallization kinetics, semicrystalline morphology, and melting behavior of miscible and immiscible polymer blends. It also includes important information from earlier books and articles and extends it to more recent achievements in this area. The miscibility of polymer blends is discussed in Section 2.2, while melting behavior and crystallization characteristics of miscible and immiscible polymer blends are included in Sections 2.3 and 2.4. Finally, compatibilized polymer blends and their crystallization behavior are discussed in Section 2.5.
2.2 Miscibility of Polymer Blends
Miscibility or compatibility is achievable when two materials can be mixed in any ratio, without the separation of the two phases. One of the important issues in polymer blends is the miscibility between the polymer components. Most polymer pairs are thermodynamically immiscible because of the positive free energy change (ΔG) during mixing. While this chapter primarily encompasses the crystallization and melting behavior of polymer blends, it is important to indicate that such properties are greatly influenced by the miscibility of the blend components. Poor adhesion at the interface of the polymer blends results in diminished mechanical properties such as lower impact resistance and elongation at break [17]. Miscibility is not the preferred state in all situations. In some cases, immiscibility is required to achieve two or multiple phases during blending.
From a thermodynamic point of view, the state of miscibility of any mixture is governed by the change in the free energy of mixing (ΔGM), which is defined as follows [18]:
where ΔGM, ΔHM, ΔSM, and T are, respectively, the Gibbs free energy change, enthalpy change, entropy change, and temperature of mixing.
Figure 2.1 shows the changes in the free energy of mixing with the composition of the overall mixture (ϕ2 is the volume fraction of component 2). Plot (a) indicates a positive free energy in entire volume fraction proportions. In case (b), the components are completely miscible in all compositions. Complete miscibility is not guaranteed by a negative free energy of mixing as shown in case (c), which shows a reversed curvature in the mid-composition range, and thus the mixture can develop an even lower free energy in this range by splitting into two phases with compositions given by the two minima. A partial miscibility is achievable in this condition.
Figure 2.1 Free energy of mixing as a function of composition for a completely immiscible (a), completely miscible (b), and partially miscible (c) binary mixtures. Reproduced from reference [20] by Paul and Barlow with permission, © 1980 Taylor & Francis.
A one-phase mixture is thermodynamically stable at fixed temperature and pressure if [18]
where ϕ is the blend composition (volume fraction). Because of the large molecular weights (small number of moles) of components in a polymer–polymer blend, the entropy term ΔSM in (Eq. 2.1) is very small and the TΔSM term is negligible. The miscibility of a polymer–polymer mixture, therefore, occurs if the heat of mixing of the components is negative (exothermic mixing). Miscibility for endothermic mixtures (ΔHM > 0) only occurs at very high temperatures. To obtain nonpolar polymer blends (ΔHM > 0) at normal temperatures, the components should have similar structures or precisely matched solubility values. To lead to exothermic heats of mixing, some specific interactions such as hydrogen bonding, acid–base, charge–transfer, ion–dipole, donor–acceptor interactions should exist among the components in a polymer blend [17,19,20]. These interdependencies are not very common for many polymers and the miscibility of polymers can only exist for a group of polymers either having precisely matched solubility values or certain specific reciprocities.
Control of the phase morphology of polymer blends is another important factor that can be used to improve the miscibility. Ratios of the phase's viscosities, interfacial properties, blend composition, and processing conditions have an influence on the size, shape, and the spatial distribution of the dispersed phases in a polymer blend [17]. An empirical relationship between the capillary number and relative melt viscosity of the dispersed and continuous phases has been presented by Wu et al. [21] in the viscoelastic systems. If the interfacial tension is lower and the viscosity ratio is closer to unity, the dispersed drops are smaller. The interfacial tension is largely controlled by the polarities of the two phases and can be varied over several orders of magnitude by using appropriate dispersants. The reduction of the interfacial tension between two components in a polymer blend can lead to the decreased particle size of the dispersed phase and therefore improve the miscibility of polymer blends. The addition of a compatibilizer and in situ compatibilization are two general methods used to decrease the interfacial tension across the interface of polymer blends. Compatibilizers are high molecular weight compounds, which act as polymeric surfactants in polymer blends. Due to the existence of the compatibilizer, the resistance against coalescence of the dispersed phase is greatly increased, and the stability of the polymer blend morphology remains the same in the different processing steps. Usually, a compatibilizer consists of one constitutive block miscible with one blend component and a second block miscible with the other blend component. These compatibilizers can be premade and added to the polymer blends or can be generated during the blending process (reactive compatibilization). A good compatibilizer reduces the interfacial energy, leads to a finer dispersion during mixing, and results in an improved interfacial adhesion. Nonbonding specific interactions such as hydrogen bonding, ion–dipole, dipole–dipole, donor–acceptor, and π–electron interactions are also useful for enhancing the compatibility of polymer blends [17].
