Bones Concepts and Polymer Properties

D.A. Ivanov , in Polymer Science: A Comprehensive Reference, 2022

1.09.ii.1 Lamellar Addiction and Thermal Properties

Linear PE is certainly the most studied semicrystalline polymer. Contrary to the simplified view of the lamellar crystal given in Figure one , the PE single crystals in solution have a 3D grade of hollow pyramids with four or half-dozen 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 loftier 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 4 (110) faces with ii 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 take been suggested, 22 such as plastic deformation without reorientation of the crystalline stems with respect to the lamellar basal plane and flattening without whatever plastic deformation leading to concatenation tilting.

An instance 23 of a solution-grown PE single crystal having the shape of a truncated lozenge is given in Figure 7 . The tilt bending of the chains in the (110) and (200) sectors was measured in an SAED experiment by tilting the unmarried crystal around its crystallographic b-axis to obtain a diffraction blueprint corresponding to the (hk0) plane (cf. Figure seven ). Information technology can seen that to recover the (hk0) reflections, the rotating angle should exist 22° in (110) sectors and xxx° in (200) sectors, which clearly shows a deviation in their microstructure. An AFM height image obtained in borer mode on a similar single crystal is given in Figure 8 . The colour code is chosen in such a way that a pocket-size deviation between the thickness of the four (110) and two (200) sectors is rendered visible. As can be seen from the AFM paradigm cross-sections traced through each of the sectors (cf. Figure 8 , bottom), this meridian difference is less than ∼4% of the lamellar thickness, which is in agreement with the dissimilar stalk inclinations in the two sectors. The crystalline stalk 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 ever oriented along the crystallographic b-axis, which proves that the collapse of the lamellar hollow pyramids is not a random procedure.

Figure vii. Meridian: electron micrograph of a PE (〈K westward〉   =   32   100, polydispersity   =   1.1) single crystal formed from a dilute solution of due north-octane at 95   °C. Bottom: electron diffraction patterns obtained at a tilt bending 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, Chiliad.; Thierry, A.; et al. Macromolecules 2003, 36, 8376. 23

Figure 8. Top: Tapping mode atomic forcefulness 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 minor pinnacle departure betwixt (110) and (200) crystal sectors. Bottom: AFM image cross-sections traced as shown in the pinnacle panel.

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

The correlation between the morphology of single crystals of PE and their thermal backdrop was in focus of many studies. Melting temperatures of 124.v   °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–xxx fabricated the in situ studies of the thermal beliefs of PE crystals feasible. Effigy ix shows AFM images of a truncated PE single crystal successively measured at 121.9, 123.9, 124.ix, and 125.9   °C. No trace of melting is observed in the AFM image recorded at 121.9   °C. At 123.nine   °C, several cracks running perpendicular to the growth faces are observed in the two (200) sectors. A smaller number of similar cracks can as well exist detected on one of the (110) faces. The (200) sectors have completely melted at 124.ix   °C and recrystallized into thicker patches, with an average thickness ranging between 30 and twoscore   nm. The thickness of the recrystallized material is comparable to the summit of the central pleat running along the crystallographic b-axis. Moreover, holes were establish 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.nine   °C.

Figure 9. Tapping manner 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 correct), 124.9   °C (bottom left), and 125.9   °C (bottom right).

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

The images bear witness that the (200) sectors indeed melt outset. Moreover, the temperature difference with the melting point of the (110) sectors is comparable to what was expected from the referenced calorimetric studies. The complimentary enthalpies σ east of the fold surface in the (110) and (200) sectors have been estimated from the Gibbs–Thomson relationship to exist 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 prove for some jagging of the (200) growth faces stems from the fact that orientation of the decorating methane series rods on PE unmarried 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 10-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 by 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 past 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 tin, therefore, differ from the experiments on single-crystal mats. Even so, 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 , South.N. Magonov , in Polymer Science: A Comprehensive Reference, 2022

2.23.three.two.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 arrangement, crystallization, melting, and processing of semicrystalline polymers. Their behavior is essentially different from that of modest molecules because long bondage are much larger than the crystal thickness and different portions of the chain tin 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 scientific discipline withal 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 evolution of new techniques such as loftier-speed calorimetry, synchrotron radiation microfocus beams, and AFM. The latter two techniques enable resolution of crystalline superstructures downward 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 angle, spherulite system, and the morphology of amorphous stage. Agreement 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 meaning 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 every bit well equally copolymers. 253 Yet, 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 procedure. 257 The existent-fourth dimension AFM studies clearly show equal thicknesses for both dominant and subsidiary crystallites, which abnegate the assumption that the first grown lamellae should be thicker than the secondary crystallites.

