Semicrystalline Polymer

Semicrystalline polymers are limited in their stiffness by the unordered material which is distributed between the folded chain crystallites, and after conventional drawing, which may allow an extension ratio of up to 10×, have extensional stiffnesses more than an order of magnitude below that for a single chain.

From: Encyclopedia of Materials: Science and Technology , 2001

Basic Concepts and Polymer Properties

D.A. Ivanov , in Polymer Science: A Comprehensive Reference, 2012

1.09.2.1 Lamellar Habit and Thermal Properties

Linear PE is certainly the most studied semicrystalline polymer. Contrary to the simplified view of the lamellar crystal given in Figure 1 , the PE single crystals in solution have a 3D form of hollow pyramids with four or six sectors. 17–20 Due to the chain tilt in the PE single crystals sedimented on solid substrate, the fold surfaces of single crystals in (110) sectors formed at high and low degree of supercooling are {(314)(110)} and {(312)(110)}, respectively. The PE single crystals grown from very dilute p-xylene solutions are delimited by four (110) faces with two truncated (200) faces appearing when the concentration or the crystallization temperature is raised. The relative importance of the (200) sector increases with crystallization temperature and polymer concentration. 21 Two modes of crystal collapse have been suggested, 22 such as plastic deformation without reorientation of the crystalline stems with respect to the lamellar basal plane and flattening without any plastic deformation leading to chain tilting.

An example 23 of a solution-grown PE single crystal having the shape of a truncated lozenge is given in Figure 7 . The tilt angle of the chains in the (110) and (200) sectors was measured in an SAED experiment by tilting the single crystal around its crystallographic b-axis to obtain a diffraction pattern corresponding to the (hk0) plane (cf. Figure 7 ). It can seen that to recover the (hk0) reflections, the rotating angle should be 22° in (110) sectors and 30° in (200) sectors, which clearly shows a difference in their microstructure. An AFM height image obtained in tapping mode on a similar single crystal is given in Figure 8 . The color code is chosen in such a way that a small difference between the thickness of the four (110) and two (200) sectors is rendered visible. As can be seen from the AFM image cross-sections traced through each of the sectors (cf. Figure 8 , bottom), this height difference is less than ∼4% of the lamellar thickness, which is in agreement with the different stem inclinations in the two sectors. The crystalline stem lengths measured along the backbone are, within experimental error, identical in both sectors. Interestingly, the central pleat resulting from the collapse of the pyramidal single crystal during sedimentation is always oriented along the crystallographic b-axis, which proves that the collapse of the lamellar hollow pyramids is not a random process.

Figure 7. Top: electron micrograph of a PE (〈M w〉   =   32   100, polydispersity   =   1.1) single crystal formed from a dilute solution of n-octane at 95   °C. Bottom: electron diffraction patterns obtained at a tilt angle of 22° in a (110) sector and at a tilt angle of 30° in a (200) sector of the single crystal shown.

 With permission from Hocquet, S.; Dosière, M.; Thierry, A.; et al. Macromolecules 2003, 36, 8376. 23

Figure 8. Top: Tapping mode atomic force microscopy (AFM) image of a single crystal of linear PE formed from a dilute solution in n-octane at 95   °C. The color code is chosen to reveal the small elevation difference between (110) and (200) crystal sectors. Bottom: AFM image cross-sections traced as shown in the top panel.

 With permission from Hocquet, S.; Dosière, M.; Thierry, A.; et al. Macromolecules 2003, 36, 8376. 23

The correlation between the morphology of single crystals of PE and their thermal properties was in focus of many studies. Melting temperatures of 124.5   °C for the (200) sectors and of 126.5   °C for the (110) sectors of PE single crystals were obtained by accurate differential scanning calorimetry (DSC) experiments coupled to electron microscopy. 24 The availability of the variable-temperature AFM 25–30 made the in situ studies of the thermal behavior of PE crystals feasible. Figure 9 shows AFM images of a truncated PE single crystal successively measured at 121.9, 123.9, 124.9, and 125.9   °C. No trace of melting is observed in the AFM image recorded at 121.9   °C. At 123.9   °C, several cracks running perpendicular to the growth faces are observed in the two (200) sectors. A smaller number of similar cracks can also be detected on one of the (110) faces. The (200) sectors have completely melted at 124.9   °C and recrystallized into thicker patches, with an average thickness ranging between 30 and 40   nm. The thickness of the recrystallized material is comparable to the height of the central pleat running along the crystallographic b-axis. Moreover, holes were found to develop in the sectors probably due to a competition between recrystallization and dewetting, which follow a similar scheme in both types of sectors. The (110) sectors are completely molten at 125.9   °C.

Figure 9. Tapping mode AFM images of a liquid phase epitaxy (LPE) single crystal formed from a dilute solution of n-octane at 95   °C taken at 121.9   °C (top left), 123.9   °C (top right), 124.9   °C (bottom left), and 125.9   °C (bottom right).

 With permission from Hocquet, S.; Dosière, M.; Thierry, A.; et al. Macromolecules 2003, 36, 8376. 23

The images show that the (200) sectors indeed melt first. Moreover, the temperature difference with the melting point of the (110) sectors is comparable to what was expected from the referenced calorimetric studies. The free enthalpies σ e of the fold surface in the (110) and (200) sectors have been estimated from the Gibbs–Thomson relationship to be 58 and 56   erg   cm−2, respectively. The closeness of the surface energy values can be due to the fact that truncated PE single crystals exhibit jagged (200) faces. 23 Supporting evidence for some jagging of the (200) growth faces stems from the fact that orientation of the decorating alkane rods on PE single crystals grown from solution is less regular in (200) than in (110) sectors. 31

The interpretation of the in situ AFM images is consistent with the observations of the temperature-dependent small-angle X-ray scattering (SAXS) patterns of PE single-crystal mats during heating. 32 The mechanisms involved in the evolution of the lamellar thickness of semicrystalline polymers have been extensively investigated in the past and were shown to depend on their initial morphology. The stacked lamellar morphology of linear ultrahigh molecular weight PE crystallized from solution exhibits on annealing a lamellar doubling, which is not accompanied by melting. 33 The lamellar thickening during annealing of solution-grown PE crystal mats could also occur without melting, depending on annealing temperature and heating rate. 34,35 The conditions of the described AFM experiment excluded any doubling of the PE lamellae and can, therefore, differ from the experiments on single-crystal mats. However, they provide a direct-space view of the reorganization and melting processes in the crystals and are therefore helpful for a general understanding of the thermal behavior of the polymer crystals.