In polyester blends, as those with poly(ethylene naphthalate) (PEN) and poly(ethylene terephthalate) (PET) where the two components are inherently immiscible, the transesterification reaction between the two polyesters can lead to the production of copolymers to improve their compatibility. Some parameters such as time and temperature are key factors that control the reactions in the melt phase. These parameters should be set at the best conditions to achieve the desired compatibilization and sufficiently high levels of transesterification reactions in the polyester blends [22,23].
2.3 Miscible Blends
Since crystallinity plays an important role in the properties of a polymer blend, understanding of the melting behavior, structure, and crystallization is required to control the properties of these materials. Optical microscopy (OM),small angle X-ray scattering (SAXS), differential scanning calorimetry (DSC), electron microscopy (SEM), and atomic force microscopy (AFM) are some techniques that have been used to study the melting behavior and crystallization of polymers.
Miscible polymer blends form a single-phase system and the properties of such materials are a combination of the properties of their pure components. If the free energy of crystallization is more negative than the free energy of the liquid–liquid mixture, crystallization occurs in a miscible blend. The crystallization and melting behavior of miscible crystallizable blends are complicated and affected by the structure and molecular weight of the components, the concentration of each component, and the intermolecular interactions between them [12,22].
The crystallization behaviors of miscible crystallizable blends containing one amorphous component are different from those of blends consisting of two crystallizable components.
The location of the amorphous component is very important. The amorphous component can be located in the interlamellar, interfibrillar, interspherulite, or a combination of two or more regions [24]. The type of segregation of the amorphous diluent plays a key role in the microstructure and crystallinity of a semicrystalline polymer blend [12]. The placement of the amorphous component in a miscible blend is affected by the size of the amorphous component (radius of gyration) and intermolecular interactions between the amorphous component and the amorphous portion of the crystallizable polymer [24].
2.3.1 Crystalline/Amorphous Polymer Blends
The glass transition (Tg) of a miscible polymer blend is generally between those of each individual component, and its crystallization occurs at a temperature between the Tg and the equilibrium melting point (Tm°). The effect of the amorphous component on the crystallization of a crystallizable, miscible, polymer blend depends on the glass transition of the amorphous component compared to that of the crystallizable one. If the Tg of the amorphous component is higher than that of the crystallizable one (as is true for the majority of polymer blends), the difference between Tm° and Tg of the blend decreases and, consequently, the tendency of the polymer blend to crystallize decreases. In the few cases where the Tg of the amorphous component is lower than that of the crystallizable component, the tendency for crystallization is increased [12,25]. Table 2.1 shows the effects of the amorphous component on the spherulite growth rate of the crystallizable component in some polymer blends.
Table 2.1 Effect of the amorphous component on the spherulite growth rate of the crystallizable component polymer blends.
Polymer blend
The effect of amorphous component on the spherulite growth rate
References
PEO/PMMA
decreases
[26]
PEO/PVAc
decreases
[27]
PVF
2
/PMMA
decreases
[28]
PVF
2
/PBSU
decreases
[29]
PCL/PVC
decreases
[30]
PCL/Phenoxy
decreases
[31]
iPS/PPO
decreases
[32]
PHB/PEC
decreases
[33]
P3HB/PVPh
decreases
[34]
P3HB/CE
decreases
[35]
P3HB/PVAc
decreases
[36]
iPS/PVME
increases
[37,38]
sPS/PVME
increases
[39]
PVF2/PBA
increases
[40]
PLLA/PBO
increases
[41]
2.3.2 Glass Transition and Melting Behavior
The glass transition temperature provides useful information on blend miscibility [42,43]. Although two Tgs are generally observed on the differential scanning calorimetry (DSC) scans of immiscible blends, only one Tg appears on that of miscible blends or copolymers. By investigating the blends of poly(2,6- dimethyl-1,4-phenylene oxide) (PPO) with poly-(styrene-co-4-chlorostyrene), Fried et al. [44] found that the width of the glass transition region (ΔTg) in which only one Tg is observed in the DSC scan can be a measure of the miscibility of the polymer blends. The ΔTg values were less than 10 °C for homopolymers or miscible blends. On the other hand, the glass transition regions were much wider (ΔTg = 30 °C) for immiscible blends. Kim et al. [45] used the same method to evaluate the PEN/poly(ether imide) (PEI) blends and found that although only one Tg was observed in the DSC scans, ΔTg values were more than 30 °C for the blends with PEI contents of 30–70 wt%.