Effigy 30. Crystallization of poly(ethylene terephthalate) was monitored past 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 notice the sectorization and make up one's mind the polymer concatenation 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 side by side reentry model. The other question is whether the chain packing and microstructure of the lamellar majority. The sensitivity of AFM to height measurements fabricated 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 side by side locations, which is consistent with the before TEM data. 259 Thickening gain gradually afterward a stepwise change at 115   °C, every bit reflected in the height histograms in Figure 31 (c). Studies of the thickening mechanisms in various polymers can exist 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. Pinnacle images of dry single crystals of PE measured at 110   °C temperature and later on 1.five   h of annealing at 115   °C. Height histograms corresponding to AFM bear witness the evolution of lamellar thickness after annealing at different temperatures.

 Reprinted with permission from Magonov, Due south. N.; Yerina, N. A.; Ungar, Chiliad.; et al. Macromolecules 2003, 36, 5637. 261 Copyright 2001 American Chemic Society.

An effective approach in elucidation of fine structural features of polymer crystals and the mechanisms of polymer crystallization is to employ model molecules. For instance, ultralong alkanes (C due north H2n+ii, north  >   150) are considered as an appropriate model for PE. Recent AFM studies have shown that the structure of unmarried crystals of C390H782 and PE is like, though the details of their thermal behavior are quite different. 261 Upon annealing, alkane crystals undergo a consummate series of transformations corresponding to stepwise unfolding from the folded-in-five conformation toward the fully extended-chain crystal, while the concatenation unfolding in PE crystals is a continuous and slower procedure. Another series of model compounds was proposed to selectively control the chain folding of polymers. 262 Structural instructions were encoded a linear courage that includes alternate crystallizable, long alkyl sequences of monodisperse sizes separated by short spacers containing side chains and acting equally 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 ∼five   nm that remains constant over a wide temperature range ( Figure 32 (b)). This arroyo allows controlling the lamellar thickness by steric interactions only, in contrast to previous attempts aimed at engineering polymer crystallization through hydrogen bonding. 263,264

Effigy 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 respective to the crystalline lamellae, wherein L a is the thickness of the amorphous layer. (b) AFM stage image of a sample measured in hard borer fashion at 85   °C. The thin white stripes represent 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, 2022

two.04.iii.1 Temperature Rising Elution Fractionation

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

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

The first stride consists of the consummate dissolution of the polymer in a adept solvent at high temperature. Afterwards 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 adjacent pace, the temperature in the column is decreased at a slow cooling charge per unit (typically i–2   °C   h−1). In this style, 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 tin 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 nearly of import ane 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). Every bit soon as the dissolution temperature of the polymer is reached, the layers deliquesce in the reverse order in which they were precipitated.

In A-TREF, the column temperature in the elution footstep is increased in a dull, abiding rate, while the polymer concentration in the eluent is monitored with an on-line mass-sensitive detector to obtain the TREF contour: the distribution of chain crystallizabilities in terms of the weight fraction of polymer eluted at each temperature. The chemic limerick distribution (CCD) and the tacticity tin be obtained from the TREF contour using a calibration curve.

In the case of P-TREF, unremarkably 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 than detailed information on the concatenation microstructure, TREF can likewise be combined with other fractionation and assay techniques, such equally GPC.

As mentioned earlier, the almost common use of TREF is for the decision 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 fabric, 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. Every bit expected, they institute out that the behavior of the copolymer is similar to that of the corresponding homopolymer if the comonomer content is depression. 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 like to an elastomeric fabric. Equally depicted in Figure 19 , TREF analysis shows that copolymers containing less than x% 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 eye range of comonomer composition.