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Polymer Characterization

S.S. Sheiko , S.N. Magonov , in Polymer Science: A Comprehensive Reference, 2012

2.23.3.2.4 Polymer crystals

Semicrystalline polymers constitute the largest fraction of industrial plastics used for fabrication of fibers, films, blends, and composites. As such, a lot of studies in the last 50 years focused on the fundamental understanding of structural organization, crystallization, melting, and processing of semicrystalline polymers. Their behavior is essentially different from that of small molecules because long chains are much larger than the crystal thickness and different portions of the chain can participate in different crystals. The fundamentals of the structure and the crystallization of polymers were established in the 1950s and 1970s, respectively; 244 however, this area of science still attracts much attention due to not only technological significance but also many unresolved questions. In recent years, research activities in this field have been intensified due to the development of new techniques such as high-speed calorimetry, synchrotron radiation microfocus beams, and AFM. The latter two techniques enable resolution of crystalline superstructures down to micrometer and nanometer range, respectively.

The morphological hierarchy of semicrystalline films, molds, and fibers is very complex, resulting in many unresolved issues including the thickness of polymer crystals, growth sectors of folded-chain lamellae, lamellar branching and bending, spherulite organization, and the morphology of amorphous phase. Understanding of the morphological hierarchy and connectivity of structural components on all levels is a key to fabricating materials with superior mechanical properties, for example, silk fibers. 245–248

It is known that kinetic trapping during crystallization of semicrystalline polymers leads to crystalline lamellae of finite thicknesses (typically 5–50   nm) with a significant portion of chain segments folding back into the crystals. 249,250 Other emerging segments accumulate in the amorphous layers as loose loops, dangling segments, or tie molecules. 251,252 From both kinetics and equilibrium points of view, there are arguments that suggest variations of crystal thickness for crystallizing homopolymers as well as copolymers. 253 However, systematic SAXS, 254,255 TEM, 256 and AFM 257 studies show that isothermal crystallization leads to crystals with uniform thickness. The only exception are polymers, for example, polyethylene (PE), with active sliding motion of chains within crystals, which leads to crystal thickening and eventually to a thickness distribution. As shown in Figure 30 , the thickness distribution of poly(ethylene terephthalate) lamellae is narrow and the mean thickness is constant throughout the whole crystallization process. 257 The real-time AFM studies clearly show equal thicknesses for both dominant and subsidiary crystallites, which refute the assumption that the first grown lamellae should be thicker than the secondary crystallites.

Figure 30. Crystallization of poly(ethylene terephthalate) was monitored by tapping-mode AFM. (a) Phase images were recorded in situ at 233   °C. Image analysis gives (b) time dependence of the mean crystal thickness and (c) thickness distribution in the final state.

 Reprinted with permission from Ivanov, D.; Amalou Z.; Magonov S. N. Macromolecules 2001, 34, 8944. 257 Copyright 2001 American Chemical Society.

Prior to AFM, TEM in combination with X-ray diffraction has been applied for in-depth examination of single PE crystals. The combination of techniques made it possible to observe the sectorization and determine the polymer chain orientation inside individual crystal sectors. However, many questions related to the organization of single crystals remain open. One of such questions concerns the structure of the lamellar surface, which is presumably formed by chain folding according to the adjacent reentry model. The other question is whether the chain packing and microstructure of the lamellar bulk. The sensitivity of AFM to height measurements made it possible to monitor chain unfolding in situ with exceptional precision. Real-time imaging of single PE crystals at elevated temperatures revealed lamellar thickening caused by unfolding of individual chains from a kinetically formed folded state to an energetically favorable extended-chain conformation. 258 Figures 31(a) and 31 (b) demonstrate that holes appear simultaneously with thickening of the adjacent locations, which is consistent with the earlier TEM data. 259 Thickening proceeds gradually after a stepwise change at 115   °C, as reflected in the height histograms in Figure 31 (c). Studies of the thickening mechanisms in various polymers can be useful in understanding the role played in these processes by the crystal/amorphous interphase 260 and the polymer nature of the reorganizing species.

Figure 31. Height images of dry single crystals of PE measured at 110   °C temperature and after 1.5   h of annealing at 115   °C. Height histograms corresponding to AFM show the evolution of lamellar thickness after annealing at different temperatures.

 Reprinted with permission from Magonov, S. N.; Yerina, N. A.; Ungar, G.; et al. Macromolecules 2003, 36, 5637. 261 Copyright 2001 American Chemical Society.

An effective approach in elucidation of fine structural features of polymer crystals and the mechanisms of polymer crystallization is to use model molecules. For example, ultralong alkanes (C n H2n+2, n  >   150) are considered as an appropriate model for PE. Recent AFM studies have shown that the structure of single crystals of C390H782 and PE is similar, though the details of their thermal behavior are quite different. 261 Upon annealing, alkane crystals undergo a complete series of transformations corresponding to stepwise unfolding from the folded-in-five conformation toward the fully extended-chain crystal, while the chain unfolding in PE crystals is a continuous and slower process. Another series of model compounds was proposed to selectively control the chain folding of polymers. 262 Structural instructions were encoded a linear backbone that includes alternating crystallizable, long alkyl sequences of monodisperse sizes separated by short spacers containing side chains and acting as stops and fold-controlling units ( Figure 32 (a)). This code translates during a crystallization process to generate a semicrystalline morphology with structure-controlled crystal thickness of ∼5   nm that remains constant over a wide temperature range ( Figure 32 (b)). This approach allows controlling the lamellar thickness by steric interactions only, in contrast to previous attempts aimed at engineering polymer crystallization through hydrogen bonding. 263,264

Figure 32. (a) Schematic illustration of the trajectory of a chain within P44/5-Prop crystals before and after crystallization. Superimposed on the right is a pictorial representation of the electron density profile corresponding to the crystalline lamellae, wherein L a is the thickness of the amorphous layer. (b) AFM phase image of a sample measured in hard tapping mode at 85   °C. The thin white stripes correspond to the crystalline lamellae grown perpendicularly to the image plane. (c) L p and L c as a function of crystallization temperature (measurements at room temperature). L c values were obtained from the fit of the SAXS.