2.3.2.1 Melting Point Depression
Because of the thermodynamically favorable interactions between the crystallizable and amorphous components in a miscible polymer blend, the melting point of the crystalline component is usually less than that of the pure polymer [12,43]. Measuring the melting point depression, therefore, provides information about interactions between polymers. According to the Flory–Huggins theory, the melting point depression under thermodynamic equilibrium can be expressed as follows [46]:
where and Tm are the equilibrium melting points of the crystallizable polymers in the bulk and the blends, respectively. Subscript 1 corresponds with the amorphous component and subscript 2 the crystallizable polymer; v represents the volume fraction; V1u and V2u are the molar volumes of the repeat units; x is the degree of polymerization; Δhu is the heat of fusion per mole of crystalline units; R is the gas constant; and χ12 is the Flory–Huggins interaction parameter. x1 and x2 are very large compared to unity in the case of polymers, and (Eq. 2.3) can be reduced to the Nishi–Wang equation [46,47]:
The authors assumed that the crystals were perfect and of finite size, and there was no recrystallization during melting. The melting point of a polymer is generally influenced not only by thermodynamic factors, but also by morphological parameters such as the crystal thickness. In miscible polymer blends, χ12 is usually negative and small [47–49].
The Flory–Huggins interaction parameter (neglecting the entropic contribution) can be defined by the following equation [48]:
where B is the interaction energy density. By substituting (Eq. 2.5) into (Eq. 2.4), we have
The interaction energy density (B) can be obtained from the slope of the plot of the left-hand side of (Eq. 2.6) as a function of v1/Tm. From (Eq. 2.4), the depression of melting point of a polymer–polymer pair is possible only when χ12 is negative, which is in agreement with Scott's condition for the miscibility of the two polymers [49]:
In some cases, (Eq. 2.4) does not fit the experimental data due to the use of observed melting temperatures instead of the thermodynamic equilibrium temperatures and the concentration dependence of the interaction parameter [12]. A modified equation has therefore been defined by other authors [50,51]:
where C is a constant related to the morphological contributions, and a is a constant from the relationship between interaction parameter and temperature.
Matkar and Kyu [52] have modified the Flory diluent theory that assumes the complete immiscibility of the solvent in the crystal. They have incorporated a crystal–solvent interaction parameter in addition to the amorphous–amorphous interaction parameter to develop a self-consistent theory for the determination of phase diagrams of a crystalline polymer solution. The changes in equilibrium melting temperatures can be explained in the context of the eutectic phase diagram approach. This has been well described by some researchers [52–55].
The equilibrium melting temperature of a polymer blend can be also determined by a plot of the experimental melting point versus the crystallization temperature (Hoffman–Weeks plot). An example of Hoffman–Weeks plots for poly(acetoxy styrene) (PAS)/poly(ethylene oxide) (PEO) blends is shown in Figure 2.2. Extrapolation from experimental data to the Tm = Tc line gives the value of . In some cases, the Hoffman–Weeks plots do not show a linearity because of the recrystallization and crystal defects in the sample [12,56].
Figure 2.2 Hoffman–Weeks plots for PAS/PEO blends. Reproduced from reference [56] by Kuo et al. with permission, © 2004 American Chemical Society.
2.3.3 Crystallization
Primary nucleation, crystal growth, and secondary crystallization are three stages, which are involved in the crystallization of polymers. In the primary nucleation process, nuclei are formed homogeneously or heterogeneously in the melt state. Heterogeneous nucleation, which is more common in polymers, occurs in the presence of impurities. After this nucleation process, crystalline lamellae develop and form three-dimensional superstructures. In general, the crystals of a polymer are spherulites, which consist of lamellar structures, growing radially from the center of the spherulites. In most cases, the crystallization will continue with the process of secondary crystallization [12,25]. A schematic drawing of a spherulite of a semicrystalline polymer is shown in Figure 2.3.
Figure 2.3 Model of spherulitic and lamellae structure. Adapted from reference [57].
The type of segregation or rejection of the amorphous component determines the morphology of polymer blends with an amorphous phase. Three different types of segregation have been considered. In the case of an interspherulitic segregation, the spherulites are imbedded in an amorphous matrix, while in the case of intraspherulitic segregation, the amorphous component is located within the spherulites, giving interlamellar and interfibrillar segregation [58,59].