Effigy xix. 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 ordinarily 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. Even so, the conventional IR detectors neglect 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 circuitous olefin copolymers. The dual-wavelength technique worked well for 2-component copolymer systems, whereas a multivariate calibration method is required for analyzing the IR spectra of more than complex multicomponent copolymer systems. Zhang investigated iii copolymer systems: ethylene–α-octene, PS-grafted ethylene–vinyl acetate, and ethylene–methyl acrylate copolymers. In addition to polymer concentration, the polymer composition (i.east., comonomer content) tin be measured by on-line FT-IR detection. This eliminates very fourth dimension- and labor-consuming TREF fraction collection equally well as the postfractionation compositional analyses by NMR and brings a benefit to the TREF analyses, especially for the circuitous olefin copolymers such every bit 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 of import 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 twenty shows a iii-dimensional plot of a Ziegler–Natta linear low-density polyethylene analyzed past ways of a combination of GPC and TREF.

Effigy 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 agile sites of a MgCl2-supported Ziegler catalyst, and to estimate the relationship betwixt polymer microtacticity and degradation rate of isotactic polypropylene (iPP). The former example revealed the conversion from high to low isospecific sites past 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 (Quaternary Edition), 1996

five.3. Semicrystalline polymers

Semicrystalline polymers found a separate course of nanostructured materials. The remarkable feature of this class of polymers is that the nanostructured morphology is ever formed if the polymers are crystallized from the melt or from solution unless crystallization occurs at high pressure or if high pressure level annealing is practical 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 bondage arrange in a periodic three-dimensional manner. The interfaces between neighboring crystals consist of both macromolecules folding back into the same crystal and necktie molecules that meander between neighboring crystals. The typical thicknesses of the crystal lamellae are in the order of x to twenty nm. These relatively small crystals thicknesses have been interpreted in terms of one of the following models. The starting time model hypothesizes the germination of the thin crystallites to result from nucleation kinetics. If the tiptop of the free 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 institute that the energy barrier of a critical nucleus consisting of extended chain molecules is larger than the bulwark pinnacle for a nucleus of folded chains. The physical reason for this energy divergence is equally 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 just one caste of freedom to reduce the energy barrier for the critical nucleus. This reduction occurs past adjusting the diameter of the needle. However, if chain folding occurs, the free energy barrier associated with the disquisitional nucleus tin be minimized by adjusting the size of the nucleus in all 3 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 concatenation-folding is based on the excess entropy associated with the folds relative to an extended-concatenation crystals. If the Gibbs gratis energies of an extended concatenation crystal and of a chain-folded crystal are compared, the concatenation folds are found to increase the internal energy of the organization. However, the chain folds also contribute to the entropy of the arrangement. Hence, at finite temperatures, a structure of lowest Gibbs free free energy is obtained, if a certain concentration of chain folds is present in the crystal. In other words, concatenation-folded crystals take a lower Gibbs costless 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 every bit a heavy line.

Polymers crystallizing from the molten state grade more than circuitous morphologies. Yet, the bones building blocks of these morphologies remain sparse lamellar crystals. Figure 28 shows spherulitic crystallization of thin molten polymer motion-picture show. 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 accept 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), loftier temperatures combined with high strain rates result in needle-similar arrangements (fig. 29b). Low temperatures and high strain rates lead to oriented micellar structures (fig. 29c). The transition betwixt 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 normally practice not form homogeneous mixtures simply separate on length scales ranging from a few nanometers to many microns depending on the thermomechanical weather condition of crystallization and the molecular construction of the costituents involved. So far the following types of nanostructured morphologies of polymer blends have been reported for blends fabricated up past one crystallizable and 1 amorphous (not-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 not-crystallizable component may be incorporated into the interlamellar regions of the spherulites of the crystallizable polymer. The spherulites are spacefilling. Blazon 3 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 2 crystallizable components, the four most frequently reported morphologies are: Type I morphology: Crystals of the 2 components are dispersed in an amorphous matrix. Type II morphology: 1 component crystallizes in a spherulitic morphology while the other crystallizes in a simpler mode eastward.chiliad., in the course of stacked crystals. Type 3 morphology: Both components exhibit a separate spherulitic structure. Type IV morphology: The two components crystallize simultaneously resulting in then-called mixed spherulites, which contain lamellae of both polymers.

Fig. 28. Brilliant field transmission electron micrograph (defocus dissimilarity) 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 night 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 ane crystallizable component has been studied for a diversity of macromolecular substances east.k., poly(ɛ-caprolactrone)/poly(vinylchloride), poly(2,6dimethyl-phenylene oxide)/isotactic polystyrene, atactic polystyrene/isotactic polystyrene blends.