 Reproduced with permission from Le Fevere De Ten Hove, C.; Penelle, J.; Ivanov, D. A.; Jonas, A. M. Nat. Mater. 2004, 3, 33. 262 Copyright Nature Publishing Group.

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Polymer Characterization

J. Eckelt , ... B.A. Wolf , in Polymer Science: A Comprehensive Reference, 2012

2.04.3.1 Temperature Rising Elution Fractionation

Semicrystalline polymers can be fractionated by means of TREF, which is based on chain crystallizabilities. It consists of two steps: crystallization and elution. The method can be operated in an analytical (A-TREF) or in a preparative mode (P-TREF). The principle of the experimental setup is the same and shown in Figure 17 . 38,39 The main differences of the two modes are the elution step and sample size as well as the used column dimensions.

Figure 17. Scheme of a typical TREF apparatus. Depending on the operation mode, the eluent leaving the column is transferred either into a detector (A-TREF) or into a fraction collector (P-TREF).

The first step consists of the complete dissolution of the polymer in a good solvent at high temperature. Afterward the homogeneous solution is introduced into a column, which contains an inert support, for example, glass beads, sea sand, steel shot, chromosorb P, or silanated silica gel. 40 In the next step, the temperature in the column is decreased at a slow cooling rate (typically 1–2   °C   h−1). In this manner, the polymer chains with the highest crystallizabilities precipitate first onto the support, followed continuously by the polymer chains with lower crystallizabilities (cf. Figure 18 ). 41 Alternatively, this crystallization step can be carried out in an automatic temperature-programmable stirred vessel, which contains the inert support. After the complete precipitation, the polymer-coated support is filled into the TREF column. The cooling rate for both operation modes (A-TREF and P-TREF) has to be slow enough to guarantee the fractionation of the polymer, as the crystallization step is the most important one in TREF.

Figure 18. Scheme of the support coated with polymer. The polymer chains with the highest crystallizabilities are located closer to the support than the polymer chains with lower crystallizabilities.

For the elution step, pure solvent is pumped through the column while the temperature increases continuously (A-TREF) or stepwise (P-TREF). As soon as the dissolution temperature of the polymer is reached, the layers dissolve in the reverse order in which they were precipitated.

In A-TREF, the column temperature in the elution step is increased in a slow, constant rate, while the polymer concentration in the eluent is monitored with an on-line mass-sensitive detector to obtain the TREF profile: the distribution of chain crystallizabilities in terms of the weight fraction of polymer eluted at each temperature. The chemical composition distribution (CCD) and the tacticity can be obtained from the TREF profile using a calibration curve.

In the case of P-TREF, normally larger columns and sample sizes are used. The temperature of the elution step is increased stepwise and all polymer eluting within a given temperature interval is recovered. This operation mode is more commonly used for preparing series of fractions that have narrower CCDs than parent samples. For more detailed information on the chain microstructure, TREF can also be combined with other fractionation and analysis techniques, such as GPC.

As mentioned before, the most common use of TREF is for the determination of the CCD of polyethylene–polypropylene copolymers. The incorporation of a second monomer into the backbone of a homopolymer has large influences on the final properties of the material, for example, crystallinity, melting and glass transition temperatures, impact resistance, and transparency. Caballero et al. 42 used TREF among other methods to investigate the influence of the chemical composition on the properties of ethylene–propylene copolymers. As expected, they found out that the behavior of the copolymer is similar to that of the corresponding homopolymer if the comonomer content is low. As the comonomer content increases, the copolymer becomes more amorphous (lower crystallization temperature and softer X-ray diffraction (XRD) signals) and easily deformable, reaching a behavior similar to an elastomeric material. As depicted in Figure 19 , TREF analysis shows that copolymers containing less than 10% and more than 80% of ethylene are semicrystalline, with elution temperatures typical for these kinds of polymers. TREF is limited to copolymers with very different comonomer contents, because the polymers are amorphous in the middle range of comonomer composition.

Figure 19. Influence of the molar ethylene content in the copolymer on TREF profiles.

 Reproduced with permission from Caballero, M. J.; Suarez, I.; Coto, B.; et al. Macromol. Symp. 2007, 257, 122. 42 Copyright Wiley-VCH Verlag GmbH & Co. KGaA.

Conventional IR detectors are commonly used as concentration detectors in TREF and the measurement is based on a single wavelength. This works well for simple copolymer systems such as ethylene–α-olefin copolymers. However, the conventional IR detectors fail for complex copolymer systems and an alternative detection method is required. Zhang 43 coupled a Fourier transform infrared (FT-IR) spectrometer with a TREF instrument to provide a tool for characterizing complex olefin copolymers. The dual-wavelength technique worked well for two-component copolymer systems, whereas a multivariate calibration method is required for analyzing the IR spectra of more complex multicomponent copolymer systems. Zhang investigated three copolymer systems: ethylene–α-octene, PS-grafted ethylene–vinyl acetate, and ethylene–methyl acrylate copolymers. In addition to polymer concentration, the polymer composition (i.e., comonomer content) can be measured by on-line FT-IR detection. This eliminates very time- and labor-consuming TREF fraction collection as well as the postfractionation compositional analyses by NMR and brings a benefit to the TREF analyses, especially for the complex olefin copolymers such as ethylene–α-olefin block copolymers and ethylene–methyl acrylate copolymers. Also, TREF/FT-IR analysis of PS-grafted ethylene–vinyl acetate provides an experimental means of measuring the grafting efficiency that is an important parameter affecting the polymer morphology and thus material properties.

