Thermal Analysis of Polymeric Materials

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Fig. 7.52

all large-amplitude motion. This combination of the various arrested equilibria gives rise to several internal variables, which are, however, often not available.

7.2.3. Melting Transitions of Block Copolymers

The chemical structure of block copolymers is given by the number of blocks, their sequence, and their length, as is discussed in Fig. 1.19 and Sect. 3.4.1. The structure of a diblock copolymer poly(styrene-block-1,4-butadiene) of type AB, for example, can have the following chemical structure:

[CH2 CH(C6H5) ]400 [CH2 CH=CH CH2 ]840

with the corresponding segment lengths of molar masses 42,300, and 45,400 Da. At sufficiently large molar mass, the segments can phase separate under the restriction that the junctions between the different repeating units define the interfaces, as is described in Sect. 5.1.11 (see also the phase diagram Fig. 7.21). The liquid-liquid phase diagram is discussed in Sect. 7.1.6, and the lamellar structure of the chosen example of poly(styrene-block-1,4-butadiene) is reproduced in the AFM picture of Fig. 5.79. The lamellar spacing is 50 to 60 nm. At the chosen large molar mass, the phase structure is over the whole temperature range that of an amphiphilic liquid crystal (see Sect. 5.5).

A microphase structure as shown in Fig. 5.79 is sufficiently large to allow molecular mobility within the different lamellae which is hindered only close to the interfaces. The glass transitions of such block copolymers are described in Sect. 7.3. Depending on the nature of the segment, crystallization is possible within the lamellar superstructure. Depending on the space available, the typically nanometer-size crystals, described in Sect. 5.2 are unaffected and experience no restriction.

748 7 Multiple Component Materials

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A well studied example is given by the poly(oxyethylene-block-styrene). In case of atactic sequences of polystyrene, only the poly(oxyethylene), POE, can crystallize. A typical morphology of the POE is shown in Fig. 5.55. Single crystals of the copolymer can be grown from a common solvent which keeps both components mobile up to the time of crystallization of the POE-component. Figure 7.53 illustrates a growth spiral out of poly(oxyethylene-block-styrene), grown at 293 K from a solution of ethylbenzene (AB diblock, 28 wt-% oxyethylene block with a molar mass of about 10,000 Da). The crystal is comparable to the lamellar crystals of Fig. 5.55, i.e., the poly(oxyethylene) crystals are chain-folded with about 2.5 nm amorphous polystyrene layers at the interfaces.

Fig. 7.53

By shortening the length of the crystallizable segment, extended chain, oncefolded, or twice-folded crystals of poly(oxyethylene), the component Ax, could be grown from the melt for poly(oxyethylene-block-oxypropylene) [25]. Table 7.1 gives a list of the crystallinity and the length of the amorphous and crystalline lengths, a, and c, respectively, for diand symmetric tri-block copolymers.

Certain conclusions can be drawn about the changes in melting temperature with the presence of the second component. The poly(oxyethylene) homopolymer segments did not reach full crystallinity. The amorphous oxypropylene component is added to the intercrystalline layers in order to account for the amorphous layer thickness. The influence of the chain folding results in an increase in the surface area, but also allows for a larger crystallinity by reducing the crystalline layer thickness, and causes an additional decrease in the melting temperature. A comparison between the AB and BAB copolymers suggests that the entropic effect of mixing on fusion is not simply proportional to the oxypropylene concentration of the block copolymer, but is also influenced by the length of the crystallizable segments, and the achieved fold characteristics.

The melting temperature, measured at the melt-end by dilatometry, is proportional to the concentration of the crystallizing units, A, whenever the ratio of the of the B- sequence length, y, to the number of B-sequences in the molecule, n + 1, is less than 2.5. For larger B-sequence lengths the concentration effect decreases, most likely due to increasing phase separation in the melt. Figure 7.54 represents a plot of the changes in melting temperature observed when the short blocks are linked to an overall higher- molar-mass copolymer By(AxBy)n. A full assessment of the data is difficult since the initial steep decrease in melting point is due to chain folding (see Table 7.1).

The overall behavior of poly(oxyethylene-block-styrene) was also studied when crystallization was carried out from cooling of the melt. In case the polystyrene phase separates before crystallization of the poly(oxyethylene) component, the glassy lamellae restrict the crystal morphology. Figure 7.55 illustrates the global phase structure as evaluated by small-angle X-ray diffraction, SAXD, electron microscopy, and standard DSC. For the chosen sequence-lengths of the components, a liquidliquid phase separation occurs on cooling at 422 K, yielding a lamellar structure as illustrated in the electron micrograph (compare with Fig. 5.40 and the phase diagram of Fig. 7.21). This lamellar structure gets fixed at the glass transition of the polystyrene component.

750 7 Multiple Component Materials

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Fig. 7.54

Fig. 7.55

The poly(oxyethylene component can, however, still crystallize at 324 or lower temperatures. Figure 7.56 illustrates such restricted crystallization. At the lowest temperature, the resulting crystals are small and are more or less randomly arranged. Between 223 and 263 K, the orientation changes and the crystallized chains are arranged parallel to the global lamellae. Above 308 K, finally, enough mobility exists to orient the chains at right angles to the glassy polystyrene lamellae.