When the diameter of gyration of an amorphous component is greater than the separation between crystalline lamellae (in highly crystalline polymer blends), the amorphous component can escape from the interlamellar zones. This leads to an interfibrillar or interspherulitic segregation. Interlamellar segregation is more favorable, if there is a good intermolecular interaction between the amorphous component and the amorphous portion of the crystalline polymer [24]. If the difference between Tg and Tc of the crystalline polymer is low, the amorphous polymer resides in interlamellar regions, while if this difference is high, some parts of the amorphous component resides in either interfibrillar or interspherulitic regions [12,24].
Type of segregation also depends on the relative values of the diffusion rate of the amorphous component and the crystallization rate of the crystallizable component. This is discussed more extensively in Section 2.3.4. The type of segregation of the amorphous component for some polymers is listed in Table 2.2 (Some data have been selected from the Polymer Blends Handbook [12] and some polymer blends are added from other articles).
Table 2.2 The type of segregation of the amorphous component in miscible polymer blends.
Polymer Blend
Amorphous Comp. (wt%)
Technique
Type of Segregation
References
PCL/CPE
0–30 CPE
SAXS, OM
interfibrillar
[60]
PCL/Phenoxy
0–50 Phenoxy
SAXS, OM
interlamellar/interfibrillar
[60]
0–10 Phenoxy
SAXS
interlamellar
[61]
PEG/PEMA
0–20 PEMA
OM
intraspherulitic
[62]
aPMMA/iPHB
0–50 aPMMA
DSC, SAXS
interlamellar
[63]
PVDF/PMMA
0 and 40 PMMA
SAXS, OM
interlamellar/interfibrillar
[64]
PVF
2
/PMMA
PMMA
SAXS, OM
interlamellar
[65]
PEO/PMMA
0–50 PMMA
SAXS
interlamellar
[66]
Cis
-PI/PVE
5–50 PVE
DSC
interlamellar
[67]
PBT/PAr
10–70 PAr
DSC, SAXS
interlamellar
[68]
PEEK/PEI
0–75 PEI
SAXS
interfibrillar/intraspherulitic
[69]
iPP/aPP
10–75 aPP
SAXS, SEM
interfibrillar/intraspherulitic
[70]
PEEK/PEI
25 and 75 PEI
TEM, OM
interfibrillar/intraspherulitic
[71]
PBT/PAr
0–80 PAr
SAXS
interfibrillar/intraspherulitic
[72]
In partially miscible polymer blends (blends with miscibility gaps in their composition ranges), two crystallization and demixing phenomena can affect the phase transitions and the morphology, depending on the amount of the crystallizable component. In crystallizable/amorphous polymer blends, if the phase separation region intersects the crystallization range, a competition between the kinetics of these two phenomena can be observed. Another phenomenon, observed in partially miscible blends, is interface crossing crystallization. In this case, there are two phases with different amounts of crystallizable component and different number density of heterogeneous nuclei in the phase separation region of partially miscible blends. The phase with a higher content of crystallizable component will crystallize first, and its crystallization will continue until the crystals reach the interface with the second phase. At this point, crystallization of the second phase will be induced, by a secondary nucleation process, on the crystals in contact with the melt [16].
2.3.4 Spherulite Growth Rate of the Crystallizable Component
An equation for the growth rate kinetics of a crystal in a homopolymer has been developed as follows by Hoffman and Lauritzen [73,74]:
where G is the growth rate of a crystal, G° is a constant that depends on the regime of crystallization, T0 is the temperature at which the required motions for the transport of molecules into the liquid–solid boundary occurs, Tc is the crystallization temperature, ΔE is the activation energy required for transferring crystals through the solid–liquid interface, and ΔF* is the energy required to form a nucleus. Two terms are involved in the growth rate kinetics of a crystal in a homopolymer: its ability to form surface nuclei and diffusion of crystalline molecules toward the crystal growth front. Based on these two opposing terms, a plot of the growth rate as a function of crystallization temperature in a homopolymer has a maximum (a bell-shape curve). The growth rate is nucleation controlled at low undercooling temperature and diffusion controlled at high undercooling temperature [12,26].
In most cases, adding an amorphous component to the polymer blend decreases the equilibrium melting temperature and, consequently, the spherulite growth rate of the crystallizable component [25]. This reduction in the equilibrium melting point occurs because of interactions between the two components [12].