Cake copolymers institute a tertiary 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 cardinal backbone of the macromolecules and/or in the class of side branches. An instance of a cake 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 arrangement 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 instance 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 airplane of the lamellae. The strain causes the lamellae to separate and so that fibres of extended tie molecules form betwixt 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 exist strained reversibly past more than 100%. The restoring force (wrinkle) 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.thou., a spherulitic microstructure, no such furnishings are noticed. In recent years, the large variety of nanostructured morphologies that may be generated for case 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 farther details, see ch. 32.

Fig. 30. Electron micrographs of the morphologies of a co-polymer consisting of polystyrene and polybutadiene blocks, as a office 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 cobweb. The dark regions are the crystallites. The pattern of brilliant 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 management (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/south. 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, 2022

v.4.ane.ii Melt Processing

Thermoplastic semicrystalline polymers exhibit a unique belongings of softening when heated in a higher place their melting point. This is usually exploited during cook processing for fabricating intercalated polymer-CNT nanocomposites. Polymers that do non deliquesce through solvents are usually processed past this technique, which involves melting the polymer and blending it with CNTs nether high shear rates to obtain well-dispersed nanocomposite blends ( Sahoo et al., 2022). The nanocomposite blends can easily be postprocessed into desired formats through rut extrusion (Cooper et al., 2002). PMMA-MWNTs and polycarbonate-MWNTs are a few well known nanocomposites that have been fabricated using this arroyo (Spitalsky et al., 2022). Industrially, these nanocomposites are like shooting fish in a barrel to produce, which just employs the Barbender twin-spiral mixer to blend a polymer cook such equally 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 gild to reduce the viscosities every bit opposed to a college CNTs loading (Sahoo et al., 2022).

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High-Temperature Engineering Thermoplastics

Vinny R. Sastri , in Plastics in Medical Devices (Second Edition), 2022

8.vi.two Properties of Polyaryletherketones

Polyaryletherketones are semicrystalline polymers and have very loftier strength, stiffness, and dimensional stability. They are also resistant to high estrus, chemicals, hydrolysis, and high-free energy radiation. Polyaryletherketones have excellent electrical backdrop over a wide range of temperatures. Carbon cobweb and glass-reinforced grades provide additional heat resistance, strength, stiffness, and wear resistance. Tabular array 8.xi gives the backdrop of unfilled PEEK, PEKK and a carbon cobweb–filled PEEK (CF-PEEK); run across Effigy viii.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. Backdrop of Polyaryletherketones

Property Unit of measurement PEEK PEKK PEKEKK 30% CF-PEEK
Density k/cc 1.31 ane.31 1.3 i.41–ane.44
Water absorption (24   h) % 0.5 < 0.two < 0.v 0.06
Drinking 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 force at break MPa ninety–110 110 115 200–220
Elongation % 20–forty 10 20 1–v
Flexural modulus GPa four.ane 4.55 four.1 13–nineteen
Impact strength, notched, 23°C J/m 55–65 69 threescore 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 eight.27. Thermal holding comparison of polyaryletherketones.

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Properties of Fluoropolymers

Sina Ebnesajjad , in Fluoroplastics (Second Edition), 2022

16.4.2 Temperature-Related Properties

Fluoropolymers are semicrystalline polymers; well-nigh 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 equally 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 add-on to the polymer sample. Table 16.23 presents these temperatures and melting points of perfluorinated and partially fluorinated fluoroplastics.

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

Resin Glass Transition Temperatures, K Melting Betoken, K
Alpha (I) Beta Gamma (2)
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 Effigy sixteen.85 provide similar information for PVDF and Tables 16.27 and xvi.28 for ETFE and ECTFE.

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

Property Temperature, °C MFR, g/x   min
xiii 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 xiv fourteen
100–150 17 18
150–210 21 22
Heat chapters PFA, J/(kg   Grand) – 1172 1172
Coefficient of linear thermal expansion of MFA, 10−v  mm/mm/°C 23–150 12–20 12–20

Table 16.25. Thermal Properties of FEP [v]

Property Temperature, °C MFR, yard/ten   min
7 3 1.5
Thermal electrical conductivity, W/(mK) 23 0.2 – –
Oestrus capacity, J/kg 23 v.1 – –
Specific rut, 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.v xiii.nine 13.five
100–150 20.8 21.2 23.iv
150–200 26.6 27.0 27.eight
Deflection temperature, °C
0.455   MPa – 77 77 74
1.820   MPa 48 48 48

Table xvi.26. Thermal Backdrop of Polyvinylidene Fluoride [fourteen]