A further progress for TREF measurements is the combination with other fractionation methods such as GPC. In this manner, polymers can be analyzed simultaneously for their distributions in chemical composition and molecular mass. 44,45 Figure 20 shows a three-dimensional plot of a Ziegler–Natta linear low-density polyethylene analyzed by means of a combination of GPC and TREF.

Figure 20. Coupled GPC-TREF diagram of a Ziegler–Natta linear low-density polyethylene.

 Reproduced with permission from Yau, W. W. Macromol. Symp. 2007, 257, 29. 44 Copyright Wiley-VCH Verlag GmbH & Co. KGaA.

Even though TREF is mainly used to determine the CCD, it can be used for other purposes also. Nakatani et al. 46 used TREF experiments to investigate the influence of extraction of an internal donor on the variation of isospecific active sites of a MgCl2-supported Ziegler catalyst, and to estimate the relationship between polymer microtacticity and degradation rate of isotactic polypropylene (iPP). The former example revealed the conversion from high to low isospecific sites by the extraction of internal donors, whereas the latter showed a negative correlation between the level of isotacticity and the degradation rate.

Recently, it has been demonstrated that high-temperature HPLC (HT-HPLC) can be used as an alternative method for the determination of the CCD of semicrystalline copolymers of ethylene and polar comonomers. 47

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MICROSTRUCTURE

H. GLEITER , in Physical Metallurgy (Fourth Edition), 1996

5.3. Semicrystalline polymers

Semicrystalline polymers constitute a separate class of nanostructured materials. The remarkable feature of this class of polymers is that the nanostructured morphology is always formed if the polymers are crystallized from the melt or from solution unless crystallization occurs at high pressure or if high pressure annealing is applied subsequent to crystallization. If a polymer is crystallized from a dilute solution, isolated single polymer crystals or multilayer structures consisting of stacks of polymer single crystals result ( fig. 27). Inside the crystals, the atoms forming the polymer chains arrange in a periodic three-dimensional fashion. The interfaces between neighboring crystals consist of both macromolecules folding back into the same crystal and tie molecules that meander between neighboring crystals. The typical thicknesses of the crystal lamellae are in the order of 10 to 20 nm. These relatively small crystals thicknesses have been interpreted in terms of one of the following models. The first model hypothesizes the formation of the thin crystallites to result from nucleation kinetics. If the height of the energy barrier for the formation of a critical nucleus of a chain-folded polymer crystal formed in a supersaturated solution is computed by means of homogeneous nucleation theory, it is found that the energy barrier of a critical nucleus consisting of extended chain molecules is larger than the barrier height for a nucleus of folded chains. The physical reason for this energy difference is as follows. Extended chain crystallization results in a needle-shaped critical nucleus, the length of which is equal to the length of the molecular chains. Hence the system is left with only one degree of freedom to reduce the energy barrier for the critical nucleus. This reduction occurs by adjusting the diameter of the needle. However, if chain folding occurs, the energy barrier associated with the critical nucleus can be minimized by adjusting the size of the nucleus in all three dimension. Detailed computations reveal that the energy barrier for chain folded nuclei is in general significantly lower than for extended chain crystallization. The second group of models for chain-folding is based on the excess entropy associated with the folds relative to an extended-chain crystals. If the Gibbs free energies of an extended chain crystal and of a chain-folded crystal are compared, the chain folds are found to increase the internal energy of the system. However, the chain folds also contribute to the entropy of the system. Hence, at finite temperatures, a structure of lowest Gibbs free energy is obtained, if a certain concentration of chain folds is present in the crystal. In other words, chain-folded crystals have a lower Gibbs free energy at finite temperatures than extended chain crystals (cf. also ch. 32, § 2.2–2.6).

Fig. 27. Schematic representation of the conformation of chain-folded polymer molecules in a semicrystalline polymer. One molecule belonging to adjacent crystals is indicated as a heavy line.

Polymers crystallizing from the molten state form more complex morphologies. However, the basic building blocks of these morphologies remain thin lamellar crystals. Figure 28 shows spherulitic crystallization of thin molten polymer film. The spherulites consist of twisted lamellae which exhibit radiating growth. If the molten thin film is strained during solidification, different morphologies may result, depending on the strain rate. However, all of these morphologies have in common that the macromolecules are more or less aligned in the straining direction. High temperatures and small strain rates favour a stacked lamellar morphology (fig. 29a), high temperatures combined with high strain rates result in needle-like arrangements (fig. 29b). Low temperatures and high strain rates lead to oriented micellar structures (fig. 29c). The transition between these morphologies is continuous and mixtures of them may also be obtained under suitable conditions (fig. 29d). The way to an additional variety of nanostructured morphologies was opened when multicomponent polymer systems, so-called polymer blends, were prepared. For thermodynamic reasons, polymer blends usually do not form homogeneous mixtures but separate on length scales ranging from a few nanometers to many microns depending on the thermomechanical conditions of crystallization and the molecular structure of the costituents involved. So far the following types of nanostructured morphologies of polymer blends have been reported for blends made up by one crystallizable and one amorphous (non-crystallizable) component: Type I morphology: The spherulites of the crystallizable component grow in a matrix mainly consisting of the noncrystallizable polymer. Type II morphology: The non-crystallizable component may be incorporated into the interlamellar regions of the spherulites of the crystallizable polymer. The spherulites are spacefilling. Type III morphology: The non-crystallizable component may be included within the spherulites of the crystallizable polymer forming domains having dimensions larger than the interlamellar spacing. For blends of two crystallizable components, the four most frequently reported morphologies are: Type I morphology: Crystals of the two components are dispersed in an amorphous matrix. Type II morphology: One component crystallizes in a spherulitic morphology while the other crystallizes in a simpler mode e.g., in the form of stacked crystals. Type III morphology: Both components exhibit a separate spherulitic structure. Type IV morphology: The two components crystallize simultaneously resulting in so-called mixed spherulites, which contain lamellae of both polymers.