Alfonso and Russel [26] have proposed a theoretical treatment for the growth rates in the miscible polymer blends containing an amorphous component. The interactions between the two components change the chemical potential of the liquid phase. These interactions affect the free energy required for forming a nucleus on the crystal surface and the mobility of both components. In addition, the amorphous component must diffuse from the growth front to the interlamellar area (segregation). There is a competition between the segregation of the amorphous component and the capability of the crystal to grow, with the crystal growth rate controlled by the slower process. One conclusion is that the crystallization growth rate depends on the molecular weight of both components. The crystallization of crystallizable component is also affected by the volume fraction of this component. Finally, the Tg of the amorphous phase has an influence on the segregation of the amorphous component at the solid–liquid interface. If the Tg of the amorphous component is higher than that of the crystalline component, the temperature range over which crystallization can occur is narrowed. Alternatively, the crystallization temperature range is widened if the Tg of the crystalline component is higher than that of the amorphous component.
Based on the above parameters and their effects, the crystallization growth rate of a polymer blend can be defined according to the following equation:
where ϕ2 is the volume fraction of the crystallizable component, Tc is the crystallization temperature, k1 is the rate of transport of the crystallizable molecule across the liquid–solid interface, and k2 is the rate of segregation of the amorphous component from the growth front. is the work required to form a nucleus on the crystal surface modified by the presence of the amorphous component. If the transport of the crystallizable component across the interface is faster than the rate of segregation (), then
Alternatively, if , then
To calculate the crystallization growth rate of a polymer blend, the values for ΔF*, k1, and k2 should be defined. For semicrystalline polymers in the presence of a low molecular weight diluent, Flory [75] has shown that
where V1u and V2u are the molar volume of the repeating unit of the amorphous component and crystallizable one; χ12 is the Flory–Huggins interaction parameter; Δhu is the enthalpy of fusion per mole of repeating unit of crystallizable component; b is the thickness of the critical nucleus; σσe is the product of the lateral and fold surface free energies; ϕ2 is the volume fraction of the crystallizable component; and parameter f is defined as follows by Hoffman and Weeks [76,77]:
In addition to the assumptions applied for the derivation of the Flory equation, χ12 and σσe are assumed to be independent of temperature and composition.
k1 is also defined as
where G0 is a constant dependent on the regime of the crystallization, ΔE is the energy of the transport, and is the value of T0 in the polymer blend.
where d is the maximum distance over which the amorphous component must diffuse away from the growth front, and D is the diffusion coefficient. D can be replaced by which is the mutual diffusion coefficient (a parameter indicating the cooperative diffusion of the amorphous and crystallizable components). can be determined experimentally or theoretically; however, the theoretical predictions by Kramer et al. [78] and Sillescu [79] are not accurate enough. Since the amorphous component diffuses in a direction normal to the growth direction, the maximum distance (d) would be half of the crystal thickness (L/2).
With the substitution of all the parameters into (Eq. 2.10), the following equation can be derived:
This can be rearranged to separate the kinetic term from the thermodynamic one:
where α is the kinetic term and contains thermodynamic parameters.
and
Since σσe are assumed to be independent of temperature and composition, a plot of α as a function of is a straight line, and σσe can be determined by the slope of this plot. Alfonso and Russell [26] have shown that deviation from a straight line can be observed when the molecular weight of the amorphous component is approximately equal to the critical molecular weight for entanglement coupling. This deviation corresponds to increasing σσe with decreasing temperature.
In (Eqs. 2.17) and (2.19), the ratio of is a modified version of the δ parameter defined by Keith and Padden [80,81]. (Equation 2.17) is derived based on the incorporation of the amorphous component in the interlamellar regions and the consistency of the growth rate on a microscopic level. It is important that the theoretical developments are more accurate in the case of low concentrations of diluent and in a temperature range near the melting point [26].
2.3.5 Overall Crystallization Kinetics
2.3.5.1 Isothermal Kinetics
The overall crystallization kinetics of polymer blends have been analyzed by the Avrami equation [82]:
where θa is the fraction of an uncrystallized material at time t, k is the Avrami constant, and n is the Avrami exponent describing the type of nucleation and the crystal growth geometry. The kinetic rate constant (k) is a function of the type of nucleation and spherulite growth rate. The kinetic parameters (n and k) can be obtained by plotting as a function of according to (Eq. 2.22):
This plot gives a straight line with the slope n and an intercept of ln (k). A typical example of an Avrami plot for the crystallization behavior of a 40/60 PEN/PET blend is illustrated in Figure 2.4. Summaries of n values and the rate constants for the crystallization behaviors of the various blends are given in Table 2.3.