Belongings Temperature, °C Value
Thermal conductivity, W/(mK)
Homopolymer
Copolymer		 23 0.101–0.125
0.19
0.17
Specific heat, kJ/(kg Grand)
Homopolymer
Copolymer		 23 0.96–one.42
0.96
one.30
Coefficient of linear thermal expansion, 10−5 mm/mm/°C
Homopolymer
Copolymer		 23 7.ii–14.four
13
fourteen–16
Rut 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   One thousand) 25 0.25
100 0.30
– –
150 0.34
300 0.38
Heat of combustion, kJ/kg – – thirteen,700 –
Coefficient of linear thermal expansion, x−5  mm/mm/°C 0–100 12.6 13.1 13.3
100–150 17.6 eighteen.5 twenty.9
150–200 22.iii 25.2 25.7
Deflection temperature, °C
0.455   MPa – 81 81 81
one.820   MPa 51 51 51

Tabular array 16.28. Thermal Properties of Ethylene Chlorotrifluoroethylene [8]

Holding Temperature, °C MFR, m/ten   min
2
Thermal conductivity, W/(m   K) xl 0.151
95 0.153
150 0.157
Specific oestrus, cal/(g   °C) 25 0.226
100 0.300
200 0.370
300 0.390
Coefficient of linear thermal expansion, x−v  mm/mm/°C −xxx to 50 8
50–85
85–125		 10
13.5
125–180 16.five
Deflection temperature, °C
0.455   MPa – 90
1.820   MPa 63

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Backdrop of Keen (Unfilled) and Filled Fluoropolymers

Sina Ebnesajjad , Pradip R. Khaladkar , in Fluoropolymer Applications in the Chemical Processing Industries, 2022

3.vii.2.2 Temperature-Related Backdrop

Fluoropolymers are semicrystalline polymers; virtually do not exhibit drinking glass transition in the conventional sense during which all crystalline structures are converted to the amorphous. The glass transitions of fluoroplastics take been described equally molecular relaxation (conformational disorder) that takes place in the amorphous stage 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. Drinking glass Transition Temperatures and Melting Points of Fluoroplastics [six,48,63–65]

Resin Glass Transition Temperatures, Grand Melting Point, Grand
Alpha (I) Beta Gamma (Two)
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 accept been listed in Tables 3.71 and three.72. Table 3.73 and Fig. 3.107 provide similar information for PVDF and Tables three.74 and 3.75 for ETFE and ECTFE.

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

Holding Temperature, °C MFR, thou/10   min
xiii two
Thermal conductivity of PFA, W/(m·K) 23 0.nineteen 0.19
Coefficient of linear thermal
Expansion of PFA (10−v  mm/mm/°C)		 23–100 fourteen 14
100–150 17 eighteen
150–210 21 22
Heat chapters of PFA, J/(kg·K) – 1172 1172
Coefficient of linear thermal
Expansion of MFA (x−five  mm/mm/°C)		 23–150 12–20 12–20

MFA, tetrafluoroethylene-perfluoromethyl vinyl ether; MFR, melt flow charge per unit; PFA, perfluoroalkoxy.

Table 3.72. Thermal Properties of FEP [48]

Belongings Temperature, °C MFR, grand/10   min
seven 3 1.v
Thermal electrical conductivity, W/(g·1000) 23 0.2 – –
Estrus capacity, J/kg 23 five.ane – –
Specific heat, kJ/(kg·G) 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.v 13.9 13.5
100–150 xx.8 21.2 23.iv
150–200 26.half dozen 27.0 27.8
Deflection temperature, °C
0.455   MPa – 77 77 74
ane.820   MPa 48 48 48

FEP, fluorinated ethylene propylene; MFR, cook flow charge per unit.

Table 3.73. Thermal Backdrop of Polyvinylidene Fluoride [fifty]

Property Temperature, °C Value
Thermal conductivity, West/(thou·One thousand) 0.101–0.125
  Homopolymer 23 0.nineteen
  Copolymer 0.17
Specific rut, kJ/(kg·K) 0.96–one.42
  Homopolymer 23 0.96
  Copolymer one.30
Coefficient of linear thermal expansion, 10−5  mm/mm/°C vii.2–14.iv
  Homopolymer 23 thirteen
  Copolymer fourteen–16
Estrus 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]

Belongings Temperature, °C MFR, g/10   min
23 seven 4
Thermal conductivity, West/(thou·Grand) 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.seven
Deflection temperature, °C
0.455 MPa – 81 81 81
ane.820   MPa 51 51 51

ETFE, ethylene tetrafluoroethylene; MFR, melt flow rate.