Fig. 28. Bright field transmission electron micrograph (defocus contrast) of a two-dimensional spherulite in isotactic polystyrene. The spatial arrangement of the lamellae formed by the folded macromolecules is indicated on the left side (Petermann [1991]).

Fig. 29. (a) Stacked lamellar morphology in polyethylene (Tem bright field). (b) Needle-like morphology in polybutene-1 (Tem bright field). (c) Oriented micellar morphology in polyethylene terephthalate (Tem dark field micrograph). (d) Shish-kebab morphology in isotactic polystyrene (Tem dark field micrograph) (Petermann [1991]).

Morphologies of lower complexity than spherulites, such as sheaves or hydrides may also be encountered. In these cases, the amorphous phase, may be arranged homogeneously or heterogeneously depending on the compatibility of the two components. The morphology of blends with one crystallizable component has been studied for a variety of macromolecular substances e.g., poly(ɛ-caprolactrone)/poly(vinylchloride), poly(2,6dimethyl-phenylene oxide)/isotactic polystyrene, atactic polystyrene/isotactic polystyrene blends.

Block copolymers constitute a third class of nanostructured polymers. All macromolecules of a block copolymers consist of two or more, chemically different sections which may be periodically or randomly arranged along the central backbone of the macromolecules and/or in the form of side branches. An example of a block copolymer are atactic polytyrene blocks alternating with blocks of polybutadiene or polyisoprene. The blocks are usually non-compatible and aggregate in separate phases on a nanometer scale. As an example for the various nanostructured morphologies possible in such systems, fig. 30 displays the morphologies formed in the system polystyrene/polybutadiene as a function of the relative polystyrene fraction. The large variety of nanostructured morphologies that may be obtained in polymers depending on the crystallization conditions and the chemical structure of the macromolecules causes the properties of polymers to vary dramatically depending on the processing conditions. An example of a polymeric material with novel properties originating from a special nanoscale microstructure is shown in figs. 31 and 32. Polyethylene with a nanostructured morphology consisting of stacked crystalline lamellae (fig. 31a) exhibits remarkable elastic properties (fig. 32) if strained in tension in the direction perpendicular to the plane of the lamellae. The strain causes the lamellae to separate so that fibres of extended tie molecules form between them (fig. 31b). Upon unloading, the surface-energy of these molecular fibres causes them to shrink and thus pull the lamellar crystals together again. In other words, one obtains a material that can be strained reversibly by more than 100%. The restoring force (contraction) of the material is driven by surface energy and hence the material may be termed surface-energy pseudoelastic. If the stacked morphology is replaced by, e.g., a spherulitic microstructure, no such effects are noticed. In recent years, the large variety of nanostructured morphologies that may be generated for example in polymer blends or block copolymers has caused a rapidly expanding research activity in this type of materials (Martuscelli et al. [1980], Petermann [1991]). For further details, see ch. 32.

Fig. 30. Electron micrographs of the morphologies of a co-polymer consisting of polystyrene and polybutadiene blocks, as a function of the fraction of polystyrene blocks. The spacial arrangements of the polystyrene and polybutadiene in the solidified polymer are indicated in the drawings above the micrographs (Petermann [1991]).

Fig. 31. (a) Defocus electron micrograph showing the stacked-lamellae structure of a polyethylene fiber. The dark regions are the crystallites. The pattern of bright lines indicates the positions of the chain folds between the lamellae (cf. fig. 27). (b) Electron micrograph of a strained polyethylene fiber (cf. fig. 31a) showing complete separation of the lamellae interconnected by fibrils. The horizontal arrow indicates the straining direction (the strain is approximately 100%).

Fig. 32. Stress-strain curve for straining and destraining (in air) of a stacked-lamellae structure of polyethylene (cf. fig. 31) at 22° C at a strain rate of 0.0005/s. In the plateau region, the deformation occurs primarily by the separation of the lamellae and the formation of the fibrils between them (cf. fig. 31b).

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Carbon Nanotubes and Their Polymer Nanocomposites

Joseph Christakiran Moses , ... Biman B. Mandal , in Nanomaterials and Polymer Nanocomposites, 2019

5.4.1.2 Melt Processing

Thermoplastic semicrystalline polymers exhibit a unique property of softening when heated above their melting point. This is usually exploited during melt processing for fabricating intercalated polymer-CNT nanocomposites. Polymers that do not dissolve through solvents are usually processed by this technique, which involves melting the polymer and blending it with CNTs under high shear rates to obtain well-dispersed nanocomposite blends ( Sahoo et al., 2010). The nanocomposite blends can easily be postprocessed into desired formats through heat extrusion (Cooper et al., 2002). PMMA-MWNTs and polycarbonate-MWNTs are a few well known nanocomposites that have been fabricated using this approach (Spitalsky et al., 2010). Industrially, these nanocomposites are easy to produce, which just employs the Barbender twin-screw mixer to blend a polymer melt such as nylon (Zhang et al., 2004a) and a low density polyethylene (LLDPE) with CNTs. It has been seen that MWNTs dispersed in LLDPE resist the thermal and oxidative deterioration with respect to pure LLDPE. However, one disadvantage of this method is the poor dispersion of CNTs in the polymer melt; hence a thorough mixing of CNTs lower-loading concentration is preferred in order to reduce the viscosities as opposed to a higher CNTs loading (Sahoo et al., 2010).

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High-Temperature Engineering Thermoplastics

Vinny R. Sastri , in Plastics in Medical Devices (Second Edition), 2014

8.6.2 Properties of Polyaryletherketones

Polyaryletherketones are semicrystalline polymers and have very high strength, stiffness, and dimensional stability. They are also resistant to high heat, chemicals, hydrolysis, and high-energy radiation. Polyaryletherketones have excellent electrical properties over a wide range of temperatures. Carbon fiber and glass-reinforced grades provide additional heat resistance, strength, stiffness, and wear resistance. Table 8.11 gives the properties of unfilled PEEK, PEKK and a carbon fiber–filled PEEK (CF-PEEK); see Figure 8.24 for acronyms. The higher aromatic content in PEKK and PEKEKK is reflected in their higher glass transition temperatures and melt temperatures compared to PEEK (Figure 8.27).