Figure 2.4 ln[−ln(θa)] as a function of ln(t) for 40/60 PEN/PET blends. Reproduced from reference [83] by Shi and Jabarin with permission, © 2001 John Wiley & Sons, Inc.
Table 2.3 Rate constant, Avrami index as function of temperature and blend composition for some miscible blends.
Polymer blend
Composition
T
c
(°C)
k (sec
−n
)
n
Reference
PCL/PVC
100/0
29.6
19.95 × 10
−3
3.5
[84]
38.0
3.88 × 10
−3
2.2
42.1
1.01 × 10
−3
2.4
90/10
28.4
13.8 × 10
−3
2.1
38.4
2.16 × 10
−3
1.7
40.4
0.87 × 10
−3
1.9
PEG/PMMA
100/0
47
1.09
2.53
[85]
54
3.52 × 10
−5
2.63
80/20
42
3.57 × 10
−2
2.45
47
4.94 × 10
−4
2.46
50
5.40 × 10
−5
2.50
PET/PEN
100/0
120
1.12 × 10
−5
2.7
[43]
140
3.08 × 10
−4
2.9
80/20
150
2.92 × 10
−6
3.8
190
7.31 × 10
−3
3.2
60/40
150
1.17 × 10
−6
3.9
170
1.99 × 10
−4
3.8
20/80
130
4.47 × 10
−4
3.9
140
6.28 × 10
−3
3.9
0/100
170
3.67 × 10
−7
3.7
200
5.36 × 10
−4
3.4
The temperature dependence of the rate constant (k) can be obtained from the Arrhenius equation [83]:
where Ea is the apparent activation energy of rate constant (k), A is the frequency factor, Tc is the crystallization temperature, and R is the gas constant. Ea can be calculated from the slope of the plot of ln (k) as a function of 1/Tc.
In most cases, the n value is a noninteger because of the assumptions applied in the Avrami model, which assumed that the shape of the growing crystal, radial density, and the rate of radial growth are constant. Other assumptions are also applied in his model such as the absence of the secondary nucleation, no overlap between the growth fronts, lack of induction time, complete crystallinity of the sample, uniqueness of the nucleation mode, and random distribution of nuclei [12,82].
A modified Avrami equation was developed by Perez-Cardenas et al. [86]. In their model, the effects of the secondary crystallization are taken into account.
where α is the weight fraction of crystallinity, and αp and αs are the weight fraction of the primary and secondary crystallization, respectively. The crystallization process can be divided into three regions that include: initial primary crystallization (region 1), primary and secondary crystallization (region 2), and secondary crystallization (region 3). The total crystallization process is described by the following equations:
where ξ is the weight fraction of the polymer crystallized by primary and secondary crystallization when the primary crystallization is finished (end of region 2); k and n are the primary crystallization parameters; and are secondary crystallization parameters; t* is the time in which pure secondary crystallization occurs (region 3) [12,86]. Experimental isothermal crystallization data can be fitted to (Eqs. 2.25) and (2.26) by using different values for the parameters in order to find the most accurate ones.
2.3.5.2 Nonisothermal Kinetics
The theory of Avrami is used for isothermal conditions. This theory has been extended by Ozawa [87] for nonisothermal conditions, assuming that the nonisothermal condition is the result of infinite isothermal processes:
where θa is the fraction of uncrystallized material at temperature T, k(T) is the cooling crystallization function, β is the cooling rate, and n is the Avrami exponent. The kinetic parameters (n and k(T)) can be obtained by plotting as a function of at a given temperature for various cooling rates.
Another approach used to describe a nonisothermal crystallization process is to apply the theory of Avrami to the results of nonisothermal crystallization thermograms by plotting as a function of ((Eq. 2.22)) at each cooling rate. Since the temperature changes constantly, n and k cannot be related to the isothermal case (the spherulite growth rate and nucleation are temperature dependent) and do not have the same physical significance as in the isothermal crystallization [88–90].
A third approach was developed by Ziabicki [91,92], describing the nonisothermal processes:
Under quasistatic conditions
where T(s) is the thermal history, s is the time required for the nucleation of crystals, t1/2 is the half time of crystallization, and n is the Avrami exponent.
Liu et al. [93,94] have presented a different kinetic equation by combining Ozawa and Avrami equations:
which can be rewritten as
where β is the cooling rate, k is the Avrami constant, k(T) is the cooling crystallization function, F(T) = [k(T)/k]1/n, with n the Avrami exponent calculated with the Ozawa method, and b is the ratio between the Avrami and Ozawa exponents. F(T) refers to the value of the cooling rate selected at unit crystallization time, when the system has a defined degree of crystallinity. A plot of log β as a function of log t at a given degree of crystallinity gives a straight line with log F(T) as the intercept and −b as the slope.