Tabular array 3.75. Thermal Properties of ECTFE [52]

Property Temperature, °C MFR, g/10   min
ii
Thermal conductivity, Due west/(m·Chiliad) 40 0.151
95 0.153
150 0.157
Specific rut, cal/(thou·°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 viii
l–85 10
85–125 xiii.5
125–180 xvi.v
Deflection temperature, °C
  0.455   MPa – xc
  1.820   MPa 63

ECTFE, ethylene chlorotrifluoroethylene; MFR, melt menses charge per unit.

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NECKING PHENOMENA AND COLD DRAWING

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

3 Experimental Procedures

Five dissimilar semicrystalline polymers were used in this study. The first was an experimental sample of isotactic polypropylene provided past the resin manufacturer. It had a viscosity average molecular weight of about 207,000 and contained 0.02 percent of stabilizer. 3 of the samples were commercial class high density linear polyethylenes having dissimilar 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 aforementioned, being in the range from 15,000 to xvi,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. Every bit-received the first iv samples were in the grade of pellets, while the UHMWPE was in the form of a fine powder.

Apartment sheets approximately 0.1 cm in thickness were prepared from each blazon of polymer by compression molding. The molding procedures differed depending upon the type of polymer, and the details of each molding operation can be plant in the following references: Polypropylene-reference [7], the three high density linear polyethylenes-reference [6], and the UHMWPE-reference [xi]. With the exception of the single step stress-relaxation experiments, the experiments were done at 23±0.five°C using a dumbbell shaped specimen cut with a die from the flat sheets. Because of the large deformations to which virtually 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 weather condition of dead loading (constant applied engineering stress). Deformation histories involving constant charge per unit of clamp separation and constant rate of loading were carried out on a servo-controlled hydraulic exam auto.

For the single footstep stress-relaxation experiments the specimens were cut with a dice 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 avert the possiblity that the zipper 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 Backdrop in Coextruded-Multilayered Films

Deepak Langhe , Michael Ponting , in Manufacturing and Novel Applications of Multilayer Polymer Films, 2022

3.iii.10 Nanoscale Solitude Effect on Mechanical Properties

Confinement of semicrystalline polymers such equally PEO, PCL limits the spherulitic crystallization in nanolayered multilayered films equally discussed in Department three.i. 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 bars morphology of PEO layers showed single lamellae with big aspect ratio, resembling single crystal structures. The result of confined structures was investigated and modeled to identify the contribution of crystalline morphology and amorphous stage on mechanical properties. PEO/EAA multilayered films with 33, 257, and 1025 alternating layers with l/l (vol./vol.) limerick 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 motion picture 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 equally shown in Figure 3.35. PEO control films exhibited breakable fracture at xiv% and EAA control showed ductile behavior with fracture strain of 340%. Since the confinement did not bear on the EAA crystallization behavior, the modulus change in EAA layers was causeless 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 charge per unit. Low temperature measurements at −10°C at 5%/min strain rate also showed a similar change equally the modulus increased from 730 ± 80 MPa to 1240 ± 50 MPa over the same thickness range. PEO layers offered a unique opportunity to modify the lamellar orientation from in-plane to on-edge by melt recrystallization approach, with selective melting of PEO layers followed by fast quenching, equally described in Section 3.1. The measured tensile modulus in on-edge lamellae was independent of PEO layer thickness. PEO/EAA model arrangement demonstrated a significant bear on 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 subtract 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 construction of PEO/EAA layers with respect to unlike crystal populations in PEO layers. In-airplane, on-edge, and mixed orientations were observed in the films at unlike 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 upwards to 100%. With further increase in strain to 400%, PEO chains became aligned with the deformation centrality by chain pull-out and recrystallization machinery. More than 85% orientation of the chains along the deformation centrality was similar to PEO-cobweb 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. Farther 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 management. 25 nm layers showed strain-induced crystallization with 40% crystals still remained in-aeroplane. The large corporeality of in-aeroplane 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 Effigy iii.36. Like to PEO, the crystal orientation changed significantly in nanolayer films of PCL, PE, sPP, and iPP as demonstrated before. 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 construction 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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