754 7 Multiple Component Materials

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Figure 7.59 illustrates the temperature-modulated DSC of the 50:50 PEBA [27]. The standard DSC trace is identical to the curve in Fig. 7.58, i.e., the crystallinity of E is 11%, and of A is 17%, as is expected for the changes in sequence lengths compared to Fig. 7.57. As in most other polymers, the heating traces show a certain amount of reversing melting which is frequency-dependent. The ether crystallinity has a melting temperature of 278 K, which compares to the pure oligomer of similar chain length (E28 2000 Da) which is about 55% and melts with a peak temperature of about 296 K. Both changes are indications of fewer and smaller (or less perfect) crystals. The amount of reversing melting for both the amide and ether crystals, normalized to crystallinity, increases with decreasing crystal perfection. The ether crystals have a higher reversibility than the pure oligomer crystals, which have practically none, since they are above the critical length of Fig. 3.91 (POTM2000) and yield an extended-chain crystal morphology on annealing.

The crystallinities of the various phase regions in the two-component systems in Figs. 7.57 to 7.61 are calculated using an adaptation of the method described in Fig. 4.80 for one component, considering the temperature dependence of all heat capacities and heats of fusion [26]. The two limiting heat capacities are the contribution of the two fractions, f(A) and f(E), multiplied with their crystalline and liquid heat capacities, Cpc and Cpa, respectively, as illustrated in Fig. 4.81. In Figs. 7.57 and 7.58 all semicrystalline heat capacities are marked, starting from the allsolid range up to the glass transition of the amorphous fraction of E, where (1 wc) of E becomes liquid. Next is the melting region of E at which the remaining wc of E also turns liquid. Shortly thereafter, at the glass transition of A the amorphous fraction of A turns liquid. Finally, beyond the melting of A, the whole sample is liquid.

Increasing the amide concentration, as seen in Figs. 7.60 and 7.61, the copolymer amide crystallinities of 35 and 37% exceed that of the corresponding nylon 12 of 30%. The copolymer melting peaks at about 436 K are about 10 K lower than seen for nylon 12. The amide crystals of the copolymers with enhanced crystallinity have less total reversibility than those of the homopolymer and the copolymers of less amide content have more [21]. One interprets these data by assuming that close to the phase boundary the ether blocks give an additional mobility to the crystallizing amide blocks. The higher reversibility of poorer crystals, in turn, seems to be a general observation for chain-folded polymers [28].

The ether-sequences, E, in Fig. 7.61 are only about 3 nm long and the crystallinity in the E-phase is only 1% due to restrictions by the surrounding glassy and 37% crystalline amide-phase, A, of 40 nm molecular chain length. The crystals seem to be located close to the phase boundary.

A different type of block copolymer can be produced by the introduction of long side chains into the polymer which by themselves may decouple and crystallize independently. Figure 7.62 displays the change of melting temperature of isotactic poly(1-alkene)s as a function of side-chain length:

[CH2 CH(Cx 2H2x 3) ]n.

These polymers are usually classified as homopolymers, but can be considered as block copolymers as soon as the side chains become sufficiently long to decouple. The polymers with short side chains have a crystal structure with a 2*3/1 helix, as

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Fig. 7.62

discussed in Sect. 5.1. 6 and 5.1.8. The side chain is part of the helix symmetry, as shown for example by the poly(1-butene) crystal structure of Fig. 5.25. As the side chain lengthens, the close packing of the side chain around the helix becomes more difficult. The polymers do not crystallize well, and the melting temperature decreases rapidly because of a decrease in heat of fusion. The minimum in the melting temperature coincides with a change in crystal structure which now permits a paraffinlike packing of the side chain and forces the backbone into a 2*4/1 helix. Each macromolecule is arranged in a plane with the side chains extending to both sides of the backbone. As a result, the melting points with the paraffin-like packing of the side-chain are identical of the paraffins with double the chain length, as can be derived from Fig. 7.62. On quenching such long-side-chain polymers from the melt, the side chains decouple and crystallize by themselves in a metastable crystal structure with a much lower melting temperature, leaving the entire backbone of the polymer in the amorphous portion. On annealing, recrystallization to crystals that include backbone and side chains is possible.

Figure 7.63 illustrates the change in melting temperatures of atactic side-chain polymers. As soon as the side chains are long enough, they crystallize, leaving the backbone in a decoupled, amorphous layer on top of the crystals. The melting temperatures are intermediate to paraffins of single and double the side-chain lengths. Figure 7.64, finally, shows the changes that occur by copolymerizing of the long sidechain polymers with short chain monomers. The poly(vinyl stearate)s show signs of phase separation and perhaps a blocky chemical structure. Poly(n-octadecyl acrylate- co-n-dodecyl acrylate)s form mixed crystals over the whole concentration region, while the three other acrylates show a reduction of the melting temperature as expected for compatible melts. Overall, these side-chain polymers behave more like decoupled small molecule systems than polymers.

756 7 Multiple Component Materials

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Fig. 7.63

Fig. 7.64

7.2.4. Melting Transitions of Regular Copolymers

The discussion of block copolymers in the last section included already a number of regular copolymers. An alternating block copolymer of type AxBy may well be considered a homopolymer, as illustrated by the example of nylon 6,6 [A is