Table 8.11. Properties of Polyaryletherketones

Property Unit PEEK PEKK PEKEKK 30% CF-PEEK
Density g/cc 1.31 1.31 1.3 1.41–1.44
Water absorption (24   h) % 0.5 < 0.2 < 0.5 0.06
Glass transition temperature °C 145 163 162 145
HDT at 0.46   MPa or 66   psi °C 160 — — —
HDT at 1.8   MPa or 264   psi °C 260–280 175 172 280–315
Melting point °C 334 360 387 340
Tensile strength at break MPa 90–110 110 115 200–220
Elongation % 20–40 10 20 1–5
Flexural modulus GPa 4.1 4.55 4.1 13–19
Impact strength, notched, 23°C J/m 55–65 69 60 54
Hardness rockwell M100 (R126) M88 — M70–M105
Processing temperature °C 345–390 345–370 375–395 350–400
Degree of crystallinity % 30–35 25–30 10–25 —

Figure 8.27. Thermal property comparison of polyaryletherketones.

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Properties of Neat (Unfilled) and Filled Fluoropolymers

Sina Ebnesajjad , Pradip R. Khaladkar , in Fluoropolymer Applications in the Chemical Processing Industries, 2018

3.7.2.2 Temperature-Related Properties

Fluoropolymers are semicrystalline polymers; most do not exhibit glass transition in the conventional sense during which all crystalline structures are converted to the amorphous. The glass transitions of fluoroplastics have been described as molecular relaxation (conformational disorder) that takes place in the amorphous phase of the polymer. These temperatures are also called second-order transitions; their value depends on the technique and the frequency of energy addition to the polymer sample. Table 3.70 presents these temperatures and melting points of perfluorinated and partially fluorinated fluoroplastics.

Table 3.70. Glass Transition Temperatures and Melting Points of Fluoroplastics [6,48,63–65]

Resin Glass Transition Temperatures, K Melting Point, K
Alpha (I) Beta Gamma (II)
PTFE 399 303 193 605
FEP 343–399 203–263 268–302 530–536
PFA 363 271 193 599
MFA – 258 – 583
PVDF 323, 373 235 203 483
ECTFE 413 363 208 536
ETFE 403 273 233 599
PCTFE 325 406 – 493

ECTFE, ethylene chlorotrifluoroethylene; ETFE, ethylene tetrafluoroethylene; FEP, fluorinated ethylene propylene; MFA, tetrafluoroethylene-perfluoromethyl vinyl ether; PCTFE, polychlorotrifluoroethylene; PFA, perfluoroalkoxy; PTFE, polytetrafluoroethylene; PVDF, polyvinylidene fluoride.

Some of the thermal properties of perfluoroalkoxy polymers (PFA and MFA) and FEP have been listed in Tables 3.71 and 3.72. Table 3.73 and Fig. 3.107 provide similar data for PVDF and Tables 3.74 and 3.75 for ETFE and ECTFE.

Table 3.71. Thermal Properties of PFA and MFA [6,47]

Property Temperature, °C MFR, g/10   min
13 2
Thermal conductivity of PFA, W/(m·K) 23 0.19 0.19
Coefficient of linear thermal
Expansion of PFA (10−5  mm/mm/°C)		 23–100 14 14
100–150 17 18
150–210 21 22
Heat capacity of PFA, J/(kg·K) – 1172 1172
Coefficient of linear thermal
Expansion of MFA (10−5  mm/mm/°C)		 23–150 12–20 12–20

MFA, tetrafluoroethylene-perfluoromethyl vinyl ether; MFR, melt flow rate; PFA, perfluoroalkoxy.

Table 3.72. Thermal Properties of FEP [48]

Property Temperature, °C MFR, g/10   min
7 3 1.5
Thermal conductivity, W/(m·K) 23 0.2 – –
Heat capacity, J/kg 23 5.1 – –
Specific heat, kJ/(kg·K) 25 0.240 0.242 0.268
100 0.266 0.267 0.294
150 0.288 0.290 0.315
Heat of combustion, kJ/kg – 5114
Coefficient of linear thermal
Expansion, 10−5  mm/mm/°C		 0–100 13.5 13.9 13.5
100–150 20.8 21.2 23.4
150–200 26.6 27.0 27.8
Deflection temperature, °C
0.455   MPa – 77 77 74
1.820   MPa 48 48 48

FEP, fluorinated ethylene propylene; MFR, melt flow rate.

Table 3.73. Thermal Properties of Polyvinylidene Fluoride [50]

Property Temperature, °C Value
Thermal conductivity, W/(m·K) 0.101–0.125
  Homopolymer 23 0.19
  Copolymer 0.17
Specific heat, kJ/(kg·K) 0.96–1.42
  Homopolymer 23 0.96
  Copolymer 1.30
Coefficient of linear thermal expansion, 10−5  mm/mm/°C 7.2–14.4
  Homopolymer 23 13
  Copolymer 14–16
Heat deflection temperature, °C at 1.820   MPa 84–118
  Homopolymer – 104–108
  Copolymer 50–72

Figure 3.107. Expansion of polyvinylidene fluoride versus temperature [50].

Table 3.74. Thermal Properties of ETFE [51]

Property Temperature, °C MFR, g/10   min
23 7 4
Thermal conductivity, W/(m·K) 23 – 0.24 –
Specific heat, kJ/(kg·K) 25 0.25
100 0.30
150 – 0.34 –
300 0.38
Heat of combustion, kJ/kg – – 13,700 –
Coefficient of linear thermal expansion, 10−5  mm/mm/°C 0–100 12.6 13.1 13.3
100–150 17.6 18.5 20.9
150–200 22.3 25.2 25.7
Deflection temperature, °C
0.455 MPa – 81 81 81
1.820   MPa 51 51 51

ETFE, ethylene tetrafluoroethylene; MFR, melt flow rate.