The nonisothermal crystallization behavior of a miscible poly(hydroxy ether of bisphenol-A) (PHE)/poly-ɛ-caprolactone (PCL) polymer blend has been investigated by de Juana et al. [88] at various cooling rates. The presence of PH reduced the overall PCL crystallization rate at a given cooling rate. Their experimental data analysis was in agreement with treatments by both Ozawa and Avrami, and agreed quite well with the theoretical results obtained using the Ziabicki method. Ozawa's treatment (the plot of ) as a function of for PHE/PCL 20/80 blend at several temperatures is shown in Figure 2.5. For all temperatures, straight lines confirmed the validity of applying Ozawa's method to PHE/PCL blends. The Avrami exponents were approximately 3, confirming a three dimensional growth process and heterogeneous nucleation.
Figure 2.5 Ozawa plot of log[−ln(1 − X(t))] as a function of log (β) for 20/80 PHE/PCL blend at various temperatures. Reproduced from reference [88] by de Juana et al. with permission, © 1996 Elsevier.
Nonisothermal crystallization behaviors of poly(trimethylene terephthalate) (PTT)/PEN blends were studied by Run et al. [95]. Plots of log β as a function of log t for 20/80 and 40/60 PTT/PEN blends are shown in Figure 2.6. Based on the log F(T) and b calculated by the authors, log F(T) values were increased with the relative crystallinity from 3.81 to 4.04 (for 20/80 PTT/PEN blend) and 3.08 to 3.45 (for 40/60 PTT/PEN blend). This indicated that a lower crystallization rate was needed to reach the given crystallinity within unit time. A small increase in parameter b was observed with increasing the crystallinity, ranging from 1.31 to 1.35 (for 20/80 PTT/PEN blend) and 1.01 to 1.09 (for 40/60 PTT/PEN blend), respectively. The results showed that with increasing PTT content, higher crystallization rate was achieved. The authors also confirmed their results with treatments by Avrami and Ozawa. They concluded that the equation proposed by Liu et al. [93] successfully describes the nonisothermal crystallization process of the binary blends for the whole crystallization process.
Figure 2.6 log β as a function of log t from the equation defined by Liu et al. [93] for (a) 20/80 and (b) 40/60 PTT/PEN blends. Reproduced from reference [95] by Run et al. with permission, © 2006 Elsevier.
2.3.6 Crystalline/Crystalline Polymer Blends
Crystalline/crystalline polymer blends have received less attention compared to the amorphous/amorphous or amorphous/crystalline systems. Since both components present in the mixture are crystallizable, the melting and crystallization behaviors of such systems are more complicated. The crystallization and morphology of these systems are related to differences in the melting points of each of component [96]. In such polymer blends, the morphology of the final mixture is also affected by thermodynamic and kinetic factors during crystallization. Eutectic or concurrent crystallization (simultaneous crystallization of the components along with fine dispersion of the crystals) is very rare in polymers. However, it may occur in low molecular weight materials. In most polymer cases, the components crystallize successively (separate crystallization), and the crystallization behavior of the first crystallized component is influenced by the other component. Since the component that crystallizes later is still in the melt state, it acts as a diluent for its crystallized partner [97]. In this case, spherulites of the component that crystallizes later will contact with the spherulites of the other component and start to grow within its spherulites. Generally, when the difference between the Tms of two components is large, the two components will crystallize separately. However, when the difference between the Tms of two components is small, they can be crystallized simultaneously [96]. Table 2.4 shows some examples of polymer blends with different types of crystallization.
Table 2.4 Types of crystallization in some crystalline/crystalline miscible blends.
Polymer Blend
Comp. (wt%)
Technique
Type of Crystallization
References
PHB/PEO
0–80 PEO
OM, DSC
Separate
[98]
PHB/POM
0–100 POM
DSC, SAXS
Separate
[99]
PEO/PES
0–100 PEO
OM, SAXS
Separate
[100]
PBT/PAr
0–100 PAr
DSC, WAXD
Separate
[101]
PVF
2
/PBA
0–100 PBA
OM, WAXD
Separate
[102]
PBSU/PEO
0–100 PEO
OM, DSC
Separate
[103]
PES/PVPh
0–100 PVPh
OM, DSC
Separate
[104]
PHB/PLLA
0–100 PLLA
OM
Simultaneous
[105]
PEC/PLLA
0–100 PLLA
DSC, SAXS
Simultaneous
[106]
PEC/PLLA
20 PLLA
AFM
Simultaneous
[107]
PES/PEO
0–100 PEO
OM
Simultaneous
[108]
PEO/PBAS
0–100 PBAS
DSC, OM
Simultaneous
[109]
PED/EVAc
0–100 EVAc
DSC, WAXD
Simultaneous
[110]
LLDPE/VLDPE
0–100 VLDPE
DSC, SAXS
Simultaneous
[111]
PEEK/PEK
0–100 PEK
DSC
Simultaneous
[112]
Qiu et al. [96] have studied a miscible poly(vinylidene fluoride) (PVDF)/poly(butylene succinate-co-butylene adipate) (PBSA) polymer blend. Since the difference in the melting points of the two components was large (165 °C and 95 °C for PVDF and PBSA, respectively), they crystallized separately. PVDF crystallized first, so PBSA which was still in the melt state acted as a diluent during the crystallization of PVDF. The crystallization of the component that crystallizes first (high temperature crystallizing component [HTC]) is a transition from the amorphous/amorphous to the amorphous/crystalline state, while the crystallization of the second component (low temperature crystallizing component [LTC]) is a transition from the amorphous/crystalline to crystalline/crystalline state.
Figure 2.7 illustrates the effect of crystallization temperature and composition on the spherulitic growth rates of PVDF (the component that is first crystallized).
Figure 2.7 Spherulitic growth rate of PVDF as a function of temperature for the PVDF/PBSA blends. Reproduced from reference [96] by Oiu et al. with permission, © 2007 American Chemical Society.
The spherulitic growth rates decrease with increasing crystallization temperature regardless of blend composition. The spherulitic growth rates also decrease with the increasing LTC content at a constant temperature. The change in this growth rate is less significant at high temperatures. This indicates that supercooling plays a key role in the morphology and growth rates of such systems. It can be concluded that factors causing reduction in the spherulitic growth rates of HTC include the depression of its Tm after blending, and the diluent role of LTC during the crystallization of HTC [96].
The spherulitic morphologies of LTC are shown in Figure 2.8. Figure 2.8a shows the spherulites of the neat PBSA crystallized at 70 °C for 20 min. The spherulites are compact, and the lamellae bundles are thick. In the case of the 20/80 blend (Figure 2.8b), PVDF is crystallized first at 120 °C for 5 min. With decreasing the crystallization temperature to 70 °C, the PBSA spherulites start to grow until impinging on other PBSA spherulites. A comparison of these two figures indicates that the nucleation of PBSA has increased in the presence of the PVDF, because the pre-existing PVDF crystals increase the nucleation of PBSA. The overall crystallization rate of the blends is larger than that of neat PBSA in this system; however, with increasing PVDF content, the overall crystallization rate decreases.
Figure 2.8 Spherulitic structures of (a) neat PBSA at 70 °C for 20 min, and (b) 20/80 PVDF/PBSA blend at 70 °C for 6 min after crystallization of PVDF at 120 °C for 5 min. Reproduced from reference [96] by Oiu et al. with permission, © 2007 American Chemical Society.
We can conclude that, in the crystalline/crystalline polymer blends, the overall crystallization of LTC is influenced by the presence of the crystals of HTC, the dilution effect, and physical hindrances in the matrix of HTC. The first parameter accelerates the overall crystallization rate, while the last two parameters slow down the rate.
Isotactic polypropylene (PP)/maleic anhydride grafted polypropylene (mPP) blends have been studied by Cho et al. [113]. Based on their results, the cooling rate has a significant effect on the DSC fusion endotherms of the blends. For a blend sample prepared at a low cooling rate such as 1 °C/min, a separate crystallization (phase separation) occurs and the DSC thermographs consist of two melting peaks (Figure 2.9). For the samples prepared at high cooling rates (40 °C/min), crystallization occurs simultaneously and a single melting peak is observed in DSC thermographs. Isothermal DSC measurements show that at a low degree of supercooling (crystallization at higher temperatures), the components crystallize separately, while at a high degree of supercooling, co-crystallization occurs and a single endotherm is observed from DSC data. This is in agreement with results from nonisothermal experiments.
Figure 2.9 DSC endotherms for 50/50 blend by nonisothermal crystallization at different cooling rates (the heating rate is 10 °C/min). Reproduced from reference [113] by Cho et al. with permission, © 1999 Elsevier.