Table 3.75. Thermal Properties of ECTFE [52]

Property Temperature, °C MFR, g/10   min
2
Thermal conductivity, W/(m·K) 40 0.151
95 0.153
150 0.157
Specific heat, cal/(g·°C) 25 0.226
100 0.300
200 0.370
300 0.390
Coefficient of linear
Thermal expansion, 10−5  mm/mm/°C		 −30 to 50 8
50–85 10
85–125 13.5
125–180 16.5
Deflection temperature, °C
  0.455   MPa – 90
  1.820   MPa 63

ECTFE, ethylene chlorotrifluoroethylene; MFR, melt flow rate.

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Properties of Fluoropolymers

Sina Ebnesajjad , in Fluoroplastics (Second Edition), 2015

16.4.2 Temperature-Related Properties

Fluoropolymers are semicrystalline polymers; most do not exhibit glass transition in the conventional sense during which all crystalline structures are converted to the amorphous. The glass transitions of fluoroplastics have been described as molecular relaxation (conformational disorder) that takes place in the amorphous phase of the polymer. These temperatures are also called second-order transitions; their value depends on the technique and the frequency of energy addition to the polymer sample. Table 16.23 presents these temperatures and melting points of perfluorinated and partially fluorinated fluoroplastics.

Table 16.23. Glass Transition Temperatures and Melting Points of Fluoroplastics [1,5,25–27]

Resin Glass Transition Temperatures, K Melting Point, K
Alpha (I) Beta Gamma (II)
PTFE 399 303 193 605
FEP 343–399 203–263 268–302 530–536
PFA 363 271 193 599
MFA – 258 – 583
PVDF 323, 373 235 203 483
ECTFE 413 363 208 536
ETFE 403 273 233 599
PCTFE 325 406 – 493

Abbreviations: PVDF, polyvinylidene fluoride; ETFE, ethylene tetrafluoroethylene; ECTFE, ethylene chlorotrifluoroethylene; PFA, perfluoroalkoxy polymer; PTFE, polytetrafluoroethylene; PCTFE, polychlorotrifluoroethene.

Some of the thermal properties of perfluoroalkoxy polymers (PFA and MFA) and FEP have been listed in Tables 16.24 and 16.25. Table 16.26 and Figure 16.85 provide similar data for PVDF and Tables 16.27 and 16.28 for ETFE and ECTFE.

Table 16.24. Thermal Properties of Perfluoroalkoxy Polymers (PFA and MFA) [1,3]

Property Temperature, °C MFR, g/10   min
13 2
Thermal conductivity of PFA, W/(m   K) 23 0.19 0.19
Coefficient of linear thermal expansion of PFA, 10−5  mm/mm/°C 23–100 14 14
100–150 17 18
150–210 21 22
Heat capacity PFA, J/(kg   K) – 1172 1172
Coefficient of linear thermal expansion of MFA, 10−5  mm/mm/°C 23–150 12–20 12–20

Table 16.25. Thermal Properties of FEP [5]

Property Temperature, °C MFR, g/10   min
7 3 1.5
Thermal conductivity, W/(mK) 23 0.2 – –
Heat capacity, J/kg 23 5.1 – –
Specific heat, kJ/(kgK) 25 0.240 0.242 0.268
100 0.266 0.267 0.294
150 0.288 0.290 0.315
Heat of combustion, kJ/kg – 5114
Coefficient of linear thermal expansion, 10−5  mm/mm/°C 0–100 13.5 13.9 13.5
100–150 20.8 21.2 23.4
150–200 26.6 27.0 27.8
Deflection temperature, °C
0.455   MPa – 77 77 74
1.820   MPa 48 48 48

Table 16.26. Thermal Properties of Polyvinylidene Fluoride [14]

Property Temperature, °C Value
Thermal conductivity, W/(mK)
Homopolymer
Copolymer		 23 0.101–0.125
0.19
0.17
Specific heat, kJ/(kg K)
Homopolymer
Copolymer		 23 0.96–1.42
0.96
1.30
Coefficient of linear thermal expansion, 10−5 mm/mm/°C
Homopolymer
Copolymer		 23 7.2–14.4
13
14–16
Heat deflection temperature, °C @ 1.820 MPa
Homopolymer
Copolymer		 84–118
104–108
50–72

Figure 16.85. Expansion of polyvinylidene fluoride versus temperature [14].

Table 16.27. Thermal Properties of Ethylene Tetrafluoroethylene [7]

Property Temperature,°C MFR, g/10   min
23 7 4
Thermal conductivity, W/(mK) 23 – 0.24 –
Specific heat, kJ/(kg   K) 25 0.25
100 0.30
– –
150 0.34
300 0.38
Heat of combustion, kJ/kg – – 13,700 –
Coefficient of linear thermal expansion, 10−5  mm/mm/°C 0–100 12.6 13.1 13.3
100–150 17.6 18.5 20.9
150–200 22.3 25.2 25.7
Deflection temperature, °C
0.455   MPa – 81 81 81
1.820   MPa 51 51 51

Table 16.28. Thermal Properties of Ethylene Chlorotrifluoroethylene [8]

Property Temperature, °C MFR, g/10   min
2
Thermal conductivity, W/(m   K) 40 0.151
95 0.153
150 0.157
Specific heat, cal/(g   °C) 25 0.226
100 0.300
200 0.370
300 0.390
Coefficient of linear thermal expansion, 10−5  mm/mm/°C −30 to 50 8
50–85
85–125		 10
13.5
125–180 16.5
Deflection temperature, °C
0.455   MPa – 90
1.820   MPa 63

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NECKING PHENOMENA AND COLD DRAWING

L.J. Zapas , J.M. Crissman , in Viscoelasticity and Rheology, 1985

3 Experimental Procedures

Five different semicrystalline polymers were used in this study. The first was an experimental sample of isotactic polypropylene provided by the resin manufacturer. It had a viscosity average molecular weight of about 207,000 and contained 0.02 percent of stabilizer. Three of the samples were commercial grade high density linear polyethylenes having different molecular weights. Their weight average molecular weights were 99,000, 160,000, and 192,000, while their number average molecular weights were very nearly the same, being in the range from 15,000 to 16,000. The fifth sample was a commercial grade linear ultra high molecular weight polyethylene (UHMWPE) which, based on the manufacturer's method of estimating molecular weight from dilute solution viscometry measurements, had a molecular weight of approximately 4×106. As-received the first four samples were in the form of pellets, while the UHMWPE was in the form of a fine powder.

Flat sheets approximately 0.1 cm in thickness were prepared from each type of polymer by compression molding. The molding procedures differed depending upon the type of polymer, and the details of each molding operation can be found in the following references: Polypropylene-reference [7], the three high density linear polyethylenes-reference [6], and the UHMWPE-reference [11]. With the exception of the single step stress-relaxation experiments, the experiments were done at 23±0.5°C using a dumbbell shaped specimen cut with a die from the flat sheets. Because of the large deformations to which most of the specimens were subjected, the strain was determined with the aid of marks placed on the specimen and a cathetometer rather than an extensometer. The creep experiments were done under conditions of dead loading (constant applied engineering stress). Deformation histories involving constant rate of clamp separation and constant rate of loading were carried out on a servo-controlled hydraulic test machine.

For the single step stress-relaxation experiments the specimens were cut with a die which conformed to the geometry of the 'T-50' bar described in ASTM D599–61 [12]. In this geometry the width of the narrow section is constant over the entire portion of the specimen exposed between the grips. In order to avoid the possiblity that the attachment of an extensometer could contribute to the premature necking or failure of the speciment the strain was determined using a cathetometer.

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Gas Transport, Mechanical, Interphase, and Interdiffusion Properties in Coextruded-Multilayered Films

Deepak Langhe , Michael Ponting , in Manufacturing and Novel Applications of Multilayer Polymer Films, 2016

3.3.10 Nanoscale Confinement Effect on Mechanical Properties

Confinement of semicrystalline polymers such as PEO, PCL limits the spherulitic crystallization in nanolayered multilayered films as discussed in Section 3.1. Coextruded films of PEO against EAA or PS systematically changed the morphology from spherulites to truncated spherulites or discoids, to oriented in-plane lamellar structure with decreasing layer thickness from micro- to nanoscale. At thicknesses close to 20 nm, the confined morphology of PEO layers showed single lamellae with large aspect ratio, resembling single crystal structures. The effect of confined structures was investigated and modeled to identify the contribution of crystalline morphology and amorphous phase on mechanical properties. PEO/EAA multilayered films with 33, 257, and 1025 alternating layers with 50/50 (vol./vol.) composition were produced with overall film thicknesses ranging from 50 μm to 130 μm. The individual PEO layer thickness changed from 45 nm to 3600 nm [50,51]. A 1025-layered film with 10/90 PEO/EAA composition with 25 nm PEO layer thickness was also produced. Stress–strain behavior of the films was measured in uniaxial tension as shown in Figure 3.35. PEO control films exhibited brittle fracture at 14% and EAA control showed ductile behavior with fracture strain of 340%. Since the confinement did not impact the EAA crystallization behavior, the modulus change in EAA layers was assumed to remain constant. The calculated PEO modulus shows a threefold increase in the modulus as compared to PEO control. With decreasing layer thickness from 3600 nm to 45 nm, the tensile modulus increased from 486 ± 84 MPa to 1450 ± 99 MPa, at room temperature measured at 100%/min strain rate. Low temperature measurements at −10°C at 5%/min strain rate also showed a similar change as the modulus increased from 730 ± 80 MPa to 1240 ± 50 MPa over the same thickness range. PEO layers offered a unique opportunity to change the lamellar orientation from in-plane to on-edge by melt recrystallization approach, with selective melting of PEO layers followed by fast quenching, as described in Section 3.1. The measured tensile modulus in on-edge lamellae was independent of PEO layer thickness. PEO/EAA model system demonstrated a significant impact of the lamellar orientation on the modulus of multilayered film [50,51].

Figure 3.35. Stress–strain curves of EAA/PEO and controls (room temperature, 100%/min). Showing (a) films retained ductility of EAA and (b) increase in the modulus and yield stress, and decrease in the yield strain with decreasing layer thickness.

Deformation mechanism of the layered composites was investigated by stretching the films to different strains, from 0% to 400% and analyzing the structure of PEO/EAA layers with respect to different crystal populations in PEO layers. In-plane, on-edge, and mixed orientations were observed in the films at different strains. In 366 and 510 nm PEO layers, isotropic orientation of PEO layers in extruded films changed to on-edge lamellae with increasing strain of up to 100%. With further increase in strain to 400%, PEO chains became aligned with the deformation axis by chain pull-out and recrystallization mechanism. More than 85% orientation of the chains along the deformation axis was similar to PEO-fiber structure. As the layer thickness decreased to 125 nm, the spherulitic morphology in thick layers changed to loosely aligned stacked lamellae, in the layer direction in extruded films. Further reduction in PEO layer thickness to 25 nm showed single PEO lamellae, resembling single-crystal structure. Stretching of the films showed nonuniform deformation (micronecking) and lamellar alignment in the deformation direction. 25 nm layers showed strain-induced crystallization with 40% crystals still remained in-plane. The large amount of in-plane crystals maintained in the 125 and 25 nm layers was different than the thick layer deformation, where almost complete orientation of chains to deformation axis was observed. AFM images of deformation behavior in PEO layers are shown in Figure 3.36. Similar to PEO, the crystal orientation changed significantly in nanolayer films of PCL, PE, sPP, and iPP as demonstrated earlier. Possibility of similar shifts in the mechanical properties of multilayer composites containing these polymers has been suggested [51].

Figure 3.36. AFM images of PEO layer structure in PEO/EAA multilayered films with increasing strain. (Top row) 510 nm PEO layers; (bottom row) 125 nm PEO layers. The layer direction is vertical in the images.

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